Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers: A Selection of Papers from the 9th International Symposium, Paris, France, 27 June 1 July 2005

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers A Selection of Papers from the 9th International Symposium, Paris, France, 27 June – 1 July 2005

Edited by Philippe Marcus Laboratoire de Physico-Chimie des Surfaces, CNRS, Ecole Nationale Supérieure de Chimie de Paris, France

Vincent Maurice Laboratoire de Physico-Chimie des Surfaces, CNRS, Ecole Nationale Supérieure de Chimie de Paris, France

Amsterdam – Boston – Heidelberg – London – New York – Oxford Paris – San Diego – San Francisco – Singapore – Sydney – Tokyo

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v

Preface This book contains a selection of papers presented at PASSIVITY-9, the 9th International Symposium on the Passivation of Metals and Semiconductors and the Properties of Thin Oxide Layers, which was held in Paris, 27 June – 1 July, 2005. One hundred and twelve peer-reviewed manuscripts have been included. The book covers all the fundamental and applied aspects of passivity and provides a relevant and updated view of the advances and new trends in the field. It is structured in ten sections: • • • • • • • • • •

Growth, (Nano)structure and Composition of Passive Films, Passivity of Semiconductors, Electronic Properties of Passive Films, Passivity Issues in Biological Systems, Passivity in High-Temperature Water, Mechanical Properties of Passive Films, Passivity Issues in Stress Corrosion Cracking and Tribocorrosion, Passivity Breakdown and Localized Corrosion, Modelling and Simulation, Surface Modifications and Inhibitors (for Improved Corrosion Resistance and/or Adhesion),

The editors would like to thank all the authors for their contributions.

Paris, 9 November 2005 The Editors Philippe Marcus, Vincent Maurice

vi

vii

Contents Preface Section A

v Growth, (Nano)structure and Composition of Passive Films

Electrochemical Properties of Fe-Cr-Mo Alloys and Fe2O3-Cr2O3-MoO2 Artificial Passivation Films in 1 M HCl K. Sugimoto, M. Saito, N. Akao and N. Hara Formation and growth processes of electrochemical passive layers (borate medium: pH 9.2) and electron stimulated oxidized films (5.10-6 Pa O2) formed on Fe-Cr alloys M. Bouttemy, M. Bertoglio and G. Lorang

1

3

9

Development and Composition of the High Temperature Oxide Film Grown on Fe-15Cr during Annealing E. Park and M. Spiegel

15

Passivity of Nickel-Containing Stainless Steels in Concentrated Sulphuric Acid M.B. Ives, Y. Li and K.S. Coley

23

An insight on the role of Nickel in the passive films generated on different stainless steels C.M. Abreu, M.J. Cristóbal, R. Losada, X.R. Nóvoa, G. Pena and M.C. Pérez

29

Passivity of Nitrogen-Bearing Stainless Steel in Acidic Solution S. Martinez, M. Metikoš-Hukovic and N. Lajci

35

Passive behaviour of stainless steels and nickel in LiBr solution at different temperatures A. Igual Muñoz, J. García Antón, J.L. Guiñón Segura and V. Pérez Herranz

41

The effect of the Cerium ion implantation in the passive films properties of a duplex stainless steel C.M. Abreu, M.J. Cristóbal, X.R. Nóvoa, G. Pena, M.C. Pérez and C. Serra

47

Use of Alloy 22 as Long-Term Radioactive Waste Containment Material O. Pensado, R. Pabalan, D. Dunn and Kuang-Tsan Chiang Effect of temperature and melt composition on the passivity of a Ni-10%Cr alloy in a molten electrolyte M. Bojinov, P. Gencheva and T. Tzvetkoff

53

59

Contents

viii Spontaneous Passivation of Amorphous Bulk Ni-Cr-Mo-Ta-Nb-P Alloys in Concentrated HCl K. Hashimoto, H. Shinomiya, A. Nakazawa and K. Asami Passivity of Fe90V10 and Fe75Cr15V10 in Alkaline Media A. Drexler and H.-H. Strehblow

65 71

Effect of anodic passivation on the corrosion behaviour of Fe-Mn-Al steels in 3.5%NaCl A.S. Hamada, L.P. Karjalainen and M.A. El-Zeky

77

Surface Characterization of 1018 Carbon Steel in Borate Medium by in-situ Electrochemical Scanning Tunneling Microscopy I. González, R. Cabrera-Sierra and N. Batina

83

Iron passivation studied by in situ Raman spectroscopy on Fe/Au(111) epitaxial films S. Joiret and P. Allongue

89

Atomic-Structure Characterization of Passive Film of Fe by Grazing Incidence X-ray Scattering at SPring-8 M. Sato, M. Kimura, M. Yamashita, H. Konishi, S. Fujimoto, Y. Tabira, T. Doi, M. Nagoshi, S. Suzuki, T. Kamimura, T. Nakayama and T. Ohtsuka

95

Electrochemical aspects of the behavior of perchlorate ions in the presence of iron group metals G.G. Láng, T.A. Rokob, M. Ujvari and G. Horányi

101

The effect of Al3+ in the passivity of iron in alkaline media containing chlorides C.M. Abreu, M.J. Cristóbal, L. Freire, X.R. Nóvoa, G. Pena and M.C. Pérez

107

Nanoscale modifications of a Ni(111) surface during nucleation and growth of the passive film A. Seyeux, V. Maurice, L.H. Klein and P. Marcus

113

Simultaneous Anodic Dissolution and Passivation of Nickel in Moderate Acid Medium J. Gregori, J.J. García-Jareño, D. Giménez-Romero and F. Vicente

119

Copper Passivity in Carbonate Base Solutions and its Application in Chemical Mechanical Planarization (CMP) E. Abelev, D. Starosvetsky, M. Auinat and Y. Ein-Eli

125

Analysis of Cu corrosion product in aqueous lithium bromide concentrated solutions M.J. Muñoz-Portero, J. García-Antón, J.L. Guiñón and V. Pérez-Herranz

131

Contents Passivity of Tin and CuSn Alloys in Alkaline Media studied by X-ray Photoelectron Spectroscopy P. Keller and H.-H. Strehblow Surface Analytical Characterization of Chromium Passivation on Tinplate R. Sandenbergh, M. Biermann and T. von Moltke

ix

137 143

To Passivate or not to Passivate, that is the Question: The Case of Barium Tin(II) Chloride Fluorides G. Dénès and A. Muntasar

149

A Thin Passivating Tin(IV) Oxide Layer on Tin(II)-Containing Fluoride Particles, or not? The M1-xSnxF2 Solid Solutions G. Dénès, E. Laou, M. Cecilia Madamba, A. Muntasar and Z. Zhu

155

Electrochemical Data About Disruption of Passivating Films. The Pb/PbSO4/H2SO4 System C. V. D’Alkaine and P. M. P. Pratta

161

Morphology, composition and structure of anodic films on binary Al-Cu alloys J. Idrac, P. Skeldon, Y. Liu, T. Hashimoto, G. Mankowski, G. Thompson and C. Blanc

167

Repassivation of aluminium during the AC-graining process by aluminium hydroxide formation Th. Dimogerontakis, H. Terryn and P. Campestrini

173

Anodization of Ti : Formation of Self-Organized Titanium Oxide Nanotube-Layers P. Schmuki, H. Tsuchiya, L. Taveira, K. Sirotna and J.M. Macak

179

Self-Organized Nanoporous Valve Metal Oxide Layers H. Tsuchiya, J. Macak, I. Sieber, L. Taveira and P. Schmuki

187

Studies on transition of titanium from active into passive state in phosphoric acid solutions E. Krasicka-Cydzik

193

The Influence of Temperature on the Passive Film Properties of ASTM Grade-7 Titanium J.J. Noël, L. Yan, D. Ofori, P. Jakupi and D.W. Shoesmith

199

Chromatic properties of anodised titanium obtained with two techniques MP. Pedeferri, B. Del Curto and P. Pedeferri

205

Contents

x Amorphous-to-Crystalline Transition of Anodic Niobia H. Habazaki, T. Ogasawara, H. Konno, K. Shimizu, K. Asami, S. Nagata, P. Skeldon and G.E. Thompson Ellipsometric Analysis of Growth Process of Oxidation Films on Magnesium and its Alloys in Neutral Aqueous Solutions N. Hara, D. Kagaya and N. Akao Corrosion Protection of Magnesium Alloys by Ce, Zr and Nb oxide layers H. Ardelean, I. Frateur and P. Marcus Section B

Passivity of Semiconductors

Passivation of (100) and (111) Silicon in KOH Solution H.G.G. Philipsen and J.J. Kelly

211

217 225

231 233

On the nature of oscillations of Open Circuit Potential of silicon immersed in CuSO4/HF solution V. Parkhutik, Y. Makushok and Y. Ogata

239

Passivation and Local Corrosion of p-Silicon in Anhydrous Organic Solutions of Chlorides U. Lelek-Borkowska and J. Banas

245

Silicon Texturing Under Negative Potential Dissolution (NPD) Conditions Y. Ein-Eli, N. Gordon and D. Starosvetsky

251

Fermi Level Pinning at n-GaAs(110) Electrodes V. Lazarescu, M. Gartner, R. Scurtu, M. Anastasescu, A. Ghita, M.F. Lazarescu and W. Schmickler

257

Composition and Growth of Thermal and Anodic Oxides on InAlP S. Kleber, M.J. Graham, S. Moisa, G.I. Sproule, X. Wu, D. Landheer, A.J. SpringThorpe, P.J. Barrios and P. Schmuki

263

High-k gate oxide formed by anodic oxidation for organic field effect transistors M. Erouel, A. Gagnaire, A.L. Deman, J. Tardy, N. Jaffrezic-Renault, Z. Sassi, J.C. Bureau and M.A. Maaref Section C

Electronic Properties of Passive Films

Photoelectrochemical analysis of the passive film formed on Fe-Cr-Ni alloys in pH 8.5 buffer solution H. J. Jang and H. S. Kwon

271

277

279

Contents

xi

Photoelectrochemical Response and Corrosion Property of Passive Films on Fe-18Cr Alloy S. Fujimoto

285

Effects of EDTA on the Electronic Properties of Passive Film on Fe-20Cr in pH 8.5 Buffer Solution E. A. Cho, S. J. Ahn and H. S. Kwon

291

Electrochemical and Mott-Schottky Behaviour of the Oxide Films on I625 Exposed to Gamma Radiation J. F. Dante, D. S. Dunn and K. T. Price

299

Oxygen Reduction on Passive Steel and Cr Rich Alloys for Concrete Reinforcement A. A. Sagüés, S. Virtanen and P. Schmuki

305

Diffusivity of point defects in the passive film on Fe S. J. Ahn and H. S. Kwon

311

Electrochemical Impedance of Thin Rust Film Fabricated Artificially M. Itagaki, H. Araki, K. Watanabe, H. Katayama and K. Noda

317

A new approach to describe the passivity of nickel and titanium oxides R. Cabrera-Sierra, I. González, J. Ávalos-Martínez, G. Vázquez and M. A. Pech-Canul

325

Relation of the Photocurrents to the Corrosion Rates of Pure Aluminum having Various Oxides M.C. Romanes, K.A. Donovan and T.D. Burleigh Growth and Characterization of Anodic Films on Al-Nb Alloys M. Santamaria, F. Di Quarto, M. Gentile, P. Skeldon and G.E. Thompson Amorphous semiconducting passive film-electrolyte junctions revisited. The influence of a non homogeneous density of state on the differential admittance behaviour of anodic a-Nb2O5 F. La Mantia, M. Santamaria and F. Di Quarto Section D

Passivity Issues in Biological Systems

Influence of Protein Adsorption on the Passivation of Dental Amalgams C.M.A. Brett, E. Jorge, C. Gouveia-Caridade and H. Dias BSA adsorption on Fe-17Cr in acid solution: electrochemical behaviour and surface composition L. Lartundo-Rojas, I. Frateur, A. Galtayries and P. Marcus

331 337

343

349 351

357

Contents

xii XPS characterisation of BSA adsorption on stainless steel S. Zanna, C. Compère and P. Marcus Study of corrosion behaviour of Ti-coated AISI 316L stainless steel in a simulated body fluid solution F. Hellal, F. Atmani, B. Malki, H. Sedjal, M. Kerkar and F. Dalard Effect of Al on the passivity of Ti base implant alloys S. Virtanen, H. Hildebrand and M. Ruczickova

365

371 377

Improved adhesion of titanium oxide film to titanium-base alloy by Ti/O compositional gradient using reactive sputter-deposition T. Sonoda, A. Watazu, K. Katou and T. Asahina

383

Liquid crystal behavior in solutions, electrode passivation, and impedance loci in four quadrants C.V. Krishnan and M. Garnett

389

Section E

Passivity in High-Temperature Water

Passive film growth and oxide layer restructuring on stainless steel in a high-temperature borate electrolyte M. Bojinov, P. Kinnunen, K. Lundgren and G. Wikmark Kinetics of passivation of nickel-base alloys (Alloy 600 and Alloy 690) in high temperature water A. Galtayries, A. Machet, P. Jolivet, P. Scott, M. Foucault, P. Combrade and P. Marcus Behaviour of oxide films in boric acid – lithium hydroxide solutions V. Ignatova, M. Vankeerberghen, S. Gavrilov, R.-W. Bosch, S. Van Dyck and S. Van den Berghe The Effect of CO2 and H2S on the Passivation of Chromium and Stainless Steels in Aqueous Solutions at Elevated Temperature and under High Pressure J. Banas, B. Mazurkiewicz, U. Lelek-Borkowska, H. Krawiec, W. Solarski and K. Kowalski Effect of Lead on Passivation of Alloy 600 Surface Z. Zhou, J. Park, J.E. Indacochea, R.W. Staehle, S. S. Hwang, N. Finnegan and R. Haasch Interaction of Oxide Layers on Structural Materials with Light Water Reactor Coolants - its influence on the mechanism of oxide growth and restructuring M. Bojinov and B. Beverskog

395

397

403

411

417

425

431

Contents Section F

xiii Mechanical Properties of Passive Films

Differences in Mechanical Properties of the Passive Metal Surfaces Obtained in Solution and Air M. Seo, D. Kawamata and M. Chiba Current Transients of Passive Iron during Micro-indentation in Solution K. Fushimi, Ken-ichi Takase and M. Seo

437

439 451

Mechanical Properties of Single-Crystal Tantalum (100) Surface Covered with Anodic Oxide Film D. Kawamata and M. Seo

457

Study of the mechanical effects on passivity breakdown by local probe techniques V. Vignal, N. Mary and R. Oltra

463

Mechanical properties of metastable r.f magnetron sputter-deposited Al1-xCux thin films M. Draissia and M.Y. Debili

469

Section G Passivity Issues in Stress Corrosion Cracking and Tribocorrosion

475

Passivity Issues in Tribocorrosion D. Landolt

477

Wear in nuclear power plants, a tribocorrosion approach A. Beaudouin, P. Combrade, D. Kaczorowski and J-P. Vernot

489

Comparison between tribocorrosion mechanisms of Stellite 6 and Zircaloy 4 in LiOH-H3BO3 solutions V.-E. Iordache, F. Wenger, P. Ponthiaux, A. Ambard, J. Peybernès and J. Vallory

495

Detrimental effect of lead on the passivity of UNS N06690 alloy B.T. Lu, J.L. Luo and Y.C. Lu

501

The role of Borides in the Passivity and SCC of Alloy 600 Y. Yi, H. Kim and J. Kim

507

Effect of Nickel and Tungsten on Repassivation Rate of Stainless Steels in Chloride Solution by Potential Step Chronoamperometry S.-Y. Kim, H. Kim and H.-S. Kwon

513

Dislocations effect on kinetic of passivation of polycristalline nickel in H2SO4 medium M. Sahal, C. Savall, J. Creus, R. Sabot and X. Feaugas

519

Contents

xiv Electrochemical potential oscillations during galvanostatic passivation of copper in NaNO2 solution and their role in TGSCC mechanism Y. Yagodzinskyy, P. Aaltonen and H. Hänninen

525

Effects of straining on oxide films and passivity of copper in nitrite solution at ambient temperature P. Aaltonen, Y. Yagodzinskyy and H. Hänninen

531

Repassivation Kinetics of Al-Alloys for Aircraft Structures J. Wloka, T. Hack and S. Virtanen Section H

Passivity Breakdown and Localized Corrosion

Localized Corrosion Growth Kinetics in AA7178 G.S. Frankel, Tsai-Shang Huang and X. Zhao

537

543 545

An in situ AFM study of the first steps of localised corrosion on a stressed 304L stainless steel in chloride media F. Martin, S. Fréchard, C. Bataillon and J. Cousty

555

Initial stage of localized corrosion in artificial pit formed on zinc coated steels by photon rupture M. Sakairi, Y. Uchida and H. Takahashi

561

Influence of the chemical dissolution of MnS inclusions on the composition of passive films and the local electrochemical behaviour of stainless steels H. Krawiec, V. Vignal, R. Oltra and O. Heintz

567

Electrochemical characterization of corrosion resistant alloys in chloride solutions F. Bolzoni, P. Fassina, G. Fumagalli and S. Goidanich

573

Effect of HCl on Pickling of 304 Stainless Steel in Iron Chloride-Based Electrolytes L.-F. Li, M. Daerden, P. Caenen and J.-P. Celis

579

Influence of thermal oxides on pitting corrosion of austenitic and duplex steels V. Alar, V. Rede, I. Juraga and B. Runje

585

Modifications in the electrochemical behaviour of SAF 2205 in alkaline media induced by Cl- ions C. M. Abreu, M.J. Cristóbal, R. Losada, X.R. Nóvoa, G. Pena and M.C. Pérez

591

Contents

xv

In-situ Raman spectroscopy study of iron and carbon steel corrosion in mineral water L. Lanarde, S. Joiret, X. Campaignolle and M. Meyer

597

Effect of Noble Element Alloying on Passivity and Passivity Breakdown of Ni Y. H. Kim, G. S. Frankel and J. C. Lippold

603

Analysis of Electrochemical Noise of Pure Aluminium in Sulphate and Molybdate Ion-Containing 0.1 M NaCl Solution S.-I. Pyun and K.-H. Na

609

Localized corrosion of 2024 alloy: structure and composition of oxide films grown on model alloys representative of the different metallurgical phases C. Blanc, G. Mankowski, C. Dufaure, C. Mijoule and Y. Kihn Passivity Breakdown of Aluminum Alloys by Trace Element Lead K. Nisancio lu, Øystein Sævik and Y. Yu Passivity Breakdown of Aluminium Alloys by Surface Enrichment of Group IIIA - VA Trace Elements B. Graver, Øystein Sævik, Y. Yu and K. Nisancioglu Electrochemical behaviour of Al and some of its alloys in chloride solutions S. Zein El Abedin and F. Endres Section I

Modelling and Simulation

Cooperative Spreading of Pit Sites as a New Explanation for Critical Threshold Potentials J.R. Scully, N.D. Budiansky, L. Organ, A.S. Mikhailov and J.L. Hudson

615 621

627 633

639

641

Atomistic simulation of the passivation of iron-chromium alloys using calculated local diffusion activation barriers B. Diawara, Yves-Alain Beh and P. Marcus

651

DFT study of the interactions of Cl- with passivated Nickel surfaces: energetic and structural aspects B. Diawara, A. Bouzoubaa, N. Pineau, C. Minot, V. Maurice and P. Marcus

659

Simulation of corrosion processes with anodic and cathodic reactions separated in space J. Stafiej, A. Taleb, C. Vautrin-Ul, A. Chaussé and J.P. Badiali

667

Contents

xvi Section J

Surface Modifications and Inhibitors (for Improved Corrosion Resistance and/or Adhesion)

On the Interaction of Organic Molecules with Metal Oxide Surfaces H. Terryn, O. Blajev, S. Van Gils, J. van den Brand, A. Ithurbide, E. Vandeweert, J. Snauwaert, A. Hubin and C. Van Haesendonck

673 675

Formation of Al-Si Composite Oxide Films on Aluminum by Electrophoretic Sol-Gel Coating / Anodizing H. Takahashi, M. Sunada, T. Kikuchi, M. Sakairi and S. Hirai

685

In-situ Raman Spectroscopy and Spectroscopic Ellipsometry study of the iron/Polypyrrole interface T. Van Schaftinghen, S. Joiret, C. Deslouis and H. Terryn

691

Spontaneous grafting of iron surfaces by reduction of aryldiazonium salts in acidic water. Applications to the inhibition of iron corrosion F. I. Podvorica, C. Combellas, M. Delamar, F. Kanoufi and J. Pinson

697

Conversion coating of zinc heptanoate in aqueous media on electrogalvanized steel S. Jacques, E. Rocca, M-J. Stébé, H. Derule, N. Genet and J. Steinmetz

703

Study of carbon steel corrosion inhibition in alkaline solution by means of EIS G. Rondelli, L. Lazzari, M. Ormellese, R. Novoa and E. Pérez

709

Nitrided/nitrocarburized and oxidized steel: Corrosion data in dependence on the N- and C-content of the ε-phase under the oxide layer U. Ebersbach

715

Passivation treatments on Zn and Zn alloys substrates involving Mo and W compounds L. Anicai, A. Petica, M. Buda and T. Visan

721

Electrochemical Properties and Applications of Sputtered Iridium Oxide Thin Films E. P. Slavcheva, U. Schnakenberg and W. Mokwa

729

Study of DSA® deactivation E. Herrera Calderon, R. Wüthrich, H. Bleuler and Ch. Comninellis

737

Index

743

Section A Growth, (Nano)structure and Composition of Passive Films

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Electrochemical Properties of Fe-Cr-Mo Alloys and Fe2O3-Cr2O3-MoO2 Artificial Passivation Films in 1 M HCl Katsuhisa Sugimotoa,b,c, Mamoru Saitoa, Noboru Akaoa , Nobuyoshi Haraa a

Department of Metallurgy, Graduate School of Engineering, Tohoku University,6-602, Aza-Aoba, Aramaki, Aoba-ku, Sendai 980-8579, Japan b Professor emeritus of Tohoku University,2-11-18, Hitokita, Taihaku-ku, Sendai 9820222, Japan c e-mail: [email protected]

Abstract - Anodic polarization curves of Fe-18Cr-xMo alloys (x = 0-10) showed that the alloys containing 5-10 % Mo suffered from no pitting in 1 M HCl. The depassivation pH of the alloy depended on the Mo content and indicated that the reductive dissolution of the passive film was suppressed by Mo. The thinning rate of Fe2O3-Cr2O3-MoO2 films as a function of potential showed that the reductive dissolution occurred on the films with low Cr and Mo cationic fractions in the potential range of pitting of the alloys. The film with the cationic fractions XCr = 0.5 and XMo = 0.1, however, suffered from no reductive dissolution in 1 M HCl. These cationic fractions coincided well with the composition of passive film on the Fe-18Cr-10Mo alloy. The reductive dissolution of locally Cr- or Mo-depleted parts in passive films on the alloys was speculated to be concerned with pit initiation in 1 M HCl. Keywords: Fe-Cr-Mo alloys, Passive films, Fe2O3-Cr2O3-MoO2 films, Pitting, HCl solution

1. Introduction In recent years, to elucidate the function of passive films on Fe-Cr-Mo alloys in the suppression of pitting, simulation experiments using Fe2O3-Cr2O3 and Fe2O3-Cr2O3-MoO2 artificial passivation films have been performed [1, 2].

4

K. Sugimoto et al.

Followings are chief results reported: (1) Fe2O3-Cr2O3 films with any Cr2O3 content show no pitting in 1 M HCl [2]. (2) Fe2O3-Cr2O3 films with low Cr2O3 content show reductive dissolution at potentials that are equivalent to pitting potentials of Fe-Cr alloys [2]. (3) Alloying MoO2 into Fe2O3-Cr2O3 films is very effective to suppress the reductive dissolution [1]. (4) Fe-Cr-Mo alloys with passive films having composition of no reductive dissolution show no pitting in 1 M HCl [1]. In result, the reductive dissolution of Cr-depleted part in passive films on Fe-Cr alloys, or Mo-depleted part in those on Fe-Cr-Mo alloys, was speculated to be concerned with the initiation of pitting [1, 2]. However, whether the alloyed Mo suppresses the reductive dissolution of passive films on Fe-Cr-Mo alloys or not is still unclear. The objective of the present study is firstly to analyze the composition of passive films formed on Fe-Cr-Mo alloys for the model of artificial passivation films, secondly to clear the effect of Mo content on the reductive dissolution of passive films on Fe-Cr-Mo alloys, thirdly to examine effects of MoO2 content on the reductive dissolution of Fe2O3-Cr2O3-MoO2 films, which simulate passive films on Fe-Cr-Mo alloys, and finally to consider collectively the pitting inhibition mechanism of alloyed Mo in 1 M HCl. 2. Experimental Solution heat-treated high purity Fe-18Cr-xMo alloys (x = 0, 2, 5 and 10 mass %) were used as specimens. The surface of specimens was finished with 1 m diamond paste polishing. Fe2O3-Cr2O3-MoO2 films about 50 nm thick with various compositions were prepared on Pt substrates by ion-beam sputter deposition technique (Tokyo Denshiyakin, DIBS-3000HC TDY-0395). The composition of the films was determined by ICPS (Seiko Instruments and Electronics, SPS-1200A). Anodic polarization curves of the alloys were measured at a scan rate of 3.8 x 10-5 V s-1 in 1 M HCl. Depassivation pH of the alloys was obtained by measuring corrosion potential vs. time curves in 1 M NaCl + x M HCl, and x M HCl with various concentrations x after the passivation of the alloys at 0.2 V in 1 M Na2SO4 for 10.8 ks. Thinning rate vs. potential curve of the Fe2O3-Cr2O3MoO2 films was measured by ellipsometry under potentiostatic control in 1 M HCl. All solutions were de-aerated by bubbling N2 gas for more than 3.6 ks. An Ag/AgCl (3.33 M KCl) reference electrode was used as a reference electrode. In the ellipsometric measurement, three-parameter ellipsometry was employed. The wavelength and the incident angle of the light were 546.1 nm and 60.00°, respectively. Three parameters, that is, the relative phase retardation, , the arctangent of the relative amplitude reduction, , and the relative reflectivity, R/R, were measured at an interval of 5 s. From these parameters obtained, the thickness, d, and the optical constant, N2 = n2 - k2i, of the films were calculated by Drude’s exact optical equations. The thinning rate

Electrochemical Properties of Fe-Cr-Mo Alloys and Fe2O3-Cr2O3-MoO 2

5

of film thickness, - d/ t, was determined from the gradient of decrease in film thickness, d, vs. time, t, curves . The in-depth composition change of passive films formed on the alloys was analyzed by angular-resolved X-ray photoelectron spectroscopy (AR-XPS) (Thermo VG Scientific Theta Probe system). The X-ray source was monochromatic Al Kg (1486.60 eV) with a spot diameter of 200 m. XPS spectra for Fe 2p, Cr 2p, Mo 3d, O 1s, and C 1s electrons were measured. The binding energy of XPS spectrum was calibrated by the C 1s peak position regarded as 285.00 eV. Quantitative amounts of Fe, Cr and Mo in the films were calculated using integral intensities of oxide peaks of these elements and their sensitivity coefficients given in the Avantage 1.68 software attached to the Theta Prove system. 3. Results and Discussion Figure 1 shows anodic polarization curves for Fe-18Cr-xMo alloys (x = 0-10) in 1 M HCl. The alloys containing 0-2 % Mo showed pitting at -0.15-0.00 V but the alloys containing 5-10 % Mo suffer from no pitting in a range of -0.200.90 V. However, Son et al. [1] reported unstable passivity for the Fe-18Cr5Mo alloy in 1M HCl. Therefore, the Fe-18Cr-10Mo alloy was selected here for the alloy without pitting susceptibility in 1 M HCl.

Fig. 1 Current density, i, as a function of potential, E, for Fe18Cr-xMo alloys (x = 0-10) anodically polarized in 1 M HCl.

Fig. 2 Change in depassivation pH, pHdp, with Mo content of alloys in HCl solutions with various concentrations.

6

K. Sugimoto et al.

The stability of passive film, however, depends on the pH of solutions. Corrosion potential, Ecorr, as a function of time, t, was measured on Fe-18CrxMo alloys (x = 0-10) with passive films formed at 0.20 V in 1 M Na2SO4 for 10.8 ks and then immersed in HCl solutions with various concentrations for 80 ks. Figure 2 shows corrosion potential, Ecorr, as a function of pH for the alloys after steady state conditions. The alloys caused depassivation and showed active potentials in solutions with low pH values after prolonged time. The depassivation pH, pHdp, depends on the Mo content of alloys and decreases with increasing Mo content. Since the depassivation is caused by the reductive dissolution of the passive film, the reductive dissolution should be suppressed by alloying Mo-oxide component into passive films.

Fig. 3 Cationic fraction as a function of detection angle, , for passive films on Fe-18Cr-10Mo alloy formed at various potentials in 1 M HCl for 10.8 ks. (a) XFe, (b) XCr, (c) XMo.

Mo content and its distribution in the passive film were examined by the ARXPS. Figure 3 shows cationic fraction as a function of detection angle, , for passive films on Fe-18Cr-10Mo alloy formed at various potentials in 1 M HCl for 10.8 ks. Each string of experimental plots exhibits the in-depth distribution of cationic fraction. In regard to Mo, the distribution of XMo is relatively uniform from the film top to the film bottom at any potential. The average values of XCr and XMo for the film formed at 0.5 V can be regarded as approximately 0.75 and 0.15, respectively. For the comparison with the data obtained by AR-XPS, the in-depth distribution of cationic fractions analyzed by AES [1] is given in Fig. 4 for passive films on Fe-18Cr-10Mo alloys formed at potentials from -0.3 V to 0.8 V in 1 M HCl. The values of XFe, XCr and XMo were evaluated by the sensitivity coefficient method. The average values of XCr and XMo for the films formed at 0.5 V are regarded as approximately 0.5, and 0.1, respectively.

Electrochemical Properties of Fe-Cr-Mo Alloys and Fe2O3-Cr2O3-MoO 2

7

Fig. 4 Cationic fraction as a function of sputtering time for passive films on Fe-18Cr-10Mo alloys formed at various potentials in 1 M HCl for 10.8 ks. Dotted lines show alloy composition.

Thinning rate of film thickness, - d/ t, as a function of potential, E, for Fe2O3-Cr2O3MoO2 films in 1 M HCl is given in Fig. 5. The films with low XCr and no XMo cause the reductive dissolution at potentials corresponding to the passivity of Fe-18Cr-xMo alloys, but those with low XCr and high XMo hardly cause the reductive dissolution. For example, the film with XCr = 0.5 and XMo = 0.0 shows reductive dissolution under 0.2V, but that with XCr = 0.5 and XMo = 0.1 shows no reductive dissolution and keeps passivity between -0.3 V and 1.0V. Namely, MoO2 alloyed in Fe2O3-Cr2O3 suppresses the reductive dissolution of the film. It should be noted that any Fe2O3-Cr2O3-MoO2 film causes no pitting dissolution at potentials under 0.9 V at which the transpassive dissolution starts.

Summarizing above results, the passive film with approximately XCr = 0.6 and XMo = 0.1 (average values of AR-XPS and AES) on Fe-18Cr-10Mo suffered from no pitting and no depassivation in 1M HCl, and the Fe2O3-Cr2O3-MoO2 film with XCr = 0.5 and XMo = 0.1 showed no reductive dissolution in 1 M HCl. Since the reductive dissolution causes the depassivation and localized depassivation leads to pitting, a passive film without reductive dissolution should suffer from no pitting in 1 M HCl. From the discussion above, the pit initiation on Fe-18Cr-xMo alloys in HCl solutions is thought as follows: In case of low Mo alloys, a local part with low XCr and low XMo is presumed to exist in a passive film with high XCr and low XMo. The local part with low XCr and low XMo should cause the reductive dissolution at low potentials in a passive region. This leads to pitting in 1 M HCl. Then, in case of high Mo alloys, a local part with comparatively low XCr

8

K. Sugimoto et al.

may exist but the XMo at the part should be higher. Such a part should cause no reductive dissolution and no pitting in 1 M HCl.

Fig. 5 Thinning rate of film thickness, - d/ t, as a function of potential, E, for Fe2O3-Cr2O3-MoO2 films in 1 M HCl at 293 K.

4. Conclusion (1) Fe-18Cr-xMo alloys containing 5-10 % Mo suffer from no pitting in 1 M HCl. (2) The depassivation pH of Fe-18Cr-xMo alloys depends on Mo content, and indicates that the reductive dissolution of passive film is suppressed by Mo. (3) The average values of XCr and XMo for the passive film formed on Fe-18Cr10Mo alloy at 0.5 V in 1 M HCl are approximately 0.6 and 0.1, respectively. (4) The Fe2O3-Cr2O3-MoO2 film with XCr = 0.5 and XMo = 0.1 shows no reductive dissolution and no pitting in 1 M HCl. (5) The passive film without reductive dissolution is considered as the film without pitting in 1 M HCl. Reference 1. M. Son, N. Akao, N. Hara and K. Sugimoto, J. Electrochem. Soc., 148 (2001) B43. 2. K. Sugimoto, M. Son, Y. Ohya, N. Akao and N. Hara, in: Corrosion ScienceA Retrospective and Current Status in Honor of Robert P. Frankenthal, Eds. G. S. Frankel, H. S. Isaacs, J. R. Scully and J. D. Sinclair, Electrochemical Society Proceedings Volume 2002-13, The Electrochemical Society, Pennington (2002) p. 289.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

9

Formation and growth processes of electrochemical passive layers (borate medium: pH 9.2) and electron stimulated oxidized films (5.10-6 Pa O2) formed on Fe-Cr alloys a

b

M. Bouttemy , M. Bertoglio and G. Lorang

a

a

Centre d’Etudes de Chimie Métallurgique- CNRS (UPR 2801), 15 rue Georges Urbain, Vitry-sur-Seine F94407 Cedex, France b Laboratoire de Génie Electrique de Paris- CNRS (Supelec), Plateau du Moulon, 11 rue Joliot Curie, Gif-sur-Yvette F91192, France

Abstract -Internal composition of passive layers (borate buffer pH 9.2, 0.3V/SCE) and Electron Stimulated Oxidized (ESO) layers formed on a series of Fe-Cr alloys (5 - 30 at%) was determined during ageing by quantitative AES depth profiling to approach and compare their film growth processes. ESO films were “in-situ” prepared by high (AES) electron irradiation under low oxygen pressures. In both cases, the film growth process seems to result of an outward cations migration (Fe >> Cr) under electric field (CabreraMott model). Slow down of growth kinetics occurs beyond a critical chromium bulk concentration of 10-15 at% and may be attributed to an improved efficiency of inner Cr2O3 layers acting like a diffusion barrier for the Fe cations transport. Keywords : Fe-Cr, Electron Stimulated Oxidation, quantitative AES depth profiling, film growth process

1. Introduction In passivated metallic alloys, the oxidized film formation stage is expected to take place instantaneously by a surface segregation of the more oxidizable alloy component, chromium in Fe-Cr alloys, induced by oxygen [1]. It would be 2+ followed by a slow growth process developing mainly with diffusing Fe and

10

M. Bouttemy et al. 3+

Fe species which accumulate as oxides and oxyhydroxides species in the outer layers. Thicker films and larger growth rates are recorded on low chromium alloy contents [2] passivated at elevated potentials [3, 4]. According to the Point Defect Model (PDM) of Macdonald et al. [5], the film growth rate would be controlled by reactions at the metal-oxide interface. In this way, the characteristics (thickness, nature and stability, hydration…) of the related internal layers – Cr oxide – may influence the mobility of the other cationic species – Fe – diffusing according to a high field mechanism towards the external surface (Mott-Cabrera) [2]. However this mechanism is often hidden or complicated by chemical reactions occurring at the film-electrolyte interface such as dissolution, rather inoperative at pH>8, dehydration of the external layers or conversion of oxides into hydroxides or oxyhydroxides into oxides [6, 7]. In other way, metallic samples exposed to high electron or photon irradiation at low oxygen pressures are known to preferentially oxidize under the beam area [8, 9]. Such damaging artifacts were reported in adsorption studies performed by AES with focused electron beam and too high current densities [10]. This so-called electron stimulated oxidation (ESO) gives enhanced oxide growth rates promoted by an increased electronic dissociation of the adsorbed oxygen molecules which, later, should react with the substrate. More recent studies privilege a mechanism by which the electron beam should create nucleation centers for the subsequent oxide film growth [11]. In-situ ESO experiments will be expected to provide an easier access to the initial oxidation stage and to the growth mechanisms than in passive layers as misleading external reactions and sample transfer operations inherent to the passivation process are eliminated. Kinetics -6 and AES depth-2 profiling experiments at different oxidation stages (5.10 Pa O2, 84 µA.cm ) will be achieved in order to determine the depth distribution of Cr and Fe contents inside the film and its growth mode. 2. Experimental Fe-Cr alloys samples (5, 10, 12.5, 15, 20, 30 at% Cr) were submitted to a preliminary electrochemical treatment in borate solutions (Na2B4O7, 10 H2O 0.073 M, H3BO3 0.05 M, pH 9.2) in order to eliminate the natural oxide film (– 1.5V/SCE, 300 s). Then passive films were grown under potentiostatic control at 0.3V/SCE in the same electrolyte during periods varying between 30 s and 112 h [2,7]. AES depth profiles were realized with a cylindrical mirror analyzer (Cameca OPC 105) and a co-axial electron gun (2.5 keV primary energy, 30° + take-off angle) in combination with Kr sputtering (differentially pumped Riber ion gun CI-50RB, 3 keV, 57° incidence angle with the sample normal). AES spectra were acquired in the derivative mode with a modulation voltage of 4V (DE/E = 0.28%). Cr LMV transition (529 eV) intensity was measured by the

Formation and growth processes of electrochemical passive layers…

11

negative part of the peak to cast off the oxygen peak overlap. Sensibility factors were formerly determined with standards samples (pure metals, alloys and oxidized species). Software based on the Hofmann Sequential Layered Sputtering model [12] allows to provide concentration profiles and film thickness data [13]. Prior to ESO experiments, the sample surface was in-situ cleaned by ion etching. Standard conditions for oxidation kinetics are achieved rastering the 2 electron beam over an area of 700´700 µm with a nominal electron beam -6 intensity of 400 ± 5 nA measured in a Faraday Cup. Oxygen pressure (5.10 Pa) is monitored with a quadrupole residual gas analyzer (Balzers QME 125). AES depth profiling was performed using the same previous procedure with a mean sputter rate of 0.088 ± 0.005 nm/min (Ta2O5 reference material). 3. Results and discussion All Fe-Cr samples were passivated at pH 9.2 at 0.3 V/SCE. XPS examinations of passive films have revealed that films are quite practically constituted of oxides with small hydration levels and weak external contribution of Fe oxy-hydroxides [2]. From XPS (and AES) depth profiling results, a simple compositional film structure is deduced where enriched internal Cr2O3 2+ 3+ layers are covered by large rich external Fe and Fe amounts. A rather diffuse interface separates each oxide layer conversely to what is usually assumed with bi-layer models in angular resolved XPS studies [4, 6]. J, Film Thickness Atomic density, at.nm-2 (atomic layer)

200

- - - - 112 hours –––– 30 seconds

Feox

100

Crox 0 20

0 0

0,1

0,2

0,3

XCr (bulk alloy), at %

Figure. 1. Effect of ageing (30 s and 112 h) on passive film contents (at.nm-2), film thicknesses (1 atomic layer: 0.215 nm) of Fe-Cr alloys passivated in borate solutions (0.3V/SCE and pH 9.2).

Quantification of AES depth profiles on passive layers provides the film contents assimilated to Cr2O3 and Fe2O3 species [13]. In Figure 1, chromium

M. Bouttemy et al.

12

and iron oxide film contents, sum of the atomic balances in each layer, were determined after 30 s and 112 h polarization time. As the chromium alloy concentration is raised up to 15 at%, oxidized chromium quantities increase and look rather stabilized beyond. At the opposite, oxidized iron amounts follow a decreasing evolution tending after 15 at% Cr towards a similar constant level. The film thickness J also continuously diminishes in pre-passive films (30 s) as in aged films (112 h). So the primordial effect of ageing is a clear increase of Feox contents and film thickness below 10-15 at% Cr (Figure 1). This threshold of the chromium concentration range is known elsewhere as a critical concentration CCr in Fe-Cr alloys which marks the passage from a corrosive alloy concentration domain to a protective domain [14]. All those characteristics well agree with the establishment of a Crox diffusion barrier acquiring a consistent thickness and an increased efficiency above CCr with ageing which affects the iron cation diffusivity and the external film growth [2, 15]. The prevalence of the internal Crox layers to reduce the Fe cations mobility under the electric field (Mott-Cabrera) inside the film may substantiate a rate limiting reaction at the metal-oxide interface according to the PDM growth model of Macdonald et al [5]. Fe-Cr 30 at %

Fe-Cr 30 at %

selective oxidation of chromium

1,1

1 preferential migration of iron

0,9

0,8

Fe and Cr film contents, at.nm-2

0

Cr

R : (ICr / IFe) / (I

200

4

5.10

-6

Pa O2 , 84 µA.cm

-2

150

3

OXYGEN

100

2

Fe (ox)

50

1

Cr (ox) 0

0,7 0

400

800

1200

-6

Oxidation time (5.10 Pa O ), min

Film thickness, nm

/ IFe0)

1,2

0 0,1

1

10

100

Oxidation time, hours

2

Figure 2. Standard ESO kinetic on a FeCr 30 at% alloy showing the R=ICr/ICr0]/[IFe/IFe0] evolution during the electron irradiation (84 µA/cm2).

Figure 3. ESO film composition versus the oxidizing time exposure.

In the literature, the Electron Stimulated Oxidation (ESO) rate is a combined function of the oxidizing pressure (CO2, H2O, O2 …) and of the electron beam current density and energy [16]. In Fe-Cr alloys, the influence of those parameters on oxidation kinetics was previously displayed [17].

13

Formation and growth processes of electrochemical passive layers…

According to Steffen et al. [18] the initial oxidation process should begin by a short oxygen adsorption stage followed by a preferential Cr oxide growth. Differentiation could be made by AES between the Fe-MVV transitions of Cr and Fe of the chemisorbed and oxidized states by convenient least squares fitting with spectral standards. Presently, the Fe-MVV transition splitting after ten minutes exposure reveals the building of a three dimensional iron oxide (44.0 and 52.0 ± 0.5 eV). The LMV transitions intensity ratio R defined in Figure 2 provides a rough estimate of the concentration Cr/Fe ratio in the film during the oxidation kinetics. In this way, the beginning of the kinetic (fig. 2) displays a chromium enrichment (R≥1) of the surface which probably results of a Cr surface segregation induced by the oxygen adsorption [18]. Conversely, after 150 min, the surface composition becomes to be enriched with iron. This seems to verify a two stages film formation process consisting, at first, in a preferential oxidation of chromium, and, secondly, in an “outward” Fe cations migration assisted by an electric field [18]. The present electric field would be applied between the external interface, negatively charged by dissociated oxygen molecules under the electron beam, and the positively charged internal interface by cationic vacancies or interstitials. Depth profiles of ESO films were achieved after different irradiation and oxidizing sequences. Like in passive layers, chromium oxide is typically accumulated in the inner layers of the film and iron oxide in the outer part. In a Fe-Cr 30 at% alloy, both iron and chromium oxides contents in the ESO film evolve with the oxygen exposure according to logarithmic time laws (Figure 3). The slope of the linear plots of Feox or Crox amounts as a function of the logarithm irradiation exposure time provides the corresponding growth rates constants kCrox and kFeox. A larger ionic mobility of iron cations – two times larger than the chromium one – is therefore measured in accordance with former literature data [18, 19]. Prevailing chromium contents in the beginning of the oxidation process support an initial preferential oxidation stage of chromium (Figure 3). Table 1. Fe and Cr oxidation rate constants in ESO films formed in the standard conditions.

k Feox (a.u.) k Crox (a.u.) kFeox/kCrox

Fe-10Cr 46.0 ± 3.0 4.9 ± 0.4 9.4

Fe-20Cr 43.0 ± 3.0 11.9 ± 0.8 3.6

Fe-30Cr 33.0 ± 1.5 15.3 ± 0.8 2.2

Similar studies were performed on Fe-Cr 10 and 20 at% alloys (Table 1). With kFeox >> kCrox, a larger ionic mobility of iron cations is established in all those Fe-Cr alloys studied. Moreover, the continuous decrease of the kFeox / kCrox ratio at rising Cr bulk alloy concentration, especially between 10 and 20 at% Cr, confirms the role of diffusion barrier effect attributed to the chromium oxide layers, as in passive films, for the Fe cations transport rate inside the film.

14

M. Bouttemy et al.

4. Conclusions In electron stimulated oxidation (ESO), the initial preferential oxidation of chromium was experimentally disclosed during oxidation kinetics. In the passivation process, this too fast stage process can only be expected from the aposteriori concentration gradients depicted in the underlying metallic alloy. Logarithmic laws (Cabrera-Mott model) can describe the film growth kinetics in ESO films. Present results confirm the higher mobility of iron cations in ESO films during the growth stage. The Cr oxide film contents prevail on iron mobility at increasing chromium alloy bulk concentration, especially beyond 15 Cr at%. This supports a diffusion barrier effect played by internal Cr2O3 layers which acquire an improved efficiency at elevated bulk chromium alloy concentration in both passive and ESO films. So formation and growth of oxide films on Fe-Cr alloys (5 to 30 at%) either passivated in borate medium (pH 9.2) or oxidized under intense electron beams are found quite related processes where ESO growth mechanisms appears a simplified version of the aqueous electrochemical passivation process by turning down hydration and dissolution reactions. Bibliography [1] D. Landolt, Surf. Interf. Anal. 15 (1990) 395. [2] M. Bouttemy, S. Réveillon, M. Bertoglio and G. Lorang, EUROCORR 2004, Nice 12-16 sept. 2004 (France), event No. 266. [3] J. Häfele, B. Heine and R. Kirchheim, Z. Metallkd 83 (1992) 395. [4] H. -W. Hoppe, S. Haupt and H. -H. Strehblow, Surf. Interface Anal. 21 (1994) 514. [5] D. -D Macdonald, J. Electrochem. Soc. 139 (1992) 3434. [6] V. Maurice, W. P. Yang, J. Electrochem. Soc. 143 (1996) 1182. [7] F. Basile, J. Bergner, C. Bombart, B. Rondot, P. Le Guevel and G. Lorang, Surf. Interface Anal. 30 (2000) 154. [8] J. M. Fontaine, O. Lee-Deacon, J. P. Duraud, S. Ichimura and C. Le Gressus, Surf. Sci. 122 (1982) 40. [9] J. P. Coad, H. E. Bishop and J. C. Rivière, Surf. Sci. 21 (1970) 253. [10] C. G. Pantano and T. E. Madey, Appl. Surf. Sci. 7 (1981) 115. [11] Wei Li, M. J. Stirniman, S. J. Sibener, Surf. Sci. 329 (1995) L593. [12] J. M. Sanz and S. Hofmann, Surf. Interface Anal. 8 (1986) 147. [13] G. Lorang, M. Da Cunha Belo and J-P. Langeron, J. Vac. Sci. Technol. A5 (1987) 1213. [14] K. Asami, K. Hashimoto, S. Shimodaira, Corr. Sci. 18 (1978) 151. [15] S. Boudin, C. Bombard, G. Lorang, M. Da Cunha Belo, « modification of passive films » Edit. P. Marcus, B. Barroux, M. Keddam, Published by the European Federation of Corrosion and the Institute of Materials No. 12 (1994) 35. [16] G. Y. Mc Daniel, S. T. Fenstermaker, D. E. Walker Jr, W. V. Lampert, S. M. Mukhopadhyay and P. H. Holloway, Surf. Sci. 445 (2000) 159. [17] M. Bouttemy and G. Lorang, to be published. [18] H. J. Steffen and S. Hofmann, Surf. Interface Anal. 19 (1992) 157. [19] G. C. Allen, S. J. Harris, J. A. Jutson and J. M. Dyke, Appl. Surf. Sci. 37 (1989) 111.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

15

Development and Composition of the High Temperature Oxide Film Grown on Fe-15Cr during Annealing Eungyeul Park, Michael Spiegel Department of Interface Engineering, Max-Planck-Institut für Eisenforschung, MaxPlanck-Straße 1, D-40237 Düsseldorf, Germany

Abstract - In this work, the morphology and elemental composition of the initial surface oxide layer of Fe-15Cr polycrystalline alloy after short-term heat treatments were investigated as a function of temperature and gas phase composition. Samples were exposed to air or nitrogen-hydrogen gas mixtures at 400, 600 and 800oC in an infrared furnace for 0 to 300 seconds, then investigated by EBSD, FE-SEM for morphological studies and XPS for the elemental depth profiling by sputter techniques. Keywords : Fe-Cr alloy, Oxidation, Annealing, Hydrogen.

1. Introduction Iron-chromium steels have been of great importance for process industries and powerplants since the development of high-temperature materials. It is generally known that the passive oxide film on the alloys is predominantly composed of trivalent Cr oxide [1-5]. Because the surface oxide composition in the beginning of oxidation is dependent on the metal composition of the surface [1,2,6], it is advantageous to modify the surface of alloys, to contain more Cr than in the bulk, in order to form rapidly a dense and thick Cr-rich oxide film in the following high-temperature applications [7,8]. In the initial stage oxidation of the surface of Fe-Cr alloys, the formation of a duplex oxide layer consisting of an outer Fe-rich oxide and an inner Cr-rich oxide, has been reported in numerous of experiments in the temperature range from room temperature to 800oC [3,7,9]. It has been shown

16

E. Park and M. Spiegel

that the enrichment of Cr on the surface at T > 400oC occurs due to the cosegregation of Cr and N or C and the selective oxidation of Cr [9-15]. However, the phenomena concerning the enrichment of Cr and the effects of hydrogen during conventional heat treatments have rarely been observed or discussed. This study investigated the surface chemistry of Fe-15at.%Cr polycrystalline alloy after short-term heat treatments at 400, 600 and 800oC, for the knowledge of the initial stage of oxidation and the effects of annealing for better understanding of the high-temperature surface phenomena and the development of the surface modifications. 2. Experimental Fe-15at.%Cr polycrystalline alloy was manufactured by the argon melting method and cut to samples of 7×7×2(mm3) in size. The surface of the samples was mechanically ground with 1000 grit of SiC paper and polished with diamond paste up to 1 m. Afterwards the samples were cleaned and degreased with ethanol in an ultrasonic bath, and then kept in a dry desiccator. The average grain size of the polished samples, measured by electron backscattering diffraction (EBSD) attached to FE-SEM, was 350±50 m. Experiments were carried out at 400, 600 and 800°C using a horizontal infrared (IR)-heating furnace (Quad Ellipse Chamber, Model 5528-10, Radiant Energy Research) with varied atmospheres and heating cycles. The reactor tube was made of quartz. The composition of inlet gas was controlled by a Mass Flow Controller (Model D-5111, Manger+Wittmann). The moisture in the inlet gas, controlled using an oxal-acid-filled moisturization, was monitored by a Moisture Analyzer (Model 8800A, Teledyne Analytical Instruments). The oxygen partial pressure of the outlet gas was measured using a fuel-cell-sensor oxygen analyzer (Model 3190, Teledyne Analytical Instruments). In oxidation experiments, atmospheric air was used. In case of annealing experiments, ultrahigh purity hydrogen and nitrogen were used. For all annealing experiments, the dew point of the inlet gas mixture was -70.0oC. The surface morphology of the samples was observed using a field emission secondary electron microscope (FE-SEM, LEO 1550VP, GEMINI). The composition of the surface and the depth profile analysis were carried out using XPS (x-ray photoelectron spectroscopy, ESCA, Quantum 2000, Physical Electronics Inc.). For each analysis, a survey spectrum was acquired before elemental depth profiling. Then Cr2p, Fe2p, O1s, C1s and N1s spectra were obtained with Ar+ sputtering. The sputtering rate was 6.30±0.02 nm/min, which was calibrated with a SiO2 monolayer on a pure Si substrate. The quantified elemental depth profiles were made by the area measurement of C1s, N1s, O1s, Fe2p 3/2 and Cr2p 3/2 using CASAXPS software (V.2.2.29). Thermodynamic calculations of the experimental systems were made using FactSageTM database.

Development and Composition of High Temperature Oxide

17

3. Results and Discussion Oxidation in air The morphology of the surfaces of samples oxidized for different times at 400oC in air is presented in Fig. 1. The images were taken from the center of grains. There was no distinguishable change in the morphology until 60 seconds exposure (Fig. 1 (a)). Then the whisker-type oxide nucleation along the point or line defects was observed on the surface of the sample oxidized for 90 seconds (Fig. 1 (b)). After 120 seconds exposure (Fig. 1 (c)), spherical-type nucleation started hindering the former whisker-type oxide. The spherical oxide grew further and became dominant with increasing exposure time, resulting in the complete coverage of the spherical-shape oxide on the surface (Fig. 1 (d)). Fig. 2 (a) and (b) present the elemental compositions of the outer 50 nm of the surface layers of the samples oxidized for 60 and 300 seconds, as obtained by XPS with Ar+ sputtering. It shows that only Fe ions took part in the oxidation resulting in the formation of an Fe2O3 layer on the surface. The oxidation proceeded with the outward diffusion of Fe ions with an increase of the thickness of the Fe2O3 layer. The interface of each sample was composed of a thin Fe-Cr oxide layer. (a) 60 sec.

(b) 90 sec.

400nm

400nm

(c) 120 sec.

(d) 300 sec.

400nm

400nm

Fig. 1. Morphology change during oxidation at 400oC

18

E. Park and M. Spiegel 100

100

O

60 40 20

Fe

(b) 300 sec.

80 atomic %

atomic %

Fe

(a) 60 sec.

80

O

60 40 20

Cr

Cr

0 0

0

10

20

30

40

50

0

10

Depth, nm

20

30

40

50

Depth, nm

Fig. 2. Elemental depth profiles of the samples oxidized at 400oC for different exposure times The morphology of the surfaces of samples oxidized for different times at 800°C are presented in Fig. 3 (a) to (d). Oxidation started with the formation of spherical-type nuclei with diameter less than 100nm (Fig. 3 (a)). The nucleation was saturated on the surface within 30 second of oxidation. With increase in time, the nuclei grew and changed to the polygonal grains (Fig. 3 (b)). Then the nucleation of a secondary oxide was observed. The secondary oxide nucleated between/on the former oxide grains and grew on their surface. It seemed that the newly formed nuclei grew or dissolved into the matrix oxide very quickly. Further oxidation resulted in the homogeneous layer composed of oxide grains with the size of 200nm (Fig. 3 (c)). The further change in the morphology with increasing oxidation time was negligible (Fig. 3 (d)). (a) 0 sec

400n (c) 120 sec

400n 400n

(b) 60 sec

400n (d) 300 sec.

400n

Fig. 3. Morphology change during oxidation at 800oC (e) 120 (f) 300

19

Development and Composition of High Temperature Oxide

Fig. 4 (a) to (d) present the elemental compositions of the outer 100nm of the surface layers of the samples, as obtained by XPS. The oxidation started with the formation of a Fe-rich mixed oxide layer (Fig. 4 (a)), representing the initial surface composition. The well-known duplex oxide layer, outer Fe-rich oxide and inner Cr-rich oxide, formed in 60 seconds (Fig. 4 (b)). The duplex oxide layer transformed to a Cr-rich mono-phase layer with increasing time (Fig. 4 (c)). The transformation combined the enrichment of Cr in the former Fe-rich oxide layer and the nucleation/growth of the Cr-rich secondary oxide layer, as observed in the morphology (Fig. 3 (d)). In the period of t = 120 to 300 sec, the Cr content in the Cr-rich mono-phase layer increased with time with negligible changes in the thickness and the morphology. It implies that in this period, the Cr diffused to the surface was consumed to reach an equilibrium phase composition in the oxide lattice, without further nucleation and growth. Except the sample exposed for 0 second at 800°C, a Cr-depletion zone was not observed, which shows the fast diffusion of Cr in the alloy. 100

100

(a) 0 sec.

60 40

O

20 0 20

40

60 40

O

60

80

Cr

0

100

0

20

Depth, nm

40

60

80

100

80

100

Depth, nm

100

100

(c) 120 sec. O

60 40

Cr

20

(d) 300 sec.

80

atomic %

80

atomic %

Fe

20

Cr 0

(b) 60 sec.

80

Fe

atomic %

atomic %

80

O

60 40

Cr

20

Fe

0

Fe

0

0

20

40

60

Depth, nm

80

100

0

20

40

60

Depth, nm

Fig. 4. Elemental depth profiles of the samples oxidized at 800oC, for different exposure times Annealing in nitrogen-hydrogen gas mixtures Fig. 5a. presents the effects of temperature on the depth profile of Cr/Fe ratio of the sample after the exposure for 60 seconds under N2-5vol%H2. The effect of annealing at 400oC was not obvious, however, further increase in the temperature increased the Cr/Fe ratio in the surface. The annealing at 800oC significantly increased the Cr/Fe ratio, as well as the thickness with higher Cr content.

20

E. Park and M. Spiegel

The Cr/Fe ratio depth profile of the samples annealed at 800oC under different hydrogen content is presented in Fig. 5b. It shows that the increase in hydrogen in the atmosphere increases the Cr content in the surface. Authors’ former works [9,15] have shown that the increase of hydrogen facilitates the diffusion of Cr through fast diffusion path in the near-surface. This phenomenon would be due to i) decarburization under dissolution of Crcarbides, especially of grain boundaries and ii) reduction of Fe oxides in the presence of hydrogen. In addition, desulfurization on the surface (iii) is also expected by the reaction H2(g) + Ssurf = H2S (g), but could not be observed since the detection of sulfur was beyond the limitation of XPS analyses in this ex-situ experiment. 10 10

400C 600C 800C Polished

Polished 100%N2 N2-1%H2 N2-5%H2 N2-10%H2

(b) Crtot /Fe tot

Crtot /Fe tot

(a) 1

0.1

1

0.1 0

10

20

30

40

50

0

10

Depth, nm

20

30

40

50

Depth, nm

Fig. 5. Effects of T (a) and H2 content (b) on the Cr/Fe depth profile

4. Conclusions The effects of short-term heat treatments at 400, 600 and 800oC on the surface composition of Fe-15at.%Cr polycrystalline alloy were investigated. The atmosphere during heat treatments varied with the gas composition, i.e. air and N2-H2. The surface of each sample was characterized using SEM and XPS technique. The surface chemistry after the heat treatments depended on the temperature and the atmosphere. Oxidation at 400oC resulted in the formation of a thin Fe2O3 oxide layer, while the formation of Cr-rich oxide layer was observed at 800oC. For the development of the surface modification, further investigation on the effects of annealing time, the diffusion of hydrogen using D2 and the effects of other minor alloying elements (Ti, Zr, etc.) is required in the future. References 1. P. Kofstad, “High Temperature Corrosion”, Elsevier, NY (1988). 2. B. Chattopadhyay and G.C. Wood, Oxid. Met., 2 (1970) 373.

Development and Composition of High Temperature Oxide 3. H.J. Mathieu and D. Landolt, Corr. Sci., 26 (1986) 547. 4. C.-O.A. Olsson and D. Landolt, Electrochemica Acta, 48 (2003) 1093. 5. E. McCafferty, Corr. Sci., 42 (2000) 1993. 6. E. McCafferty, Corr. Sci., 44 (2002) 1409. 7. C.P. Jensen, D.F. Mitchell and M.J. Graham, Corr. Sci., 22 (1982) 1125. 8. I. Saeki, H. Konno and R. Furuichi, Corr. Sci., 38 (1996) 19. 9. E. Park, B. Huning and M. Spiegel, Appl. Surf. Sci. (2005), in publication. 10. C. Leygraf, G. Hultquist and S. Ekelund, Surf. Sci., 46 (1974) 157. 11. H.J. Grabke, R. Dennert and B. Wagemann, Oxid. Met. 47 (1997) 495. 12. E. Clauberg, C. Uebing and H.J. Grabke, Surf. Sci., 433-435 (1999) 617. 13. H.J. Grabke, V. Leroy and H. Viefhaus, ISIJ Int., 35 (1995) 95. 14. H.J. Grabke, W. Paulitschke, G. Tauber and H. Viefhaus, Surf. Sci., 63 (1977) 377. 15. E. Park, B. Huning and H.J. Grabke and M. Spiegel, Defect and Diffusion Forum, 237-240 (2005) 928.

21

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

23

Passivity of Nickel-Containing Stainless Steels in Concentrated Sulphuric Acid M.B. Ives a, Y. Li a, K.S. Coley a a

Walter W. Smeltzer Corrosion Laboratory, McMaster University, Hamilton, Ontario, L8S 4L7, Canada. ( e-mail: [email protected])

Abstract - The spontaneous periodical oscillation of corrosion potential between the active and passive states, typical of nickel-containing stainless steels in concentrated (93.5%) sulphuric acid is used to determine the nature of passivity in such systems. Detailed analysis of the electrochemical kinetics of the oscillating potential in combination with (ex-situ) surface analysis has indicated the presence of nickel sulphides with a limited potential range of stability. A minimum critical coverage of sulphide on the surface is required for the mixed potential to be shifted into the passive range. The sulphide catalyses the cathodic reaction thereby raising the potential. However, when the coverage exceeds a critical value, the sulphides effectively block the formation of a passive film. Consequently, there is a critical range of nickel content required in these alloys for their effective use in sulphuric acid production plants. Keywords: sulphuric acid, nickel sulphide staibility, potential oscillation, exchange current density.

1. The Potential Oscillation Phenomenon It has been well-documented [1,2,3] that at appropriate temperatures and acid concentrations, many stainless steels when immersed in concentrated (>90%) sulphuric acid exhibit a regular periodic oscillation of their “open circuit” (corrosion) potential. Kish et al. [4] investigated this phenomenon in the stainless steel S30403, identifying the importance of nickel in the alloy, and Li et al. [5] investigated the temperature and flow rate dependence of the major features of the potential oscillations.

24

M.B. Ives et al.

Figure 1 shows a corrosion potential time plot for a typical oscillating situation, and Figure 2 a high resolution plot of the “spike” portion of the oscillation. While the sample is at the spike potential, it is actively corroding, therefore, the total corrosion rate is directly proportional to the frequency of the oscillation, as was shown by Chang [2].

Figure 1 Typical potential oscillation for S30403 in 93.5 wt% sulphuric acid at 60EC.

Figure 2 High resolution view of active “spike”

The experiments described here were made on a rotating cylinder electrode in a typical three-electrode cell, employing a mercurous sulphate electrode (MSE). Recording the oscillation characteristics as a function of temperature and electrode rotation rate [5] indicated that while the oscillation period is transportcontrolled, the width of the active spike is purely activation-controlled. Furthermore, during the spike, a dark cloud may be visually observed above the sample, which disappears when the potential shifts back into the (“peak”) passive region. This has lead to the view that during the active spike the acid in

Passivity of Nickel-Containing Stainless Steels in Concentrated Sulphuric Acid

25

the vicinity of the surface is supersaturated with corrosion product, which subsequently precipitates out. It was proposed by Kish [4] that this salt is nickel sulphide. The oscillating sequence is modeled by assuming that the exchange current density for the acid reduction reaction is greater on a surface containing sulphide than on a bare steel surface. Figure 3 shows a mixed-potential diagram which demonstrates this effect. On first immersion into the acid, the cathodic reaction would provide a mixed potential in which the steel is active. On precipitation of the nickel salt, the cathodic reaction is catalyzed, increasing the exchange current density and a cathodic E-log i plot which produces a mixed potential in the passive range. However, at the higher passive potential, it is assumed that the salt phase is no longer thermodynamically stable. Consequently, it starts to dissolve and reduces the effective surface area over which the cathodic reaction is catalysed.

Figure 3. Schematic mixed-potential diagram indicating the cathodic reactions for bare steel and for the steel with a nickel salt on its surface.

The mixed potential thereby is lowered, corresponding to the continuous decrease in potential observed after the “peak” in the oscillations. Eventually, the steel surface returns to its original state, and the mixed potential drops into the active “spike”. The process may then be repeated indefinitely. To better understand these stages involving the repeated formation and removal of the sulphide film, additional electrochemical polarization experiments were

26

M.B. Ives et al.

performed [6]. When the sample is pre-polarized under potentiostatic control at a series of potentials between the “peak” (at 0.0VMSE) and “spike” (at -0.6VMSE ) potentials different subsequent behaviour was observed. In brief, when the pre-polarization was made for 1h at higher than -0.3VMSE the corrosion potential continued to oscillate when the applied potential was removed. But when held at potentials lower than -0.3VMSE the samples remained active, eventually approaching a steady potential of -0.4VMSE. It is concluded that -0.3VMSE represents the limit of potential stability for the salt phase at this acid concentration and temperature. When held at the higher potentials, the salt phase is de-stabilized, permitting the return of the active “spike”. However, stabilizing the salt phase by pre-polarization at the lower potentials can lead to complete coverage of the surface, thereby blocking the formation of a passive film despite the catalyzing influence of the salt. These latter samples were heavily corroded due to their continuous active condition. This situation would correspond to the mixed potential diagram shown in Figure 4.

Figure 4. Schematic mixed-potential diagram for a situation wherein the formation of a passive film is blocked.

A more detailed investigation of the pre-polarization treatments [6] has indicated that there is a critical time for polarization at -0.6VMSE to produce sufficient salt phase to block passivation and cause a high rate of corrosion. Specifically, for a rotating cylinder electrode at 1000 rpm in 93.5 wt% acid at 60EC the critical time is between 56s and 112s. This must be the time required to grow a salt film sufficient to block the formation of the passive film. 2. Identification of Surface Phase(s) Limited success has been achieved in identifying the surface phase(s) responsible for the repeated active-passive transitions in this system. Ideally,

Passivity of Nickel-Containing Stainless Steels in Concentrated Sulphuric Acid

27

such analysis should involve in-situ measurement, due to the hygroscopic nature of sulphuric acid and the problems of exposing samples removed from the acid into humid air. While in-situ Raman spectroscopy is being investigated, analysis to date has only been possible using ex-situ X-ray photo-electron spectroscopy (XPS) on samples removed from acid, rinsed, immersed in hexane and transferred into the spectrometer through a controlled-atmosphere glove box. From the observations of the XPS analyses it is clear that the number and abundance of sulphide phases present varies with the progress of the potential oscillation. Figure 5 shows the sulphur 2p spectra obtained at three points on an oscillation. The peaks correspond to a series of sulphides and polysulphides in

Figure 5. XPS S-2p spectra from (a) active, (b) passive, and (c) transition stages of an oscillation; (d) schematic of the sampling points.

the region of 163 eV and sulphates around 169 eV. The latter do not change significantly through the cycle, but the amount of sulphides decreases successively in going from the active “spike”, to the passive “peak”, and to the “transition” just before the potential drops back to the active spike (Fig 5d). This is consistent with the proposed model for the manner in which the sulphide is deposited on and dissolved from the surface determines the potential oscillation. While the details of the deposit are not germane to the argument, it

28

M.B. Ives et al.

is clear from Figure 5 that there are a number of sulphides and polysulphides involved in the process. The XPS measurements also provide information on the chemical state of the major cations -- iron, chromium and nickel [7] -- and confirms that the important precipitating phases are nickel sulphides. 3. Conclusions and Consequences for Alloy Development It is now clear that the nickel content of the stainless steel is primarily responsible for the oscillation phenomenon which keeps the alloy in the passive range for extended times…. but not continuously. However, if the nickel content is too great, there is a danger of blocking the passive film and causing high corrosion rates (see Figure 4.) From an alloy development viewpoint, there is a critical range of alloy nickel content to minimize corrosion of these alloys in concentrated sulphuric acid. It would be instructive to determine the influence of the nickel/chromium ratio in the steel, with a view to obtaining the ratio which minimizes the fraction of total time spent in the active spike. The acid production industry would benefit greatly by such alloy development, thereby removing the necessity to protect these alloys in service through the application of anodic protection [8]. References 1. R.M Kain, P.E. Morris, Paper No. 149, Corrosion/76, NACE, Houston, TX, 1976 2. Y.S. Chang, “Periodic Active-Passive Corrosion Behaviour”, PhD Dissertation, Cambridge University, 1984 3. Michael H.W. Renner, “Corrosion Behaviour of Stainless Steel and Nickel Alloys in Hot Concentrated Static and Flowing Sulphuric Acid”, Ph.D. Dissertation, University of Teeside, 1991 4. J.R. Kish, M.B. Ives and J. Rodda, Corrosion Science, 45 (2003) 1571 5. Y. Li, M.B. Ives, K.S. Coley, J.R. Rodda, Corrosion Science, 46 (2004) 1969 6. Y. Li, M.B. Ives, K.S. Coley, “Corrosion Potential Oscillations of Stainless Steel in Concentrated Sulphuric Acid: 1 Electrochemical Aspects.”, submitted to Corrosion Science, March 2005. 7. Y. Li, M.B. Ives, K.S. Coley, “Corrosion Potential Oscillations of Stainless Steel in Concentrated Sulphuric Acid: 2 Surface Analysis.”, to be submitted to Corrosion Science 8. D. Fyfe, D. Sanz, F.W.S. Jones, G.M. Cameron, “Anodic Protection of Sulphuric Acid Plant Cooling Equipment” CORROSION/75. Paper no. 63 (NACE, Houston, TX) 1975

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

29

An insight on the role of Nickel in the passive films generated on different stainless steels C.M. Abreu, M. J. Cristóbal, R. Losada, X. R. Nóvoa, G. Pena, M. C. Pérez* E. T. S. E. I. e M., Campus Universitario, Universidade de Vigo, 36310 Vigo (Spain) *Corresponding author: Tel.: +34986812213, email: [email protected]

Abstract - The conductivity of air-formed films on Fe-Cr-Ni alloys is studied in alkaline solution using Impedance Spectroscopy. Increasing Ni content in the alloy decreases conductivity and produces thinner films. Air-formed films are not homogeneous from the electrical point of view and two well differentiated conducting paths appear in parallel to dielectric islands.

Keywords: Fe-Cr-Ni alloys, Stainless steel, High frequency IS, air-formed films, alkaline medium

1. Introduction The increasing use of highly alloyed steels and nickel-base alloys for applications in very aggressive environments has encouraged the studies on the effect of alloying elements in passive layers’ stability. But, despite the efficient analytical techniques applied in those studies, the role of Ni, Mo, and other elements continues to be a matter of debate [1, 2]. One important issue in corrosion of steel rebars in concrete is the on-set of possible galvanic effects when stainless steels are coupled to carbon steel. In this sense, the oxygen reduction current (an aspect of major importance in galvanic corrosion) was demonstrated to be highly dependent of the chemical composition and microstructure of the material [3] and, thus, on passive film conducting properties. Previous studies [3, 4] suggest that the electronic conducting properties of passive films formed on stainless steels in alkaline media can be detected in the high frequency region of the impedance spectra. The present work is focussed in analysing the Ni influence on the growth and conducting properties of the oxide layers on the different alloys.

30

C. M. Abreu et al.

2. Experimental Part Table 1 summarises the chemical compositions of the four materials studied. Specimens of 1.1x1.1 cm were cut, wet ground, then polished to 6µm diamond paste and ultrasonically cleaned just before the immersion in the test solution. The exposed area was limited by an O-ring defining 0.28 cm2 active surface. Cyclic voltammetry (CV) was performed in a conventional three electrode cell arrangement where the working electrode was the selected steel; a Pt mesh was used as counter electrode and an Hg/HgO 0.1M KOH as reference electrode. The experiments were carried out at room temperature in NaOH 0.1M. An AUTOLAB 30 Potentiostat (from EcoChemie, NL) was used for electrochemical measurements. The electrode potential was scanned from hydrogen to oxygen evolution reactions at dE/dt=1 mVs-1. Table"1. Chemical composition of all studied materials (% weight)." C

Cr

Si

Ni

Mn

Mo

P

Cu

S

Fe

AISI 430

0.046

16.14

0.36

0.16

0.43

0.03

0.018

0.09

0.01

Bal.

AISI 304L

0.023

18.22

0.34

8.58

1.79

0.43

0.026

0.29

0.001

Bal.

AISI 316

0.031

16.85

0.45

11.75

1.66

2.10

0.020

0.31

0.001

Bal

Ni Alloy

0.07

28.79

0.6

42.36

1.53

3.56

0.04

0.03

Bal.

Before the CV, the high frequency impedance range was examined in detail using impedance spectroscopy (IS). Measurements were performed in a twoidentical-electrode cell, without DC polarisation, in the frequency range 40MHz to 100Hz using a HP 4194A Impedance analyzer [3]. 3. Results and discussion The passive layers developed on air for the different stainless steels (AISI 430, AISI 304L, AISI 316 and Ni base alloy) have been analysed by our group in previous works [3-7]. In general, the oxide layers of these stainless steels and Ni-base alloy can be considered as a Fe-Cr oxide, with decreasing Fe/Cr ratio towards the oxide/metal interface. Ni2+ is not detected throughout the films of the studied materials. Differences between films concern the Cr/Fe ratio, which increases with the nickel content in the alloy. This effect more important in the passive layer developed on the nickel-base alloy, which is mainly chromium oxide (Cr2O3) containing only small amount of Fe-oxide (Fe2+ and Fe3+). The topotactic character of surface coverage by an insulating phase (Cr2O3) and a conducting one (Fe3O4) was already postulated for AISI 316 steel and Ni base alloy [7, 8]. A closer look to the problem is presented here with an extensive use of IS. Figure 1 shows the initial impedance behaviour corresponding to AISI 430. Apparently, the Nyquist diagram corresponds to a pure capacitive

31

An insight on the role of Nickel in the passive films

behaviour. When the same measurement is performed on the other materials, a capacitive arc appears in the scanned frequency domain. The associated resistance decreases with immersion time and finally the Nyquist impedance plot becomes capacitive. Figures 2, 3 and 4 summarise this evolution with time for AISI 304L, AISI 316 and the Ni-base alloy. The elapsed time to reach the capacitive behaviour increases with Ni content in the alloy and is 0.5 h, 96 h and 273 h, respectively. 250

300

0.5

2

-Imaginary part/ kW.cm

- Imaginary part/ kW.cm

2

0.4

250 200 150 100 1 kHz

50

200

0.3 0.2

150 t=0h 100

50

10 kHz

0

0

50

1000 0h 24 h 164 h

100 Hz

2

0.1

0.1

0.2

1 MHz 10 kHz

250

10 kHz

0.3

100 Hz 100 Hz 100 Hz

0

250

500

750

Real Part / k W. cm

1000

2

Figure 3: Evolution of the HF impedance spectra with immersion time for AISI 316 in NaOH 0.1 M.

- Imaginary part/ kW.cm

- Imaginary part/ kW.cm

2

96 h

750

0.0

200

250

Figure 2: Evolution of the HF impedance spectra with immersion time for AISI 304 in NaOH 0.1M.

1000

10 MHz

150 2

Figure 1: High frequency impedance spectra for AISI 430 stainless steel in NaOH 0.1 M.

500

100

Real Part/ k W. cm

2

0h 3h 24 h 90 h

100 Hz

10 MHz

100 150 200 250 300

0.2

1 MHz 0.0 0.0 0.1 0.2 0.3 0.4 0.5

50

Real Part/ W. cm

0

0.1

100 kHz

10 MHz

0 0

t=0.5 h

100 Hz

100 Hz

750

0.4 372 h

100 Hz

0.3 0.2 0.1 1 kHz

500

0.0 0.0

40 MHz

0.1

0.2

0.3

0.4

1 kHz

250

0

100 Hz

10 MHz

0

250

500

750

1000

2

Real Part / k W. cm

Figure 4: Evolution of the HF impedance spectra with immersion time for Ni base alloy in NaOH 0.1 M.

32

C. M. Abreu et al.

From the apex frequency of the capacitive arc shown in Figure 2 a capacitance value close to 80 pF.cm-2 can be obtained, which can be associated to the dielectric behaviour of chromium oxide islands in the passive film. The “loss resistance”, about 200 kΩ.cm2 in Figure 2 could be associated to an iron-rich oxide phase [4]. Nevertheless, this apparently single capacitive arc flattens (Figures 3 and 4) when Ni content in the alloy increases, which reveals the presence of more than one time constant in this frequency region of the impedance spectra. XPS depth profiles (not shown here) demonstrate that air formed passive films become thinner as Ni content in the alloy increases (films thickness ranges between about 10 nm for AISI 430 and 1.5 nm for the Ni-base alloy). Those data correlate well with impedance spectra presented in Figures 2 to 4: Film growth is limited by film conductivity which, in turn, determines film resistance to electrolyte attack. The dissimilar passive layer structure between the ferritic steel and those containing nickel is evident also in voltammetry behaviour. Figure 5 depicts the first voltammetric cycle for the studied alloys. The main difference concerns Cr3+ oxidation domain (between about +0.1 and +0.6V). The high intensity peak at +0.45V corresponds to the AISI 430 alloy and can be considered as proportional to the amount of Cr3+ present in the conducting phase (as substituted magnetite). The peak at +0.25 V will correspond to Cr3+ oxidation at boundaries of Cr2O3 islands. The surface of those boundaries seems to be inversely related to Ni content in the alloy (and thus to film thickness).

Current Density / mA cm

-2

150 100

AISI 430 AISI 304L AISI 316 Ni-Base Alloy

50 0 -50 -100 -150 -1.4 -1.2 -1.0 -0.8 -0.6 -0.4 -0.2 0.0

0.2

0.4

0.6

Potential / V (vs. Hg/HgO, 0.1M KOH) Figure 5: First cyclic voltammograms obtained for the studied materials in 0.1 M NaOH solution.

With the above discussed data we can state that the passive film contains insulating Cr2O3 islands and varying ratios of conducting phases. The Cole-Cole analysis of impedance spectra once achieved the steady state in solution can help in clarifying this idea. The data to be analysed correspond to

33

An insight on the role of Nickel in the passive films

those obtained at 0.5h, 96h and 372h (inserts of Figures 2, 3 and 4). Those data are transformed into equivalent capacitance using equation 1 to obtain in Fig. 6. In Eqn. 1, C(w) is the complex capacitance, w = (–1)1/2, RHFL represents the high frequency limit (10 MHz) of the impedance Z(w). C(w ) =

("Eqn. 1" )

1 jw (Z(w) - R HFL )

Capacitance dispersion is clearly seen in Figure 6 all over the scanned frequency range, showing two well differentiated time constants (one in the 100 kHz to 10kHz range and the other one close to 1kHz) with a low frequency limit close to 6 mF cm–2 for AISI 430, 5mF cm–2 for Ni base-alloy and 4 mF cm–2 for both austenitic steels. Thus, the dispersion of electrical properties all throughout the passive films becomes evident. The Cole–Cole spectra have been modelled using the equivalent circuit depicted in Fig.7, to which Eqn. 2 corresponds. The best fitting parameters obtained for the studied materials are summarised in Table 2. The high frequency limit, C1, is poorly defined and highly dependent on RHFL, so an average value from Figures 2 to 4 has been fixed in the fitting procedure. 4

5

4

3

2

1 kHz

1

0

1

2

3

4

5

Re (Capacitance) / mF cm

C1

3

2

R2

C2

R3

C3

1 100 kHz

100Hz

10MHz

0

Experimental Fitting

-2

AISI 430 AISI 304L AISI 316 Ni base alloy

-Im (Capacitance) / mF cm

-Im (Capacitance) / mF cm

-2

6

10 kHz

1 kHz

10 MHz

6

0

-2

0

100Hz

1

2

3

4 -2

Re (Capacitance) / mF cm

Figure 6: Cole–Cole plot obtained from the Figure 7: Experimental and best fitting data impedance data depicted in the inset of Figures 2, corresponding to the Cole–Cole spectra for 3 and 4. AISI 316 at 96 h immersion in NaOH 0.1M.

C(w ) = C1 +

C3 C2 + a 2 b2 (1 + ( jw R 2 C2 ) ) (1 + ( jw R 3 C3 )a3 ) b3

("Eqn. 2." )

Besides C1 parameter, that will be characteristic of the Cr2O3 islands, two parallel paths for electronic conduction are clearly differentiated. With capacitance values around 2 mF.cm–2, the charge carrier concentration results in the order of 1020 cm–3, value typically found for semiconductors [4]. The conduction path at higher frequency (R2C2) shows charge carrier concentration

34

C. M. Abreu et al.

decreasing as Ni-content increases, while low frequency one (R3C3) shows higher carrier concentration for the two extreme Ni compositions. Thus probably the former path can be associated to magnetite (oxidised by Ni2+ [8]) and the latter to grain boundaries conduction. Table 2. Best fitting parameters for data in Figure 6 using equation 2. C2 / mF.cm-2 3.22

C3 / mF.cm-2 2.84

R2 / Wcm2 15

R3 / Wcm2 124

g2

g3

AISI 430

C1/ pF.cm-2 70

0.99

0.97

0.99

0.99

AISI 304L

70

2.54

1.40

2.65

142

0.97

0.77

1

1

AISI 316

70

2.40

1.60

2.33

355

1

0.81

1

0.60

Ni Alloy

70

1.20

4.60

4.22

267

0.89

0.64

0.99

0.58

2

3

4. Conclusion The above discussed results allow concluding that the conductivity of airformed films on Fe-Cr-Ni alloys decreases as Ni content increases. Less conducting films are thinner and stand longer in alkaline media. These films are not homogeneous from the electrical point of view. Two well differentiated conducting paths appear in parallel to dielectric islands. Acknowledgement The authors whish to acknowledge the Spanish “Ministerio de Educación y Ciencia” for financial support under project MAT2004-06435-C02-01. 5. References 1. M. Bojinov, I. Betova, G. Fabricius, T. Laitinen, R. Raicheff, T. Saario, Corros. Sci., 41 (1999) 1557. 2. T. Nishimura, T. Kodama, Corros. Sci., 45 (2003) 1073. 3. C. M. Abreu, M. J. Cristóbal, M. F. Montemor, X. R. Nóvoa, G. Pena and M. C. Pérez, Electrochim. Acta, 47( 2002) 2271. 4. C.M. Abreu, M.J. Cristóbal, R. Losada, X.R. Nóvoa, G. Pena, M.C. Pérez, J. Electroanal. Chem., 572 (2004) 335. 5. C.M. Abreu, M.J. Cristóbal, X.R. Nóvoa, G. Pena, M.C. Pérez, Electrochim. Acta, 47 (2002) 2215. 6. C.M. Abreu, M.J. Cristóbal, R. Losada, X.R. Nóvoa, G. Pena, M.C. Pérez, Electrochim. Acta, 49 (2004) 3049. 7. C.M. Abreu, M.J. Cristóbal, X.R. Nóvoa, G. Pena, M.C. Pérez, Surf. Coat. Tech., 158-159 (2004) 582. 8. C.M. Abreu, M.J. Cristóbal, R. Losada, X.R. Nóvoa, G. Pena, M.C. Pérez, Electrochim. Acta, submitted.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

35

Passivity of Nitrogen-Bearing Stainless Steel in Acidic Solution Sanja Martineza, Mirjana Metikoš-Hukovića, Nushe Lajcib a

University of Zagreb, Faculty of Chemical Engineering and Technology, Marulićev Trg 19, P. O. Box 177, HR-10000 Zagreb, Croatia, Phone +385(1)4597116, Fax +385(1)4597139, E-mail [email protected] b University of Prishtina, Mining and Metallurgical Faculty, str Lead Factory. 38000, Mitrovica, Kosova +33744221107, E-mail [email protected]

Abstract - The mechanism and kinetics of the passivation process, the mechanisms of charge transfer and potential distribution in the system NTR50 steel½passive film½0.5 M H2SO4, were studied using transient electrochemical techniques in combination with cyclic voltammetry and the earlier results of XPS analysis. It has been established that the point defect model, that takes into account the interfacial kinetic effects, presents the best fit to the experimental data yielding physically reasonable parameters. The model has been applied to quasi-stationary experimental data to determine: (i) the transfer coefficient for the metal dissolution reaction through the barrier film, a11 = 0.017 (ii) the proportionality coefficient between the barrier film½solution potential and the applied potential, a = 0.446 and (iii) the electric field strenght, e = 4.4 ´106 V cm-1. Cation vacancies, VCr3- represent the main charge carriers under quasi – stationary conditions of formation and growth of the barrier Cr2O3 film. The barrier layer on NTR50 stainless steel is essentially a cation conductor. The outward migrating species are presumed to be the Fe3+ ions and the molybdenum Mo4+ and Mo6+ ions, in concordance with the film composition comprised of an inner Cr2O3 barrier layer and a Fe-rich outer hydrated layer. Keywords: stainless steel, nitrogen-bearing, point defect model

36

S. Martinez et al.

1. Introduction Passive films represent a topic of special interest today, though electrochemists have been concerned with passivity for almost 200 years [1]. Nowadays oxide films play an increasing role in microelectronics and micromechanics. Due to their wide application in materials science, thin film technology and corrosion, oxide films and passive films have been discussed at the ISE meetings [2 and ref. therein]. The influence of nitrogen as an alloying element on corrosion of austenitic steel was previously investigated in various aggressive electrolytes [3]. It was found that austenitic steel NTR50 with interstitially dissolved nitrogen in the austenite microstructure passivates spontaneously in 0.5 M H2SO4 and 1M NaCl solutions. The anodic polarization of nitrogen bearing austenitic stainless steel (NTR 50) was found to proceed in essentially two stages, with an initial formation of pure Cr2O3 film. According to the XPS data, other alloying elements accumulate underneath the Cr (III) oxide film during a steady – state enrichment of the alloy at the alloy½film interface via solid-state relations. Subsequently, during anodic polarization, iron and molybdenum ions get incorporated into the anodic Cr (III) oxide film in their alloy proportions according to the lattice conservative solid-state reactions. It was suggested that nitrogen addition enhances passivity of austenitic stainless steel via enhancement of anodic segregation process. Segregation is believed to be tied to the formation of an intermetllic interphase with alloying additions such as Mo and Ni, beneath the oxide. The aim of this work is to improve the understanding of passivity phenomena at the atomic level of new class of engineering materials – nitrogen alloyed austenitic steel NTR50 by studying the oxide growth and steady state. 2. Experimental Nitrogen bearing austenitic stainless steel NTR50 was used as a test specimen. Measurements were preformed in PAR cell and the potentials in the text refer to the SCE scale. Prior to measurements, samples were abraded with wet SiC emery paper down to grit size 1000 and were polished successively with alumina suspensions of granulation ranging from 1 mm to 0.05 mm and degreased in isopropanol steam, rinsed in doubly distilled water, ultrasonically cleaned, and dried in air. All experiments were carried out at room temperature (22 °C) under continuous deaeration with high purity nitrogen in 0.5 M H2SO4 solution, prepared from analytical grade acid using doubly distilled water. Electrochemical measurements were done using a PAR&EGG (model 273) potentiostat / galvanostat. Prior to all measurements, the sample was cathodically prepolarized at -0.7 V for 60 s. Potentiostatic current transients during film formation were recorded for 900 s at a step of 0.1 V from 0.3 to 0.9 V. The potential steps were taken, from -0.65 V to the formation potential,

Passivity of nitrogen-bearing stainless steel in acidic solution

37

immediately following the pretreatment. Short-term galvanostatic transients were recorded for 3 s also at a step of 0.1 V, from 0.3 to 0.9 V. The square wave current signals of amplitudes ranging from 10 to 100 mA cm-2 were applied starting from the current value obtained after stabilization at the corresponding formation potential. After each signal the system was left polarized at formation potential until current recovered approximately to the starting value. Galvanostatic charging curves were recorded during film formation for 300 s at each step of 10 mA cm-2 from 30 to 100 mA cm-2. 3. Results and discussion Study of the oxide growth and steady state and application of various models to the experimental data is a common method of investigation of protective properties of passive films and processes that lead to their formation. Two theories, the high-field model (HFM), and the variants thereof, and the point defect model (PDM) have been advanced to account for the properties of thin anodic passive films [4].

Figure 1. a) galvanostatic charging curves at formation current densities of 30, 35, 40, 50, 60, 70, 80 and 100 mA cm-2, respectively represented by curves 1 – 8 and b) potentiostatic current transients for various potentials applied immediately after cathodic pretreatment at -0.7 V for 60 s. Vext equals 0.3, 0.4, 0.5, 0.6, 0.7, 0.8 V, respectively represented by curves 1 – 6. The measurements were recorded on NTR50 steel at 22 °C in 0.5 M H2SO4.

For that purpose, galvanostatic charging curves (Figure 1a) and potentiostatic transients (Figure 1b) have been recorded in the passive potential range of the NTR50 - 0.5 M H2SO4 system. Application of the HFM to galvanostatic charging curves according to ref. [5] yielded the electric field across the passive layer ~ 107 V cm-1, the value above those expected to cause dielectric breakdown. Such result is often calculated for very thin passive films [4] as the HFM does not limit the field strength. A more physically realistic description should be called for.

38

S. Martinez et al.

In the presently investigated system, a linear variation of the logarithm of the steady state current density iss with the applied voltage Vext has been observed and is consistent with the PDM model [4]. The experimental data shown in Figure 2a were fitted to the PDM relation:

¶ ln ( iss ) ¶Vext = a11adg

(1)

where a represents the proportionality coefficient between the barrier film½solution potential difference ff/s and the applied potential Vext. a11 is the transfer coefficient for the metal dissolution reaction through the barrier film. The correlation coefficient of the straight line in Figure 2a equals 0.997. From its slope the product aa11 was found to be equal to 0.0083.

Figure 2. Dependence of the a) quasi-steady state passive current density and b) oxide thickness deduced from the cyclic voltammetry measurements, on the passive film formation potential for NTR50 steel at 22 °C in 0.5 M H2SO4.

The above results may be linked to the electronic properties of the barrier layer. The PDM attributes the linear increase in the logarithm of iss with Vext to the predominantly p-type semi conducting mechanism [5]. This finding is concordant with the view that points to the generation of soluble products as a main process occurring on the Fe-Cr-Mo alloys in passive the region [6]. This has been mostly attributed to the selective dissolution of Fe from the alloys whereby Fe(III) is assumed to substitute Cr(III) in hexagonal a-Cr2O3 lattice via a solid-state reaction Fem+VCr3-®Fe(III)Cr+3e-. Molybdenum also contributes to the steady state current trough the reactions Mom+VCr3-®Mo(IV)Cr+4e- and Mo(IV)Cr ®Mo(VI)Cr+2e-. The substitution of Fe3+ and Mo4+ (Mo6+) ions inside mobile cation vacancies of the barrier film, their pairing and formation of neutral complexes may cause a drastic drop in the diffusivity and concentration of cation vacancies. The increase in the electronic conductivity of the film, enhances the passivity of steel, while film growth and dissolution occur very slowly.

Passivity of nitrogen-bearing stainless steel in acidic solution

39

The steady state barrier film thickness Lss also varies linearly with the applied voltage Vext (Figure 2b) and according to the PDM model:

( ¶Lss

¶Vext )pH,C

M h+

= (1 - a ) / Î

(2)

where e is the electric field. The unknowns contained in Eqs. (1) and (2) are: a, a11 and e. Apparently, more unknowns than observables exist, so that steady state properties cannot be used to obtain an unequivocal set of model parameters unless some of the parameters can be determined in an independent experiment [8]. For the purpose of determination of the transfer coefficient a11 in an independent experiment, has to be conducted. When a potential, denoted by Vext, is applied at the oxidized metal surface, it is distributed to yield potential drops at different parts of the system i.e. Vext = fm/f +ff + ff/s +fR where fm/f and ff/s are the potential drops at the metal½barrier film and barrier film½solution interfaces, respectively, ff is the potential drop in the barrier film and fR is the potential drop at the solution/reference electrode interface. Taking into account Lss being a linear function of Vext and a being constant and knowing that at steady state e = Dff / DLss, Eq. (2) may be rewritten in the following manner:

Î= ⎡⎣

(1 - a )Vext

Lss ⎤⎦ pH,C

Md +

= ff

Lss

(3)

It is apparent from the above relation that when Eq. (3) applies, the whole change of (1-a)Vext potential difference is reflected trough the change in ff, while the potential drop fm/f (as well as fR) remains constant.

Figure 3. a) short-term galvanostatic transients obtained for NTR50 steel electrode (initially at quasi-steady state at the formation potential of 0.4 V), after application of square-wave current signals of amplitudes from 10 to 100 mA cm-2 every 10 mA cm-2 (respectively represented by curves 1-10) and b) Tafel lines deduced from short-term galvanostatic transients for the film initially stabilized at 0.4, 0.5, 0.6 V respectively represented by curves 1 – 3.

40

S. Martinez et al.

Providing that under specific experimental conditions, Lss of the film formed at particular Vform could be kept constant, while Vext is changed, the transport coefficient a11 of dissolution reaction could be measured. The above condition is assumed to apply when short term square wave potentiostatic or galvanostatic pulses are applied to the electrode at which the oxide layer has been previously formed at Vform [9]. The initial polarization data may be determined by back extrapolation from such curves and they obey the Tafel law with the slope depending on Vform (Figure 3a). The logarithm of the current density is, therefore, linearly dependent upon Vext with the Tafel slope, ba = (a11d g )-1 (Figure 3b). The average value of a11 = 0.017 has been computed from ba measured at formation potentials between 0.4 and 0.6 V. Furthermore, a may now be computed. The value of a = 0.446 is in a good agreement with the literature data [7]. Once the parameter a has been calculated, the electric field may be computed from the known dLss/dVext dependence. The electric field calculated from Eq. (2) equals 4.44 ´ 106 V cm-1. This value lies within the range of most commonly obtained field strengths of Fe-Cr-Mo alloys (4 – 6 V cm-1) [8]. 4. Conclusions The PDM has been applied to the results of the transient electrochemical measurements and the results deduced from cyclic voltammetry preformed on NTR50 steel at 22 °C in 0.5 M H2SO4. A predominantly p-type semi conducting mechanism has been found to be operative and cation vacancies, VCr3- represent the main charge carriers under quasi – stationary conditions of formation and growth of the barrier Cr2O3 film. The conclusions are consistent with the earlier results of XPS analysis as well as with the principal observations concerning the passivity of Fe-Cr-Mo alloys. References 1. Uhlig H., in: R.P.Frankenthal, J. Kruger (Eds.), Passivity of Metals, The Electrochemistry Society, The Corrosion Monograph Series, Princeton, NJ, 1978, p.1. 2. Schultze J.W., Lohrengel M.M., Electrochim. Acta, 45 (2000) 2499. 3. Vehovar L., Vehovar A., Metikoš-Huković M., Tandler M., Materials And CorrosionWerkstoffe Und Korrosion, 53 (2002) 316-327. 4. Sikora E, Macdonald D.D., Solid State Ionics, 94 (1997) 141-150. 5. Ammar I.A., Darwish S., Khalil M. W., Z. Werkstoffech. 12 (1981) 421-431. 6. M. Bojinov, G. Fabricius, T. Laitinen, K Makela, T. Saario and G. Sundholm, Electrochim. Acta. 46 (2001) 1339. 7. D. Hamm, C.-O. A. Olsson and D. Landolt, Corros. Sci. 44 (2002) 1009. 8. D. D. Macdonald, S. R. Biaggio and H. Song, J. Electrochem. Soc. 139 (1992) 170. 9. T. Hurlen, Electrochim. Acta. 39 (1994) 1359.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

41

Passive behaviour of stainless steels and nickel in LiBr solution at different temperatures A. Igual Muñoz, J. García Antón, J.L. Guiñón Segura, V. Pérez Herranz Departamento de Ingeniería Química y Nuclear. ETSI Industriales. Universidad Politécnica de Valencia. P. O. Box 22012, E-46071 Valencia. Spain Tel. +34 963877632 Fax. +34 963877639 e-mail: [email protected]

Abstract - The aim of the present work was to study the passive behavior of different stainless steels in commercial LiBr heavy brine solution (850 g/L, which contains chromate as inhibitor), at different temperatures (25, 50, and 80ºC). LiBr solutions are the most common medium in the refrigeration systems which uses the absorption technology. The materials tested were stainless steels EN 1.4591, with nominally (wt%) 33 Cr, 32 Fe, 31 Ni; EN 1.4562, an iron-nickelchromium-molybdenum alloy with nitrogen addition; and two austenitic stainless steels EN 1.4311 and EN 1.4429. Pure nickel was also tested to study its influence in the alloys. Passive behaviour was estimated from the cyclic potentiodynamic polarization curves, analysing the pitting potentials, passive current densities and repassivating properties at different temperatures. Keywords: Stainless steels; Lithium Bromide; Repassivation; Inhibitor; Temperature.

1. Introduction Absorption technology is one of the best alternatives to compression cooling in terms of energy diversification and environmental protection. Recently, the triple-effect absorption cycle has attracted much interest to replace the conventional machine as the more efficient one [1]. The conventional LiBr/H2O solution is known to cause a serious corrosion problem to the metal part of the machine operating at high temperature (specially the generator). Nowadays, the most

A. Igual Muñoz et al.

42

commonly used material for absorption machines is type SS 316L [2-4]. However, failures in the form of pitting and crevice corrosion are sometimes observed in the hottest parts of the. With the advances in materials, newly developed alloys with greater corrosion resistance and better mechanical properties have been introduced. Austenitic and duplex stainless steels, and nickelbased alloys are now available for absorption machines applications. The objective of the present work was to study the influence of temperature on the pitting and general corrosion behaviour of several high-alloyed stainless steels and nickel in a commercial LiBr solution. The aim of the evaluation of their relative resistance to pitting and general corrosion is to provide a reliable basis in the selection of structural materials for absorption machines.

2. Experimental The materials tested and their chemical composition (%wt) were summarized in Table 1. Table 1. Electrochemical data of the materials tested in LiBr solution at different temperatures. Material (EN)

Fe

1.4591

31.43 0.005

0.017 1.5

1.4562

32.2

0.007

0.01

14311

Bal

0.015

14429

Bal

0.019

Nickel

C

P

Mn

S

Si

Ti

Cu

Cr

Ni

N

Mo

0.002

0.1

-

1.21

26.75

31.85

0.193

6.6

0.004

0.31

-

0.58

32.85

30.95

0.39

1.7

0.033 1.383

0.026

0.379

0.002

0.35

18.55

8.798

0.135

0.41

0.027 1.575

0.001

0.27

0.005

0.267

17.81

11.64

0.16

2.83

0.54

100

The solution used was a commercial LiBr solution, from FMC Corporation Lithium Division (U.S.A.), which has a concentration of 850 g/L LiBr and contains Li2CrO4 (4.3 g/L) and LiOH (0.08 g/L) as additives. It is generally accepted [5] that chromate is an effective corrosion inhibitor which can passivate metals by forming a monoatomic or polyatomic oxide film on the electrode surface, not only for iron and steels but also for many other metals and alloys. The tests were conducted in a three-electrode cell held in a water bath at a constant temperature. The working temperatures were kept constant during the whole tests at 25ºC, 50ºC and 80ºC.The electrolyte was continuously purged with purified nitrogen gas. Before each polarization, the probe was immersed in the test solution for 1 h at the open-circuit potential (OCP).. The sample was polarized anodically at a scan rate of 0.5 mV/s from –1000 mV to the potential at which the current density reached 10 mA/cm2. At this potential, scanning was reversed. The potential was measured with reference to an Ag/AgCl 3M KCl electrode. The counter electrode was a platinum wire. Corrosion current density (icorr) and corrosion potential (Ecorr) were obtained from the curves using the Tafel slopes. The potential at which the current density exceeded 100 mA/cm2 was defined as the pitting potential (Ep). Passive current density (ip) and repassivation potentials (Erp) were calculated.

Passive behaviour of SS and Nickel in LiBr solutions at different temperatures

43

3. Results and data analysis Figure 1 shows the cyclic curves of all the materials at 25ºC. All alloys exhibit hysteresis, characteristic of passivation breakdown on the upward sweep and repassivation at or near their corrosion potential Ecorr., which indicates that the repassivation of an existing pit is difficult. The larger the hysteresis loop, the more difficult repassivation will be. 5

LOG lil (m A/cm2)

EN 14591

EN 14311

4 3

EN 14429 2 Ni 1 0

-1 EN 14562 -2 -3 -1000

-800

-600

-400

-200

0

200

400

POTENTIAL (mV vs Ag/AgCl)

Figure 1. Cyclic potentiodynamic curves for all the tested materials in a commercial LiBr solution at 25ºC.

(a) 270 mV

(b) 310 mV

(c) 80 mV (reverse scan)

(d) 75mV

(e) 110 mV

(f) 180 mV

(g) 51.0 mV

(h) 63.7 mV

(i) 70.4 mV

Figure 2. Images of EN 14562 (a), (b) and (c); Nickel (d), (e) and (f) and EN 14429 (g), (h) and (i) at different moments of the test in commercial LiBr solution at 25ºC.

44

A. Igual Muñoz et al.

At 25 ºC all the materials were spontaneously passive. This is in accordance with their high corrosion resistance. Pitting and repassivation potentials are higher for EN 14591 and EN 14562, which show lower passive current density. But both alloys always reached the highest current after the scan was reversed, from 25 ºC to 80 ºC. Images of the electrodes during the scan were acquired using an optical device coupled to the electrochemical system. Figure 2 shows the morphology of the corrosion of two stainless steels and the nickel. It is possible to observe the different way corrosion advance: for the stainless steels, corrosion product catalizes de corrosion process so that the metallic surface degrades along the wetted zones; by other way, nickel corrodes thorough more generalized way although this way corresponds to multiple localized sites around the nickel surface. At 50 ºC cyclic scan exhibited the same tendency, typical of a passive material: the anodic current density maintains a value of » 1 mA/cm2 up to a potential between » 100 and 300 mV depending on the material, then an abrupt increase of the current density is observed, Figure 3. 5 EN 14591

4 EN 14311

LOG lil (mA/cm2)

3 2

Ni

1

EN 14562

0 -1

EN 14429

-2 -3 -4 -1000

-800

-600

-400

-200

0

200

400

POTENTIAL (mV vs Ag/AgCl)

Figure 3. Cyclic potentiodynamic curves for all the materials tested in a commercial LiBr solution at 50ºC. 6 EN 14591 5 EN 14311

LOG lil (mA/cm2)

4

EN 14429

3 2

Ni 1 0

EN 14562

-1 -2 -3 -1000

-800

-600

-400

-200

0

200

POTENTIAL (mV vs Ag/AgCl)

Figure 4. Cyclic potentiodynamic curves for all the materials tested in a commercial LiBr solution at 80ºC.

Passive behaviour of SS and Nickel in LiBr solutions at different temperatures

45

When temperature increased above 50 ºC, Figure 4, the hysteresis loop of all the alloys notably increased, which means that they reduce their repassivating nature, in fact EN 14429 and EN 14311 did not repassivate at this temperature.

3.1. Corrosion current and corrosion potential Corrosion potentials and current values are summarized in Figure 5. 5

EN 14591 EN 14562 Nickel EN 14311 EN 14429

-200 -300 -400 -500 -600 -700 -800 20

30

40

50

60

70

80

90

TEMPERATURE (ºC)

CORROSION CURRENT DENSITY 2 (mA/cm )

CORROSION POTENTIAL (mV)

-100

EN 14591 EN 14562 Nickel EN 14311 EN 14429

4.5 4 3.5 3 2.5 2 1.5 1 0.5 0

20

40

60

80

TEMPERATURE (ºC)

(a)

(b)

Figure 5. Corrosion potentials (a) and corrosion current densities (b) of the three materials in a commercial LiBr solution at different temperatures. Corrosion potentials did not present any clear trend with temperature, while corrosion current values increased in all the alloys when temperature increases. The highest corrosion current was shown by EN 14562. However the ability of generating a passive film makes this alloy very suitable for absorption systems applications. EN 14429 corroded at the lowest corrosion rates from 25ºC to 80 ºC. Only above 50 ºC icorr significantly increases with temperature, which could also be due to the lost of chromate efficiency at those higher temperatures.

3.2. Resistance to localized corrosion These alloys are less pitting corrosion resistant and they show diminution in repassivation properties as temperature increases. A linear trend is observable in the range of temperatures 2580ºC, in the case of EN 14562 pitting potential decreases at 3.2 mV/ºC. The tendency presented by Erp also decreased with temperature; this trend was also observable in the difference Ep- Ecorr. This fact was more noticeable for pure nickel above 50 C. The passivation range did not narrow significantly with temperature for EN14562, which also presented the widest passive region at all temperatures (except at 25ºC). With respect to the passive current, EN 14591 showed the lowest values at all the temperatures under study, while EN 14311 exhibited the highest passive current in most of the tests. Thus in the cases studied, the lowest pitting potential means a higher passive current in the passive region. The protective properties of the passive films decrease as the temperature of the electrolyte increases.

46

A. Igual Muñoz et al.

Table 2. Electrochemical data of the materials tested in LiBr solution at different temperatures.

25ºC

50ºC

80ºC

Materials

Ep (mV)

ip (mA/cm2) Erp (mV) Ep-Erp (mV) Ep-Ecorr (mV)

EN 14591

274.9

0.95

20.9

254

978.1

EN 14562

322.0

1.17

118.7

203.3

825.2

Ni

302.7

1.76

70.3

232.4

876.1

EN 14311

-40

4.6

-240

200

390

EN 14429

56

0.8

-67

164

506

EN 14591

126.3

1.13

-46.4

172.7

855.7

EN 14562

241.7

1.59

28.62

213.1

864.9

Ni

172.2

1.52

-55.2

227.4

652.5

EN 14311

-95.2

5

-267

172

76

EN 14429

40

1.2

-160

120

269

EN 14591

49.8

2.1

-210.1

259.9

745.9

EN 14562

154.7

3.7

-105.2

259.9

862.6

Ni

-51.4

6.5

-164.2

112.8

584.0

EN 14311

-163.2

-

-

-

-

EN 14429

21.18

2.2

-

-

171

4. Conclusions The results of cyclic polarization studies indicate excellent corrosion resistance by stainless steels in LiBr solution at room temperature and the absence of big differences at temperatures above 50 ºC. The pitting corrosion resistance of all alloys decreases with increasing temperature. Molybdenum plays an important role in the localized resistance of the alloys and in this sense EN 14562 is the most suitable alloy for the design in the absorption machines. EN 14311 and EN 14429 do not repassivate at 80 ºC. However, the hysteresis loop in the cyclic scans became bigger with temperature for the rest of materials. Acknowledgements: To MCYT (PPQ2002-04445-C02-01) for the financial support of this research and to Dr. M. Asunción Jaime for her translation assistance.

References 1. 2. 3. 4. 5.

J.S. Kim; F. Ziegler; H. Lee Applied Thermal Engineering 22 (2002) 295-308 J.L. Guiñón; J. Garcia-Anton; V. Perez-Herranz; G. Lacoste. Corrosion 50(3) (1994) 240. A.Igual Muñoz;J.García Antón;J.L.Guiñón;V.Pérez Herranz. Corrosion. 59(7) (2003) 606. A.Igual Muñoz;J.García Antón;J.L.Guiñón;V.Pérez Herranz. Corrosion. 58(12) (2002) 995 A.Al-Odwani; J. Carew; M. Al-Tabtabaei; A. Al-Hijji. Desalination 135 (2001) 99-110.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Published by Elsevier B.V.

47

The effect of the Cerium ion implantation in the passive films properties of a duplex stainless steel C.M.Abreu, M.J. Cristóbal*, X.R Nóvoa, G.Pena, M.C. Pérez, C. Serra ET.S.E.I. e M. Campus Universitario, Universidade de Vigo 36310 Vigo (Spain) *Corresponding author: e-mail: [email protected]

Abstract - The effect of Ce+ implantation on the electrochemical behaviour of SAF 2205 DSS is studied in alkaline medium. XPS studies have detected the implanted cerium as Ce3+ throughout the oxide films. The peak current of the magnetite formation peak is directly related to the passive film thickness. Nevertheless, the Cr3+ oxidation process is not affected by cerium implantation, which suggests cerium incorporation in the iron spinel.

Keywords: Cerium Implantation, Duplex Stainless Steel, SAF 2205, XPS, alkaline medium 1. Introduction Active elements with high oxygen affinity, such as yttrium, cerium, lanthanum and other rare earths, are known as elements that improve the oxidation behaviour of “high temperature” alloys, whose protection derives from the formation of Cr2O3 scales [1-3]. However, concerning the effect of those elements on electrochemical corrosion of stainless steels few papers have been found in the reviewed literature [4-6]. One way to incorporate the active elements to the metallic matrix is using ion implantation. This technique gives the unique possibility of introducing controlled concentration of an element in a thin surface layer, typically up to 100 nm thick. The few works related with Ce-implanted steels observed cathodic reaction inhibition that was attributed to the formation of a stable cerium oxide [4]. The present investigation analyses the effect of Ce+ implantation in the development of oxide layers on a duplex stainless steel, SAF 2205.

48

C. M. Abreu et al.

2. Experimental Samples of SAF 2205 (22.38 wt.% Cr, 5.39 wt.% Ni, 2.94 wt.% Mo, 1.68 wt.% Mn, 0.51 wt.% Si, 0.019 wt.% C, Fe balance) with dimension of 1.1 x 1.1 cm2 were used. Before implantation, the samples were ground with silicon carbide paper up to 600 grit, then mechanically polished with diamond paste to 6 mm finish. Implantation of one coupon face, at the nominal dose of 1x1017 Ce+ ions.cm-2, was undertaken using 150 keV acceleration voltage. Electrochemical experiments were performed at 30°C in pre-deaerated NaOH 0.1M solutions and under continuous Ar bubbling, in a conventional threeelectrode cell, where the working electrode was the selected steel (0.5 cm2 of test area). A Pt mesh was used as large area counter electrode and a Hg/HgO 0.1M KOH as the reference electrode. An AUTOLAB 30 Potentiostat (EcoChemie) was used for cyclic voltammetry. The oxide films are formed by scanning the potential from hydrogen to oxygen evolution reactions. In order to establish the evolution of the voltammetric curves and to generate films thick enough to be characterized, each sample was cycled eight times at 1mVs-1 scan rate. The films were chemical characterized by X-ray photoelectron spectroscopy (XPS) using a VG ESCALAB 250iXL spectrometer. The XPS data were collected using monochromatic Al Ka radiation at 1486.92 eV, and at 20 eV constant analyser pass energy. Depth Profile experiments were performed using an EX05 Ar+ Ion Gun at 3 kV. A survey spectrum was first recorded to identify all elements present at the surface, then high resolution spectra of the following regions were recorded: Fe 2p, Cr 2p, Ni 2p, Mo 3d, Ce 3d, and O 1s. Sputtering rate values were determined using perfilometry measurements. 3. Results 3.1. Cyclic voltammetry measurements Figure 1 depicts the last polarization curve corresponding to eighth cycle for both materials. The main characteristics of the three potential regions, appearing in this voltammograms, have been previously described [7] and can be summarised as follows: The first potential region (from hydrogen evolution to about -0.6V) is mainly related to iron redox processes: Fe/Fe2+ and Fe2+/Fe3O4. Region II corresponds to the passivity region, from –0.6 to 0V, and finally, in Region III (between about 0V and 0.6V), two main processes take place: Cr3+/Cr6+ and Ni2+/Ni3+. Making a comparison for the unimplanted steel voltammogram and the corresponding Ce-implanted, it is possible to note an important change in their morphology. Peak currents are lower for the implanted steel in region I (that corresponds to Fe redox processes). However, an increasing in the passivity current (Region II) is registered for the implanted steel due to Ce3+/Ce4+ redox process. As it concerns Region III, the peak

49

The effect of the Cerium ion implantation

assigned to the Cr3+/Cr6+ oxidation shows small density increase current for implanted steel. The same was observed for the current at more anodic potentials. The anodic current increasing observed at +0.5V for the implanted steel could be attributed to an activation of the O2 evolution process, because the signals assigned to Ni redox process do not change significantly. 100 Region I

Region II

Region III

+2

Current Density / mA cm

-2

Fe /Fe3O4 50 3+

Cr /Cr

0

2+

6+

Ni /Ni 3+

Ce /Ce

4+

om Zo

600

-50

400

200

SAF 2205 Ce-implanted SAF 2205

0 0.40

-100 -1.4

3+

-1.2

-1.0

-0.8

-0.6

-0.4

-0.2

0.0

0.45

0.2

0.50

0.4

0.6

Potential / V vs. Hg/HgO

Figure 1: Cyclic voltammograms obtained in NaOH 0.1M solution for SAF 2205 and Ce-implanted SAF 2205. Potential range: -1.4 to +0.5 V (vs. Hg/HgO). dE/dt=1 mVs-1

3.2. X-ray photoelectron spectroscopy (XPS) In order to clarify the elemental composition of the electrochemical formed passive film on duplex stainless steels in NaOH medium XPS depth profiles were acquired. For the passive layers develop after polarization, XPS measurements show that main constituents are: OH-, O2-, Fe2+, Fe3+ and Cr3+. The curve fitting of the spectra has been published in a previous work [7]. Fig. 2 shows the XPS depth profiles of the oxide films developed on unimplanted SAF 2205 after electrochemical measurements. The obtained results (Fig. 2) indicate that mainly Fe oxide (Fe2+ and Fe3+) was formed containing small amount of Cr3+ through the oxide layer, although slightly concentrate at the oxide/metal interface. The thickness of the oxide layer, about 95 nm, is taken at the end of the oxygen profile. There is no evidence of the presence of Mo species in the passive layer, and Ni2+ is present only as traces in the hydroxide outer part. The two oxidation states of iron correspond to a Fe3O4 or a mixture of FeO and Fe2O3. Since Fe3O4, an inverse spinel, is structurally similar to a mixture of FeO and Fe2O3 [8], no difference between the two oxides is measurable with XPS.

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Moreover, it is important to note that it is not possible to determine whether the oxide film is a mixture of two oxides (Cr2O3 + Fe3O4) or a mixed ironchromium spinel in which chromium replaces some of the iron positions. 70

2-

O OH Fe Metal Cr Metal Ni Metal Mo Metal Fe Oxide Cr Oxide

60

Atomic Conc. %

50

40

30

20

10

0 0

10

20

30

40

50

60

70

80

90

Depth (nm)

Figure 2: XPS depth profile obtained for unimplanted SAF 2205 after eight cycles in NaOH 0.1M solution

Cerium implantation in SAF 2205 penetrates » 60 nm, as evidence XPS depth profile, as it can be observed in Fig. 3. The high-resolution spectrum of Ce 3d shows a complex structure (Fig. 4a and b). The peaks show spin-orbit split into a doublet (3d5/2 and 3d3/2) with each doublet showing further structure due to final state effects. Taking account the reference spectra of CeO4 [9] and Ce2O3 [10], the XPS spectra obtained for the cerium implanted steel can be identified as Ce3+. Quantification has been done using 3d3/2 ionisation (mean peak at 899.9 eV and satellite peak at 904.2 eV), because iron-Auger peak at 888.0 eV interfere on the quantification using the 3d5/2 ionisation. The oxidation state of cerium ions are the same on the layer developed after polarization. One of the most important effects of cerium implantation is the reduction of the film thickness formed under cyclic polarization (about 50 nm). The presence of cerium does not seem to affect the shape of the element depth profile for iron ions and the Cr/Fe ratio; however, the distribution of Cr3+ throughout a film is a little bit different showing enrichment on the most outer part of the oxide layer. The XPS results cannot establish whether the cerium is incorporated to the

The effect of the Cerium ion implantation

51

passive layer as an independent oxide (Ce2O3) or as a mixed Fe-Ce oxide, probably a Fe-rich spinel with Ce3+ incorporated. 70

60 2-

O OH FeMetal CrMetal NiMetal FeOxide CrOxide CeOxide

Atomic Conc. %

50

40

30

20

10

0 0

10

20

30

40

50

60

70

80

Depth (nm )

Figure 3: XPS depth profile obtained for SAF 2205 implanted with cerium after eight cycles in NaOH 0.1M solution

B

KCS

A

Binding Energy/ eV

Binding Energy/ eV

Figure 4: A) High resolution XPS for Ce 3d emission through the passive film on SAF 2205. B) Identification of the peaks at Ce 3d emission.

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4. Discussion The aforementioned results demonstrate that the change in the electrochemical behaviour of duplex stainless steel is due to the modifications that cerium implantation promotes in the passive layer and in his formation. The XPS results can not establish whether the cerium is incorporated to the passive layer as an independent oxide or as a mixed Fe-Ce oxide. However, as Fig. 1 shows big effect on iron redox processes and almost no modification of the Cr3+ oxidation peak, the most likely hypotheses is that cerium is incorporated to the iron oxide lattice. In any case, the presence of Ce induces an increase in the protective character of the passive film. This fact is confirmed by the diminution in current density of iron redox processes in the corresponding voltammograms (Fig. 1). It can be noted the good agreement between the reduction in film thickness and the decrease in magnetite peak in the implanted sample after cycling. 5. Conclusions In the present work cerium implantation has been applied to modify the surface of a duplex stainless steel in order to improve its corrosion resistance. The surface characterization techniques have detected the implanted cerium as Ce3+ throughout the oxide films. The peak current of the magnetite formation peak is directly related to the passive film thickness. Nevertheless, the Cr3+ oxidation process is not affected by cerium implantation, which suggests cerium incorporation in the iron spinel. References 1. M. F. Stroosnijder, M. J. Cristóbal, and J. D. Sunderkötter, Mater. Sci. Forum, 251–254 (1997) 254 2. K. Przybylski and G. J. Yurek, Mater. Sci. Forum, 43 (1989) 1. 3. J. Pérez, M. J. Cristóbal, M. P. Hierro, G. Arnau and J. Botella, Oxid. Met., 54 (2000) 87. 4. Y. C. Lu and M.B. Ives, Corros. Sci., 34 (1993) 1773. 5. Y. C. Lu, M. B. Ives, Corros. Sci., 37 (1) (1995) 145. 6. S. Virtanen, M.B. Ives, G.I. Sproule, P. Schmuki and M.J. Graham, Corros. Sci., 39 (1011) (1997) 1987. 7. C. M. Abreu, M. J. Cristóbal , R. Losada, X. R. Nóvoa, G. Pena, M. C. Pérez, Electrochimica Acta 49 (2004) 3049-3056 8. C. Giacovazzo, Fundamentals of Crystallography, 3rd edn., Oxford University Press, New York, (1995) 442. 9. C. D. Wagner, Practical Surface Analysis, Vol. 1., 2nd Edn, J. Wiley and Sons, London, 1990. 10. J.Z. Shyu, K. Otto, W.L.H. Watikins, G.W Graham, R.K. Belitz, H.S. Gandhi, J. Catal., 114 (1988) 23.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Use of Alloy 22 as Long-Term Radioactive Waste Containment Material Osvaldo Pensado, Roberto Pabalan, Darrell Dunn, and Kuang-Tsan Chiang Center for Nuclear Waste Regulatory Analyses (CNWRA), 6220 Culebra Rd. San Antonio, TX, 78238-5166, USA, [email protected]

Abstract - An approach to estimate upper probability bounds for localized corrosion initiation in Alloy 22 as function of water chemistries relevant to a potential geologic repository for high-level radioactive waste is discussed. The approach is used to compare the effect of fabrication processes (e.g., welding) accounting for various chemical components that promote or inhibit localized corrosion. Keywords: Ni-Cr-Mo alloy, passive dissolution, localized corrosion, radioactive waste, risk assessment

1. Introduction The waste package design for the potential geologic repository for high-level radioactive waste at Yucca Mountain in the United States includes an outer container made of Alloy 22 (Ni-22Cr-13.5Mo-3W-4Fe), for corrosion resistance, and an inner container made of Type 316 nuclear grade stainless steel, for structural support. The evaluation of Alloy 22 oxide stability in the repository setting is unique because of the long time frames involved and its dependence on changing environments. No natural processes, aside from the establishment of adverse environmental conditions, are envisioned capable of inducing passivity breakdown for extended periods. Therefore, Alloy 22 passivity is dependent on the quantity and chemistry of the waters that could contact waste packages. The repository design includes self supported titanium shields to protect waste packages from seeping water. For reasons discussed elsewhere [1, 2], passivity breakdown (in the form of crevice corrosion) could only occur when (i) seeping water directly contacts waste packages and (ii) dynamic evaporation causes the development of brines. A simplified model to

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develop upper bounds for the probability of passivity breakdown, in the form of crevice corrosion, is discussed accounting for distributions of brine compositions after water evaporation, and for the different susceptibility of mill-annealed and weld zones. 2. Approach A mechanistic equation for the computation of the corrosion potential, Ecorr, as a function of pH, temperature, and the anodic current density associated with the passive dissolution of Alloy 22 has been developed elsewhere [2]. The equation is consistent with experimental values and trends (e.g., the corrosion potential decreases with increasing values of pH and temperature). Studies indicate that localized corrosion is initiated when the corrosion potential is above a critical potential; and that fabrication processes such as welding can cause Alloy 22 to become more susceptible to localized corrosion [3]. The crevice corrosion repassivation potential, Ercrev, can define the critical potential. Ercrev is a function of temperature and chloride concentration [3]. Oxyanions such as nitrate, carbonate, and sulfate are effective inhibitors of localized corrosion in Alloy 22 [4]. The effect of nitrate has been modeled as an increase in the repassivation potential as a function of the oxyanion to chloride concentration ratio [2]. This functional dependence has been extended to account for the possible simultaneous presence of multiple inhibiting oxyanions [5]. To estimate the probability of localized corrosion of Alloy 22 under the dynamic evaporation conditions stated in the introduction, values from distribution functions relevant to the repository setting of pH and anion concentrations (chloride, nitrate, carbonate, and sulfate) were randomly sampled to determine values of Ecorr and Ercrev (the latter for both mill-annealed and weld areas). The conditional probability of localized corrosion, Pc, was estimated by counting the fraction of times Ecorr exceeds Ercrev. Assuming that Pw is the probability of brine formation on the waste package, the localized corrosion probability can be estimated as Pw´Pc. The approach to estimate distributions of water compositions is described below. Thermodynamic simulations of the chemical evolution of in-drift waters resulting from evaporation of seepage waters were conducted. The simulations allowed determination of the types of brines that may form and the ranges of brine chemistry that may exist in emplacement drifts. The thermodynamic calculations were supplemented by an alternative approach based on the concept of chemical divide developed by Hardie and Eugster [6]. In the chemical divide concept, the chemical types of brines and salt minerals that form upon evaporation of natural waters are determined by early precipitation of insoluble minerals. Natural waters at Yucca Mountain are considered to evolve into three types of brines upon evaporation: (i) calcium-chloride, (ii) neutral, and (iii) alkaline. The evaporation simulations were conducted using thermodynamic

Use of Alloy 22 as Long-Term Radioactive Waste Containment Material

55

equilibrium software [5]. Twenty-nine water compositions were used as input to the evaporation simulation, selected to represent the broad composition range of the more than 150 samples of Yucca Mountain unsaturated zone porewater reported by Yang et al. [7, 8, 9]. The evaporation simulations were done at 110 °C [230 °F] and at constant pressures of 0.85 or 1 bar [12.3 or 14.5 psi]. The likelihood of localized corrosion increases with increasing temperature. At temperatures higher than 110 °C [230 °F], water from seepage may not contact engineered barrier materials because of thermal mobilization of water away from drifts. Thus, consideration of this reference temperature is appropriate to estimate upper probability bounds for localized corrosion. The thermodynamic simulation of seepage water evaporation used only a subset of the unsaturated zone porewater chemistry data reported by Yang et al. [7, 8, 9]. To provide a basis for estimating the frequency of occurrence of the three brine types and their respective chemical characteristics, the full set of Yang et al. [7, 8, 9] data on unsaturated zone porewater compositions was used with the chemical divide concept of Hardie and Eugster [6]. From 156 compositions, 8, 24, and 68 percent resulted into calcium chloride, neutral, and alkaline type brines, respectively. Distribution functions were numerically derived by interpolated sampling of empirical distribution functions for each brine type, and combining the three populations into a single population, preserving the brine type frequency (i.e., in a population with 10,000 stochastic vectors {pH, [Cl-], [NO3-], [CO32-]+[HCO3-], [SO42-]}, 800 were calcium chloride; 2,400, neutral; and 6,800, alkaline brine type of vectors). The rank-correlation matrix of the 10,000 vector population was computed. A dominant correlation between pH and the carbonate concentration (equal to 0.9) indicated that highest carbonate concentrations were exhibited in alkaline brines. A negative correlation of -0.8 between chloride and pH mirrored the fact that acidic calcium chloride brines are highly concentrated in chloride. 3. Results and Discussion Sampling from empirical distribution functions (10,000 samples) for the anions and pH while preserving the relevant rank correlations, and computing the corresponding values for the corrosion potential, Ecorr, and critical potentials, Ercrev, for each sample, the conditional distribution functions in Figure 1 were derived. Figure 1 (b) indicates that in 26 percent of the samples the corrosion potential exceeded the critical potential in thermally aged material (simulating welds), and only in 3 percent of the samples for mill-annealed material. It is therefore interpreted that Pc equals 0.26 along weld areas, and 0.03 on the rest of the body. The probability for Alloy 22 to exhibit localized corrosion equals Pw´Pc, where Pw is the probability of brine formation on the waste package. If (i) drip shields perform their water-diverting function, or (ii) fail at later times when

56

O. Pensado et al.

development of brines is not longer possible (i.e., when the seepage rate exceeds the evaporation rate), or (iii) there is no water seepage during the thermal pulse, Pw would equal zero. In any of these three cases, localized corrosion would not be exhibited and passive dissolution would be the dominant corrosion mode (sustained by moisture in the environment or condensed water). A more detailed assessment is needed to estimate the probability Pw.

(a)

(b)

1

1 Thermally aged 0.9

0.8

E corr

0.8 Conditional Complementary CDF

Conditional Cumulative Distribution Function

0.9

0.7

E rcrev (thermally aged)

0.6 0.5 0.4 0.3

0.7 0.6 Mill annealed

0.4 0.3

0.1

E rcrev (mill annealed)

3%

0

0 -200

26%

0.2

0.2 0.1

-400

0.5

0

200 400 Potential, mVSHE

600

800

1000

-1500

-1000

-500 0 500 E corr - E rcrev , mV

1000

1500

Figure 1. (a) Conditional distribution functions for the corrosion potential and critical potential for localized corrosion resulting from the adopted stochastic sampling approach. (b) Conditional complementary cumulative distribution function for the difference between the corrosion potential and repassivation potential.

In addition to the dependence on Pw, other caveats to this simplified analysis are offered. There exists evidence that localized corrosion propagation could stifle, even in strongly oxidizing conditions [5]. Stifling may limit the extent of crevice attack. Localized corrosion experiments are performed in bulk solutions. If brines develop in the repository, localized corrosion may be activated under a limited supply of water. The stifling mechanism is likely to be more effective in solutions of limited volume due to local saturation with corrosion products. A second caveat is that the kind of localized corrosion that could affect Alloy 22 is crevice corrosion. Formation of crevices on waste packages may be limited to the contact area with the waste package support system. Additional crevice sites could form from direct contact between the waste package and drip shield caused by drift degradation and rockfall. Therefore, the probability for waste packages to be affected by localized corrosion is also conditional on the probability of formation of crevice sites. If crevice corrosion affects a waste package, the corroded area will be limited and the degraded waste package could still offer protection against radionuclide release. 4. Conclusions The passive dissolution can only be affected by processes inducing compositional changes to the alloy at some significant depth. No processes are

Use of Alloy 22 as Long-Term Radioactive Waste Containment Material

57

currently envisioned that could accomplish such a change. The persistence of Alloy 22 passivity in the repository setting is mainly dependent on the quantity and chemistry of water that could contact this alloy. An approach to estimate the conditional probability for waste packages to be affected by localized corrosion was discussed accounting for possible brine chemistries developed after water evaporation. The approach did not consider protection from additional engineered and natural barriers, localized corrosion stifling, nor crevice site requirements. The analysis accounted for the effect of various chemical components that promote (chloride) or inhibit (nitrate, carbonate, and sulfate) localized corrosion in Alloy 22, and for fabrication processes. Notice of Disclaimer This paper was prepared to document work performed by the CNWRA for the Nuclear Regulatory Commission (NRC) under Contract No. NRC-02-02-012. The activities reported here were performed on behalf of the NRC Office of Nuclear Material Safety and Safeguards, Division of High Level Waste Repository Safety. This paper is an independent product of the CNWRA and does not necessarily reflect the view or regulatory position of the NRC. References [1] L. Browning, R. Fedors, L. Yang, O. Pensado, R. Pabalan, C. Manepally, and B. Leslie, Mat. Res. Soc. Symp. Proc. 824 (2004) 417-424. [2] D.S. Dunn, O. Pensado, and G.A. Cragnolino, “Performance Assessment of Alloy 22 as a Waste Package Outer Barrier” in Proceedings of the CORROSION 2005 Conference, paper 588. NACE International, Houston, Texas (2005). [3] D.S. Dunn, L. Yang, Y.M. Pan, and G.A. Cragnolino, “Localized Corrosion Susceptibility of Alloy 22” in Proceedings of the CORROSION 2003 Conference, paper 697. NACE International, Houston, Texas (2003). [4] D.S. Dunn, L. Yang, C. Wu, and G.A. Cragnolino, Mat. Res. Soc. Symp. Proc. 824 (2004) 3338. [5] D.S. Dunn, O. Pensado, Y.-M. Pan, R.T. Pabalan, L. Yang, X. He, and K.T. Chiang. “Passive and Localized Corrosion of Alloy 22—Modeling and Experiments.” CNWRA 2005-02, San Antonio, Texas (2005). [6] L.A. Hardie and H.P. Eugster, Mineralogical Society of America, Special Paper No. 3 (1970) 273-290. [7] I.C. Yang, Z.E. Peterman, and K.M. Scofield, Journal of Contaminant Hydrology, 1878 (2003) 1-20. [8] I.C. Yang, P. Yu, G.W. Rattray, J.S. Ferarese, and J.N. Ryan, “U.S. Geological Survey WaterResources Investigations Report 98-4132,” Denver, Colorado (1998). [9] I.C. Yang, G.W. Rattray, and Y. Pei, “U.S. Geological Survey Water-Resources Investigations Report 96-4058,” Denver, Colorado (1996).

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

59

Effect of temperature and melt composition on the passivity of a Ni-10%Cr alloy in a molten electrolyte Martin Bojinov, Petia Gencheva, Tzvety Tzvetkoff Department of Physical Chemistry, University of Chemical Technology and Metallurgy, 1756 Sofia, Bulgaria, e-mail [email protected] Abstract - The effect of temperature and carbonate addition on the mechanism of growth and the composition of anodic oxide films on a Ni–10 wt.% Cr alloy in molten NaOH was studied. Impedance measurements were performed in the passive potential region to investigate the conductivity mechanism of the oxide films as depending on temperature and melt composition. The surface composition of the oxides has been estimated by X-ray photoelectron spectroscopy and their in-depth composition by Auger electron spectroscopy. The main passivation product on the Ni surface was found to be a non-stoichiometric NiO doped with Cr close to the metal/film, and Na or Li at the film/melt interface. The transport of nickel cation vacancies through the barrier sublayer was demonstrated to be the rate-limiting step of the overall process in the passive region. Keywords: Ni-Cr alloy; Molten electrolyte; Anodic oxide film; Electrochemical Impedance Spectroscopy; Surface analysis; Solid-state ion transport; Electronic defect; Kinetic model

1. Introduction In a previous paper [1], we presented voltammetric and impedance results on the anodic behaviour of two nickel-based alloys in a NaOH melt at 470 °C coupled to the analysis of the oxides by X-ray photoelectron spectroscopy (XPS) and Auger electron spectroscopy (AES). The conduction mechanism in the oxide was modelled on the basis of previous approaches for films formed in high-temperature gas systems [2,3] and high-temperature aqueous electrolytes [4-6]. In the present paper, electrochemical and surface analytical results on Ni10%Cr in a hydroxide-based melt are discussed in terms of the effects of temperature and carbonate addition on the thickness, composition, electric and electrochemical properties of the oxides. An attempt is made to summarise the electrochemical and surface analytical results into a qualitative picture of the conduction mechanism of the anodic passive films.

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2. Experimental A Ni–10%Cr alloy (9.96 wt% Cr, 0.22 wt% Fe, 0.50 wt% Si, bal. Ni) was used as working electrode, pure Ni served as a counter, whereas pure Mo (E= -0.9 ± 0.02 V vs. Pt) was employed as a pseudo-reference electrode. Experiments were performed in NaOH and NaOH-Li2CO3 (1:1) melts deaerated with N2 (99.999%) for 2 h prior to and during the experiments. An Autolab PGSTAT 30 with a FRA2 module (EcoChemie, the Netherlands) was used. Voltammetric curves were registered at a sweep rate of 1 mVs-1. EIS measurements were performed at 0.1 V intervals from the corrosion potential up to 1.0 V in a frequency range 0.02 Hz - 30 kHz and an ac amplitude 15 mV(rms). XPS analyses were carried out using an ESCALAB Mk II (VG Scientific) with a base pressure of ~10-7 Pa, a Mg Ka (1253.6 eV) source and a pass energy of 20 eV. The AES spectra were recorded with a Riber LAS 3000 in ultrahigh vacuum conditions (2 ´10 –8 Pa). The conversion of the sputtering time to depth has been performed vs. a SiO2 standard (sputtering rate 4.4 nm min-1). 3. Results and Discussion

a b

Fig.1. Current density vs. potential curves for Ni–10 % Cr in molten NaOH at different temperatures (a) and in the two studied melts at 470 °C ( b). Sweep rate 1 mV s-1.

The effect of temperature on the current vs. potential curves is illustrated in Fig.1a. The active-to-passive transition (0.2–0.5 V), the passive range (0.5–0.9 V) and the transpassive range related to the decomposition of the melt are distinguished. The passive current increases slowly with potential indicating a solid-state transformation of the oxide. The width of the passive region is the largest at 370 °C, and current densities in it are lower at 370 and 420 °C than at 470 °C (Fig.1a). The effect of carbonate addition is illustrated in Fig.1b in which the results obtained in NaOH and NaOH–Li2CO3 (1:1) at 470 °C are compared. The passive currents are smaller and the onset of the transpassive range is located at a higher potential in the mixed melt, which could be explained by the smaller tendency of dissolution of Cr as chromate in the more acidic melt [7,8].

Effect of temperature and melt composition on the passivity of…

61

b

a

c

Fig. 2. Impedance spectra of Ni-10%Cr in molten NaOH at 370 °C (a) and 470 °C (b) and the corresponding spectra in molten NaOH–Li2CO3 (1:1) at 470 °C (c). Impedance magnitude plot (left) and phase angle plot (right). Points – experimental values, solid lines – best-fit calculation.

The impedance magnitudes at low frequencies (Fig. 2a-c), decrease with increasing potential in the passive range in accordance to the increase in the currents (Fig.1). The characteristic frequency of the high-frequency time constant (at 0.1–1 kHz, associated with the electronic properties of the oxide [1]) also decreases with increasing potential, indicating p-type semiconductor properties. The low-frequency time constant (at 0.05–1 Hz) represents the solidstate transport in the barrier part of the oxide [1]. The characteristic frequency of this time constant increases with temperature. The impedance magnitude at low frequencies is higher in the carbonate-containing melt (Fig. 2c) in accordance with the observed trends in the passive current densities. Fig. 3a-b shows detailed XPS spectra of Ni2p, Cr2p and O1s at the surface of the oxides formed at 0.5 V in molten NaOH at 370 and 420 °C. Both Ni2p and O1s peaks are split to low- and high-energy contributions, i.e. the composition of the oxide corresponds to non-stoichiometric and/or partly hydrated NiO containing a certain amount of Cr3+. Fig. 4a-c shows the Auger depth profiles measured after 1 h oxidation of the alloy at 0.5 V in molten NaOH at 370 °C (a) and 470 °C (b), as well as in NaOH–Li2CO3 (c) at 470 °C. The oxide thickness, estimated by the depth at which the intensity of the O signal has dropped to half of its maximum value, increases with temperature. The thickness of the film in the carbonatecontaining melt is higher than that in NaOH (c), the values for Ni–10%Cr in

M. Bojinov et al.

62

NaOH–Li2CO3 electrolyte being comparable to those for Ni–3%Al-2%Mn in NaOH [1]. The depth profiles indicate a bilayer structure of the oxide, the outer part being a mixed Ni-Na oxide and the inner part a Cr-substituted NiO. 8000 533.0

370 °C, 0.5 V

7000

6000

Intensity / a.u.

861.7

3000 2000 1000 864

860

856

150

100

0 595

852

585

580

575

570

8000

b

862.0

3000 2000 1000 0 868

864

860

856

852

Binding Energy / eV

300 587.0

200

100

590

585

580

575

532

420 °C, 0.5 V 533.2

6000

530

528

530.0

531.4

4000

2000

0

0 595

534

Binding Energy / eV

576.5

420 °C, 0.5 V

856.2

Intensity / a.u.

Intensity / a.u.

4000

0

590

400

854.7

5000

3000

Binding Energy / eV

420 °C, 0.5 V

6000

4000

1000

Binding Energy / eV 7000

5000

2000

50

Intensity / a.u.

0 868

531.4

576.6

586.7

a

4000

6000

370 °C, 0.5 V

200

Intensity / a.u.

5000

Intensity / a.u.

530.0

854.4 370 °C, 0.5 V 855.9

534

570

532

530

528

Binding Energy / eV

Binding Energy / eV

Fig. 3. Detailed XPS of Ni, Cr, and O at the Ni-10%Cr surface oxidised for 1 h at 370 and 420 °C in molten NaOH at 0.5 V.

For the interpretation of the impedance data we have used a transfer function based on the MCM [5,6]: 1 ⎧⎡ ⎛ ⎪⎪ ⎢ p ⎜ 1 + jwr d ee 0 e p Z = Rohm + ⎨ ⎢ ln ⎜ ⎪ ⎢ jw C ⎜ 1 + jwr d ee 0 ⎝ ⎩⎪ ⎣

-1

-1 ⎫ ⎞⎤ -1 ⎡ tanh( jwt )0.5 ⎤ ⎪⎪ ⎟⎥ ⎟ ⎥ + ⎢ Rt + s ( jw )0.5 ⎥ ⎬ ⎣ ⎦ ⎪ ⎟⎥ ⎠⎦ ⎭⎪

(1)

where Rohm is the electrolyte resistance, Rt the charge transfer resistance at the film / melt interface, the first member in the brackets is the capacitance of a semiconductor layer with spatially and energetically variable defect density, whereas the second is the faradaic impedance of generation, transport and consumption of ionic point defects. For the case in which cation vacancies act as predominant current carriers (typical for nickel oxide [9]), the following expressions for the parameters in (1) have been derived [5,6]: p –1 = [const + 2(F/RT) (1-a)E ] rd = {(F De / RT) [kg exp(bg aE) / 2

0

(2) kc0]

}

-1

(3)

Effect of temperature and melt composition on the passivity of…

63

s = RT / {F2 (32DV)1/2 [kg0 exp(bg aE) / kc0](1-a)

(4)

t = [2(FE /RT) DV]

(5)

2

-1

a

0

20

40

60

80

100

Depth / nm

120

0

50

100

150

Auger intensity / a.u.

200

300

350

c

B C Na Cr

100

250

Depth / nm

NaOH-Li2 CO3 470 °C, 0.5 V

0

200

Auger intensity / a.u.

O Ni Na Cr

300

400

500

Auger intensity / a.u.

Auger intensity / a.u.

370 °C, 0.5 V

b

Auger intensity / a.u.

Auger intensity / a.u.

470 °C, 0.5 V O Ni Na Cr

600

Depth / nm

Fig. 4. AES depth profiles for the oxide on Ni–10% Cr for 1 h at 0.5 V in molten NaOH at 370 °C (a), 470 °C(b) and NaOH- Li2CO3 at 470 °C (c).

In these expressions, C is the depletion layer capacitance, E is the field strength, a is the polarisability of the film / melt interface, De is the hole diffusivity, DV the diffusivity of cation vacancies, kg0 and kc0 are the rate constants of generation / consumption of cation vacancies, bg is an exponential factor and e the dielectric constant. Fitted spectra are traced in Fig. 2 and demonstrate the ability of the transfer function described by eqs. (1)–(5) to account for the experimental data. The dependences of the parameters C, p, rd e and s on potential for all the studied are shown in Fig. 5. The depletion layer capacitance decreases with increasing temperature, i.e., the number of defects decreases. A further decrease in C is observed in molten NaOH–Li2CO3 at 470 °C, which points to a less defective layer formed in that melt. The parameter p-1 increases linearly with potential in agreement with eq. (2). The polarisability of the film/melt solution interface increases with temperature, i.e. the potential drop within the film correspondingly decreases. The quasi-exponential decrease of he product rde and the Warburg constant s with increasing potential and the decrease of these two parameters with temperature can be understood in terms of eqs. (3)–(4).

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3.5

26

3.0

25

a

370 °C 420 °C 470 °C 470 °C 2.0 NaOH-Li2CO3

21

1.5

20

1.0

19 0.5

p-1

22

1

1

b 0.1

370 °C 420 °C 470 °C 470 °C NaOH-Li2 CO3

18 0.4

0.5

0.6

0.7

0.8

E / V vs. Mo

0.9

1.0

17

0.01

0.4

0.5

0.6

0.7

0.8

s / kWcm s

-1

23

2 -0.5

24

2.5

0.0

10

10

rde / MW cm

3

-2

10 C / (m Fcm )

-2 -2

The larger values of rd e and s in the NaOH–Li2CO3 melt can be due to lower kg0 , De and DV values in the carbonate-containing melt.

0.1 0.9

1.0

E / V vs. Mo

Fig. 5. Parameters of the transfer function used to fit the impedance data as depending on the potential and the composition of the molten electrolyte.(a) C-2 and p-1 vs. E, (b) rd e and s vs. E.

4. Conclusion In the present paper, an attempt is made to characterise the surface film formed on a Ni-10%Cr alloys in molten NaOH and NaOH-Li2CO3 in terms of the relationship between surface and in-depth composition and conduction mechanism. To rationalise the correlation between composition and conduction mechanism, a quantitative model approach based on the MCM is used for the interpretation of the experimental impedance spectra. The picture of the anodic film that emerges from the present study is that of a non-stoichiometric layer with appreciable ionic and electronic conductivity, the main current carriers being identified as cation vacancies and electron holes. It can be also tentatively concluded that the surface film is composed of two layers, an inner barrier-like layer which grows by a solid-state mechanism and an outer layer growing by a dissolution-precipitation mechanism. References 1. 2. 3. 4. 5. 6.

P. Gencheva, Tz. Tzvetkoff and M. Bojinov, Appl. Surf. Sci. 241 (2005) 458. W. S. Epling, G.B. Hoflund, Thin Solid Films 292 (1997) 236. R. Peraldi, D. Monceau, B. Pieraggi, Oxid. Met. 58 (2002) 249. T.M. Angeliu, P.L. Andresen, M.L. Pollick, Corros. 53 (1997) 114. M. Bojinov, P. Kinnunen, G. Sundholm, Corros. 59 (2003) 91. B. Beverskog, M. Bojinov, P. Kinnunen, T. Laitinen, K. Mäkelä, T. Saario, Corros. Sci. 44 (2002) 1923. 7. T. Ishitsuka, K. Nose, Mater. Corros. 51 (2000) 177. 8. R.A. Rapp, Mater. Sci. Eng. 87 (1987) 319. 9. T. Bak, J. Nowotny, C.C. Sorrell, Key Eng. Mater. 125 (1997) 1.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

65

Spontaneous Passivation of Amorphous Bulk NiCr-Mo-Ta-Nb-P Alloys in Concentrated HCl Koji Hashimotoa*, Hiroyuki Shinomiyaa, Akito Nakazawaa and Katsuhiko Asamib a

Tohoku Institute of Technology, Sendai, 982-8588 Japan; Phone: +81-22-229-1151; Fax: +81-247-8150; E-mail: [email protected] b Institute for Materials Research, Tohoku University , Sendai, 981-8577 Japan

Abstract - Amorphous alloy rods of 1 mm diameter such as Ni-(1-3)at%Cr-(1-8)at%Mo22at%Ta-14at%Nb-4at%P were prepared by copper-mold casting. Most of them were spontaneously passive and immune to corrosion in 6 and 12 M HCl. Their air-formed films were stable in these hydrochloric acids without significant change in the composition after immersion for 168 h. The films seemed to be composed of the outer triple oxyhydroxide of Cr3+, Ta5+ and Nb5+ and the inner MoO2. Keywords : Spontaneous passivation, amorphous bulk Ni-Cr-Mo-Ta-Nb-P alloys, concentrated HCl, bi-layer film

1. Introduction Spontaneous passivation is prerequisite for the use of metallic materials. Various amorphous alloys have high corrosion resistance due to spontaneous passivation in aggressive environments. The use of corrosion-resistant amorphous alloys is, however, restricted by their limited thickness, and the formation of amorphous bulk alloys is eagerly requested for application to extremely aggressive environments such as concentrated HCl where no conventional crystalline alloys can be used. It was found in the mid 1990s that some amorphous alloys became the supercooled liquid state just below the crystallization temperature, where the amorphous alloy powders were able to be consolidated in the form of amorphous bulk alloys [1,2]. These procedures are, however, too laborious for trial to form amorphous bulk alloys. More simple copper-mold casting sometimes gave rise to the formation of amorphous alloy rods with a diameter of 1 mm or more [3], although necessary are not only rods but also sheets, tube, etc which can be formed by consolidation of amorphous alloy powder. Since sputter-deposited binary amorphous Cr-Ta alloys have the highest corrosion resistance among known metallic materials in concentrated hydrochloric acids [4], we tried to prepare amorphous bulk alloys containing chromium and tantalum and

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succeeded using a copper-mold casting in preparing amorphous bulk Ni-20at%(P-B) alloys with chromium, tantalum and/or molybdenum which were spontaneously passive in concentrated hydrochloric acids [5]. When their temperature intervals of the supercooled liquid state were wider than 50 K at the heating rate of 20 K min-1 and when amorphous alloy rods of 1 mm diameter were formed by copper-mold casting, hot rolling consolidation of their amorphous alloy powders in the supercooled liquid state resulted in the formation of amorphous bulk alloy sheets of a few mm thickness [6,7] . For other amorphous alloy series with high crystallization temperatures, Ni40.5at% Nb and Ni-36.5at% Ta alloys are the eutectic mixtures and become amorphous by rapid quenching from the liquids state. The amorphous Ni-Ta alloys are particularly attractive because of their high crystallization temperatures exceeding 980 K. For further high corrosion resistance and for the formation of amorphous bulk alloys the addition of chromium and molybdenum to Ni-Ta-Nb alloys is required. The addition of small amount of phosphorus will also be effective for the formation of amorphous bulk alloys. In the present work an attempt was made to tailor corrosion-resistant amorphous bulk Ni-Cr-Mo-Ta-Nb-P alloys by copper-mold casting.

2. Experimental Procedures Crystalline alloy ingots were prepared by argon arc melting of commercial electrolytic nickel, electrolytic chromium, 99.9% pure molybdenum, tantalum and niobium, and nickel phosphide (Ni-14.17wt%P). These ingots were used for coppermold casting to form rod-shaped alloys of 1 mm diameter. The structure of specimens prepared was identified by X-ray diffraction with Cu Ka radiation. Prior to immersion tests and electrochemical measurements the surface of specimens was polished in cyclohexane with silicon carbide paper up to No. 1500, degreased in acetone and dried in air. Immersion tests and electrochemical measurements were carried out in 6 and 9Mo-1Cr 8Mo-1Cr

Diffraction Intensity / Arbitraey Unit

Diffraction Intensity / Arbitrary Unit

1Mo-4Cr 1Mo-3Cr 2Mo-2Cr 1Mo-2Cr 1Mo-1Cr

1Mo-0Cr 30

40 50 60 70 2q / Degree (Cu Ka)

80

7Mo-1Cr 6Mo-1Cr 5Mo-1Cr 4Mo-1Cr 3Mo-1Cr 2Mo-1Cr 0Mo-1Cr 30

40 50 60 70 2q / Degree (Cu Ka )

80

Fig. 1 XRD patterns of Ni-Mo-Cr-22Ta-14Nb-4P alloy rods of 1 mm diameter.

12 M HCl solutions open to air at 303 K. Corrosion rates were estimated from the

Spontaneous Passivation of Amorphous Bulk Ni-Cr-Mo-Ta-Nb-P Alloys...

67

weight loss measured by a microbalance after immersion for 168 h. The length of specimen used for the immersion test was about 16 mm and their densities were about 10.2 Mgm-3. Thus the weight loss of 1 x 10-6 g by immersion for 168 h corresponded to the corrosion rate of about 1 x 10-4 mmy-1 which was determined as the detectable limit. The corrosion loss was measured 2-4 times for each specimen. Potentiodynamic polarization curves were measured with a potential sweep rate of 0.5 mVs-1. The potential was swept from the open-circuit potential to anodic and cathodic directions using different specimens after immersion in hydrochloric acids for about 10 min. The surfaces before and after immersion were analyzed by X-ray photoelectron spectroscopy using Mg Ka excitation. 3. Results and Discussion Fig. 1 shows X-ray diffraction patterns. The proper addition of both chromium and molybdenum leads to the formation of the amorphous single phase showing halo patterns. Unless both molybdenum and chromium are added amorphous single phase alloys cannot be formed, but the alloys consist of minor crystalline phases in the amorphous matrix. The excess addition of chromium such as 4 at% to Ni-Cr-1Mo22Ta-14Nb-4P alloy results in precipitation of nanocrystalline phases in the amorphous matrix. For Ni-1Cr-Mo-22Ta-14Nb-4P alloys, amorphous single phase alloys are formed in a considerably wide concentration range of molybdenum from 1 to 8 at%. Table 1 shows the corrosion rates measured in 6 and 12 M HCl for various Ni-

Table 1 Corrosion rates of copper-mold cast Ni-Mo-Cr-Ta-Nb-P alloys of 1 mm diameter in HCl solutions at 303 K. Corrosion Rate / mmy-1 Alloy 6 M HCl 12 M HCl Ni-1Mo-1Cr-20Ta-18Nb-4P < 1 x 10-4 < 1 x 10-4 -4 Ni-2Mo-2Cr-22Ta-14Nb-4P < 1 x 10 < 1 x 10-4 -4 Ni-1Mo-1Cr-22Ta-16Nb-4P < 1 x 10 < 1 x 10-4 -4 Ni-1Mo-1Cr-22Ta-14Nb-4P < 1 x 10 3.5 x 10-4 -4 Ni-1Mo-2Cr-22Ta-14Nb-4P < 1 x 10 < 1 x 10-4 -4 Ni-1Mo-3Cr-22Ta-14Nb-4P < 1 x 10 < 1 x 10-4 Ni-0Mo-1Cr-22Ta-14Nb-4P 4.7 x 10-4 < 1 x 10-4 (including minor crystalline phases) Ni-2Mo-1Cr-22Ta-14Nb-4P < 1 x 10-4 2.0 x 10-4 Ni-6Mo-1Cr-22Ta-14Nb-4P < 1 x 10-4 3.0 x 10-4 Ni-9Mo-1Cr-22Ta-14Nb-4P 6.9 x 10-4 < 1 x 10-4 (including minor crystalline phases) Ni-1Mo-4Cr-22Ta-14Nb-7P < 1 x 10-4 < 1 x 10-4 -4 Ni-2Mo-3Cr-22Ta-14Nb-4P < 1 x 10 < 1 x 10-4 -4 Ni-1Mo-1Cr-19Ta-19Nb-3P < 1 x 10 1.7 x 10-4 -4 Ni-2Mo-2Cr-18Ta-18Nb-3P < 1 x 10 1.8 x 10-4

K. Hashimoto et al.

68

Cr-Mo-Ta-Nb-P alloy rods of 1 mm diameter. The addition of chromium, molybdenum and sufficient amount of tantalum such as 20 at% or more to amorphous single phase alloys is necessary for immunity to corrosion in 6 and 12 M HCl. In particular, an increase in chromium content is effective in enhancing the corrosion resistance as seen from the examples that the Ni-1Mo-2Cr-22Ta-14Nb-4P alloy is stable while the Ni1Mo-1Cr-22Ta-14Nb-4P alloy suffers detectable corrosion loss in 12 M HCl. If the tantalum contents are not sufficiently high such as 18-19 at%, the alloys suffer corrosion in aggressive 12 M HCl even if the alloys consist of an amorphous single phase. In this manner, the presence of chromium, molybdenum and proper amounts of tantalum in addition to the formation of amorphous single phase is prerequisite in providing the sufficiently high corrosion resistance in aggressive concentrated hydrochloric acids. The excess addition of molybdenum is detrimental and the alloy suffers detectable corrosion loss in 12 M HCl. Fig. 2 shows potentiodynamic polarization curves of representative alloys measured in 12 M HCl. At potentials lower than the open circuit potential of about 0.1 V, the cathodic current for reduction of oxygen and proton appears. Because of high cathodic current and because of low anodic current these alloys are spontaneously passive in the aggressive solution. Even for spontaneously passive amorphous alloys the excess addition of molybdenum is apt to increase the anodic current density since the open circuit potential is higher than the potential for the transpassive dissolution of tetravalent molybdenum as hexavalent molybdenum unless the tetravalent molybdenum is protected by the outer layer of the passive film. The presence of nanocrystalline phases also increases the anodic current density. Fig. 3 shows an example of the analytical results for the surface film by X-ray photoelectron spectroscopy. Air exposure of the alloy forms the film in which the 101

Current Density / Am-2

Ni-9Mo-1Cr-22Ta-14Nb-4P with nanocrystalline phases 100

Ni-0Mo-1Cr-22Ta-14Nb-4P with nanocrystalline phases

10-1 10-2 Ni-6Mo-1Cr-22Ta-14Nb-4P 10-3

Ni-2Mo-1Cr-22Ta-14Nb-4P

~

-0.5

0

0.5 1 Potential / V vs. Ag/AgCl

1.5

~ ~

2

Fig. 2 Potentiodynamic polarization curves of Ni-Mo-Cr-22Ta-14Nb-4P alloy rods of 1 mm diameter measured in 12 M HCl.

Spontaneous Passivation of Amorphous Bulk Ni-Cr-Mo-Ta-Nb-P Alloys...

69

corrosion-resistant elements are significantly concentrated. In particular, the chromium content exceeds 20% of cations in spite of 1 or 2 at% in the alloys. Because of the high stability of the air-formed film, spontaneous passivation in concentrated hydrochloric acids occurs without appreciable change in the film composition. A high concentration of chromium together with tantalum, niobium and molybdenum is responsible for immunity to corrosion in 12 M HCl due to spontaneous passivation. According to the study of binary Cr-Ta and Cr-Nb alloys, the passive film was not the heterogeneous oxyhydroxide mixtures of chromium and tantalum or niobium but homogenous double oxyhydroxide of chromium and tantalum or niobium [8]. Thus chromium, tantalum and niobium in the surface films on the Ni-Cr-Mo-Ta-Nb-P alloys seemed to form triple oxyhydroxide of Cr3+, Ta5+ and Nb5+. The presence of the triple oxyhydroxide with significantly high stability in concentrated hydrochloric acids is partly responsible for spontaneous passivation in addition to the beneficial role of molybdenum, although it was difficult to obtain the precise analytical results for molybdenum because the Mo 3 d spectra were superimposed upon the Ta 4d signal. According to the analysis of the spontaneously passive films on Fe-Cr-Mo-P-C [9], Mo-Cr [10] and Mo-valve metal [11], in which the molybdenum content was widely changed, mostly tetravalent molybdenum existed in the inner part of the passive film under the protection by the outer film. On the basis of these results the spontaneously passive films on the amorphous alloy rods seemed to consist of the bi-layer structure of the outer Cr1-x-yTaxNbyOzOH3+2(x+y-z) and the inner MoO2. The both layer should act as the barrier for diffusion of cations and oxygen. In particular, the inner MoO2 layer was protected by the outer layer from oxidation to Mo6+.

Cationic Fraction

0.8

Ni-1Mo-2Cr-22Ta-14Mo-4P

0.8

Ta

0.6

0.4

0.2

0

1

Cationic Fraction

1

Nb

Ni-2Mo-1Cr-22Ta-14Mo-4P

Ta

0.6

0.4 Nb 0.2 Cr

Cr Mo P

Before immersion

0

After After immersion immersion in 6 M HCl in 12 M HCl

Mo P

Before immersion

After After immersion immersion in 6 M HCl in 12 M HCl

Fig. 3 Cationic fractions in surface films on Ni-Mo-Cr-22Ta-14Nb-4P alloy rods of 1 mm diameter before and after immersion in 6 and 12 M HCl for 168 h.

70

K. Hashimoto et al.

4. Conclusions Some Ni-Cr-Mo-Ta-Nb-P alloys prepared by copper-mold casting consisted of the amorphous single phase in a limited composition range. The addition of both chromium and molybdenum to Ni-Ta-Nb-P alloys was necessary for the formation of amorphous single phase. The molybdenum content range amorphizable for Ni-1Cr-Mo22Ta-14Nb-4P alloys was considerably wide from 1 to 8 at% but the chromium content was restricted to 1-3 at% in Ni-Cr-1Mo-22Ta-14Nb-4P alloys . After immersion for 168 h in 6 M HCl at 303 K the corrosion weight loss for these alloys could not be detected even for some alloys containing minor crystalline phases in the amorphous matrix. If the Ni-Cr-Mo-22Ta-14Nb-5P alloys consisted of an amorphous single phase and contained chromium, molybdenum and 20 at% or more tantalum, the corrosion weight loss was not detected after immersion for 168 h in 12 M HCl at 303 K. These corrosion-resistant amorphous alloys were spontaneously passive in 6 and 12 M HCl. XPS analysis revealed that spontaneous passivation was due to the presence of the air-formed film consisting of tantalum, niobium, a high concentration of chromium and a small amount of molybdenum and being stable in concentrated hydrochloric acids.

References [1] A. Kato, T. Suganuma, H. Horikiri, Y. Kawamura, A. Inoue, T. Masumoto, Mater. Sci. Eng., A179/A180 (1994) 112. [2] Y. Kawamura, H. Kato, A. Inoue, T. Masumoto, Appl. Phys. Lett., 67 (1995) 2008. [3] A. Inoue, Mater. Sci. Eng., A267 (1999) 171. [4] J. H. Kim, E. Akiyama, H. Habazaki, A. Kawashima, K. Asami and K. Hashimoto, Corros. Sci., 34 (1993) 1947. [5] e.g. H. Katagiri, S. Meguro, M. Yamasaki, H. Habazaki, T. Sato, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci., 43 (2001)183, and H. Habazaki, T. Sato, A. Kawashima, K. Asami, K. Hashimoto, Mater. Sci. Eng., A304-306 (2001) 696. [6] H. Habazaki, H. Ukai, K. Izumiya and K. Hashimoto, Mater. Sci. Eng., A318 (2001) 77. [7] H. Habazaki, Y. Naruse, H. Konno, H. Ukai, K. Izumiya and K. Hashimoto, J. D. Sinclair, E. Kalman, M. W. Kendig, W. Plieth, W. H. Smyrl, (eds.) Corrosion and Corrosion Protection, the Electrochemical Society PV 2001- 22 (2001) 130. [8] J. H. Kim, E. Akiyama, H. Habazaki, A. Kawashima, K. Asami and K. Hashimoto, Corros. Sci., 36 (1994) 511. [9] M.-W. Tan, E. Akiyama, A. Kawashima, K. Asami and K. Hashimoto, Corros. Sci., 37 (1995) 1289. [10] P.-Y. Park, E. Akiyama, A. Kawashima, K. Asami and K. Hashimoto, Corros. Sci., 37 (1995) 1843. [11] P. Y. Park, E. Akiyama, H. Habazaki, A. Kawashima, K. Asami and K. Hashimoto, Corros. Sci., 37 (1995) 307.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

71

Passivity of Fe90V10 and Fe75Cr15V10 in Alkaline Media Andreas Drexler,a Hans-Henning Strehblowa a

Heinrich-Heine-Universitaet, Universitaetsstr. 1, 40225 Duesseldorf, Germany, [email protected], [email protected]

Abstract - The formation of passive films on Fe90V10 and Fe75Cr15V10 in 1.0 M NaOH is examined by combined electrochemical and surface analytical studies. The electrochemical behaviour was investigated by potentiodynamic polarization curves and the composition of the passive layer has been studied with X-ray Photoelectron Spectroscopy (XPS) with a systematic variation of the electrochemical parameters. Fe90V10 forms a passive layer with an outer Fe(OH)2 layer whereas Fe(III) accumulates within the inner part of the film. Vanadium does not contribute to the outer parts of the film due to its dissolution as vanadate, however, it is accumulated within the inner part of the layer. Fe75Cr15V10 forms a duplex layer with a sequence Fe-oxide / Cr-V-oxide / bulk-metal and thus an accumulation of the Cr and V cations within the centre of the film. Keywords: Vanadium, Iron-Vanadium-Alloy, Iron-Chromium-Vanadium-Alloy, Passivity, XPS

1. Electrochemical studies Fig. 1a depicts a typical polarisation curve of Fe90V10 in 1.0 M NaOH. Transpassive dissolution as vanadate (VO43-) starts at E = 0.1 V with two wide anodic peaks VA I and VA II which are found in the first scan only. In addition a strong anodic peak FA II and a related cathodic peak FC II is found for the alloy at E = -0.55 V and -0.85 V respectively, which is attributed to the oxidation of Fe(II) to Fe(III) and the reverse reaction. These peaks get larger with the number of scans as for pure Fe and other Fe-alloys [1]-[3]. The small and broad peak FA III at E = -0.30 V corresponds to the oxidation of the outer Fe(OH)2 layer to Fe2O3.

72

A. Drexler and H.-H. Strehblow

The polarization curve of Fe75Cr15V10 in 1.0 M NaOH is shown in Fig 1b. The oxidation of Fe(II) to Fe(III) and the reverse reaction occurs at potentials of E = -0.40 V (FA II) and E = -0.7 V (FC II) respectively, with related current peaks. Two different processes are superimposed in the potential range between E = 0.2 and E = 0.9 V. Vanadium is oxidized with an anodic peak VA I which is found in the first scan only. In addition the peak CA I which is found in the following scans which is assigned to the oxidation of Cr3+ to Cr6+.

Fig. 1: Cyclic voltamograms in 1.0 M NaOH. a) Fe90V10 (dE/dt = 20 mV/s), b) Fe75Cr15V10 (dE/dt) = 30 mV/s)

2. Data analysis After background subtraction according to the method of Shirley [4] the XPS signals V 2p3/2, Fe 2p3/2, Cr 2p3/2, and O 1s were separated into their contributions of the different species with their specific oxidation states by a fit procedure as described in ref. [1] applying software which has been developed in the group of the authors. The quantitative evaluation of the different species within the passive layers was performed on the basis of characterized standards. These standards are described by one or more Gauss-Lorenzian functions with fixed parameters and a tail function. XPS data evaluation of an actual specimen involves a superposition of these standard signals keeping their form and energy value but with a variation of their height, to obtain a close fit of the calculated to the measured signals of the actual samples. The data evaluation on the basis of standards permits a reliable separation of the signals into the peaks of the contributing species, which may be close to each other. XPS data evaluation and the related specimen preparation has been applied successfully to the study of many other systems [5]-[9]. The separation and quantitative evaluation of the XPS-signals in Fe(II) and Fe(III) contents has been previously performed successfully with the same methods which have been applied in this work. XP-spectra of electrochemically passivated Fe90V10 show two oxide species in the V 2p3/2 signal with binding energies at EB = 514.1 eV and EB = 515.7 eV.

73

Passivity of Fe90V10 and Fe75Cr15V10 in alkaline media

Due to these values the oxidation state of vanadium can be at most V(III). Higher oxidation states should cause higher binding energies. For example the binding energy of V2O5 is approximately EB = 517.4 eV [10], [11], however in this energy range there are no peaks detectable. The two vanadium species can be assigned in good agreement with the literature [12], [13] to V(OH)3 (EB = 514.1 eV) and V2O3 (EB = 515.7 eV). The binding energies of the standard signals for all species for this work are summarized in table 1. Table 1: Binding energies in eV for the XP-standards of the metals and cations of this work Fe(0)

Fe(II)

Fe(III)

V(0)

V(OH)3

V2O3

Cr(0)

Cr(OH)3

Cr2O3

706,6

709,5

710,6

512,5

514,1

515,7

574,2

576,7

577,1

3. XPS studies In order to understand the surface chemistry including the details of the chemical composition of the passive layers, Fe90V10 samples have been prepared in 1.0 M NaOH for 300 s with a systematic variation of the potential in the range of E = -1.0 V to E = +0.4 V with a subsequent investigation by XPS.

Fig. 2: change of the composition of the passive layer with the passivation potential EP formed on Fe90V10 in 1.0 M NaOH (pH 13.9) for 300 s deduced from XP spectra a) cationic fraction X, b) cationic fraction for vanadium species only Y

The quantitative evaluation of the XPS-signals of Fe90V10 yields the composition of the passive layer presented in Fig. 2. Fig. 2a shows the cationic fraction X of Eqn. (1) of the layer in dependence on the passivation potential Ep. X(Fe(III)) = n(Fe(III)) / [ n(Fe(III)) + n(Fe(II)) + n(V(OH)3) + n(V2O3) ] (1) It can be subdivided into two potential ranges. At potentials below E < -0.2 V the passive layer contains predominantly Fe(II). The contribution of Fe(III) stays on a constant low level of X(Fe(III)) < 0.2. X(Fe(III)) increases with a step at E = -0.3 V to E = -0.2 V and then linearly for E > -0.2 V up to X(Fe(III)) > 0.8. This behaviour corresponds to that of pure iron, which has been described in ref. [2], [3]. The total content of vanadium cations in the passive

74

A. Drexler and H.-H. Strehblow

layer X(Vox) of Eqn. (2) remains constant for the whole potential range X(Vox) < 0.10. X(Vox)=[n(V(OH)3)+n(V2O3)] / [n(Fe(III))+n(Fe(II))+n(V(OH)3)+n(V2O3)] (2) A more detailed evaluation of both vanadium species considers that the cationic fractions of the vanadium cations only Y of Eqn. (3) change similarly with potential as X(Fe(II)) and X(Fe(III)) respectively. Fig. 2b presents Y for the passive layer on Fe90V10 in the potential range E = -1.0 V to E = +0.4 V. Y(V2O3) = n(V2O3) / [n(V(OH)3) + n(V2O3) ]

(2)

For E ≤ -0.2 V the V(OH)3 content is potential independent and stays on a constant level Y(V(OH)3) = 0.23 with the related value Y(V2O3) = 0.77. For E > -0.2 V Y(V2O3) increases linearly with potential with a corresponding decrease of Y(V(OH)3) which is in good agreement with the behaviour of the iron cations as shown in fig. 2a. Apparently there is a close link between the V(OH)3 to V2O3 reaction and the Fe(II) to Fe(III) oxidation. To investigate the structure of the passive layer in more detail AngularResolved-XPS (AR-XPS) measurements have been performed. It allows the distinction of a predominant location of a species within the inner or outer part of the layer and avoids artefacts like Fe(III) to Fe(II) reduction by preferential sputtering of oxygen during depth profiling.

Fig. 3: AR-XPS of Fe90V10 passivated in 1 M NaOH (pH 13.9), Ep = +0.40 V, tp = 300s a) Fe(II) / Fe(III), b) OH- / O2-, c) Feox / Vox

The intensity ratio of the Fe 2p3/2 signal for the iron cations Fe(II) and Fe(III) in dependence on the take off angle Q after film formation at Ep = +0.40 V for tp = 300 s are depicted in fig 3a. Its increase with increasing Q shows the accumulation of Fe(II) at the layer/electrolyte interface whereas Fe(III)-cations are located within the inner part of the film. Previous studies have pointed out that iron and iron-alloys form an inner Fe(II) and an outer Fe(III) layer [1], [2] i.e. just the opposite situation. The analogue evaluation of the intensity ratio of the OH- / O2- part of the O 1s signal is shown in fig. 3b. This result shows an enrichment of hydroxide at the outer film/electrolyte interface. Fig. 3c depicts the intensity ratio of the Fe 2p3/2 and the V 2p3/2 signal as a function of Q [Feox = Fe(II) + Fe(III); Vox = V2O3 + V(OH)3]. The increase of the iron cations with Q

Passivity of Fe90V10 and Fe75Cr15V10 in alkaline media

75

illustrate the accumulation of iron cations at the outer film/electrolyte interface whereas the vanadium cations are accumulated at the inner film-alloy interface. To determine the thickness of the passive layer, sputter depth profiles have been performed relative to sputter profiles of electrochemically grown Ta2O5 standards [14]. Fig. 4a presents a typical sputter depth profile for Fe75Cr15V10 passivated at Ep = 0.0 V (SHE) for tp = 300 s at E = 0.0 V in 1.0 M NaOH as an example. The oxide/metal interface is defined by the drop of the mole fraction of the oxygen below 0.05, which corresponds to 5%.

Fig. 4: XPS sputter depth profiles of Fe75Cr15V10 passivated in 1 M NaOH (tp = 300 s, Ep = 0.0V) a) molar fraction b) cationic fraction (Feox = Fe(II) + Fe(III), Vox = V(OH)3 + V2O3, Crox = Cr(OH)3 + Cr2O3).

The composition of the anodic oxide on the ternary alloy and its change with depth has been also deduced from the XPS sputter depth profile (Fig. 4b). It shows an enrichment of vanadium- and chromium-oxides within the inner part of the passive layer of 40% and 60% resp. whereas the surface contains 5% and 25% only. In contrast the iron oxides accumulate at the outer phase boundary. The alloy forms in 1.0 M NaOH at E = 0.0 V a 3.7 nm thick duplex layer with an outer iron rich film and an inner vanadium-chromium rich mixed oxide.

Fig 5: Fe 2p3/2, Cr 2p3/2 and V2p3/2 signals for a passive layer on Fe75C15V10 formed at Ep = 0.80 V for 300 s in 1.0 M NaOH in dependence on the sputter depth

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This film-structure is even more pronounced at higher potentials. Fig. 5 depicts the Fe 2p3/2, Cr 2p3/2 and V2p3/2 signals of Fe75Cr15V10 in dependence on the sputter depth and the binding energy at E = 0.80 V in a 3-dimensional presentation. The cations with a higher valency are located at higher binding energies. One discerns from Fig. 5 very easily, that the Fe oxides accumulate at the outer electrolyte/oxide phase boundary. At E = 0.80 V the Fe oxide layer is about 4 nm thick and turns to a Cr-V oxide layer located at a depth of 4 to 8 nm. 4. Summary and Conclusion Electrochemical passive layers on Fe90V10 and Fe75V10Cr15 were prepared in 1.0 M NaOH with systematic variation of the passivation potential and were subsequently examined with XPS. Fe90V10 forms a passive layer which consists mainly of Fe cations. Fe(II) forms the outer Fe(OH)2 layer whereas Fe(III) accumulates within the inner part. Although time dependent XPS measurements of oxide formation show the continuous presence of a Fe(OH)2 film it cannot excluded that it is formed by precipitation from dissolved Fe(II) ions. V does not contribute to the outer parts of the film due to its dissolution as vanadate into the electrolyte. Fe75Cr15V10 forms a duplex layer with a sequence Feoxide / Cr-V-oxide / bulk-metal and thus an accumulation of the Cr and V cations within the centre of the film. 5. Acknowledgement Financial support of the work by the Deutsche Forschungsgemeinschaft for Project Str 200/23-1 is gratefully acknowledged. 6. References 1. H.-H. Strehblow S. Haupt, C. Calinski, U. Collisi, H. W. Hoppe, H. D. Speckmann, Surf. Interface Anal., 9 (1986) 357 2. S. Haupt and H.-H. Strehblow, Langmuir, 3 (1987), 873 3. S. Haupt and H.-H. Strehblow, Corrosion Science, 29 (1989) 163 4. D. A. Shirley, Physical Review, 5 (1972) 4709 5. H.-H. Strehblow in R. C. Alkire, D. M. Kolb (eds.), Advances of electrochemical science and engineering, Vol. 8, Whiley-VCH, Weinheim, Germany (2003) 271 6. D. Schaepers, H.-H. Strehblow, Journal of the Electrochemical Society 142 (1995) 2210 7. D. Schaepers, H.-H. Strehblow, Corrosion Science, 39 (1997) 2193 8. P. Keller, H.-H. Strehblow, Corrosion Science 46 (2004) 1939 9. C. Schmidt, H.-H. Strehblow, Journal of the Electrochemical Society, 145 (1998) 834 10. G. C. Bond, S. Flamerz, Applied Catalysis, 46 (1989) 89 11. V. I. Nefedov et. al., J. Elec. Spec. Rel. Phen., 10 (1977) 121 12. B. Horvath et. al., Z. Anorg. Allg. Chem., 483 (1981) 181 13. R. J. Colton, A. M. Guzman, J. W. Rabalais, Journal of Applied Physics, 49 (1978) 409 14. J.M. Sanz, S. Hofmann, Surface and Interface Analysis, 5 (1983) 210

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Effect of anodic passivation on the corrosion behaviour of Fe-Mn-Al steels in 3.5%NaCl A.S. Hamada,a L.P. Karjalainen,a M.A. El-Zekyb a

Dept. of Mechanical Engineering, University of Oulu, 90014 Oulu, Finland Dept. of Metallurgy & Materials Eng., Faculty of Petroleum & Mining Eng., Suez Canal University, Suez, Egypt [email protected] b

Abstract - The effects of anodic passivation in 30%HNO3 on the structure of the passive film and on the corrosion behaviour of three Fe-Mn-Al steels in 3.5%NaCl solution were studied by electrochemical measurements, X-ray diffraction and scanning electron microscopy. In general, the corrosion resistance of the Fe-30Mn-4.5Al-4Cr (in wt-%) alloy, characterized by the polarization curves, was found to be superior to that of Fe-24Mn-5.7Al and Fe-23Mn-8Al, as passivated in 30%HNO3 or in the unpassivated condition. Green Rust was observed as the main product of anodic dissolution of the unpassivated Fe-24Mn-5.7Al and Fe-23Mn-8Al alloys. Prolonged anodic ageing in the anodic passive regime in 30%HNO3 induced a thick, protective and stable passive film that enhanced the corrosion resistance in 3.5%NaCl. The highest corrosion resistance of Fe-30Mn-4.5Al-4Cr can be attributed to the enrichment of Al2O3 and Cr2O3 and the depletion of Fe and Mn oxides in the passive film. Keywords: Fe-Mn-Al steels, Cr alloying, anodic passivation, passive film, oxide layer

1. Introduction Fe–Mn–Al alloys with 20–30 wt-% Mn and 4–20 wt-% Al can be single-phase austenitic, ferritic or duplex-phase steels [1]. Austenitic grades have gained attention as potential substitutes for conventional austenitic Cr-Ni stainless steels [2-6]. In recent years interest towards these steels has increased even more, for some austenitic compositions possess the property of twinninginduced plasticity that increases their ductility, energy absorption and toughness [7-9]. Such steels can find applications in the automotive industry, for instance.

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There are several papers on the electrochemical polarization behaviour of FeMn-Al alloys. The authors have reported that these steels do not passivate in aqueous solutions of 3.5%NaCl, 10%HCl or 10%HNO3 and that they are susceptible to pitting corrosion [10-14]. In the present study, a potential method for improving the corrosion resistance of Fe-Mn-Al alloys by modifying the composition and the constitution of the surface layer has been investigated. This consists of reducing the surface concentrations of elements that decrease the corrosion resistance and enriching elements that improve the corrosion resistance by means of electrochemical anodic polarization in an oxidizing electrolytic solution. 2. Experimental Three high-Mn Fe-Mn-Al type steels were prepared by an induction melting. Their chemical compositions are given in Table I. After homogenization at 1100°C for 24 h, the cast alloys were hot rolled into strips of 1 mm thickness. Table I: Chemical compositions of the investigated Fe-Mn-Al alloys (wt-%).

Alloy code Fe-25Mn-5Al Fe-25Mn-8Al Fe-30Mn-4Al-4Cr

C 0.22 0.20 0.27

Mn 24 23 30

Al 5.7 8.4 4.5

Cr 0.08 0.08 4.14

Si 0.48 0.49 0.43

Fe Bal. Bal. Bal.

For polarization tests, a conventional three-electrode, single compartment, cylindrical glass cell was used with graphite as the counter electrode. All the potentials were recorded with respect to the saturated calomel electrode (SCE). Passive films were analyzed by X-ray diffraction (XRD) and scanning electron microscopy (SEM) equipped with an energy dispersive spectrometer (EDS). 3. Results and discussion The microstructures of the solution treated Fe-25Mn-5Al and Fe-30Mn-4Al-4Cr consisted equiaxed austenitic grains with the coarse grain size of 180 mm. In Fe25Mn-8Al, both the austenite and ferrite phases were present, as seen in Fig. 1. The volume fraction of the ferrite phase was about 55%, as determined by the point-counting method. XRD analysis of the tested alloys confirmed the presence of those phases in the microstructure. Potentiodynamic polarization curves for the three Fe-Mn-Al steels in the 30%HNO3 aqueous solution are displayed in Fig. 2. The polarization curve of Fe-30Mn-4Al-4Cr presents a stable passive region at a high positive potential and the low passivation current density of 8 mAcm-2.

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(a)

(b)

A F

F

A

Figure 1. Optical micrographs of the solution-treated alloys, (a) Fe-25Mn-5Al, (b) Fe-25Mn-8Al. A = austenite; F = ferrite phase. For Fe-25Mn-8Al, the polarization curve shows the occurrence of passivation at a higher current density of about 22 mAcm-2. Furthermore, oscillations in the current density of the order of one mAcm-2 are present at the onset of passivation and beside the Flade potential. These oscillations are known as the oscillatory phenomenon [15]. This is now observed for the first time for Fe-MnAl type steels, although there are several investigations on the corrosion behaviour of these alloys [10-14]. As shown by the polarization curve, Fe25Mn-5Al has the same electrochemical behaviour as that of Fe-25Mn-8Al in 30%HNO3. In addition, the current oscillations are observed in the vicinity of the Flade potential.

Figure 2. Polarization curves of the Fe-Mn-Al steels in 30%HNO3. Concluding from the high frequency oscillations the passivation of the Fe25Mn-8Al and Fe-25Mn-5Al steels is a result from the formation and dissolution of salt films on the surface, not from the spontaneous formation of

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an oxide film as in the Fe-30Mn-4Al-4Cr [15]. The relaxation oscillations in these two steels are characterized by the quasistationary phase with a negligible dissolution rate and a gradual relaxation to the passive state. The efficient passivation of Fe-30Mn-4Al-4Cr is due to the presence of Al and Cr, which induce the formation of a passive film with a higher stability than that formed in the steels without Cr. The self-passivation, a spontaneous activepassive transition, of Fe-Mn-Al steels containing Al and Al with Cr can be attributed to the autocatalytic nature of HNO3 reduction [16, 17]. Zhang et al. [18] reported that the stability of Mn is low in Fe-Mn base alloys containing Al and Cr. Therefore, Mn readily forms oxides and it is preferentially dissolved at the oxide/electrolyte interface. Accordingly, this autocatalytic reduction induced the preferential dissolution of Mn and Fe and the enrichment of Al and Cr that formed the passive film. The polarization curves for the steels in an aerated 3.5%NaCl solution before the anodic passivation had a similar shape but showed different values of the electrochemical parameters. Under the conditions evaluated here, oxygen reduction is the prevailing cathodic reaction. With increasing potential in the anodic dissolution regime, the dissolution of the alloying elements of Mn and Al as well as matrix Fe would occur. As the Fe(II) ions move away they meet hydroxide ions and produce Fe(II) hydroxide. The visual inspection of the corrosion solutions during the anodic polarization of Fe-25Mn-5Al and Fe25Mn-8Al steels revealed a film of the green rust (GR) suspension. GR is an intermediate compound between Fe(II) hydroxide and Fe(III) oxyhydroxide formed by the oxidation of Fe(OH)2 aqueous suspensions. The occurrence of GR in the course of corrosion of iron and steel under aerobic and anaerobic conditions has commonly been reported [19]. GR(Cl-) was observed here as the main product of the first stage of the corrosion process of the alloys without chromium in 3.5%NaCl solution, but not in the Cr-bearing steel. Table II shows the composition of the corrosion products of the investigated steels after filtration. The chemical balance of the reaction leading to the formation of GR(Cl-), Fe(II)3Fe(III)(OH)8Cl(2H2O), has been described in Ref. [19]. The order of the corrosion current in 3.5%NaCl for the present steels is Fe-25Mn5Al > Fe-25Mn-8Al > Fe-30Mn-4Al-4Cr. Table II: The wet analysis of the corrosion products of the alloys in 3.5%NaCl solution.

Compound Cr-free alloys Cr-bearing alloy

Fe2O3 18.70 66.30

MnO 6.84 8.56

GR(Cl-) Bal. -

FeO Bal.

In the following experiments, samples of the alloys were passivated anodically in a 30% HNO3 solution at +800 mV, i.e. aged at a constant potential

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corresponding to the median passive region of the alloys, for various times of 15, 60 and 300 min. Subsequently, the effects of anodic ageing time on the corrosion potential Ecorr and the corrosion current Icorr in the 3.5%NaCl solution were studied by recording the electrochemical polarization of the alloys. The results revealed that the prolonged anodic passivation time up to 300 min increased the stability and improved the corrosion resistance of the passive film in 3.5%NaCl. The most pronounced effect was obtained for Fe-30Mn-4Al-4Cr, as seen in Fig. 3. The positive influence of short passivation times of 15 min and 60 min was not significant for Fe-25Mn-5Al and Fe-25Mn-8Al. However, even these short times had obviously beneficial effects on Ecorr and Icorr of Fe30Mn-4Al-4Cr, for no GR(Cl-) was formed during the polarization measurements on this steel. Neither was it formed in the passivated Fe-25Mn5Al nor Fe-25Mn-8Al after the longest ageing time.

Figure 3. Effect of anodic passivation time on Icorr and Ecorr in 3.5%NaCl.

Figure 4. XRD of the passive film formed on Fe-30Mn-4Al-4Cr after 5 h ageing.

The structure of the passive film of Fe-30Mn-4Al-4Cr was analyzed by SEM and XRD after the electrochemical treatments. The thickness of the passive film increased with increasing passivation time from 90 nm at 15 min to 460 nm at 5 h. This suggests that the formation of passive film takes place by the nucleation and growth mechanism. The XRD examination, as shown in Fig. 4, revealed the presence of Al, Cr, Mn and Fe oxides in the passive film that formed during the 5 h ageing.

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4. Conclusions Polarization measurements and XRD analysis indicate that the rather good passivation of Fe-30Mn-4Al-4Cr steel is due to the presence of Al and Cr. The high passivity of Al and Cr induces the formation of a passive film that is more stable than that formed in the steels without Cr. The passive film on this steel mainly consists of the oxides of Al, Cr, Mn, and Fe. The formation of both Al and Cr oxides takes place by a direct nucleation and growth mechanism on the steel surface, where the enrichment of Al and Cr and the depletion of Fe and Mn have occurred. A prolonged anodic ageing time in 30% HNO3 aqueous solution can result in the modification of the passive layer, which has a beneficial influence on the stability of the passive film and consequently on the corrosion resistance of the steel. Hence, anodic passivation ageing of passive films in an oxidizing electrolytic solution can be suggested as a surface modification method for improving the corrosion resistance of Fe-Mn-Al and Fe-Mn-Al-Cr type steels. References [1] [2] [3] [4] [5] [6] [7] [8] [9]

K. Sato, K. Tanaka, Y. Inoue, ISIJ Int. 29 (1989) 788 J. Charles, A. Berghezan, A. Lutts, P.L. Dancoisne, Metal Progr. (1981) 71 Y.G. Kim, Y.S. Park, J.K. Han, Metal. Trans. A 16A (1985) 1689 Y.G. Kim, J.K. Han, E.W. Lee, Metall. Trans. A 17A (1986) 2097 V.C. Lins, M.A. Freitas, E.M. Paula e Silva, Corr. Sci. 46 (2004) 1895 A. Dias, V. de Freitas Cunha Lins, Corr. Sci. 40 (1998) 271 S. Allain, J.P. Chateau, O. Bouaziz, Steel Research 73 (2002) 299 G. Frommeyer, U. Brux and P. Neumann, ISIJ Int. 43 (2003) 438 S. Vercammen, B. Blanpain, B.C.D. Cooman, P. Wollants, Acta Mat. 52 (2004) 2005 [10] L.W. Liu, C.J. Chen, J. Zhejiang Univ. 35 (2001) 360 [11] S.T. Shih, C.Y. Tai, T.P. Perng, Corrosion 49 (1993) 130 [12] Y.S. Zhang, X.M. Zhu, Corr. Sci. 41 (1999) 1817 [13] M. Ruscak, T.P. Perng, Corrosion 51 (1995) 738 [14] Y.J. Gau, J.K. Wu, Corrosion Prev. & Contr. 44 (1997) 56 [15] S.G. Corcoran, K. Sieradzki, J. Electrochem. Soc. 139 (1992) 1568 [16] V.P. Razygraev, R.S. Balovneva, E.Y. Ponomareva, M.V. Lebedeva, Prot. Met. 26 (1990) 43 [17] D.G. Kolman, D.K. Ford, D.P. Butt, T.O. Nelson, Corr. Sci. 39 (1997) (1997) 2067 [18] Y.S. Zhang, X.M. Zhu, M. Liu, R.X. Che, Appl. Surf. Sci. 222 (2004) 89 [19] Ph. Refait, O. Benali, M. Abdelmoula, J.M. Genin, Corr. Sci. 45 (2003) 2435.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Surface Characterization of 1018 Carbon Steel in Borate Medium by in-situ Electrochemical Scanning Tunneling Microscopy Ignacio Gonzáleza,*, Román Cabrera-Sierraa,b, Nikola Batinaa a

Universidad Autónoma Metropolitana-Iztapalapa. Departamento de Química. Apartado Postal 55-534. 09340 México, D.F. (MEXICO). [email protected] b Escuela Superior de Ingeniería Química e Industrias Extractivas (ESIQIE-IPN) Academia de Química Analítica, Edificio Z5. A.P:75-874, C.P. 07338, México, D.F. (MEXICO).

Abstract - Microstructures of low carbon steel are ferrite (Fe-a) and pearlite (alternate mixture of Fe-a and Fe3C) and each one has its own oxidation mechanism. These two phases were identified using in-situ electrochemical scanning tunneling microscopy (ECSTM). Real time images were obtained during the immersion of 1018 carbon steel probes in 0.642 M H3BO3 and 0.1 M NaOH, pH = 7.8. Two different corrosion mechanisms were identified and correlated with the observed surface changes. Keywords : ECSTM, Corrosion, Carbon Steel, Pearlite, Ferrite .

1. Introduction Carbon steel corrosion in aqueous media is one of the most studied systems for its impact on the chemical industry. Usually, the carbon steel - aqueous media interface has been studied in-situ using different electrochemical and ex-situ using different spectroscopic techniques [1-8]. In particular, the surface composition of the low carbon steel is strongly related to its oxidation processes. In dependence on the fabrication process, a carbon steel surface has different phases such as ferrite, Fea (iron rich phase), and pearlite (sequential arrangement of ferrite and iron carbide, Fe3C) [9]. Usually, corrosion studies of a steel interface in an aqueous medium do not consider separate oxidation mechanism for each phase, thereby obtaining only global information on the

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mechanism relative to that interface [1-8]. This work focus on determination and characterization of corrosion of ferrite and pearlite on a steel sample, in the borate medium, using an in-situ Electrochemical Scanning Tunneling Microscopy (ECSTM). The real time monitoring of changes in the steel surface topography, allows to understand each phase oxidation as well as their influence on the global carbon steel corrosion process. 2. Results and Discussion The experimental conditions and methodology of the ECSTM image recording process were described in details in our previous work [10]. It is important to indicate that the experimental methodology proposed in this work for surface characterization of 1018 carbon steel in borate medium using in-situ ECSTM technique is different from that reported in other works [11-13]. In the literature, such characterization is carried out after cleaning the surface of corrosion products formed during its preparation (through an imposed cathodic potential). However, here it is proposed that the surface is studied with no prior electrochemical treatment. Moreover, this work focuses on studying the corrosion process of steel in aqueous medium using the time of immersion (at early stages of corrosion) in order to differentiate the reactivity of distinct microstructures on the basis of oxidation mechanism of each steel phase. Nevertheless, it is important to indicate that such goal can be achieved only if the corrosion process images are obtained at short immersion times, since longer times make the process visualization difficult. Mechanism of the imaging of iron oxides is still not clear for us. We believe the iron oxides are conductives specially magnetite. However we also believed that specially important for imaging obtention is the thickness of the oxide films, which is not defined in our work. After a systematic analysis of the major part of the sample surface, at different immersion time in a borate medium, two different types of surface with characteristic structures were identified (pearlite and ferrite) and selected as model surfaces for further studies of the corrosion behavior. A set of images presented in Figures 1 (a-b), is related to characterization of the pearlite phase on the 1018 carbon steel surface using the ECSTM technique, in a borate medium. The key issue was an identification of the periodical arrangement (linear rows) characteristic for the pearlite phase, of the carbon steel [14,15]. The surface topography (surface features) observed in our ECSTM images is very similar to surface microstructure seen by Scanning Electron Microscopy (SEM), reported and describe in details in previous works [14,15]. In agreement with this previous work, we also found that distance between the observed pearlite rows in SEM (0.095 mm) and ECSTM images (0.1 mm) is identical. As a further step in our study, the part of the sample surface identified as a pearlite phase, was a subject of continuos monitoring during the oxidation progress as a function of the immersion time (characterization time). Figure 1b shows the progress of the oxidation of the pearlite phase after 84 min of immersion in the

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borate solution. The sample surface is covered by granular features which size varies from 0.4 to 1 mm. Pearlite

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Figure 1. In-situ ECSTM images of 1018 carbon steel immersed in 0.642 M H3BO3 and 0.1 M NaOH, pH = 7.8, after different immersion time. Pearlite phase: a) t =11.5 min (Z scale: 0 – 27.5 nm), b) t = 84 min (Z scale: 0 – 22.5 nm ). Ferrite phase: c) t = 23 min (Z scale: 0 – 20 nm), d) t = 36 min (Z scale: 0 – 22.5 nm). All images were recorded at the Open Circuit Potential (-0.125 V), with a bias potential of -0.180 V, a tip potential of –0.305 V and tip current of 2 - 3 nA.

The oxide growth occurs along the preferential directions determinate by the texture and the microstructure of the previous pearlite layer. It persists even at the larger time of immersion with more progress of oxidation. The width and the distance between the original pearlite rows increase as a function of the immersion time. The observed surface topography changes a typical for low oxidative properties of the pearlite phase of the steel surface. Distinguished surface topography without linear alignment, with larger granular features (1-2

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mm) and with significantly higher magnitude of the surface perturbations during the oxidation process, was observed for ferrite (Fe-a) phase (Figs. 1 (c-d)). Note that this microstructure is very different from pearlite. Images at Figs. 1 (a-d) shows that oxide morphology depends on the characteristic and composition of the steel surface (pearlite vs. ferrite phase). However, at open circuit potential (OCP) conditions, after few first minutes of immersion, the sample surface is completely covered by the oxide film. In accordance to the literature [16-21], in case of both phases, it is an internally formed oxide film (barrier layer with passive properties [11,12,22]) known as magnetite. Furthermore this internally formed iron oxide film will be compared with the externally formed oxide film, grown by applying external oxidation potentials. Images at Figs 1b and 1d, recorded at –0.125 V vs. NHE (OCP) serve as the reference for external oxide growth monitoring over the inner oxide film of pearlite and ferrite phases, respectively. Figure 2a shows typical image of oxidized pearlite surface obtained by imposing a potential of – 0.120 V vs. NHE (oxidation overpotential of 5mV) after the period of 33.33 minutes. The original laminar arrangement (pearlite rows) is still visible, as in Figs.1 a and b. The surface feature becomes slightly smaller, which could be first sign of the surface oxide dissolution, and possibly formation of a passive film formed externally (maghemite). This is in agreement with results reported for iron oxidation in a borate medium [11,12,22]. Figure 2b shows a set of images obtained by imposing an overpotential of 65 mV after 3.33 minutes to the pearlite phase in a borate medium. Detail analysis shows that rows are higher and slightly wider, as result of a preferential growth of the new oxide at the top of the pearlite/magnetite lines. Contrasting to pearlite, the further oxidation of ferrite phase promotes development of different topography (Figs. 2 c and d). Surface is covered by large oxide grains (1.5 – 2 mm) not observed before in the case of any stage of ferrite oxidation (Figs. 1c and 1d). After oxidizing the ferrite phase up to 33.33 minutes, deterioration of the oxide film is visible due to appearance of the nanometric size pits. We suppose that it is result of rather complex mechanism of formation of the external oxide film, which is associated to maghemite. However, at this point we could not speculate more about mechanisms involved in this process. Here, we just could conclude that observed changes in the surface topography differs significantly for different phases: pearlite and ferrite in case of formation of internal and external iron oxide layers. In other words ECSTM appears to be a useful tool for in-situ monitoring and characterization of the steel surface oxidation. 3. Conclusion Using ECSTM it was successfully characterized the state of the electrode surface, identified phases at the steel sample, determinate surface morphology characteristics and by systematic monitoring of topography change, it was

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possible to determinate and differentiate oxidation mechanisms of different steel phases (pearlite vs. ferrite). In particular it is interesting a visualization of development of internal and external oxide layers, which open new possibilities in studies of the oxidation processes at the micro and nanoscale level, at in-situ conditions. Pearlite

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Figure 2. Typical ECSTM images (6mm x 6mm) recorded after applying different oxidation overpotentials to the different carbon steel phases in contact with the borate medium (pH=7.8). Pearlite, a) overpotential 5 mV after 33.33 minutes (Z scale: 0 – 25 nm, tip potential –0.300 V), b) overpotential 65 mV after 3.33 minutes (Z scale: 0 – 50 nm, tip potential –0.120V). Ferrite, c) overpotential 150 mV after 3.33 minutes (Z scale: 0 – 27 nm, tip potential –0.155 V), d) overpotential 170 mV after 33.33 minutes (Z scale: 0 - 45 nm, tip potential –0.135 V). The usual imaging conditions were a bias potential of -0.180 V and tip current of 2 - 3 nA.

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Acknowledgments: The authors are grateful for the financial support of CONACyT through the project 47162 and IMP, project FIES-98-100-I. References [1] T.A. Ramanarayanan and S.N. Smith. Corrosion. 46 (1990), 66. [2] H. Huang and W.J.D. Shaw. Corrosion. 48 (1992), 931. [3]H. Vedage, T.A. Ramanarayanan, J.D. Munford and S.N. Smith. Corrosion. 49 (1993), 114. [4] H. Shutt and P.R. Rhodes. Corrosion. 52 (1996), 947. [5] Y.F. Cheng and J.L. Luo. Electrochim. Acta. 44 (1999), 2947. [6] Y.F. Cheng and J.L. Luo. Appl. Surf. Sci. 167 (2000), 113. [7] A. Groysman and N. Erdman. Corrosion. 56 (2000), 1266. [8] D.A Lopez, S.N. Simison and S.R. de Sánchez. Electrochim. Acta. 48 (2003), 845. [9] J. R. Davies, ed. Properties and Selection: Iron and Steel. Vol. 1; Materials Characterization. Vol. 10. American Society of Metals (ASM), Ohio, USA. [10] R. Cabrera-Sierra, N. Batina and I. González. Materials Chemistry and Physics. Submitted. [11] I. Diéz-Pérez, P. Gorostiza, F. Sanz and C. Müller, J. Electrochem. Soc. 148 (2001), B307. [12] J. Li, D. J. Maier, J. Electroanal. Chem. 454 (1998), 53. [13] R.C. Bhardwaj, A. González-Martín, and J. O’M. Bockris, J. Electrochem. Soc. 138 (1991), 1901. [14] D.G. Enos and J. R. Scully, Metall. Mater. Trans. A, 33A (2002), 1151. [15] S. W. Thompson and P. R. Howell, J. Mat. Sci. Lett., 17 (1998), 869. [16] C.Y. Chao, L.F. Lin, and D.D. Macdonald, J. Electrochem. Soc. 128 (1981), 1187. [17] D. D. Macdonald and M. Urquidi-Macdonald. J. Electrochem. Soc. 137 (1990), 2395 [18] D. D. Macdonald, Sonia R. Biaggio, and Herking Song. J. Electrochem. Soc. 139 (1992), 170. [19] L. Zhang, D.D. Macdonald, E. Sikora, and J. Sikora. J. Electrochem. Soc. 145 (1998), 898. [20] J. Liu and D. D. Macdonald. J. Electrochem. Soc. 148 (2001), B425. [21] D.D. Macdonald, K.M. Ismail, and E. Sikora. J. Electrochem. Soc. 145 (1998), 3141. [22] M. Nagayama and M. Cohen, J. Electrochem. Soc. 109 (1962), 781.

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Iron passivation studied by in situ Raman spectroscopy on Fe/Au(111) epitaxial films S. Joiret a, P. Allongueb a

Laboratoire des Interface et Systèmes Electrochimiques, UPR 15 CNRS Université P & M Curie, 4 Place Jussieu, 75005 Paris b

Laboratoire de Physique de la Matière Condensée, CNRS - UMR 7643, Ecole Polytechnique, 91128 Palaiseau, France

Abstract - New substrates giving enhanced Raman spectra have been used to follow the reduction of iron passive layer grown in neutral (pH 8.4; borate) or alkaline (pH 12.7; NaOH) solutions. A thin (20 monolayers) of epitaxial bcc iron layer is electrodeposited on Au (111) , 30 nm thick evaporated film on mica. The electrochemical behaviour of these films is analogous to the one of a massive iron electrode.Depending on the gold evaporation conditions an electromagnetic enhancement of the iron oxide related Raman signal is obatined. This property has been used to record the variation of the Raman signal during the passive film reduction in potentiodynamic mode. Special attention has been paid to the key role of magnetite formation during the reduction of the passive film. In borate solution the film is essentially of a spinel structure in its whole existence range while in basic solution the presence of an outer layer, still under investigation, is evidenced. Keywords: Iron passivity, SERS, Spinel structure, borate solution 1. Introduction The growth mechanism, structure and chemical composition of the passive film on iron in borate solution have been investigated for several decades [1,3]. As to structural models of the passive film, Heusler et al. [1] postulated very early that

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the anodic oxide presents a structure with a miscibility between Fe3O4 layer (inner) and g-Fe2O3 (outer) layer. A second popular model considered the passive layer as an intern anhydrous oxide phase covered with a hydrated hydro / oxyhydroxide phase [2]. The recent structural studies performed on welldefined iron substrates suggest that the oxide film is a compact spinel structure with no foreign species inside nor hydroxide or oxyhydroxide atop. The LAMM phase found on iron single crystals [3] is a cation deficient spinel oxide structure which resembles that of g-Fe2O3 but with randomly distributed cations vacancies and octahedral Fe(III) in interstitials while the oxygen lattice is fully occupied. The same film structure, ruling out the formation of g-Fe2O3, was found from in situ Raman on Fe(110)/Au(111) epitaxial layers [4] with no additional layer. The above brief review outlines a long lasting controversy between the duplex models and the formation of a homogenous phase. The fact that in situ SERS [4] confirmed the LAMM model [3] constituted the starting point of this work because this proves that we can safely characterize in situ and in real time the passive film chemistry and structure. Using the same ultrathin Fe(110)/Au(111) as electrodes as for SERS we have identified, for the first time, the final reduction step of the passive film in borate buffer solutions. This study has been extended to basic solution where the existence of an outer layer, different from passive film itself has been checked. 2. Experimental Fe thin film electrodes, hereafter designated as Fe/Au(111), were obtained by electrodepositing iron on a SERS-active 30nm thick highly textured Au(111) films obtained by evaporation on a freshly mica substrate [5]. After flame annealing the film consists of large flat top grains onto which the 22x√3 Au(111) reconstruction is visible by STM. A conventional three-electrode electrochemical cell was connected to a potentiostat (PGSTAT12, Eco Chemie, Utrecht, The Netherlands) with a Pt wire as counter electrode and a saturated mercurous sulfate electrode (MSE) as reference. Iron was electro-deposited under potential control right after the flame annealing at –1.5 V/MSE from a 1 mM FeSO4 solution in 10-2 M K2SO4 + 10-4 M KCl + 10-3 M H2SO4 [6]. Passivation was studied in a borate buffer solution of pH 8.4 and NaOH 0.1M pH 12.7. All solutions were prepared with pro analysis purity grade reagents (Merck, Germany) and 18 MW·cm Milli-Q water. In what follows, all potentials are quoted versus NHE (reversible hydrogen electrode in the same solution).

Iron passivation

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3. Results and discussion 3.1. Passivation in the borate solution Prior to use, the freshly prepared Fe/Au electrode was pre-polarized at –0.45 V for 5 minutes. The anodic oxide was formed by stepping the sample potential into the passive region ( +1.0 V < UA < +1.2 V) instead of slowly ramping the potential anodically. This procedure is avoiding iron dissolution and ensures the growth of a compact passive film with no formation of oxyhydroxide [7]. After passivation (nominally during 2 min), the electrode potential was scanned negatively towards the cathodic limit UC = –0.45 V at a rate of 5 mV/s. . In what follows the different peaks will be characterized by their potential U(I’) and charge Q(I’) for peak I’ etc .To estimate the life-time of a given Fe/Au electrode, the above described process was repeated several times on a Fe(20 ML)/Au electrode. Fig. 1a evidences some changes during successive cathodic scans: after 3 complete passivation / reduction cycles Q(I’) is divided by a factor 2 and complete disappearance of peak I’ is achieved after 6 cycles. Given the initial iron thickness (20 ML) we conclude that about 3-4 equivalent iron monolayers are dissolved per electrode cycle.

Fig 1: Voltammograms of Fe(20ML)/Au(111) electrodes in the borate solution : (a) effect of scan number ; (b) Influence of stopping the sweep of potential at potentials indicated by arrows.

The effect of holding the potential at different stages of the reduction scan is displayed in Fig. 1b where the reference reduction scan is also shown (continuous sweep at 5 mV/s). Stopping the scan at 0.5 V for 20 min suppresses the peaks II’a and II’b but has no impact on peak I’. A control experiment (not shown), in which the potential is held at 0.7 V for 20 min, keeps intact both the peaks II’ and peak I’. Stopping for 15 min the reductive scan between peaks I’ and II’ (U = 0.1 V) decreases of Q(I’) by a factor 10 (dashed line).

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3.2. Passivation in 0.1 M NaOH Cyclic voltammograms of a Fe(20ML)/Au(111) electrode in 0.1M NaOH are displayed in Fig. 2. The same peak I' is evidenced with a charge twice as large as the one measured in the borate solution. A new pair of quasi-reversible peaks labeled III and III' is growing upon continuous cycling. The Fe(20ML)/Au can be cycled 10 times before Q(I') begins to diminish and 20 cycles are necessary to let disappear peak I’ while peak III reaches a maximum. The increase of peak III is correlated with the process on peak I' as reverting the voltammograms before I' evolution suppress peak III.

Fig 2: Cyclic voltammograms of Fe(20ML)/Au(111) in 0.1M NaOH.

3.3. In situ Raman spectroscopy The Raman spectra shown in Fig. 3a and 3b were recorded in the borate solution at high potential (a) and between peak II' and I' (b) in 3s only thanks to the SERS active gold substrate.As discussed in [4], the two lines in spectrum 3a (670 and 720 cm-1) are characteristic of an oxidized spinel (LAMM phase [3] ). Spectrum 3b is characteristic of magnetite Fe3O4(111) which is formed from U < U(II’) [4]. No spectrum could be resolved at potentials negative of peak I'. The potential dependence of spectra (Fig. 3c) recorded in real time during the cathodic sweep of potential (5 mV/s) evidences that the intensity of the peak at 720 cm-1 is vanishing from peak II’ which a direct manifestation of the above mentioned transition LAMM å magnetite transition (filled circles). The quenching of the magnetite related Raman intensity (integration range 630 to 710 cm-1, open circles) is closely correlated with peak I’. There is also a slow decrease of the peak intensity for U(I’) < U < U(II’).

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Fig 3: In situ Raman of Fe(20ML)/Au(111) in borate solution : (a) and (b) spectra at 1.2 and 0V (exposure time 3 s). Two lines at 670 and 720 cm-1 and one line at 670 cm-1 are respectively identified. (c) Potential dependence of the 720 cm-1 line intensity (see (a)) (filled circles) and 680 cm-1 line intensity (see (b)) (open circles). The solid line recalls the I – V sweep of potential.

Passivation in the basic solution at high anodic potential leads to a different Raman spectrum. In Fig 4a a spinel related peak is however clearly identified after correction from a wide asymmetric background signal, which line shape is weakly potential dependent (Fig. 4b). The potential dependence of the spinel related peak (Fig 4c) shows that the signal does not return to zero for U < U(I’).

Fig 4: In situ Raman spectra of Fe(20ML)/Au(111) in basic solution : (a) as measured at 1.2V/SHE and after background correction; (b) after 20 cycles at potentials slightly negative (1) and positive (2) of peak III (c) Potential dependence of spinel related peak intensity (square) with current versus potential The solid line is the voltammogram.

3.4. Electrochemical model From the above results a common model can be proposed for the reduction of the iron passive film either at neutral or basic pH. In both cases, one starts with a spinel structure (LAMM phase [3]) with a fully occupied O-lattice and cationic vacancies. From a stoechiometry viewpoint, the LAMM phase may be described as “Fe20O32”.The following reaction occur at different potentials:

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Fe20O32 + 4Fe + 2O2 + 4H2O + 8e- å Fe24O32+ 8OH(1) 0 Fe3O4 (s) + 8H2O + 8e ® 3Fe (s) + 8OH + 4H2 (2) Fe3O4 (s) + 4H2O + 2e® 3 Fe(OH)2 (aqueous or solid) + 2OH- (3) Fe2+(OH-,H2O) ® Fe3+(OH-,H2O) + e(4) Reaction (1) stands for the transition LAMM å magnetite at peak II’. It coincides with the onset of oxygen reduction and accounts for the conservation of the O sub-lattice. It also involves the oxidation and the incorporation of Fe atoms (from the iron substrate) to fill the vacancies existing in the LAMM phase, in agreement with the mixed ionic-electronic conduction at this potential [8]. Reactions (2) describes the reduction of magnetite into metallic iron at peak I’ at pH 8.4 (total disappearance of the Raman signal). Reaction (3) accounts for the reductive dissolution of magnetite when U(I’) < U < U(II’) in agreement with the slow decrease of the Raman signal (Fig. 3c). At more elevated pH, the dissolution products become insoluble. Reaction (3) was rewritten as reaction (4) to express that a sort of "solid" highly hydrated and amorphous adlayer forms atop the spinel layer. 4. Conclusion The use of epitaxial iron films deposited on SERS active Au(111) layer allows the real time monitoring of passive layer composition. The present data show that the passive layer formed at high potential in neutral or basic solution undergoes a phase transition into magnetite at peak II’. At pH 8.4, the magnetite layer can be either chemically dissolved or electrochemically reduced into metallic iron (at peak I’). In the basic solution, these reactions produce also partially insoluble ferrous species, which are forming a hydrated layer on the surface. References 1. K.E. Heusler, K.G.Weil, K.F. Bonhoeffer, Z. Phys. Chem. N. F., 15, 149 (1958) 2. R. W. Revie, B. G. Baker,J. O'M. Bockris, J. Electrochem. Soc.,122, 1460 (1975) 3. A. J. Davenport, L. J. Oblonsky, M. P. Ryan, M. F. Toney, J. Electrochem. Soc., 147, 2162 (2000) 4. P. Allongue and S. Joiret, Phys. Rev. B, 71, 115407/1 (2005). 5. P. Allongue, L. Cagnon, C. Gomes, A. Gündel, V. Costa, Surf. Sci. 557, 41 (2004) 6. L. Cagnon, T. Devolder, R. Cortès, A. Morrone, J. E. Scmidt, C. Chappert and P. Allongue, Phys. Rev. B 63,104419 (2001) 7. S.Virtanen, P. Schmuki, A. J Davenport,.; C. M Vitus,. J. Electrochem. Soc. 144, 198 (1997) 8. M. Bojinov, T. Laitinen, K. Mäkelä, T. Saario, J. Electrochem. Soc., 148, B243-B250 (2001).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Atomic-Structure Characterization of Passive Film of Fe by Grazing Incidence X-ray Scattering at SPring-8 Masugu Satoa, Masao Kimurab, Masato Yamashitac, Hiroyuki Konishid, Shinji Fujimotoe, Yasunori Tabiraf, Takashi Doig, Masayasu Nagoshih, Shigeru Suzukii, Takayuki Kamimurag, Takenori Nakayamaj, and Toshiaki Ohtsukak a

Japan Synchrotron Radiation Research Institute, 1-1-1 Kouto, Mikazuki, Sayo, Hyogo, Japan b Nippon Steel Corporation, 20-1 Shintomi, Futtsu, Chiba, Japan c University of Hyogo, 2167 Shosha, Himeji, Hyogo, Japan d Japan Atomic Energy Research Institute, 1-1-1 Kouto, Mikazuki, Sayo, Hyogo, Japan e Osaka University, 2-1, Yamada-oka, Suita, Japan f Mitsui Mining & Smelting Co., Ltd., 1333-2 Haraichi, Ageo, Saitama, Japan g Sumitomo Metal Industries, Ltd., 1-8 Fuso-cho, Amagasaki, Hyogo, Japan h JFE Steel Corporation, 1 Kokan-cho, Fukuyama, Hiroshima, Japan i Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai, Miyagi, Japan j Kobe Steel, Ltd., 1-5-5 Takatsukadai, Nishi-ku, Kobe k Hokkaido University, 8 Kita-13-jo-nishi, Kita-ku, Sapporo, Hokkaido, Japan Abstract - We investigated the atomic structure of the passive films on the porycrystalline substrates of pure iron by grazing incidence X-ray scattering, using synchrotron radiation at SPring-8. The X-ray scatteing data clearly showed the dependence of the environment on the crystalicity of the passive film. The atomic radial distribution function derived from the data of the passive film, which was formed by anodic polarization in a borate buffer solution, indicated that its atomic structure had the characteristics of a spinel-type iron oxide. Keyword: pure iron, passive film, grazing incidence X-ray scattering, synchrotron radiation

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1. Introduction The origin of the high corrosion resistance of stainless steel is attributed to a very thin passive film on its surface, stabilized by specific alloying elements, such as Cr, Ni, etc. The stabilization mechanism has not yet been elucidated, because the structure of the passive film is not clear. The reason for this is that the passive films are too thin (a few nm) for a standard X-ray diffraction study. Previously, Davenport et al. [1] performed a structural analysis of passive films on single crystalline substrates of pure iron with grazing incidence X-ray scattering (GIXS). They succeeded in obtaining precise data of the X-ray scattering from the passive films and concluded that the atomic structure of the passive films on their pure iron specimens was of a spinel-type iron oxide, like Fe3O4 or -Fe2O3. They called the structure as LAMM. The reason for their success was that the X-ray diffraction signals of the passive films from their specimens were easily separated from those of the single crystalline substrates, because the passive films on the single crystalline substrates were highly oriented. However, the single crystalline specimens of the material for practical uses were not easily obtained. Therefore, the developmental research on steels for practical use need an X-ray scattering measurement method that can be applied to passive films on polycrystalline substrates. We examined the capability of the GIXS for the investigation of the passive films on polycrystalline substrates of pure iron. 2. Experiment We tried to derive the atomic radial distribution function (RDF) from the GIXS data [2] to obtain the structural information of the passive films. We needed to reduce the background noise due to the diffraction from polycrystalline substrates by controlling the penetration depth of the incident X-ray beam to the substrates, because this analysis needs to obtain highly precise data of the X-ray scattering profile. In GIXS, we can control the penetration depth by controlling the glancing angle of the X-ray beam. If the glancing angle can be controlled precisely enough to cause a total reflection, the penetration depth should be suppressed to the order of thickness of the passive films (e.g.: a few nm). We employed the polycrystalline substrates of pure iron, mirror-finished by lapping, for the precise control of the glancing angle. In addition, we prepared two kinds of specimens in order to examine the dependence of the atomic structure on the environment; (a) as-polished with an air-formed film and (b) electrochemically passivated film. The as-polished specimen was exposed to the atmosphere after being polished by lapping. The electrochemically-passivated film was prepared as follows. The specimen was immersed in a borate buffer solution of pH= 8.4, then polarized at –1000 mVAg/AgCl for 30 min, in order to reduce preformed film. The applied potential was switched to +700 mVAg/AgCl, and then kept for 30 min

97

Atomic-Structure Characterization of Passive Film of Fe

at the potential to form a passive film. The thickness of the passive films was estimated at about 3 nm by X-ray reflectivity measurements. The experiments were carried out using the multi-axis diffractmeter at BL46XU of SPring-8. The light source of this beamline is an undulator. The energy of the X-ray beam we used was 12 KeV. We fixed the glancing angles at 0.1 degrees, which is below the critical angle of the total X-ray reflection of the pure iron substrate ( c= about 0.25 degrees). In order to achieve a high S/N ratio of data, we set the samples in a He chamber for prevention of the background noise from air scattering. (a) Naturally formed film (b) Electrochemically Passivated film

3. Results Intensity (a.u.)

(110)

Iron substrate( a-phase) (200) (211) (220)

10

20

30

50

40

(310)

70

60

2q(degree)

Fig.1 X-ray scattering data from surface of specimens (a) (dashed line) and (b) (solid line). (b)Electrochemically Passivated Film

Data Form Factor BG

Intensity(a.u.)

Figure 1 shows the X-ray scattering profiles from the specimens (a) and (b). Although the sharp diffraction peaks from the substrates were not perfectly suppressed, we succeeded in clearly observing X-ray scattering peaks from the passive films. Comparing the data, the profile of the X-ray scattering intensities from the specimen (b) film is broader than that of specimen (a). It indicated that the crystallinity of the electrochemically passivated film was lower than that of the film on the as-polished specimen. The reason for this outcome was considered to be that the film on the as-polished specimen was formed by precipitation in the solution used for the polishing. Therefore, the aspolished specimen was regarded as unsuitable for the study on the intrinsic characteristics of the passive film. We decided to analyze the RDF of specimen (b). Figure 2 shows the X-ray scattering data of specimen (b) (electrochemically-passivated film) as the function of wave vector Q, which is corrected by the instrumental geometrical factor

0

2

4

-1

6

8

10

Q(Å )

Fig.2 X-ray scattering data of passive film of iron formed by anodic polarization (specimen (b)) corrected by subtraction of diffraction peaks from substrate. The broken line and the hatched area are estimated form factor and background, respectively.

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and the subtraction of the diffraction from the substrate. The scattering intensities are written by the function as in equation (1). Intensity (Q ) = A × S (Q ) × ∑ f i 2 (Q ) + Background

,

(1)

QI(Q)

where A is the scaling parameter, fi2(Q) is the atomic form factor per unit of composition, and S(Q) is the structure factor. The form factor was approximated, assuming the composition of the film as Fe1.9±0.2O3 on the basis of the structure model reported by Davenport et al. [1], because we have no information on the composition of the film. The function of the background was assumued to be BiQi+C1exp(-C2Q) (i=0~3). The coefficients of A, Bi, C1 and C2 in equation (1) were estimated by the least (b)Electrocehmically Passivated Film square optimization. The form 4 factor weighted by A and the background estimated by the 2 above procedure were shown as 0 the broken line and the hatched area in Fig. 2, respectively. -2 Figure 3 shows the interference function I(Q) (= S(Q)-1 ) -4 weighted by Q which was derived by equation (1). Taking 10 4 8 0 2 6 -1 into consideration that this film Q(Å ) has no prefered orientation, its Fig. 3 Interference function weighted by wave RDF, D(r), can be derived from vector, QI(Q) estimated from X-ray scattering data the Fourier transform of Q·I(Q) of electrochemically passivated film (specimen (b)) with equation (2).

∑f ∑Z × f 2

D(r ) = 4p r 2 r 0 +

i

i

2 e

×

2r

p

∫

¥ 0

Q × I (Q) × sin(Qr )dQ ,

(2)

where Zi is the total number of electrons per unit of composition, fe is the mean atomic form factor per electron (fe = fi / Zi), and 0 is the mean electron density. We discuss the structure of the films by the differential radial distribution function (dRDF) as follows, H(r) = D(r)-4

r2

0

, (3)

which is the differential between RDF and 0, because we have no information of 0. Figure 4 shows the dRDF derived by equation (2). The dRDF has 3 peaks at r = about 2 Å (peak A), 3 Å (peak B), and 3.5 Å (peak C), indicating the existence of the nearest, the second nearest, and the third nearest neighboring

99

Atomic-Structure Characterization of Passive Film of Fe

atomic pairs, respectively. The peaks in the region below r = about 1.5 Å (the shaded area) are regarded as the noises of the estimation of S(Q).

H(r) (electron/Å)

200

(b)Electrochemically Passivated Film B

A1

100

C

A2

A 0

-100

-200

0

1

2

3

4

r(Å)

Fig. 4 Differential radial distribution function derived by Fourier transformation of QI(Q) of electrochemically passivated film (specimen (b)).

4. Discussion As shown in Fig. 4, the first peak A of the dRDF for the passive film of specimen (b) can be divided into two peaks at r = about 1.9 Å (peak A1) and 2.2 Å (peak A2). The atomic structure of the spinel-type iron oxide (e.g.: Fe3O4, Fe2O3) is constituted by two kinds of units, octahedrons and tetrahedrons centered by Fe atoms surrounded by O atoms. The Fe sites of these units are 16d site (FeO) and 8a site (FeT) in Fd 3 m , respectively. Therefore, it is characterized by the existence of two kinds of Fe-O bonds. Peaks A1 and A2 of the dRDF from specimen (b) can be reflected in these characteristics of the spinel-type structure. Figure 5 compares the dRDF of specimen (b) with the simulation of the dRDFs for Fe3O4 and the LAMM model proposed by Davenport et al. [1]. For these simulations, we approximated the peak widths of the dRDF at 0.4 Å (FWHM). In the dRDF of Fe3O4, one can see two peaks corresponding to the two kinds of Fe-O bonds, FeO-O and FeT-O at r = about 1.75 Å and 2.1 Å (pointed by arrows in Fig. 5). The LAMM model has the modified spinel structure, which has vacancies at the FeT and FeO sites, and the Fe atoms partially occupying the octahedral interstitital sites (16c sites in Fd 3 m : FeI). The occupancies of Fe sites are 66± 10% for FeT, 80± 10% for FeO and 12± 4% for FeI, respectively. The distances of the atomic pairs of the FeI site (FeI-FeI and FeI-O) are indicated by the arrows in the simulated dRDF of the LAMM in Fig. 5. The dRDF of the LAMM resembles that of Fe3O4 in spite of a little difference in the ratio of their peak intensities. Comparing the dRDF of specimen (b) with these simulated dRDFs, the dRDFs of Fe3O4 and the LAMM have peaks corresponding to peaks A1, A2, B, and C of specimen (b). This feature suggests that the structure of the passive film of specimen (b) has

M. Sato et al.

100

5. Conclusion

1000

Fe3O4 simulation 800

FeO-O

H(r) (electron/Å)

the characteristics of a spineltype structure. However, the ratios of the peak intensities of the dRDFs were quite different from each other. If these differences are intrinsic, it possibly suggests the difference in the occupancies of the Fe-sites originated in the differences in the sample conditions, for example, the experimental condition (exsitu vs. in-situ). We cannot examin the validity of this suggestion in this paper, because the S/N ratio of the X-ray scattering data was not sufficient enough to quantitatively discuss the amplitude of the dRDF as shown in Fig. 2.

600

400

FeT-O

FeI-O

FeI-Fe I

200

LAMM simulation A1

A2

B

0

C

Experimental data

-200

(b)Electrochemically passivated film

2.0

r(Å)

3.0

4.0

Fig. 5 Results from simulations of dRDFs for Fe3O4 and LAMM compared with experimental results of electrochemically passivated film.

We succeeded in analyzing the atomic structures of passive films on pure iron (about 3 nm thick) by the GIXS measurements. The dRDF derived from the data clearly indicated the characteristics of a spinel-type structure, which has two kinds of Fe sites. However, there was not enough quality GIXS data for a quantitative analysis on the evidence of the vacancies of Fe sites and the interstitial Fe sites, which were proposed by Davenport et al. [1] Improvements in this technique for a quantitative analysis will require, (1) information on the composition and density of the passive film, and (2) the achievement of a high S/N quality of the GIXS data. We will carry on with the improvements in the structural analysis method of the passive films by GIXS, combining them with other techniques. References 1. A. J. Davenport, L. J. Oblonsky, M. P. Ryan, and M. F. Toney, J. Electrochem. Soc. 146, 2162, (2000) 2. M. Sato, T. Matsunaga, T. Kouzaki, and N. Yamada, Mat. Res. Soc. Symp. Proc. 803, 245

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

101

Electrochemical aspects of the behavior of perchlorate ions in the presence of iron group metals G.G. Lánga, T.A. Rokoba, M. Ujvaria , G. Horányib a

Department of Physical Chemistry, Eötvös Loránd University, PO Box 32, Budapest 112, H-1518, Hungary, e-mail: [email protected] b Research Institute of Materials and Environmental Chemistry, Chemical Research Center, Hungarian Academy of Sciences, PO Box 17. Budapest, H-1525, Hungary Abstract: The problem of the stability of perchlorate ions in the presence of iron group metals is discussed. It is suggested that electrochemists should consider the reduction process as important evidence that should be taken into account. It has been shown that the mechanical properties of the passive layer on nickel are influenced by the presence of Cl- ions, presumably because chloride ions incorporated in the passive film facilitate the formation of a more ordered structure. Keywords : Perchlorate ions; Iron Group Metals; Reduction; Chloride ions; Film Stress

1. Introduction From thermodynamic point of view ClO -4 ions at solution/metal electrode interfaces, mainly in acidic medium, should be instable against reductive attacks in a wide potential range [1]. In contrast to this, it is a general view in the electrochemical literature that ClO -4 ions are very resistant to reduction. Therefore, perchloric acid and sodium perchlorate are widely used as supporting electrolytes in electrochemical studies with various electrodes. Among these investigations reports concerning dissolution, deposition, passivation and corrosion of iron group metals can also be found. (See refs [2-4] and literature cited therein.) However, it has been already demonstrated that the reduction of ClO -4 ions takes place during the corrosion of Co, Ni and Fe in HClO4 solutions [5-12]. A survey of the literature can be found in reviews [5,13].

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Besides the significance of perchlorate reduction for the fundamental electrochemistry nowadays a very practical reason, the so-called perchlorate contamination challenge, came into foreground orienting the attention towards the reductive elimination of perchlorate ions [14-16]. It follows from these considerations that the problem of perchlorate reduction constitutes an intersection of several branches of sciences: electrochemistry (anodic dissolution of metals, discharge of protons); corrosion science (coupling of cathodic and anodic processes, passivity of metals); catalysis (electrocatalysis); heterogeneous chemical processes. 2. Experimental Analytical grade Fe, Ni and Co metal powders (Fluka), and metal chips (Aldrich) were used for radiotracer adsorption studies and for analytical measurements, Fe, Co and Ni foils (Aldrich) and rods (Johnson Matthey) for the corrosion potential measurements. Fresh bottles of reagent grade perchloric acid, and ‘‘Millipore’’ grade water was used to prepare all solutions. For the preparation of solutions containing sodium perchlorate reagent grade perchloric acid (Merck), and ‘‘Suprapur’’ sodium-hydroxide monohydrate (Merck) or sodium perchlorate monohydrate was used. The formation of Cl - ions during the corrosion of metals was followed by potentiometric titration using a silver chloride electrode (detection limit: 2´10-5 mol dm-3 Cl - ) [8,9]. The measurement of the Cl - concentration was combined with the determination of the amount of metal consumed, measuring the mass of the metal before and after the corrosion. In addition, in the case of Co the Co2+ concentration in the solution was determined with a spectrophotometer at l = 511 nm. In some cases the volume of H2 gas evolved during the dissolution process was also measured with the help of a gas burette. The principle of the radiotracer method used is the measurement of radiation intensity originating from labeled adsorbed species on a powdered metal layer sprinkled on a thin plastic foil that serves simultaneously as the window for radiation measurement. The measurements were carried out at ambient temperature in an Ar atmosphere. 36Cl labeled HCl was used for study of the chloride adsorption [17,18]. The changes of the surface stress, D , in a thin metal plate (serving as a working electrode) was estimated from the changes in the bending of the plate (only one side of the plate was in contact with the electrolyte solution, the other side of it was covered with a nonconducting film). For the calculation of D the changes of the radius of curvature of the plate D (1 R) should be known. The values of D (1 R) = D ki (ki is a constant) can be calculated if the changes D of the deflection angle ( ) of a laser beam mirrored by the metal are measured by using an appropriate experimental setup [19,20].

Electrochemical aspects of the behavior of perchlorate ions

103

3. Interaction of ClO -4 ions with iron group metals From electrochemical point of view the interaction of a metal with dissolved perchlorate ions should be considered as a corrosion process involving three charge transfer processes Me ® Me 2+ + 2e (1)

H + + e- ® 1 2 H 2 ClO -4 + 8H + + 8e - ® Cl - + 4H 2 O

(2)

(3) This last step should be a very complex one composed of several elementary steps. The occurrence of the reduction process is well reflected by the changes of rest potential values at a fixed pH with increasing perchlorate concentration. This is demonstrated for Co in Fig. 1.

Fig.1 Mixed (corrosion) potential of Co measured in solutions: (1) 0.1 mol dm-3 HClO4 (2) 1 mol dm-3 HClO4 (3) 0.1 mol dm-3 HClO4 + 0.9 mol dm-3 NaClO4 (4) 0.1 mol dm-3 HClO4 + 2.9 mol dm-3 NaClO4 (5) 3 mol dm-3 HClO4 (6) 1 mol dm-3 HClO4 + 2 mol dm-3 NaClO4. The potential values indicated in the figure correspond to the maxima of the rest potential vs. time curves as shown in the insert.

It was found that the kinetics of the anodic dissolution of iron, cobalt and nickel are quite similar and a common mechanistic picture can be used for the interpretation of phenomena occurring with these metals. In Table 1 some data for metal-perchlorate (perchloric acid) interaction are summarized in order to estimate the relative role of ClO -4 reduction in the overall dissolution process of Ni, Fe and Co. These data were obtained using metal powders thus they

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G.G. Láng et al.

should be considered as average values as during the dissolution process the parameters of the systems are changing continuously. -

Table 1 Interaction of Ni, Fe, Co with ClO 4 ions at T = 25°C Electrolyte solution

Ni

Fe

Co

nA / mol

nB / mol (amount of Cl- formed during the dissolution of metal)

nB nA

Conversion* /%

cHClO4 /

cNaClO4 /

-3

mol·dm

mol·dm-3

(amount corresponding to the mass of dissolved metal)

3

0

3.57´10-2

7.5´10-5

2.1´10-3

0.84

1

0

1.26´10

-2

-5

-3

0.80

0.1

0.9

2.6´10

2.0´10

4.48´10-4

2.92´10-5

6.52´10-2

26.1

-3

-5

-2

1

0

1.27´10

0.1

1.9

11.73´10-4

1

0

9.48´10

-4

5.75´10

4.53´10

18.1

9.54´10-5

8.1´10-2

32.5

-5

-2

20.1

4.95´10

5.2´10

-

* The percentage of dissolved metal involved in ClO 4 reduction.

In experiments concerning the adsorption of labeled Cl - ions on powdered Fe in 0.5 M NaClO 4 solution at low pH values, for instance, pH » 2 (by addition of HClO4) the count-rate vs. time curves (curves 1, 2 and 3 in Fig. 2) going through maximum were obtained. The decrease in the count-rate means that no equilibrium or steady state is attained with respect to the surface concentration of the isotope emitting the radiation. This phenomenon can be explained by the production of inactive Cl - ions (i.e. by the decrease of specific activity of Cl - species). The displacement of the active species by non-active ones occurs immediately at the surface and this could be the very reason for the rapid decrease of the count-rate in the case of curve 1. As it can be seen in Fig. 3 the mechanical properties of the passive layer on nickel are also influenced by the presence of Cl - ions. The shape of the voltstressogram recorded in 0.1 M perchloric acid solution (curve 1) changes substantially if the solution contains 10-4 M Cl - ions (curve 2). The increase in the deflection is more pronounced at higher Cl - concentrations. It is known from the literature, that chloride ions are incorporated in the passive film on nickel, if the film is formed in Cl - -containing solutions [21,22]. Since the change of the stress is positive (tensile stress), it can be assumed, that Cl - ions incorporated in the passive film facilitate the formation of a more ordered structure.

Electrochemical aspects of the behavior of perchlorate ions

105

Fig.2 Count-rate vs. time curves obtained following the addition of 0.25 ml 1 mol·dm-3 HClO4 to a solution of 0.5 mol·dm-3 NaClO4 +1·10-4 mol·dm-3 labeled Cl- in contact with 0.5 g Fe powder. (1) first run, (2) second run, (3) third run.

Fig.3 Changes in the radius of curvature of a Ni plate (one side of it covered with a Teflon® foil) and the current during potential sweep in a 0.1 mol·dm-3 perchloric acid solution (curves 1 and 1A), in the same solution containing 10-4 mol·dm-3 Cl- ions (2, 2A), and in the same solution containing 2.5·10-4 mol·dm-3 Cl- ions (3, 3A). Sweep rate: 3 mV·s-1.

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4. Conclusions The reduction of perchlorate ions to Cl- ions in contact with iron group metals is so pronounced that it seems inevitable to revise many statements in text-books and monographs dealing with systems containing iron group metals and perchlorates. On the basis of film stress measurements it has been shown that the mechanical properties of the passive layer on nickel are influenced by the presence of Cl- ions. It can be assumed, that chloride ions incorporated in the passive film facilitate the formation of a more ordered structure. ACKNOWLEDGEMENT Financial support from the Hungarian Scientific Research Fund is acknowledged (Grants OTKA T037588, M042115; T045888).

References 1. G. Charlot; A. Collumeau, M.J.C. Marchon, Oxidation-Reduction Potentials of

Inorganic Substances in Aqueous Solution, Butterworths, London, 1971. 2. F. Zucchi, M. Fonsati, G. Trabanelli, J. Appl. Electrochem. 28 (1998) 441. 3. S.Y. Zhao, S.H . Chen, H.Y. Ma, D.G. Li, F.J. Kong, J. Appl. Electrochem. 32

(2002) 231. 4. H.Y. Ma, G.Q. Li, S.C. Chen, S.Y. Zhao, X.L. Cheng, Corros. Sci. 44 (2002) 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.

1177. G.G. Láng, G. Horányi, J. Electroanal. Chem. 552 (2003) 197. A.M. Lecco, V. Canić, Glasnik Hem. Drustva Beograd 14 (1949) 249. V. Canić, Glasnik Hem. Drustva Beograd 16 (1951) 13. M. Ujvári, G. Láng, G. Horányi, J. Appl. Electrochem.31 (2001) 1171. M. Ujvári, G. Láng, G. Horányi, J. Appl. Electrochem.32 (2002) 581. M. Ujvári, G. Láng, G. Horányi, J. Appl. Electrochem. 32 (2002) 1403. G. Láng, M. Ujvári, G. Horányi, Corrosion Sci. 45 (2002) 1. G.G. Láng, A. Vrabecz, G. Horányi, Electrochem. Comm. 5 (2003) 609. G. Horányi, in Catalysis; J.J. Spivey, Ed.; A Specialist Periodical Report; The Royal Society of Chemistry, Cambridge, GB. 1996; Vol. 12, pp. 254-301. Perchlorate in the Environment; E.T. Urbansky, Ed.; Kluwer Academic Publisher, Dordrecht, 2001. E.T. Urbansky, Environ. Sci. Pollut. Res. 9 (2002) 187. E.T. Urbansky, M.R. Schock, J. Environ. Manage. 56 (1999) 79. G. Horányi, Radiotracer studies of adsorption at electrode surface, in: A. Hubbard, Ed., Encyclopedia of Surface and Colloid Science, Marcel Dekker, New York, 2002, pp. 4423–4437. G. Horányi, Corros. Sci. 46 (2004) 1741. G.G. Láng, K. Ueno, M. Ujvári, M. Seo, J. Phys. Chem. B 104 (2000) 2785. G.G. Láng, M. Seo, J. Electroanal. Chem. 490 (2000) 98.

18. 19. 20. 21. J.M. Herbelin, N. Barbouth, P. Marcus J Electrochem. Soc 137 (1990) 3410. 22. P. Marcus, J.M. Herbelin Corros. Sci. 34 (1993) 1123.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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The effect of Al3+ in the passivity of iron in alkaline media containing chlorides C.M. Abreu, M.J. Cristóbal, L. Freire, X.R. Nóvoa*, G.Pena, M.C. Pérez Universidade de Vigo, E. T. S. E. I., Campus Universitario, 36310 Vigo. Spain *Corresponding author: Tel.: +34986812213, email: [email protected]

Abstract - The behaviour of iron in alkaline media containing chlorides is studied in presence and absence of aluminate anions. The presence of aluminium in solution inhibits pitting of iron at least up to Cl- to OH- ratio equal to 5. This effect is due to lower cathodic reactivity. Keywords: iron corrosion, alkaline medium, pitting inhibition, aluminates, voltammetry 1. Introduction The Bayer Process of alumina (Al2O3) production from bauxite ore generates an alkaline iron-rich waste called red mud (RM). Those wastes have been employed for making bricks, for land recovering (due to its alkalinity), and also as absorber for heavy metallic cations [1] and as steel corrosion inhibitors [2]. The encouraging results obtained for iron corrosion inhibition in RM solutions containing chlorides leads to study the inhibiting effect of the slurry’s components. The speciation of RM’s shows Bayer-sodalite content in the range 10-20%. This is an aluminium compound in which aluminium is present as Al2O3. So, the aim of the present paper is to study the effect of soluble aluminium in the corrosion behaviour of iron in alkaline media. Aluminium is soluble in acid media as Al3+ and in alkaline media as aluminate (complexed Al3+, either as the hydrated Al(OH)-4 or anhydrous AlO -2 forms). The minimum solubility occurs at pH 5, for more alkaline pH values the dominant form is aluminate. So, the amount of aluminates in solution is pH dependent and ranges between the limiting values obtained from equations 1 and 2, that correspond to the more soluble form Al(OH)3 and more insoluble one, hydrargillite (Al2O3.3H2O), respectively[3]. Log( Al(OH) -4 )= -10.64 + pH

(1)

Log( AlO -2 )= -14.6 + pH (2)

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2. Experimental Solutions were prepared by dissolving Al foils in NaOH 0.1M to reach the solubility limit, that was about 13.5 mg.L-1. So, the employed solutions were 0.5mM in Al3+. The tested material was pure iron (99.9%) obtained from Goodfellow as rods (f= 5 mm). The exposed area was the cross section (0.2 cm2), freshly polished and rinsed surface. Polishing was done using SiC paper down to 1200 grade. The testing solutions were 0.1M NaOH as reference, and 0.1M NaOH solutions containing Al3+ and different additions of Cl- (as NaCl). Solutions were prepared the day of testing (to minimise carbonation) using distilled water and reagent grade chemicals. The electrochemical cell was a conventional three electrode arrangement: the tested sample as working electrode, platinum gauze as counter electrode, and an Hg/HgO 0.1M KOH as reference electrode. When chloride ions were added to the solutions, a saturated calomel electrode (SCE) was employed. All electrode potentials in the text are referred to SCE. The electrochemical tests were performed immediately after surface preparation at room temperature and open to the laboratory atmosphere, using an AUTOLAB 30 Potentiostat (from Ecochemie, NL). Different electrochemical techniques were employed: Potentiometry, Cyclic Voltammetry and Electrochemical Impedance Spectroscopy. The potential region for cyclic Voltammetry tests was scanned at 1 mV s-1 scan rate, which allows relaxation of the redox process taking place in the passive layer [4,5]. The electrochemical generated film was studied by Optical Microscopy, SEM, and XPS. 3. Results and discussion 3.1. Cyclic voltammetry Figure 1 corresponds to the first cycle of the voltammogram for the iron electrode in a blank solution (0.1 M NaOH) and the same solution added Al3+. It can be noted that the blank solution shows the usual shape [4] with the activity peak at about -1V corresponding to the Fe/Fe2+ process and the magnetite formation peak at about -0.8V. Nevertheless, in presence of aluminates both peaks vanish, the current of passivity is smaller, and the potential of O2 evolution shifts anodically. It seems that aluminates block to some extent the iron surface. In Figure 2 chlorides were added to the solution to reach Cl- to OH- ratio = 2. It can be noted that again the activity peak current is lower for the aluminatescontaining solution. Moreover, in the passivity region, current is steady in the later case while in the former current transients typical of pitting/repassivation events are visible. Thus, again aluminates seem to be effective in blocking the metallic surface, in this case preventing pitting events.

The effect of Al3+ in the passivity of iron in alkaline media containing chlorides

109

In Figure 3 chloride concentration was increased to reach Cl- to OH- = 5. Now the iron passive film breaks down at 0.15V and the reverse curve shows a large current loop typical of pit growth. Nevertheless, in presence of aluminates current remains steady and the reverse curve corresponds to lower current values than those observed for the forward scan. This behaviour is typical of passive film thickening during the forward scan.

Figure 1: Comparison of the first voltammetric cycles obtained for iron electrode in 0.1M NaOH + 0.1M NaCl solution with and without 0.5 mM Al3+. Scan rate = 1 mV.s-1.

Figure 2: Comparison of the first voltammetric cycles obtained for iron electrode in 0.1M NaOH + 0.2M NaCl solution with and without 0.5 mM Al3+. Scan rate = 1 mV.s-1.

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Another interesting feature to be discussed is the free corrosion potential evolution. Ecorr shifts from −0.216 V to −0.339 V when chloride concentration is increased from 0.2M to 0.5M. In presence of Al3+ potential shifts cathodically too but being always lower than that corresponding to the solutions without Al3+. Ecorr in presence of aluminates are −0.309 V and −0.400 V, respectively for the solutions 0.2M and 0.5M in NaCl. Potential shifts in the cathodic direction in presence of chlorides probably due to activation of the anodic process. Nevertheless, the fact that corrosion potentials are lower in presence of Al3+ suggests lower local oxygen concentration due to surface blockage. Lower electrode potential corresponds to higher resistance of passive steel to chloride attack [6], probably due to the higher hydration degree of the passive film [4], that decreases locally the Cl− to OH− ratio.

Figure 3: Comparison of the first voltammetric cycles obtained for iron electrode in 0.1M NaOH + 0.5M NaCl solution with and without 0.5 mM Al3+. Scan rate = 1 mV.s−1.

3.2. Electrochemical Impedance Spectroscopy Figure 4 corresponds to the EIS data obtained at the respective corrosion potential for the solutions containing NaCl 0.5M before visible pitting. The experimental data have been modeled using an equivalent circuit with two hierarchically distributed RC time constants. The corresponding impedance is given in equation 3: the meaning of the model parameters has been discussed elsewhere [4, 7] and is as follows: Re, the high frequency resistance, corresponds to the electrolyte resistance; C1 is associated to the double layer capacitance and R1 to the charge transfer resistance. The impedance Z2 is associated to a redox process occurring in the passive layer (between magnetite and iron (III) oxides). The α parameters account for the Cole−Cole dispersion

The effect of Al3+ in the passivity of iron in alkaline media containing chlorides

111

of time constants. This model allows good quality fitting of experimental data using a simplex method of 2 minimisation [8]. The best fitting parameters for the data depicted in Figure 4 are given in Table 1. R1 R2 (3) Z (w ) = R e + ; Z2 (w ) = a1 ( jw R1C1 ) + 1/(1 + (Z2 (w ) / R1 )) 1 + ( jw R2C2 )a2 The phase angle depicted in Figure 4B shows visually that the main difference between the presence or not of aluminates in the solution concerns the low frequency domain, that related to redox activity in the passive film. 35 NaOH 0.1M + NaCl 0.5M 3+ NaOH 0.1M + NaCl 0.5M + Al 0.5mM

30

A -Imaginary Part/ kW.cm

2

25 20

0.01 Hz

15 10

0.1 Hz 0.1 Hz 0.01 Hz

5

1 mHz

0

0

5

10

15

20

25

Real Part / k W. cm

30

35

2

100 80

B |Z|/ kW cm

2

60

1 40

0.1

20

0.01

1E-3 -4 10

NaOH 0.1M + NaCl 0.5M 3+ NaOH 0.1M + NaCl 0.5M + Al 0.5mM

10

-3

10

-2

10

-1

10

0

10

1

10

2

10

3

10

4

-Phase angle/ degree

10

0

10

5

Frequency /Hz

Figure 4: Nyquist (A) and Bode (B) plots of the impedance obtained for iron electrode in 0.1M NaOH + 0.5M NaCl solution with and without 0.5 mM Al3+. Spectra obtained at the corresponding Ecorr: (-0.34 and -0.40) Vvs. SCE, in absence and presence of Al3+, respectively.

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The Nyquist plot (Figure 4A) reveals also the great difference in overall impedance induced by the presence of aluminates. This difference is mainly due to the redox process in the oxides layer, as revealed by R2 values in Table 1. Figures 2 and 3 show that in presence of chlorides pit initiates and then an increase in surface roughness shall be expected, which will increase the double layer capacitance. Nevertheless, the double layer capacitance increases in presence of aluminates (see C1 in Table 1) which can suggest charge accumulation at the interface due to adsorption of aluminates that will hinder mass and charge transfer. Moreover, C2 is much smaller for the solution containing aluminates, which suggests thinner passive film than for the solution without them. The idea of adsorption of aluminates is supported by SEM examination and XPS depth profiles that show neither aluminium incorporation into the passive film nor structural differences with reference samples. Table 1: Best fitting parameters corresponding to data in Figure 4 using equation 3 as model function.

R1/ Re/ C1/ W.cm2 kW.cm2 mF.cm-2 NaOH 0.1M + NaCl 0.5M NaOH 0.1M + NaCl 0.5M + Al3+ 0.5mM

a1

R1/ C2/ kW.cm2 mF.cm-2

a2

3.2

7.6

43.1

0.9

16.6

1.9

0.6

3.1

0.3

53.7

0.9

92.5

0.3

0.5

4. Conclusion The above discussed results allow concluding that the presence of Al3+ is able to inhibit pitting of iron in chlorinated alkaline media. These preliminary results suggest that inhibition is due to adsorption of aluminates blocking the interface Acknowledgement The authors whish to acknowledge the Spanish “Ministerio de Educación y Ciencia” for financial support under project MAT2004-06435-C02-01. References [1] E. López, B. Soto, M. Arias, A. Núñez, A., D. Rubinos, M.T. Barral, Water Res. 32(1998) 1314. [2] B. Díaz, S. Soiret, M. Keddam, X.R. Nóvoa, M.C. Pérez, H. Takenouti, Electrochim. Acta 49 (2004) 3039. [3] M. Pourbaix, Atlas d’équilibres électrochimiques, Gauthier-Villars Ed., Paris, 1963, p. 170. [4] S. Soiret, M. Keddam, X.R. Nóvoa, M.C. Pérez, C. Rangel, H. Takenouti, Cem. Concr. Composites, 24 (2002) 7. [5] X.R. Nóvoa, M.C. Pérez, Corros. Rev., 23 (2005) 195. [6]C. Alonso, M. Castellote, C. Andrade, Electrochim. Acta 47 (2002) 3469. [7] C. Andrade, M. Keddam, X.R. Nóvoa, M.C. Pérez, C.M. Rangel, H. Takenouti, Electromchim. Acta 46(2001) 3905. [8] C. Abreu, M. Izquierdo, M. Keddam, X.R. Nóvoa, H. Takenouti, Electromchim. Acta 41(1996) 2405.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Nanoscale modifications of a Ni(111) surface during nucleation and growth of the passive film A. Seyeux, V. Maurice, L. H. Klein, P. Marcus Laboratoire de Physico-Chimie des Surfaces, CNRS / ENSCP (UMR # 7045), Ecole Nationale Supérieure de Chimie de Paris, 11 rue P. & M. Curie, 75005 Paris, France ; [email protected]

Abstract - EC-STM data on the nucleation and growth of the passive film on Ni(111) in 0.1 M NaOH + x M NaCl (0 x 1) are reported. Nucleation at low over-saturation potentials is characterized by the absence of dissolution and the preferential growth of 2D nuclei of the passivating oxide at the surface defects (i.e. step edges). This process is not modified by the presence of Cl- despite indications of a competitive adsorption mechanism between Cl- and OH-. At increased over-saturation, the growth of a crystalline 3D hydroxide layer with a typical facetted topography is observed in the outer part of the passive film. In the presence of Cl-, a nanogranular film with an increased density of grain boundaries is formed, indicating a marked inhibiting effect of Cl- on the coalescence and crystallization of the passive film. Keywords : growth mechanism, structure, nickel, oxide/hydroxide, chlorides, EC-STM

1. Introduction The crystalline structure of the passive film formed on nickel is well documented [1-11]. In acidic solution, the NiO inner layer of the duplex film has been observed to grow along the [111] direction in anti-parallel epitaxy on Ni(111). The (111) orientation of NiO is stabilized by surface hydroxylation in the outer layer. The surface of the inner NiO(111) layer is facetted as a result of a tilt of the oxide lattice with respect to the substrate lattice [3,4,6,8,9]. In alkaline solutions, a crystalline and facetted lattice of hexagonal symmetry has been observed on Ni(100) [5] and Ni(111) [10,11], and assigned to the growth of the 3D outer layer of -Ni(OH)2(0001) in tilted epitaxy [11].

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Nucleation and growth of the passive film have been much less studied. This is due to the intense dissolution occuring in acid solution at low-oversaturation that strongly modifies the substrate structure at the sub-nanometer scale and prevents its detailed study. In alkaline solution, where dissolution is minimized, a distorted structure, observed prior to the growth of the 3D Ni(OH)2(0001) hexagonal phase [5,11], has been assigned to a strained 2D hydroxide layer formed in the initial stages of growth [11]. Here we report new data on the initial stages of growth of the passive film obtained in 0.1 M NaOH(aq.). The (sub-)nanometer scale structural modifications of a Ni(111) surface were investigated with in situ Electrochemical Scanning Tunneling Microscopy (ECSTM). The effect of chlorides has been studied. 2. Experimental Ni(111) single-crystals were used. The surface was mechanically and electrochemically polished, and subsequently annealed at 1000°C in a flow of ultra-pure hydrogen at atmospheric pressure. The sample was then transferred in air to the electrochemical cell of the STM. The measurements were performed in 0.1 M NaOH(aq.) solutions, without or with Cl- (0.05 M, 0.1 M and 1 M NaCl), prepared from ultra-pure chemicals and Millipore water (resistivity > 18 M .cm). The EC-STM instrument was a Molecular Imaging system. The tungsten tips were prepared from a 0.3 mm diameter wire by electrochemical etching and isolated with Apiezon wax. All images were obtained in the topographic (constant current) mode. After cathodic reduction of the air-formed oxide to produce an oxide-free metallic surface, the potential was increased anodically to values corresponding to the onset or center part of the AI anodic peak corresponding to the Ni(0) to Ni(II) oxidation reaction [11]. 3. Results and Discussion Nucleation was investigated by characterizing the surface modifications occurring at the onset of the AI anodic peak. Recent measurements, performed in 1 mM NaOH(aq.), pH~11, E = -300mV/SHE [11], have shown that dissolution occurs preferentially at step edges and slowly consumes the terraces by a step flow process. This is the predominating process modifying the surface immediately after anodic polarization. No ordered superstructure was observed on the terraces, despite indications of the possible specific adsorption of hydroxide species. The subsequent development, preferentially localized at and near step edges, of 2D nanograins was assigned to the nucleation of the passive film. Increasing the potential causes the formation of 2D islands of nuclei of the passivating oxide, progressively covering the entire surface, and forming a crystalline 2D layer assigned to Ni(OH)2(0001) in strained epitaxy with the substrate [11].

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Nanoscale modifications of a Ni(111) surface…

At higher OH- concentration (0.1 M NaOH(aq.), pH~13, E = -550mV/ SHE), the mechanism is slightly different. Dissolution of the surface is not observed immediately after anodic polarization, indicating the blocking of this reaction at the step edges. After prolonged polarization, the development of islands with an hexagonal lattice with a parameter of 0.49±0.04 nm corresponding to a (2x2) superstructure, was observed on the terraces (Fig. 1a), indicating the formation of an ordered adlayer. These differences possibly reflect the effect of the higher concentration of OH- species on their adsorption at step edges and on the terraces. Figure 1b illustrates the development of nanograins (marked G) preferentially located at step edges. The grains are ~1 nm across and appear higher than the step edges (0.25 nm), quite similarly to the observations made previously for 1 mM NaOH(aq.) [11]. Consistently, these nanograins are assigned to nuclei of the passive film, likely consisting of Ni(II) hydroxide species.

G G G G

a

b

Figure 1 : EC-STM images of the Ni(111) surface recorded at –550mV/SHE in 0.1 M NaOH. (a) 5 nm x 5 nm, Etip=-150mV/SHE, Itip=1.3 nA, t=2600 s; (b) 50 nm x 50 nm, Etip=-950mV/SHE, Itip=1.3 nA, t=2800 s.

The growth of the passive film was investigated by characterizing the surface modifications induced by polarization in the center of the AI anodic peak. Fig.2 illustrates the progressive development of a facetted topography in 0.1 M NaOH(aq.). As in 1 mM NaOH(aq.), this is characteristic of the complete crystallization of the outer hydroxide part when a 3D passive film is formed. 3D growth is consistent with the fact that the substrate terraces can no longer be observed after passivation (Fig. 2b), in contrast with the observations made when a 2D layer is formed (Fig. 2a). The surface lattice could not be resolved due to the lack of resolution. At potentials higher than the AI anodic peak, it became extremely difficult to obtain stable images due to the decrease of the conductivity of the surface. These limitations of the STM imaging are assigned to an increase of the thickness of the passive film with increasing potential,

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already observed by XPS [12], causing a decrease of the tunneling probability of the electrons through the growing semi-conducting film from/or to the tip, via the surface states at the passive film/electrolyte interface.

a

b

Figure 2 : EC-STM images of the Ni(111) surface recorded at –450mV/SHE in 0.1 M NaOH. (a) 67 nm x 67 nm, Etip=-850mV/SHE, Itip=1.3 nA, t=40 s; (b) 100 nm x 100 nm, Etip=-850mV/SHE, Itip=1.4 nA, t=3800 s.

(2x2)

Figure 3 : EC-STM image of the Ni(111) surface recorded at –550mV/SHE in 10-1M NaOH + 10-1M NaCl (17 nm x 17 nm; Etip=-1200mV/SHE, Itip= 1.8 nA, t=2040 s).

In the presence of chloride ions, the formation of the passive film by a nucleation and growth mechanism is also observed. Figure 3 shows a typical surface topography obtained at the onset of the AI anodic peak for a concentration ratio [Cl-]/[OH-]=1. The step flow caused by preferential removal of the atoms at the step edges is not observed, indicating a blocked mechanism like in the absence of Cl-. The formation of (2x2) islands is observed on the terraces. This structure is similar to that observed in the absence of Cl- ions, which supports its assignment to the specific adsorption of OH- species on the surface. For [Cl-]/[OH-]=10, this superstructure could not be observed prior to the growth of the passive film. Although a lack of resolution of the imaging in

117

Nanoscale modifications of a Ni(111) surface…

this measurement cannot be fully excluded, the effect of the increasing [Cl]/[OH-] ratio suggests a competitive adsorption mechanism between hydroxides, resulting in the formation of the (2x2) ad-islands, and chlorides, forming a nonordered adlayer, as shown by the absence of superstructure on the terraces. No significant effect of the presence of Cl- ions was observed on the nucleation stage of the passive film. Irrespective of the Cl - concentration (in the investigated range), nucleation of the passive film was observed at potentials corresponding to the onset of the AI anodic peak. It was characterized by the growth of 2D nanograins forming small 2D islands with increasing time or potential. This result suggests the absence of blocking effect of chlorides on the nucleation of the passive film.

a

b

Figure 4 : EC-STM images of the Ni(111) surface recorded at –300mV/SHE in 10-1M NaOH + 10-1M NaCl. (a) 140 nm x 140 nm, Etip=-1100mV/SHE, Itip=2nA, t=800 s; (b) 140 nm x 140 nm, Etip=-600mV/SHE, Itip=2nA, t=1200 s.

In contrast, the addition of chloride in the solution affects significantly the 3D growth of the passive film. Indeed, the faceting of the surface, characteristic of the growth of a 3D crystalline film, observed at a potential corresponding to the top of the AI anodic wave (E=-450mV/SHE in 10-1M NaOH) in the absence of Cl- ions, was not observed in equivalent polarization conditions in the presence of chlorides. For [Cl-]/[OH-] = 0.5, the growth of the nanogranular and nonordered passive film is observed at E=-400mV/SHE. At E=-260mV/SHE, just above the AI anodic peak, the surface faceting characterizing the crystallization of the passive film is observed only locally. Polarization at higher anodic potentials (E=-50mV/SHE) is necessary to produce the crystalline 3D films with a typical facetted topography. For [Cl-]/[OH-] > 0.5, surface faceting was not observed even at higher potentials. Figure 4 shows images obtained on Ni(111) in 10-1M NaOH + 10-1M NaCl at E = -300 mV/SHE, i.e. above the AI anodic peak. The surface is roughened with formation of uncoalesced grain aggregates separated by local

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depressions. The difference of apparent height between the grain aggregates and the deepest depressions can reach 10 nm. This high value suggests that localized attacks of the metal have been initiated as the result of a local increase of the dissolution in the passive state. An inhomogeneous thickening of the passive film possibly also contributes to the surface roughening (thicker passive films have been reported to form in the presence of Cl- [13]). Thus it is observed that chlorides inhibit the coalescence of the nuclei of passive film, and coalescence is necessary for the formation of an homogeneous and crystalline film. As a result, a nanogranular film with a high density of grain boundaries is formed. 4. Conclusion The initial stages of growth of the passive film on Ni(111) in 0.1 M NaOH(aq.) have been studied with EC-STM. The nucleation stage is characterized by the absence of dissolution of the metal surface at low over-saturation potentials corresponding to the initiation of the Ni(0) to Ni(II) reaction. A (2x2) superstructure corresponding to the formation of adsorbed islands of OH- is observed. The growth of 2D nuclei is observed preferentially at the defects (i.e. step edges) of the metal surface. At higher over-saturation, the faceting of the passivated surface is observed. It is assigned to the formation of a crystalline 3D hydroxide layer. No marked effect of chlorides on the nucleation stage have been observed at low over-saturation except for the absence of ordered superstructure at [Cl-]/[OH-]>1, suggesting a competitive adsorption mechanism for OH- and Cl-. At higher-oversaturation, a marked inhibiting effect of chlorides on coalescence and crystallization produces a nanogranular film with an increase density of grain boundaries. References 1. J. Oudar & P. Marcus. Appl. Surf. Sci. 3 (1979) 48. 2. R. Cortes, M. Froment, A. Hugot-Legoff & S. Joiret, Corros. Sci. 31 (1990) 121. 3. V. Maurice, H. Talah & P. Marcus, Surf. Sci. 284 (1993) L431. 4. V. Maurice, H. Talah & P. Marcus, Surf. Sci. 304 (1994) 98. 5. S. -L Yau, F.-R. F. Fan, T. P. Moffat & A. J. Bard, J. Phys. Chem. 98 (1994) 5493. 6. T. Suzuki, T. Yamada & K. Itaya, J. Phys. Chem. 100 (1996) 8954. 7. D. Zuili, V. Maurice & P. Marcus, J. Electrochem. Soc. 147 (2000) 1393. 8. O. M. Magnussen, J. Scherer, B. M. Ocko & R. J. Behm, J. Phys. Chem. B 104 (2000) 1222. 9. J. Scherer, B.M. Ocko & O.M. Magnussen, Electrochim. Acta 48 (2003) 1169. 10. N. Hirai, H. Okada & S. Hara, Transaction JIM 44 (2003) 727. 11. A. Seyeux, V. Maurice, L. H. Klein & P. Marcus, J. Solid State Electrochem. 9 (2005) 337. 12. H.-W. Hoppe & H.-H. Strehblow, Surf. Interf. Anal. 14 (1989) 121. 13. P. Marcus & J.-M. Herbelin, Corrosion Sci. 34 (1993) 1123.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Simultaneous Anodic Dissolution and Passivation of Nickel in Moderate Acid Medium Joan Gregori, José Juan García-Jareño, David Giménez-Romero, Francisco Vicente* Department of Physical Chemistry, University of Valencia, C/ Dr. Moliner 50, 46100 Burjassot (Spain). *Corresponding autor; E-mail: [email protected]. Phone: 34963543022, Fax: 34963544564

Abstract - The EQCM results show that nickel electrodissolution and nickel passivation occur simultaneously in a sulphate acid media of pH = 3.5. Mass balances have been done from the instantaneous F(dm/dQ) function. The fitting of the experimental i = f(E) and –dm/dt = g(E) curves to the theoretical equations allow to obtain information about the kinetic parameters and the molecular mass of the species involved in the electrochemical processes. Keywords: Nickel; Anodic dissolution; Deposition; Passivation; EQCM

1. Introduction. It is commonly accepted that nickel electrodissolution proceeds according to two consecutive single-electron transfer steps [1-4]. But, in the case of nickel, several points are yet unclear due to the strong tendency for self-passivation of nickel [5]. This tendency should prove more pronounced in a weakly acid or alkaline media [6-7] and it is affected drastically to a great extent by the anions [8] and other experimental conditions [2,9]. These factors affect the recorded cyclic voltammograms for a nickel electrode to a great extent, where a well defined anodic peak represents the active-passive transition. At a relatively weakly acid pH, nickel electrodissolution and nickel passivation simultaneously take place at the potential range which corresponds to the anodic peak defined in the voltammetric curves. Electrochemical Quartz Crystal Microbalance

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(EQCM) measurements could allow to separate both processes since they imply opposite mass changes. The aim of this work is to perform an EQCM study of nickel electrodissolution at moderately acid medium. Also a quantitative study of the active-passive transition has been performed based on EQCM results in this potential range. 2. Experimental. All the experiments have been carried out in a typical three electrodes cell. The potential was measured versus a Hg/Hg2SO4/K2SO4(sat.) (saturated sulphate electrode, SSE) reference electrode. A platinum sheet of a relative large area (A = 2 cm2) was used as an auxiliary electrode. Solutions were prepared from 10-3 M NiSO4 (Scharlau, a.g.), 10-2 M H3BO3 (R.P. Normapur, a.g.), 10-4 M H2SO4 (Merk, a.g.), pH = 3.6, with distilled and double deionised water (MilliQ). In EQCM experiments the working electrodes have been made from a quartz sheet embedded between two pieces of gold connected to a resonance circuit. The resonance frequency of the quartz at air was 6 MHz. One of the pieces of gold acts as electrical surfaces in contact with the electrolyte. The electrical area was 0.228 cm2 and the mass sensitive area was 0.196 cm2. The potential sweep has been carried out in the [200,-1500] mV (vs SSE) potential range at 20 mV/s. Before starting the potential sweep the potential has been kept at 200 mV during 2 minutes. The microbalance was an UPR15/RT0100 (UPR of the CNRS). The resonance frequency of the quartz has been measured with a frequency-meter Fluke PM6685. The current in the auxiliary electrode has been measured with a multimeter Keithley PM2000. The potential has been applied with a Potentiostat 263A EG&G PAR. All the system has been controlled with a GPIB board. In all the experiments the temperature was T = 298 K. All the measurements have been performed under inert atmosphere. EQCM measurements in combination with voltammetry allow to define the F dm dQ function [10] or instantaneous mass/charge ratio as:

F dm dQ = ∑n i MWi ni

(1)

where the summatory is extended to all the processes which involves ni n i = 1 . The function electrons. ni is the charge fraction due to process i, and F dm dQ can be determined numerically at each applied potential. Equation (1) provides in situ information on the species which participate in the electrochemical reaction at each applied potential.

∑

3. Results and discussion. The voltammogram and the mass changes for the microbalance electrode in the [0.2,-1.5] V potential window are represented in figure 1.

Simultaneous Anodic Dissolution and Passivation of Nickel

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Figure 1 Voltammogram and mass variations in the potential range [200,-1500] mV for a nickel electrodeposition/stripping experiment. Scan rate 20 mV/s. 10-2 M H3BO3, 10-4 M H2SO4, 10-3 M NiSO4. T = 298 K. pH = 3.7.

In the [-1,-1.5] V and [-1.5,-1.25] V potential ranges, in the forward scan and in the backward one respectively, a considerable mass increase takes place. This mass increase is associated with nickel electrodeposition process. Moreover a well defined cathodic peak, peak I in figure 1, can be seen. This peak is not clearly associated with the electrodeposition process since it is not accompanied by an appreciable mass variation. In the [-0.675,-0.2] V potential interval a considerable mass decrease is observed. This mass decrease is accompanied by a broad anodic peak, peak II. Consequently this mass decrease is associated with electrodeposited nickel electrodissolution. The mass decrease and the calculated charge in nickel electrodissolution potential range are Dmdiss = 0.33 mg·cm-2 and DQ = 1.58 mC·cm-2, respectively. Then, the experimental mass/charge ratio value is F Dm /DQ = -20 g·mol-1. This value does not agree with the theoretical one F Dm /D Q th = -29 g·mol-1 expected for an electrodissolution process according such as Ni ® Ni2+ + 2e-. The reason for this is a side reaction which proceeds together with nickel electrodissolution: nickel passivation. Moreover a considerable mass gain of Dmgain = Dmdep Dmdiss = 0.12 mg·cm-2 is measured on the electrode surface after a complete cycle is finished. This mass gain and the discrepancies between F Dm /DQ values can be explained if a partial passivation of the nickel deposit takes place. It is considered that passivation proceeds according to Ni + 2OH- ® Ni(OH)2 + 2e- [11], although in passivity potential range a bilayer structure for passive film should be considered with an inner thin crystalline NiO layer and an outer thicker amorphous Ni(OH)2 one [12]. From F Dm /DQ definition the charge fraction due to nickel electrodissolution, gdiss, and passivation, gpass, can be calculated as:

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F Dm / DQ|exp = g diss F Dm / DQ| diss + g pass F Dm / DQ| pass

(2)

A value of gpass = 0.2 is obtained. Then the total charge associated with the passivation process is DQpass = 0.3 mC·cm-2 to which corresponds a mass increase due to OH- deposition of 0.05 mg·cm-2 in good agreement with the value 0.04 mg·cm-2 calculated from the Dmgain value. The analysis of the instantaneous mass/charge ratio, as represented in figure 2, shows that initially nickel electrodissolution is the main process, but at more anodic potentials F dm /dQ experimental values decreases as the passivation process becomes more important. All these experimental findings indicate that at the potential range of anodic peak II two electrochemical processes take place: k1 2+ Ni ¾¾® Ni aq + 2e -

(3)

k2 Ni + 2OH - ¾¾® Ni (OH ) 2 + 2e -

(4)

Figure 2. Instantaneous F(dm/dq) mass/charge ratio (full and open circles) and voltammetric curve (continuous line). Same experimental conditions than in figure 1.

The charge and mass balances at the electrode surface are, respectively:

i F / 2 FA = k1q 0 + k 2q 0

(5)

dq 0 / dt = -k 2q 0

(6)

where q0 is the surface concentration of nickel metal sites, k1 and k2 are the rate constants for reaction (3) and (4), respectively, where it is supposed that

Simultaneous Anodic Dissolution and Passivation of Nickel

123

ki=k0iebiE. In a linear sweep experiment the applied potential changes with time in such a way that E = Ei + vt, where v is the scan rate and Ei is the initial potential. Our attention is focused in the anodic peak II potential range. The integration of equation (6) with an appropriate initial condition (q0 = q0 for t = ti(Ei), where q0 is the maximum concentration of nickel sites on the electrode surface, and Ei sufficientely cathodic) allows to obtain an expression for the faradaic current: q0

∫q

E

dq 0 /q 0 = - ∫ k 2 dt ® q 0 = q 0 e - k 2 / b2n

(7)

i F / 2 FA = k1q 0 + k 2q 0 = (k1 + k 2 )q 0 e - k 2 / b2n

(8)

0

Ei

Moreover the mass variation rate can be expressed as [11]:

- dm / dt = (-2 M OH k 2 + M Ni k1 )q 0

(9)

where MOH- and MNi are OH- (supposedly responsible of passivation) and Ni molecular weights respectively.

Figure 3a. Experimental(circles) and fitted curves(continuous line) for measured current intensity at the nickel electrodissolution potential range. Figure 3b. Experimental and fitted curve for mass change rate numerically determined from measured mass changes in figure 1 at the nickel electrodissolution potential range. In both cases same experimental conditions as in figure 1.

The fitting of the experimental i = f(E) and –dm/dt = g(E) curves to the theoretical ones defined by equations (8) and (9) allows to obtain information about the kinetic parameters and the molecular mass of the species involved in the electrochemical processes. The fitted curves can be seen in figure 3a and figure 3b for the f(E) and g(E) curves, respectively. For this purpose a value of q0 = 3.7·10-9 mol·cm-2, determined from nickel density, is considered. In this

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way it is obtained that b1 = 13 V-1, b2 = 11 V-1, MOH- = 17 g·mol-1, and MNi = 68 g·mol-1. The obtained value for MOH- is close to that expected from theory. 4. Conclusion. Nickel electrodissolution and nickel passivation take place simultaneously at the potential range which corresponds to the defined anodic peak on the voltammetric curves at a relatively moderate acid pH. The use of the instantaneous mass/charge ratio analysis can allow to separate nickel electrodissolution and nickel passivation current components from the total measured charge. This analysis seems to indicate that in this case Ni(OH)2 is the species responsible for passivation. The theoretical expressions for intensity current and mass rate variations are deduced for a simplified electrodissolution/passivation mechanism. The fitting to the experimental curves allow to obtain a mass of MOH- = 17 g·mol-1 for the species responsible for passivation, in good agreement with the expected value. 5. Acknowledgements. This work has been supported by CTQ 2004-08026/BQU. J. Gregori acknowledges a Fellowship from the Spanish Education Ministery (FPU program). J.J. García-Jareño acknowledges his position the (Program “Ramón y Cajal”) to the Spanish Science and Technology Ministery. D. Giménez-Romero acknowledges a fellowship from the Generalitat Valenciana. 6. References. 1. J. Gregori, J.J. García-Jareño, D. Giménez and F. Vicente J. Solid State Electr. 9 (2005) 83. 2. M.R. Barbosa, S.G. Real, J.R. Vilche and A.J. Arvía J. Electrochem. Soc. 135 (1988) 1077. 3. A. Jouanneau, M. Keddam and M.C. Petit Electrochim. Acta 21 (1976) 287. 4. M. Itagaki, H. Nakazawa, K. Watanabe and K. Noda Corros. Sci. 39 (1997) 901. 5. M. Keddam, “Anodic Dissolution”, in: Corrosion Mechanism in Theory and Practice, Ed. P. Marcus and J. Oudar, Marcell Dekker, New York (1995) p.55. 6. C.V. D’Alkaine and M.A. Santanna J. Electroanal. Chem. 457 (1998) 5. 7. C.V. D’Alkaine and M.A. Santanna J. Electroanal. Chem. 457 (1998) 13. 8. R.D. Armstrong and H.R. Thirsk Electrochim. Acta 17 (1972) 171. 9. M.R. Barbosa, J.A. Bastos, J.J. Gacía-Jareño and F.Vicente Electrochim. Acta 44 (1998) 957. 10. D. Giménez-Romero, J.J. García-Jareño and F. Vicente J. Electroanal. Chem. 558 (2003) 25. 11. C. Gabrielli and M. Keddam Electrochim. Acta 41 (1996) 957. 12. P. Marcus, J. Oudar, I. Olefjord J. Microsc. Spectrosc. Electron. 4 (1979) 63.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Copper Passivity in Carbonate Base Solutions and its Application in Chemical Mechanical Planarization (CMP) E. Abelev, D. Starosvetsky, M. Auinat and Y. Ein-Eli Department of Materials Engineering, Technion-Israel Institute of Technology, Haifa 32000, Israel e-mail address: [email protected]

Abstract - Copper is fully passivated in sulfate solutions containing potassium carbonate. The potential range of copper passivity strongly depends on the relationship between sulfate and carbonate concentrations. At potentials above the passivity range copper suffers from localized attack by pitting corrosion. Copper passivity is more pronounced in solutions containing higher carbonate content. The increase in carbonate concentration shifts the breakdown potential towards positive anodic potentials while decreasing the anodic currents values in the region of passivity. In a solution containing carbonate copper passivity was detected in a wide potential range between OCP (~ 異0.15 VSCE) and ~1.0 VSCE. Increasing the sulfate concentration has the opposite effect on copper passivity than carbonate does. Keywords: copper, extended passivity, alkaline, carbonate solution

1. Introduction In the last decade the interest in studying the electrochemical behavior of copper in different electrolytes systems is significantly increased due to the role of copper in the microelectronic industry [1-5]. Copper is the metal of choice, replacing aluminum in integrated circuit (IC) interconnects. Copper dual damascene technology includes two electrochemical steps: electroplating and chemical-mechanical polishing (CMP) [1]. The main purpose of applying the CMP process is to remove copper overburden from interconnects patterning through rapid dissolution and global planarization of the wafer surface. CMP process is usually conducted with polishing pads in slurry.

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An effective CMP can be obtained once the slurry is fulfilling the following requirements [1]: (1) the planarized metal is fully passivated in the slurry of choice. (2) The dissolution of the planarized metal occurs only at the activated upper surface sites. (3) Exposed and activated sites should be rapidly repassivated. Thus, the necessary condition for CMP slurry is to provide a rapid self-passivity to the planarized metal. Thermodynamic calculations indicate that copper passivity in aqueous solutions can be achieved in a pH range of 7–13 and at potentials region above 0.58 – 0.0592(pH) VSHE [8]. The oxides covering the copper surface in this region are mainly Cu2O, CuO and Cu(OH)2 [8-21]. In the present work we evaluate a chemical composition of solutions, which can provide copper passivity in a wide potential range. For this purpose we studied potassium carbonate based solutions. The formation of protective films containing carbonates and hydroxide copper species (CuCO3·Cu(OH)2, 2CuCO3·Cu(OH)2 and CuCO3(OH)2異2 was reported [17, 21]. Copper passivity was reported by Tromans and Sun [17] in sulfate solution with a pH of 10.8, which was controlled by carbonate buffer (HCO3異/ CO32異). Our work is focused on identification of copper passivity parameters such as potential range and current characteristics. 2. Experimental A pencil-type specimen made of 5 mm diameter copper rod (99.995 wt.%) mounted in a room temperature-curing-epoxy was used in the electrochemical measurements. After polishing to 1200 grit alumina, the samples were carefully degreased with acetone and water rinsed. The electrochemical measurements were conducted with 273A EG&G potentiostat in 500 mL three electrode electrochemical cell equipped with a reference saturated calomel electrode (SCE) and Pt-wire counter electrode. A saturated calomel reference electrode was installed in solution through a Luggin-Habber capillary tip assembly. The studied solutions were prepared with DI water with addition of analytical grade chemicals, such as K2CO3, Na2SO4, and KOH. All chemicals were purchased from Aldrich Chemicals and were used without any further purification. 3. Results and Discussion Figure 1 presents anodic potentiodynamic curve obtained from Cu polarization in 1 gr/l Na2SO4 and cyclic polarization curves obtained from Cu in 1 gr/l Na2SO4 containing different concentrations of K2CO3. Positive potential sweep was applied from a potential slightly below OCP. Scan direction was reversed at potentials where current density reached a current density value of 1 mA/cm2.

Copper Passivity in Carbonate Base Solutions and its Application in CMP

1.0

1gr/l Na 2 SO 4 1gr/l Na 2 SO 4 + [K 2 CO 3 ]

0.8

Potential (VSCE)

127

0.6

0.5 gr/l, pH 10.5 1 gr/l, pH10.7 2 gr/l, 10.9 4 gr/l, 11.1

0.4 0.2 0.0 -0.2 -8 10

10

-7

10

-6

10

-5

10

-4

10

-3

10

-2

2

Current (A/cm )

Figure 1: Potentiodynamic curves of copper in 1 gr/l Na2SO4 solutions without K2CO3 and with different concentrations of K2CO3.

The anodic current gradually increases in the carbonate-free solution containing 1 gr/l Na2SO4, indicating active Cu dissolution (Fig.1). The addition of potassium carbonate to the sulfate solution significantly affects the electrochemical behavior of copper. Small current peak and narrow region of copper passivity up to 0.14 V (SCE) can be seen in the anodic curve obtained at the positive potential scan with the addition of 0.5 gr/l K2CO3 to the sulfate solution. At potential above 0.14V, a breakdown is detected and anodic current rapidly increased up to 1 mA/cm2 during further positive potential shift. The marked hysteresis, characteristic for pitting corrosion, can be detected between anodic curves obtained at positive potential scan and a back scan (Fig.1). The breakdown potential significantly increased with increase in K2CO3 concentration. In sulfate solutions containing 1, 2 and 4 gr/l K2CO3 passivity breakdowns were measured at 0.77 V, 0.86 V and 0.93 V, respectively. It should be noted that current densities values measured from polarizing copper in the carbonate containing solutions in the region of copper passivity were very small, indicating a strong Cu passivity. Figure 2 presents anodic curve of Cu obtained at positive potential scan and back scan in 1 gr/l Na2SO4 solution containing 4 gr/l K2CO3. The reverse scans were applied at potentials related to different potential regions of anodic characteristic. As can be seen, at potentials below 0.75 V the anodic current gradually decreased subsequent to a reverse scan, indicating copper passivity. At a potential of 0.96 V a breakdown of copper passivity occurred. Within a further anodic potential shift the current rapidly increased. A large hysteresis was detected once the reverse scan was applied at potentials above the breakdown potential (0.96 V). Similar passivation effect of carbonate was obtained in 10 gr/l Na2SO4 solution. Figure 3 presents cyclic polarization curves of Cu polarized in a free-carbonate solution containing 10 gr/l Na2SO4 and in solutions

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containing 10 gr/l Na2SO4 and different K2CO3 concentrations. Positive potential sweep was applied from a potential slightly below OCP and the reverse scan direction was applied at potentials where current density reached 1 mA/cm2. R everse scan (V SC E ):

Potential (VSCE)

0.9

- 0.1

0.3 0.6 0.75 1

0.6 0.3 0.0

10

-9

10

-8

-7

-6

-5

10 10 10 10 2 Current (A/cm )

-4

10

-3

Figure 2: Potentiodynamic curves obtained from polarization at positive potential scans and back scan in 1 gr/l Na2SO4 containing 4gr/l K2CO3. The reverse scans were applied at the following potentials: - 0.1, 0.3, 0.6, 0.75 and 1 V.

Potential (VSCE)

As can be seen in Figure 3, Cu is actively dissolved in carbonate-free solution (containing only sulfate). The addition of carbonates to the solution results in copper passivation. At small K2CO3 concentration, copper is passivated. However, the region of passivity was very narrow and a breakdown was detected at 0.06 V. At further potentiodynamic shift of the applied potential rapidly the current increased up to 1 mA/cm2 and a large hysteresis between positive and back scan curves can be clearly seen. 1.0

10 gl Na2SO4 K2CO3 concentration:

0.8

1 gr/l, pH 10.9 2 gr/l, pH 11.0 4 gr/l, pH 11.1

0.6

(a)

KOH concentration: 0.1 gr/l, pH 10.9 0.2 gr/l, pH 11.2 0.3 gr/l, pH 11.4

(b)

0.4 0.2 0.0 -0.2 -9

10

-8

10

-7

-6

-5

-4

-3

-2

10 10 10 10 10 10 2 Current Density (A/cm )

-7

10

-6

10

-5

-4

-3

-2

10 10 10 10 2 Current Density (A/cm )

Figure 3: Potentiodynamic curves of copper in 10 gr/l Na2SO4 solutions without and with different concentrations of K2CO3 (a) and KOH (b).

Copper Passivity in Carbonate Base Solutions and its Application in CMP

129

Increase in K2CO3 concentration results in a more pronounced copper passivity. One can see that the region of passivity was extended to 0.3 V in solution containing 2 gr/l K2CO3 while a breakdown of passivity was detected at this potential. In a solution containing 4gr/l K2CO3 the breakdown was detected at much more positive potential (close to 0.7 V). Comparing these results with those obtained in 1 gr/l K2CO3 we can conclude that the effect of carbonate concentration on copper passivity appeared much weaker at higher sulfate concentration. It was shown previously that the addition of carbonate significantly increases the pH value of sulfate solution. This in turn would affect the copper passivity. In order to clarify the role of pH and CO32- concentration on copper passivity potentiodynamic experiments were conducted in 10 gr/l Na2SO4 solutions containing different concentrations of KOH (Fig. 3b). In these experiments we have set the KOH concentrations, which provide the same pH values as in the case of K2CO3 addition. As can be seen in Figure 3b, the effect of KOH addition on copper passivity was much weaker compared with K2CO3 addition. 4. Conclusion Copper is fully passivated in sulfate solutions containing potassium carbonate. The potential range of copper passivity strongly depends on the relationship between sulfate and carbonate concentrations. At potentials above the passivity range copper suffers from an attack by pitting corrosion. Copper passivity is more pronounced in solutions containing higher carbonate content. The increase in carbonate concentration shifts the breakdown potential towards anodic potentials positive while decreasing the anodic currents values in the region of passivity. In a solution containing 1 gr/l Na2SO4 with 4 gr/l K2CO3 copper passivity was detected in a wide potential range between OCP (~ -0.15 VSCE) and ~1.0 VSCE. Increasing the sulfate concentration has the opposite effect on copper passivity than carbonate does. References 1.J.M. Steigerwald, S.P. Murarka, R.J. Gutmann, ”Chemical Mechanical Planarization of Microelectronic Materials.”, 1997. 2.S. Lakshminarayanan, J. Steigerwald, D. T. Price, M Bourgeois, T. P. Chow, R. J. Gutmann, and S. P. Murarka, IEEE Electron Device Lett.,15, 307 (1994). 3.M. B. Small and D. J. Pearson, IBM J. Res. Develop., 34, 858 (1990). 4.P. Singer, Semiconductor International, 90, June 1998. 5.R. Lui, Solid State Electronics, 43, 1003 (1999). 6.G.B. Shinn, V. Korthuis, A.M. Wilson. Handbook of semiconductor manufacturing technology, Ed. Y. Nishi, R. Doering, MARCEL DEKKER, Inc. N.Y. 2000, 415. 7.F.B. Kaufman, D.B. Thompson, R.E. Broadie, J. Electrochem.Soc.,138(11), 3460 (1991).

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8.M. Pourbaix,”Atlas of Electrochemical Equilibria in Aqueous Solutions”, 2nd US ed. NACE,Houston,TX,1974. 9.V. Maurice, H.H. Strehblow, P. Marcus, Surface Science, 458, 185 (2000). 10. J.Kunze, V.Maurice, L.H.Klein, H-H. Strehblow, P. Marcus, Corrosion Science, 46 (2004) 245264. 11. J.Kunze, V.Maurice, L.H.Klein, H-H. Strehblow, P. Marcus, J. Phys,Chem. B, 105 (2001) 42634269. 12.J.Kunze, V.Maurice, L.H.Klein, H-H. Strehblow, P. Marcus, J. Electroanal. Chem., 554-555 (2003) 113-125. 13.S.M. Wilhelm,Y. Tanizawa,C. Liu,N. Hackerman, Corrosion Science, 22(8), 791(1982). 14.D.D. MacDonald, J. Electrochem. Soc., 121, 208 (1966). 15.W. Kautec, J.G. Gordon, J. Electrochem. Soc., 137(9), 2672 (1990). 16.H. H.Strehblow, B. Titze, Electrochemica Acta, 25, 839 (1980). 17.D. Tromans, R. Sun, J. Electrochem. Soc., 139(7), 1945 (1992). 18.H.Y.H. Chan, C.G. Takoudis, M.J. Weaver, J. Phys. Chem. B, 103, 357 (1999). 19.R.L. Deutscher, R. Woods, J. Appl. Electrochem., 16, 413 (1986). 20.J. Ambrose, R.J. Barradas, D.W. Shoesmith, J .Electroanal. Chem., 47, 47 (1973). 21.J. Ambrose, R.J. Barradas, D.W. Shoesmith, J. Electroanal. Chem., 47, 65 (1973). 22.M. Perez-Sanchez, M.Barrera, S.Gonzalez, R.M. Souto, Electrochimica Acta, 35(9),1337 (1990).

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Analysis of Cu corrosion product in aqueous lithium bromide concentrated solutions M.J. Muñoz-Portero, J. García-Antón, J.L. Guiñón, V. Pérez-Herranz Departamento de Ingeniería Química y Nuclear. Universidad Politécnica de Valencia P.O. Box 22012, E-46071 Valencia (Spain) E-mail: [email protected] Abstract : The corrosion product of copper produced by lithium bromide (LiBr) concentrated solutions of 400 g/L and 700 g/L (4.61 M and 8.06 M) at room temperature, 70 ºC, and 110 ºC was studied using scanning electron microscopy (SEM). The corrosion product formed in 400 g/L LiBr solutions presented a crystalline and granular morphology, believed to be associated with a mixture of CuBr and CuO-Cu(OH)2. However, the corrosion product formed in 700 g/L LiBr solutions presented a gelatinous and amorphous morphology, thought to be associated with a mixture of CuBr and CuBr2·3Cu(OH)2. Keywords : Lithium bromide, copper, corrosion product, scanning electron microscopy, temperature

1. Introduction Concentrated solutions of lithium bromide (LiBr) are widely used as absorbents in refrigeration technology [1,2]. However, LiBr solutions can cause serious corrosion problems in the structural materials that make up components of refrigeration systems and heat exchangers at absorption plants. The aim of the present work was the study of the corrosion product of copper produced by LiBr concentrated solutions of 400 g/L and 700 g/L (4.61 M and 8.06 M) at different temperatures (room temperature, 70 ºC, and 110 ºC) using scanning electron microscopy (SEM). 2. Experimental procedure Copper cylindrical probes (99.9 % purity) of 1 cm in diameter and 5 cm high (with a total surface of 17.28 cm2) were used in the study of corrosion product.

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Previously the copper probes were cleaned with ethanol, rinsed with distilled water, and dried with air, according to ASTM G 1-03 [3]. Tests were carried out maintaining the probes in aqueous solutions of 400 g/L and 700 g/L (4.61 M and 8.06 M) LiBr at different working temperatures: room temperature, 70 ºC, and 110 ºC. These temperatures represented the values of the temperatures reached in the absorber and generator of the absorption machine that uses LiBr. The temperatures of 70 ºC and 110 ºC were maintained with a thermostatic bath model “Precisterm” of P SELECTA, with a capacity of 12 L, which contained ethylenglycol (C2H6O2). LiBr solutions were prepared from a LiBr (98 wt %) of PANREAC. The corrosion products of copper were observed with a scanning electron microscope model JSM 6300 JEOL, which had associated a microanalysis equipment. Copper probes were immersed in 60 mL of pre-nitrogenous 400 g/L and 700 g/L LiBr solutions at room temperature, 70 ºC, and 110 ºC into a cylindrical glass bottle, according to ASTM G 31-72 (2004) [4]. The specimens were maintained for one year. The copper probes were removed from the cylindrical glass bottle and were drained off. The copper probes were observed by SEM to study the morphology of the corrosion products adhered to them [3]. A new copper probe was also observed with the microscope to compare the aspect of the corroded and noncorroded copper. 3. Results and Discussion Figures 1 and 2 show the SEM micrographs of the copper probes after one year approximately of exposure in pre-nitrogenous 400 g/L (Figure 1) and 700 g/L (Figure 2) LiBr solutions at room temperature, 70 ºC, and 110 ºC. Figures 1(a) and 2(a) show the micrographs of a noncorroded copper probe, which are used as reference for the study of the corrosion product. For copper probes exposed in 400 g/L and 700 g/L LiBr solutions, a blue color corrosion product was obtained at the three temperatures, the amount of corrosion product obtained in 700 g/L LiBr solution being greater than in 400 g/L LiBr solution. The amount of corrosion product obtained was not enough for the characterization by X-ray power difraction (XRD). In the analysis of corrosion product by energy dispersive X-ray (EDX), three elements were detected: copper, oxygen, and bromine. However, the composition can not be determined because it was not possible to distinguish between the copper in the corrosion product and the copper in the electrode. Therefore, the predictions of compound formed was realized using the Pourbaix diagrams for copper in 400 g/L and 700 g/L LiBr solutions at 25 ºC (Figure 3) [5,6], in similar way as in previous works [7]. A line at the appropiate pH has been marked in the Pourbaix diagrams for copper in 400 g/L (pH = 6.80) and 700 g/L (pH = 5.65) LiBr solutions at 25 ºC to facilitate the identification of the compounds.

Analysis of Cu corrosion product in aqueous lithium bromide...

mm 5050mm

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mm 5050mm

(a) mm 5050mm

(b) mm 5050mm

(c)

(d)

Figure 1. SEM micrograph (1000 X) of copper (a) before corrosion, and after corrosion in 400 g/L LiBr solutions at (b) room temperature, (c) 70 ºC, and (d) 110 ºC.

The comparison of the micrographs of the corrosion product in 400 g/L LiBr solutions at room temperature, 70 ºC, and 110 ºC (Figure 1) shows that the increase of the temperature generated a greater amount of corrosion product adhered on the surface of the copper probes, which presented a crystalline and granular morphology. At room temperature, only isolated crystals were formed on the surface of the copper probe. However, the amount of crystals increased with the temperature, which did not cover totally the surface of the copper probes and allowed the interaction of the copper with the corrosive medium. According to the Pourbaix diagram for copper in 400 g/L solutions at 25 ºC (Figure 3(a)), the corrosion product initially formed in 400 g/L LiBr solutions should be copper (I) bromide (CuBr), although it could be oxidized with time, forming copper (II) oxide (CuO) and copper (II) hydroxide (Cu(OH)2).

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mm 5050mm

mm 5050mm

(a) 5050 mm

(b) 50 mm

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Figure 2. SEM micrograph (1000 X) of copper (a) before corrosion, and after corrosion in 700 g/L LiBr solutions at (b) room temperature, (c) 70 ºC, and (d) 110 ºC.

The comparison of the micrographs of the corrosion product in 700 g/L LiBr solutions at room temperature, 70 ºC, and 110 ºC (Figure 2) shows that when increasing the temperature the amount of corrosion product adhered on the surface of the copper probes increased, in a similar way as what happened in 400 g/L LiBr solutions. However, in 700 g/L LiBr solutions the corrosion product presented a gelatinous and amorphous morphology, which explained the adherent and protective nature of the corrosion product. At room temperature the corrosion product partially covered the surface of the copper probe. At 70 ºC and 110 ºC the corrosion product totally covered the surface of the copper probes, forming a compact and little porous film that protected it against corrosion. According to the Pourbaix diagram for copper in 700 g/L LiBr solutions at 25 ºC (Figure 3(b)), the corrosion product formed initially should be also CuBr, but in this case CuBr could be oxidized to form copper trioxybromide (CuBr2·3Cu(OH)2).

Analysis of Cu corrosion product in aqueous lithium bromide...

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Figure 3. Pourbaix diagrams for Cu-Br -H2O system at 25 ºC with LiBr concentrations of (a) 400 g/L and (b) 700 g/L.

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4. Conclusions In the SEM study of the corrosion product formed after immersing copper probes in pre-nitrogenous 400 g/L and 700 g/L LiBr solutions at room temperature, 70 ºC, and 110 ºC, the corrosion product on the surface of the copper probes increased with the temperature, the corrosion product in 700 g/L LiBr solutions being greater than in 400 g/L LiBr solutions. The corrosion product formed in 400 g/L LiBr solutions presented a crystalline and granular morphology, believed to be associated with a mixture of CuBr and CuOCu(OH)2, which did not totally cover the copper probes and allowed the evolution of the corrosion process. The corrosion product formed in 700 g/L LiBr solutions presented a gelatinous and amorphous morphology, thought to be associated with a mixture of CuBr and CuBr2·3Cu(OH)2, which formed an adherent and protective film on the surface of the copper probes. Acknowledgments We wish to express our gratitude for the support of this work by the DGSIC. (Convention No. PPQ2002-04445-C02-01), and to Dr. M. Asunción Jaime for her translation assistance.

References 1. J.W. Furlong, The Air Pollution Consultant 11/12 (1994) 1.12. 2. S-F. Lee, and S.A. Sherif, ASHRAE Transactions 105, Part 1 (1999) 1256. 3. ASTM G 1-03, “Standard Practice for Preparing, Cleaning, and Evaluating Corrosion Test Specimens” (West Conshohocken, PA: ASTM International, 2003). 4. ASTM G 31-72 (2004), “Standard Practice for Laboratory Immersion Corrosion Testing of Metals” (West Conshohocken, PA: ASTM International, 2004). 5. M.J. Muñoz-Portero, J. García-Antón, J.L. Guiñón, and V. Pérez-Herranz, “Pourbaix diagrams for copper in aqueous lithium bromide concentrated solutions”, 15th International Corrosion Congress, paper no. 464, Granada, Spain (2002) p. 454. 6. M.J. Muñoz-Portero, J. García-Antón, J.L. Guiñón, and V. Pérez-Herranz, Corrosion 60, 8 (2004) 749. 7. M.J. Muñoz-Portero, J. García-Antón, J.L. Guiñón, and V. Pérez-Herranz, Corrosion 61, 5 (2005) 464.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Passivity of Tin and CuSn Alloys in Alkaline Media studied by X-ray Photoelectron Spectroscopy Petra Keller*, Hans-Henning Strehblow Heinrich-Heine-Universtät, Universitätsstr. 1, D-40225 Düsseldorf, Germany * [email protected]

Abstract - Passive layers on tin, CuSn4 and CuSn19 were prepared by electrochemical oxidation in 0.1 N KOH under potentiostatic control with systematic variation of the relevant parameters, such as potential and time, and examined by X-ray photoelectron spectroscopy (XPS). The oxidation of the sputter cleaned samples was performed in an closed system without exposure to laboratory air. The quantitative evalution of the XPS signals of Sn 3d5/2, Cu 2p3/2, O 1s and the Auger signal Cu L3MM was performed with a fitting procedure with standard spectra. It includes the determination of the composition, i.e. the cationic and anionic fractions of the layer. Sputter depth profiles and angle-resolved measurements allow an insight into the layer structure and the calculation of the layer thickness. Keynotes: tin, CuSn-alloys, XPS, passivity, electrochemistry

1. Introduction Since many centuries tin and bronces are important materials with many practical applications because of their good corrosion resistance. In alkaline solutions the anodic passivation of tin is a complex, multi-step process which is determined by the formation of an oxide layer at the metal surface. The composition and the thickness of this passive film changes continuously with the electrode potential and the pH. But there is some disagreement in literature about the composition of the film and the mechanism of its formation [1-7]. The passivation behaviour, the composition, structure and the electronic properties of copper in alkaline solutions are very well documented [8,9], however, the knowledge about anodic films on CuSn is still poor [10-12]. The present

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investigation gives informations about the potential and time dependent formation of passivating layers on Sn and CuSn alloys in 0.1M KOH with XPS. 2. Experimental The electrochemically grown passive films on tin, CuSn4 and CuSn19 were studied by XPS in a commercial ultra high vacuum (UHV) spectrometer (VG ESCALAB 200X) using Al K -radiation. The take-off angle of the emitted photo electrons was adjusted usually to 15° with respect to the surface normal. A specifically designed electrochemical chamber [13] flanged to the spectrometer permits specimen preparation under argon atmosphere and sample transfer into the UHV without contamination. Argon ion sputtering (Penning source, Specs PS IQP 10/63) was applied to clean the sample surfaces before the preparation started and to investigate the thickness and composition of the prepared passive layers by sputter depth profiles. The depth scale was calibrated relative to anodically formed Ta2O5 layers of known thickness. The settings of the ion source (540V, 4.5 mA) yield a sputter rate of 1.8nm/min. The polished and cleaned samples were passivated at various potentials for 5min in 0.1M KOH and for various times at specified potentials. All potential values are given relative to the standard hydrogen electrode (SHE). After passivation the samples were rinsed within the electrochemical chamber with deionised water, blown dry with argon and transferred into the analyser chamber of the spectrometer without exposition to the atmosphere. All solutions were prepared with chemically pure substances (p.a.) and deionised water (Millipore system). The XPS signals Sn 3d5/2, O 1s, Cu 2p3/2, and the X-ray induced Auger signal Cu L3MM were evaluated on the basis of standard spectra to yield the qualitative and quantitative composition of the passive layer. The binding energies of the tin standards were taken from XP-spectra of sputter-cleaned metal samples (Sn EB= 484.7eV) and of commercial powders (SnO EB = 485.6eV; SnO2 EB = 486.3eV). The shift of the binding energy of the XPS signals of 0.8eV for the tin oxide species allow to differentiate between Sn(II) and Sn(IV) as was found by other groups [1,20-22]. The parameters for the Cu standards were obtained from anodic layers on pure copper electrodes [15-16]. The peak areas of the different signals were corrected by the corresponding photoionisation cross section according to Scofield [17] and used to evaluated the cationic fractions. X-ray induced Auger Electron Spectroscopy was applied to distinguish between Cu(I) and Cu(II) due to their small chemical shift for XPS [15]. 3. Results and Discussion Fig.1 depicts the potentiodynamic polarisation curve of pure tin with two anodic peaks ASn1 and ASn2 at E = -0.8 and E = -0.5 V (SHE) respectively, which agree reasonably well with the thermodynamic data of the Pourbaix diagram for the formation of soluble HSnO2- and SnO32- [18]. A large passive range follows

Passivity of Sn and CuSn Alloys in Alkaline Media studied by XPS

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between E = -0.3 V and 1.2 V with i = 50 Acm-2 and a secondary passive range up to 2.1 V with an increased current density of ca 0.8 mAcm-2. At E > 2.1 V follows a steep current increase attributed to oxygen evolution. The cathodic peak CSn at E = -1.0V is comparatively small. The integration of the peak yields a charge of 2.9 mCcm-2, the calculated layer thickness for SnO2 reduction amounts 1.6 nm. XPS studies in combination with potentiostatic transients yield a complete reduction of passive layer on Sn at E = -1.0 V. However also partial dissolution of the anodic film for these conditions is suggested.

Fig.1: potentiodynamic polarisation curve of a) Sn, b) CuSn4 and c) CuSn19 subsequent scans in 0.1N KOH

Tin specimens have been prepared as described above with a systematic variation of the potential within the whole range beginning at values below the anodic peaks A1 and A2 up to those in the secondary passive range and were subsequently investigated by XPS. Only small amounts of SnO / Sn(OH)2 are found in the potential range of E -0.9V to E = -0.3V corresponding to ca. one monolayer of oxide. For E > -0.3V the Sn(IV) oxide signal appears and the metal signal disappears due to the growth of the oxide film. In this passive potential range of E = - 0.2V to E = 1.4V the layer thickness increases linearly with the potential. At the surface the film contains a mixture of Sn(II) and Sn(IV)-compounds, i.e. SnO/Sn(OH)2 and SnO2/Sn(OH)4, with pure SnO2 in the inner part, which increases with the potential. Close to the metal surface SnO is found again with an constant thickness of 1nm [19]. At potentials in the secondary passive range with E > 1.4V (SHE) to E = 2.0 V (SHE) only SnO2 was detected at the top of the layer with a significant increase of its thickness, SnO was detected at the oxide/metal/interface by sputter depth profiles. Time dependent measurements from 2s up to 90 min at a potential of E = -0.3V(SHE) shows a constant monolayer of tin oxides and no development of the passive layer thickness or composition. A growth of the inner SnO2 layer thickness with time was observed at a potential of E = 0.5V [19]. The described results for the passive layers on pure tin and the well documented results for pure copper [8,9,15-16] are the basis for the following experiments with Cu/Sn alloys and their interpretation. The potentiodynamic polarisation curves of CuSn4 and CuSn19 are shown in Figs. 1b,c. The oxygen evolution is shifted down to E = 0.8 V for both alloys. CuSn4 behaves almost like pure

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cationic fraction

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copper: Two anodic peaks ACu1 and ACu2 appear at E = -0.2 V and E = 0.05 V and two cathodic peaks CCu1 and CCu2 at E = -0.3 and E = -0.6 V which are related to the formation of Cu(I)- and Cu(II) -oxide and their reduction. The peaks for tin oxidation remain small and increase with the tin content of the alloy. For CuSn19 the characteristic oxidation peaks of tin ASn1 and ASn2 are well developed at E = -0.8V and E = -0.5V. Repeated scans display an increase of the negative tin CSn peak at E = -0.8 V. The cathodic peak of the Cu(II) to Cu(I)-reduction CCu2 at E = -0.3V is well developed. Only a small Cu(I)reduction peak appears at E = -0.6V and E = -0.8V with an obvious splitting into two contributions. The quantitative evaluation of the XPS-signals yields the composition of the passive layer which is presented in figures 2 and 3.

Sn(ox) Cu(ox)

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Fig.2: Potential dependent composition of passive layer on CuSn4 in 0.1N KOH (tp = 5min)

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Fig.3: Potential dependent composition of passive layer on CuSn19 in 0.1N KOH (tp = 5min)

The passive film on CuSn4 has basically the same composition and structure as that on pure copper: At a potential of E = -0.3V only Cu2O is present with a small SnO fraction. In the potential range of E = 0.2V to E = 0.8V the well known duplex structure of the copper passive film was found with a Cu2O / CuO, Cu(OH)2 sequence. The XPS sputter depth profiles for CuSn4 passivated in 0.1M KOH for 5min show the dominating presence of copper oxides and a small content of tin oxides in the middle of the film [19]. For CuSn19 the composition of the passive layer differs significantly due to the increased influence of the less noble metal tin (Fig.3). Up to E = -0.3V a monolayer of SnO / Sn(OH)4 covers the surface. In the passive range of E = 0.2V to E = 0.8V the film consists of SnO / SnO2 / Cu2O / CuO,Cu(OH)2. The XPS sputter depth profile of CuSn19 passivated for 5min in 0.1M KOH

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Passivity of Sn and CuSn Alloys in Alkaline Media studied by XPS

yields a very thin outer layer of copper oxides on top of a SnO2 film. The total thickness of the passive layer increases linearly with potential for both alloys with almost the same values [19]. To study the development of the anodic layer and its change with time, CuSn4 and CuSn19 specimens were prepared at the potential E = 0.5V for different passivation times. The layer on CuSn4 (Fig.4) shows a constantly low content of tin oxides within the film at any time. The dominating duplex layer Cu2 O/CuO, Cu(OH)2 was found for short passivation times up to 5 min. The CuO / Cu2O ratio increases for t > 5min due

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to the oxidation of Cu O to CuO. Fig 4: Time dependent composition of the passive layer on CuSn4 in 0.1N KOH (Ep = 0.5V) SnO SnO2 Cu2O CuO Cu(OH)2

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Fig 5: Time dependent composition of the passive layer on CuSn19 in 0.1N KOH (E p = 0.5V)

The time dependent growth of the oxide layer on CuSn19 alloy (Fig.5) shows a strong influence of the alloying element tin. The reaction at the surface starts with oxidation of the less noble component tin. After 2s the passive film SnO / SnO2 / Cu2O / CuO, Cu(OH)2 is observed with a high content of copper oxide at the electrolyte/oxide interface. A closer look to the valency of Cu ions within the oxides shows the increase of the Cu(II) species with passivation time. At 10 min the amount of copper oxides within the layer decreases. It is suggested that islands of Sn-oxides grow with time and finally form a continuous film. This tin oxide layer prevents further oxidation of copper which finally accumulates at the metal surface underneath the tin oxide. Sputter depth profiles indicates an increase of the tin oxide content inside the layer up to 75at% while the growth of the layer thickness slows down with time [19].

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4. Conclusions The formation and chemical composition of passivating layers on Sn and CuSn alloys have been studied in 0.1M KOH with electrochemical methods and X-ray Photoelectron Spectroscopy. The passive layer on tin consists of a Sn(II)-oxide / hydroxide film in contact with the electrolyte followed by an inner part of Sn(IV) and a Sn(II) fraction at the oxide/metal interface. Time dependent measurements show a growth of the inner SnO2 layer. The composition of the passive layers on CuSn4 and CuSn19 is significantly different. On CuSn4 alloys the well known duplex structure of copper oxides Cu2O / CuO, Cu(OH)2 is formed with a small tin oxide fraction inside the layer. The tin oxide content remains constant with time. For CuSn19 the layer composition is significantly different due to the increased influence of the less noble metal tin. In the passive potential range a minor contribution of copper oxide is found in the outer part of the layer. Tin is accumulated in the inner part of the protective oxide film which contains mainly Sn(IV). 5. References 1. R.O. Ansell, T. Dickinson, A.F. Povey, P.M.A.Sherwood, J. Electrochem. Soc., 124 No.9 (1977) 1360 2. T.D. Burleigh, H, Gerischer, J. Electrochem. Soc. 135 No. 12, (1988) 2938 3. M. Meticos-Hukovic, A. Resetic, V. Gvozdic, Electrochimica Acta, 40 No.11, (1995) 1777 4. J-M. Themlin, M. Chtaib, L. Henrard, Physical Review B, Vol 46 No 4, (1992) 5. M. Metikos-Hukovic, S. Omanovic, A. Jukic, Electrochimica Acta, 45 (1999), 977-986 6. M. Seruga, M. Metikos-Hukovic, T. Valla, M. Milun, H. Hoffschultz, K. Wandelt, J. of Electroanalytical Chemistry, 407 (1996) 83-89 7. P.E. Alvarez, C.A. Gervasi, Corrosion science 46 (2004) 91-107 8. U. Collisi, H.-H. Strehblow, J. Electroanal. Chem., 284 (1990) 385-401 9. H.-H. Strehblow, H.-D. Speckmann, Werkstoffe und Korrosion 35 (1984) 512-519 10. T.N. Vorobyova, V.P. Bobrovskaya, V.V. Sviridov, Metal Finishing, (1997) p. 14 11. M: Wadsak; T. Aastrup, I. Odnevall Wallinder, C. Leygraf, and M. Schreiner, Corrosion science 44 (2002) pp.791-802 12. I. Mabille, A. Bertrand, E.M.M. Sutter, C. Fiaud, Corrosion Science 45 (2003) 855-866 13. S. Haupt, C. Calinski, U. Collisi, H.W. Hoppe, H.D. Speckmann, H.-H. Strehblow Surf. Interface Anal. (1986) 357 14. D.A. Shirley, Phys. Rev. B, 5 (1972) 4709 15. P.Druska, H.-H. Strehblow, Surface aund Interface Anal. 23, (1995), 440-450 16. H.H. Speckmann, S. Haupt, H.-H. Strehblow Surf.and Interf. Anal. 11 (1988), 148-155 17. J.H. Scofield, J. Electron Spectrosc, 8 (1976) 129 18. M. Pourbaix, “Atlas of Electrochemical Equilibria in Aqueous Solutions” Perganom Press, Oxford, (1966), 475-484 19. P. Keller, H.-H. Strehblow, to be published 20. J.H. Thomas, S.P. Sharma, J. Vac. Sci. Technol. Vol. 14, No. 5, (1977) , 1168-1172 21. J.-M. Themlin, M. Chtaib, L. Henrard, Physical Review B, 46, No 4, (1992) 2460-2466 22. L. Köver, Z. Kovacs, R. Sanjines, G. Moretti, Surf. and Interf. Anal., 23, (1995), 461-466

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Surface Analytical Characterization of Chromium Passivation on Tinplate Roelf Sandenbergh, Marius Biermann, Thomas von Moltke Department of Materials Science and Metallurgical Engineering, University of Pretoria, Pretoria, 0002, South Africa; [email protected]

Abstract - X-ray Photo-electron Spectroscopy was used to characterise chromium passivation layers formed on tinplate during dip and cathodic dichromate (CDC) passivation treatments. The passive layers consisted of chromium hydroxides and/or oxides for the CDC and dip treatments respectively with tin oxides below it. No chromium or tin metal was detected in the passive layers. The poorer lacquer adhesion on CDC treated material is probably related to the presence of chromium hydroxides in these passive layers. Keywords: chromium, passivation, tinplate, lacquer adhesion

1. Introduction Chromium passivation on tinplate is used to protect the tin plated surface from oxidation when exposed to the atmosphere, during curing of lacquer and when exposed to sulphur containing food. Tinplate is typically produced on continuous production lines consisting of cleaning, tin electroplating, tin reflowing and chromium passivation steps. The chromium passivation is applied by exposure of the tinplate to an aqueous sodium dichromate solution as a simple dip, i.e. dip passivation, or with the application of cathodic polarization of the tinplate, i.e. CDC passivation. During passivation the soluble Cr(VI) in solution is reduced to insoluble Cr(III) species, or possibly Cr(O), by oxidation of the tin to tin oxide or by cathodic polarization. Tin lines typically operate at speeds of 400 – 450 m/minute resulting in exposure times to the passivation treatment of approximately 1s, and producing chromium contents on the surface of the tinplate ranging from 1 - 3, and 5 - 10 mg/m2 for the dip and CDC

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treatments respectively. While the corrosion and oxidation resistance of the tinplate increase with increasing chromium applied, lacquer adhesion may be decreased by higher chromium applications [1]. In this work the nature and composition of the passive layers produced by the different passivation treatments were investigated by XPS. 2. Experimental Standard production tinplate samples were used and surface characterization of the samples were done using X-Ray Photo Electron Spectroscopy (XPS) using a Physical Electronics model 5400, equipped with both an Al/Mg dual x-ray source and an Al monochromator. The instrument was operated at a base pressure < 4x10-9 Torr and had a lateral resolution of 10mm. 3. Results and Discussion The XPS survey analyses of tinplate passivated by dip and CDC treatments in sodium dichromate are shown in figure 1 and indicate that there is more chromium present on the CDC treated sample, as would be expected based on electrochemical measurements [2]. Multiplex analysis of the Sn 3d peaks, shown in figure 2, indicate that the tin surface was fully oxidized during the dip process, but also that some tin oxidation occurred, even under the cathodic polarization applied during the CDC process. The tin metal peak was shown to occur beneath the oxides by angle resolved analysis, and thus originated from the tin substrate and not from reduced tin oxide on top of the tin oxide. The XPS multiplex surveys of the Cr 2p peak for the dip and CDC treated tinplate, shown in figure 3, indicate the presence of mainly Cr2O3 in the case of the dip passivation, and of Cr2O3 and Cr(OH)3 for the CDC passivation, and also some residual Cr(VI) due to incomplete removal of the sodium dichromate solution used for the passivation. The surveys also clearly indicate that no chromium metal is formed during the CDC treatment. This was further confirmed with angle resolved XPS analysis of the Sn 3d and Cr 2p peaks on CDC passivated tinplate, shown in figure 4, where photoelectrons from deeper in the layer are detected at larger take off angles, confirming that no chromium metal is present in the passive layer. It was also confirmed in related work that the chromium passive layer is indeed continuous and amorphous in nature [3]. 4. Conclusions The present results confirm that the chromium passivation layer on tinplate is formed by the reduction of soluble chromium (VI) species to insoluble chromium (III) hydroxide and/or oxides, but not chromium metal, during both the dip and CDC treatments. Oxidation of the tin prior to chromium passivation

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would thus limit the ability of the surface to reduce the chromium in the case of the dip treatments, and thus the efficiency with which the passive layer is formed. It is interesting to note that the tin surface is also partially oxidised even for the CDC treatment. The poorer lacquer adhesion of the CDC passivated tinplate is probably related to the formation of chromium hydroxides, stimulated by the alkaline conditions formed under cathodic polarization.

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References 1. E.J. Helwig, 6th International Tinplate Conference, ITRI, 1996, 96-103. 2. M.C. Biermann, MSc Thesis, University of Pretoria, 2005. 3. A.S. Tuling, M.C. Biermann and T.vS. von Moltke, Proc. Microscopy Society of Southern Africa, Vol. 34, 2004, 26.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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To Passivate or not to Passivate, that is the Question: The Case of Barium Tin(II) Chloride Fluorides Georges Dénès and Abdualhafed Muntasar Laboratory of Solid State Chemistry and Mössbauer spectroscopy, Department of Chemistry and Biochemistry ,Concordia University, Montréal, Québec, Canada, [email protected]

Abstract - Barium tin(II) chloride fluorides were found to have a diversified behavior with respect to oxidation. Three types of situations were observed: (i) slow “fluoridelike” oxidation, (ii) moderate speed of oxidation, (iii) “fast” oxidation. The two latter categories were found only for the Ba1-xSnxCl1+yF1-y solid solution, that has the BaClF structure, with full Ba/Sn disorder, and also disorder between y Cl and (1-y) F (for y>0), or between –y F and (1+y) Cl (for y<0). A combination of X-ray diffraction and Mössbauer spectroscopy has shown that, in the barium tin(II) chloride fluorides, the key criteria that determine the kinetics of oxidation are (i) the type of bonding (covalent or ionic), (ii) the tin(II) sublattice strength, and (iii) the method of preparation. Keywords: Barium tin(II) chloride fluorides, oxidation, passivation, Mössbauer spectroscopy

1. Introduction The phenomenon of passivation by a thin surface layer of oxide seems to work well in tin(II) fluoride containing materials, that show a weak tin(IV) Mössbauer signal attributable to tin(IV) oxide, which seems to remain constant over many years at ambient temperature [1]. The SnO2 passivating layer is too thin to be observed by X-ray diffraction. The hypothesis of a thin passivating layer of SnO2 was confirmed by the fact its Mössbauer signal could not be detected on a single crystal (a plate thin enough to allow the -ray beam through it) when it is easily detected on a powdered sample [2]. This is not surprising

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since the surface/bulk ratio is negligible for a single crystal, and it increases in importance when the particle size decreases. The crystal structure of BaClF (PbClF type) is a tetragonal distortion of that of fluoride type BaF2, with anionic order. In BaF2, the unique cation site is in a MF8 cubic environment. The structure of BaClF is derived from that of BaF2, by replacing each second sheet of fluoride ions parallel to (a,b) by a layer of chloride ions [3]. The chloride ions are much larger than the fluoride ions, they cannot fit in the same plane, and they shift by (¼,¼,±z). This results in each barium ion being in a distorted monocapped square antiprism, BaF4 ClCl 4/ , whereby the Ba-Cl apical bond goes through the center of the Cl 4/ square. Recently, it was found in our laboratory that three new compounds, stoichiometric BaSn2Cl2F4 and BaSnClF3·0.8H2O, and non-stoichiometric Ba1-xSnxCl1+yF1-y, a doubly disordered solid solution that has the BaClF structure, could be obtained by precipitation when solutions of SnF2 and a of BaCl2·2H2O are mixed together [4-6]. A similar, although not quite identical, Ba1-xSnxCl1+yF1-y solid solution, is obtained by heating together a mixture of appropriate stoichiometry of SnF2, BaF2 and BaCl2 [5]. 2. Experimental Sample preparation was carried out by precipitation or in dry conditions at high temperatures according to methods that have already been reported [4] and were studied by X-ray powder diffraction and Mössbauer spectroscopy using methods that have been described in details by us elsewhere [6]. 3. Results and Discussion 3.1. Passivation of Stoichiometric Barium Tin(II) Chloride Fluorides Like in the case of fluorides, X-ray diffraction of stoichiometric BaSn2Cl2F4 and BaSnClF3.0.8H2O show no peak of SnO2, however, also like in the case of fluorides, their Mössbauer spectrum contains a broad peak at ca. 0 mm/s characteristic of tin(IV) coordinated by oxygen (fig. 1). Therefore, it is obvious that the chloride fluorides are passivated by a thin layer of SnO2, in the same way as for the fluorides. The stronger relative tin(IV) peak intensity of BaSn2Cl2F4 (fig. 1) does not necessarily mean that this compound contains a larger amount of tin(IV). It could also be due to a smaller recoil-free fraction of its tin(II), therefore located in a softer sublattice. This could be due to its Cl:F ratio of 2:4, compared to 1:3 in BaSnClF3.0.8H2O, since Sn-Cl bonds are weaker than Sn-F bonds. The SnO2 signal does not increase on aging, therefore passivation works as efficiently as for the fluorides.

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3.2. Poorer passivation of the Ba1-xSnxCl1+yF1-x solid solution For the Ba1-xSnxCl1+yF1-x solid solution, the situation is quite different from the above stoichiometric phases, and it changes with the method of preparation and the chemical composition. For the precipitated solid solution, the relative tin(IV) Mössbauer signal at a young age(a few months old) is similar to that of fluorides and covalently bonded tin(II) chloride fluorides (fig. 2). However, on aging in air at ambient conditions, the relative tin(IV) signal increases more than for the stoichiometric compounds. Therefore, passivation does not work as

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efficiently. For Ba1-xSnxCl1+yF1-y prepared by the dry method, the oxidation rate is a function of x and y. The tin(IV) Mössbauer line intensity increases much faster at low x, and for high x values at highly positive y values (fig. 3). This

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Figure 3: 119Sn Mössbauer spectra of Ba1-xSnxCl1+yF1-y at RT, prepared by direct reactions under dry conditions (freshly prepared and aged): (a1 & a2) AM-1180, x = 0.050, y = 0.10 [freshly prepared (a1) and aged for 32 months (a2)], (b1 & b2): AM-315, x = 0.121, y = 0.000 [freshly prepared (b1) and aged for 36.7 months (b2)], (c1 & c2): AM-1155, x = 0.225, y = 0.200 [freshly prepared (c1) and aged 31.0 months (c2)]

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tin(II) is mostly in the form of Sn2+ stannous ions (single Mössbauer line at ca. 4 mm.s-1 (figs. 3a1 & 3a2)) [4]. Since Sn2+ is smaller than Ba2+, it binds weakly to the anions surrounding it and is therefore an easy target for oxidation. At high x, in addition to Sn2+, clusters of covalently bonded Sn(II) linked by F bridges are present [4, 6], that give a single Sn2+ Mössbauer line at ca. 4 mm.s-1 and a covalent Sn(II) doublet (figs. 3b1 & 3b2)). For y<0, i.e. F/Cl>1, the excess fluorine maximizes the number of Sn(II)-F bonds, that are strongly covalent and therefore more difficult to break, resulting in a lower oxidation rate, like in the case of stoichiometric Ba/Sn(II) chloride fluorides, while the ionic (Sn2+) fraction oxidizes faster, resulting in a total oxidation rate higher than for the stoichiometric phases. At high x and y values (high Sn and Cl), a very weak line at ca. 4 mm.s-1 (Sn2+) is observed, showing a low recoil-free fraction, therefore a very weak tin(II) sublattice [4, 6]. Such a weakly bonded tin can be expected to yield easily to oxidation, and indeed, a fast growth of the relative tin(IV) signal is observed (figs. 3c1 & 3c2), not detected by diffraction. 4. CONCLUSION New barium tin(II) chloride fluorides experience a wide range of sensitivity to oxidation by atmospheric oxygen. “To passivate or not to passivate” is a question of bond strength: covalently bonded tin(II) provides a reasonably strong lattice and passivates well. Ionic Sn2+ oxidizes faster, particularly in the BaClF structure, mostly because the dimension of the anionic sublattice is determined by the size of the Ba2+ ions, which results in an oversized site for lodging the Sn2+ ions, which therefore bind weakly to the anions. Acknowledgement: This work was made possible by the support of Concordia University and the Natural Science and Engineering Research Council of Canada. Grateful thanks are also due to the Procter and Gamble Co. (Mason, Ohio) for supporting our Mössbauer laboratory. References 1. 2. 3. 4. 5. 6.

G. Dénès and E. Laou, Hyperf. Interact. 92, 1013 (1994). T. Birchall, G. Dénès, K. Ruebenbauer and J. Pannetier, J. Chem. Soc. (A), Dalton Trans., 2296 (1981). J. Flahaut, J. Solid State Chem .9, 124 (1974). A. Muntasar and G. Dénès, In: I. Ortalli (ed.), Proc. Of the Internat. Conf. on the Applications of the Mössbauer Effect, ICAME-95, Vol. 50, 1996, p. 127. G. Dénès and A. Muntasar, Hyperf. Interact.153, 119 (2004). A. Muntasar, PhD thesis, Concordia University, Québec, Canada (2002).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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A Thin Passivating Tin(IV) Oxide Layer on Tin(II)Containing Fluoride Particles, or not? The M1-xSnxF2 Solid Solutions Georges Dénès, Eva Laou, M. Cecilia Madamba, Abdualhafed Muntasar, and Zhimeng Zhu Laboratory of Solid State Chemistry and Mössbauer Spectroscopy, Department of Chemistry and Biochemistry ,Concordia University, Montréal, Québec, Canada, [email protected]

Abstract - Tin-119 Mössbauer spectroscopy shows that all powdered tin(II) compounds give a small tin(IV) signal that has been attributed to the presence of a very thin layer of SnO2 on the surface of the particles, that passivates against further oxidation, and is undetectable by X-ray diffraction. In a new development of this work, efficient passivation against further oxidation of the M1-xSnxF2 (M = Ca and Pb) solid solution poor in tin was found to be as high as for phases rich in tin, even though the amount of tin is way too low to be able to produce full coating of the particles by a thin layer of SnO2. This observation required developing a new model to explain the sluggish oxidation of phases poor in tin, based on bonding type and strength. Keywords: alkaline earth tin(II) fluorides, order-disorder, passivation, Mössbauer spectroscopy

1. Introduction In previous works, the spontaneous oxidation of SnF2 and MSnF4 in air at ambient conditions was studied in our laboratory [1, 2]. It was observed that the 119Sn Mössbauer spectrum of all powdered samples contains, in addition to the tin(II) lines in the 2-4 mm/s range, a weak peak at ca. 0 mm/s that can be attributed to tin(IV) coordinated by oxygen, most likely SnO2 stannic oxide. The tin(IV) oxide phase cannot be detected by X-ray diffraction, because the crystallites are likely to be smaller than the ca. 30Å critical size required to give coherent scattering. The tin(IV) Mössbauer signal intensity does not increase

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with further aging, thereby demonstrating that the remaining tin(II) is well passivated, and this has been accounted for by the presence of SnO2, which is therefore thought of forming an uninterrupted layer that prevents contacts between atmospheric oxygen and the bulk of the particles. The work presented here on passivation in samples very poor in tin will challenge this interpretation. Our earlier work showed that tin-119 Mössbauer spectroscopy is an excellent method for detecting minute amounts of tin(IV) oxide in tin(II) compounds [3]. The two main reasons are (i) the large velocity difference between Sn(II) and Sn(IV) lines makes them well separated, and (ii) the large recoilless fraction difference between tin(II) halides and tin(IV) oxides makes traces of tin(IV) oxide being easily detected. Consequently, a very small amount of SnO2 in a tin(II) halide will give a subspectrum disproportionably higher than that expected from its small content in the sample. Mössbauer spectroscopy is therefore capable of detecting traces of SnO2, regardless of crystallinity. Large alkaline earth metal fluorides MF2 (M = Ca-Ba) and -PbF2 crystallize in the fluorite type structure (space group Fm3m), whereby the metal ion is in a cubic coordination (fig 1a) [4]. In contrast, stannous fluoride SnF2 has a molecular structure, with a large degree of covalency in the Sn-F bonds [5]. The numerous materials we have prepared in the MF2/SnF2 systems result from an unusual combination of covalently bonded tin(II) with the MF2 ionic fluoritetype structure. The resulting combination gives structures that can be ordered or disordered. The best understood ordered structure is that of -PbSnF4 and BaSnF4 [6]. Several disordered MF2/SnF2 compounds have been prepared and characterized, all in our laboratory: (i) stoichiometric PbSn4F10, high temperature -PbSnF4, and µ -MSnF4 (M = Ba and Pb) obtained by ballmilling, and (ii) the non-stoichiometric M1-xSnxF2 solid solution (M = Ca and Pb), all of which have a cubic unit-cell with full M/Sn disorder. However, a Mössbauer spectroscopic study showed that, despite the M/Sn disorder, tin does not have ionic bonding like the M2+ ions, and instead, it forms covalent bonding like in the ordered phases. Tin(II) is in a SnF4E pseudo-square pyramidal coordination similar to the case of BaSnF4, without the axial Sn-F bond, with the tin-lone pair axis pointing randomly towards the three main axes a, b and c of the unit-cell. 2. Experimental Sample preparation was carried out by precipitation (Ca1-xSnxF2, 0.10≤x≤0.15, and Ba1-xSnxCl1+yF1-y, 0≤x≤0.25, -0.15≤y≤0.25) or in dry conditions (Pb1-xSnxF2, 0≤x≤0.50) at high temperatures according to methods that have already been reported [2] and were characterized by X-ray powder diffraction and Mössbauer spectroscopy using methods that have been described elsewhere [2].

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Figure 1: Projection of the structures of: (a) BaSnF4 (g-PbSnF4 type), and (b) BaF2

3. Results and Discussion The X-ray powder diffraction pattern of the M1-xSnxF2 solid solution (M = Ca and Pb) contains only the peaks characteristic of the F lattice of the fluorite-type structure (fig. 2A). The absence of superstructure reflections or extra lines characteristics of a lattice distortion shows that the unit-cell is still cubic and is not a multiple of that observed for MF2 fluorite-type structures (CaF2 and PbF2), therefore the two metals M and Sn are fully disordered on the unique metal site of the fluorite-type structure in the Fm3m space group. The same can be said, for the same reasons, for Ba and Sn, and for Cl and F (on the F site for y>0 or on the Cl site for y<0), for the Ba1-xSnxCl1+yF1-y solid solution in the P4/nmm tetragonal space group of the BaClF structure (fig. 2B). Since the main conditions for forming solid solutions are that the two atoms that are disordered (i) have similar sizes, and (ii) form the same type of bonding, one might expect that tin(II) forms ionic bonding to F and Cl like Ca, Pb and Ba in the M1-xSnxF2 and Ba1-xSnxCl1+yF1-y solid solutions. Such an assumption is difficult to verify by X-ray diffraction, because of the very high scattering

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power of the metals in comparison to that of fluorine. Fortunately, for tincontaining compounds, 119Sn Mössbauer spectroscopy provides a simple and convenient method for probing the tin(II) bonding method: (i) ionic tin(II) (Sn2+ stannous ion) has a non-stereoactive (spherical) lone pair, and gives a single line at ca. 4 mm.s-1; (ii) covalent tin(II) has a stereoactive (axial) lone pair and as a result, a highly distorted coordination, which result in a smaller isomer shift h and a large quadrupole doublet (ca. 1.5 mm.s-1) and an isomer shift of ca. 3.3 mm.s-1 (figs. 2C(a) & 3(b)), both of which are characteristic of covalently bonded tin(II); (b) in contrast, Ba1−xSnxCl1+yF1-y gives a Mössbauer single line at ca. 4.1 mm.s-1 (fig. 2C(b)) that is typical of ionic bonding with a nonstereoactive lone pair. X-ray powder diffraction of the solid solutions gives

(hence a smaller quadrupole (A) isomer shift h ) and a large (B) efg at tin (hence a large(C) splitting ). Figure 2C shows clearly that the two types of solid solutions have different tin(II) bonding types: (a) M1-xSnxF2 gives a large quadrupole doublet -1 powder diffraction pattern of (a) -PbF2, (b)-1Ca1-xSnxF2 (x = 0.27), (B) X-ray Figure (A) X-ray ) and an isomer shift of ca. 3.3 mm.s (figs. 2C(a) & 3(b)), both (ca. 1.52: mm.s powder diffraction pattern of (a) BaClF, (b) Ba1-xSnxCl1+yF1-y (x = 0.15 , y = 0.094), (C) 119Sn of which are characteristic of covalently bonded tin(II); (b) in contrast, ambient temperature Mössbauer spectrum of (a) Ca1-xSnxF2 (x = 0.27), (b) ) Ba1-x-1SnxCl1+yF1-y (x = Cl1+yF1-y gives a Mössbauer single line at ca. 4.1 mm.s (fig. 2C(b)) Ba 1−x,Sn 0.15 y =x 0.094) that is typical of ionic bonding with a non-stereoactive lone pair. X-ray powder no SnO2 that could indicate that some of the tin has oxidized on prolonged

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contact with air. Figures 2C(a) and 3 show that there is no significant tin(IV) Mössbauer line near velocity 0 mm.s-1, just a very weak signal, barely detectable to the untrained eye, showing that oxidation to tin(IV) is very minimal, no more than a trace, in the M1-xSnxF2 solid solutions, similar to the On the other hand, it is clear from stoichiometric phase -PbSnF4 (fig. 3). figure 2C(b) that a non-negligible broad tin(IV) line is observed for Ba1−xSnxCl1+yF1-y. These results require revisiting the mechanism of passivity in stannous tin compounds that was proposed earlier [1, 2]. Obviously, the model based on the presence of a thin protective layer of tin(IV) oxide on the surface of the particles cannot apply to solid solutions containing an amount of tin insufficient to cover entirely the particles, such as M1-xSnxF2 for x = 0.10. In addition, a reliable model must explain why passivation is less efficient for Ba1−xSnxCl1+yF1-y, even though it is as rich in tin as Pb0.90Sn0.10F2 that is better passivated. Even though both contain probably an insufficient amount of tin to provide enough SnO2 to make a uniform passivating layer on the surface of the solid particles, both are rather well passivated, although unequally, with the fluorides being clearly more protected than the chloride-fluoride. First, one must assume that the activation energy required to break the tin(II) bonding from the fluoride or chloride-fluoride lattice is sufficiently highly to make the kinetics of oxidation slow. In addition, the difference between the fluoride and

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chloride-fluoride solid solutions can be understood by their different type of bonding. In a fluoride matrix, tin(II) is hybridized and forms covalent bonding, therefore the activation energy required to oxidize tin must be able to destroy the hybridization and break the Sn-F covalent bonds, hence a very efficient protection. On the other hand, in the Ba1-xSnxCl1+yF1-y solid solution, tin(II) forms ionic bonds (single Mössbauer line at ca. 4 mm.s-1), it is not hybridized, and the activation energy required to donate the bare lone pair to oxygen and oxidizing tin is lower. 4. Conclusion The present work has shown that the passivation of fluoride or chloride-fluoride solid solutions poor in tin against the oxidation of tin(II) to tin(IV) cannot be accounted for by the formation of a protective layer of tin(IV) oxide at the surface of the particles. Passivation is explained by kinetic inertness due to the activation energy required to break existing bonds, and the difference between fluorides and chloride-fluorides by the tin orbital hybridization of the fluorides (covalent bonding) that requires a higher activation energy of oxidation than the ionic bonding of the chloride-fluorides. Acknowledgement: This work was made possible by the support of Concordia University and the Natural Science and Engineering Research Council of Canada. Grateful thanks are also due to the Procter and Gamble Co. (Mason, Ohio) for supporting our Mössbauer laboratory. References 1. G. Dénès and E. Laou, Hyperf. Interact. 92, 1013 (1994). 2. G. Dénès , E. Laou, M. C. Madamba and A. Muntasar, In: M. B. Ives, J. L. Luo and J. Rodda (eds), Passivity in Metals and Semiconductors, Proc. of the 8th Internat. Symposium, Electrochemical Society Proceedings, Vol. 99(42), The Electrochemical Society, Pennington, NJ, 2001, p. 233. 3. T. Birchall, G. Dénès, K. Ruebenbauer and J. Pannetier, J. Chem. Soc. (A), Dalton Trans., 1831 (1981). 4. A. R. West, Solid State Chemistry and its Applications, Wiley, New York, pp. 230-242 (1984). 5. G. Dénès, J. Pannetier, J. Lucas and J. Y. Le Marouille, J. Solid State Chem. 30, 335 (1979). 6. T. Birchall, G. Dénès, K. Ruebenbauer and J. Pannetier, Hyperf. Inter. Soc. 29, 1331 (1986).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Electrochemical Data About Disruption of Passivating Films. The Pb/PbSO4/H2SO4 System. C. V. D’Alkainea *, P. M. P. Prattaa a

Group of Electrochemistry and Polymers, DQ-UFSCar, CP 676, São Carlos (SP), 13565-905, Brazil. [email protected]

Abstract - The paper discusses experimental electrochemical data to support the idea that passivating films, under some circumstances, can suffer disruption, giving rise to disrupted films, attached by superficial tension forces, on the remaining continuous non disrupted films, glued on the metal surfaces. The case of Pb on H2SO4, forming PbSO4 films, is analyzed. 1.Introduction The disruption problem in passivating films can appear after the formation of few monolayers. It is related with the interruption of the forming electrical field (galvanostatically, potentiostatically or voltammetrically) or with its reduction. In the literature, in multiple cases, the formed films correspond to higher thicker films, where the disruption, if possible, it has already happened. The films discussed here correspond to the initial states of those thicker films. The present paper will give attention to electrochemical facts which signalize the disruption, principally when the reduction process of the film is taken into account. To do that, it has been selected the case of PbSO4 formation on lead in H2SO4 solutions. 2. The Disrupted Film Model. Previous Experimental Facts. For the growth of a passivating film in thickness, it is necessary to apply high electrical fields, to move the ions through it. These fields stabilize the films in their actual thickness through the electrostriction forces [1]. When these high fields are reduced, inverted or interrupted, the disruption forces produced by the fact that, in general, the molar partial volumes of the films are higher than those of the metals, become higher than that of the electrostriction forces [2]. The films can be then partially disrupted, giving rise to disrupted external films

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attached by superficial tension forces to the remaining glued films. The remaining glued films are glued on the metal surface because, at any interface, there is an inner potential difference which produces always a remaining field inside the remaining film. This field gives rise to a remaining stabilizing electrostriction force. Several are the experimental facts pointing out to this disruption process when it is looked for systems in which the metal ions are highly insoluble and then, where the dissolution of the disrupted films is impossible, even when they can slowly recrystallize, due to the small initial sizes of the particles. One demonstration of the disruption has been done through the combination of cycle voltammetry and ellipsometric measurements, for the case of Fe in alkaline solutions [3]. In this case, when the measurements are made at potentials anodic to the oxidation peaks for the film formation, they present two films. An inner one with the dielectric constant near of that of an iron oxide, and an external one with a dielectric constant near, but lower than that of the water. On the other side, when they are made at potentials cathodic to the reduction peaks, it is observed only the external film. The inner film, when detected, presents always the same thickness, but the external film thickness increases with the increase of the number of cycle voltamograms. The inner film is, evidently, the glued continuous one and the external film is the disrupted one. The disruption is also related with the higher amounts of anodic charge densities (qa) when compared with the cathodic ones (qc), for convenient intervals of potential sweeping, in solutions in which the metal cations are highly insoluble. If not, where have gone the excess of qa? Finally, the disruption can be related with the fact that all these kinds of systems in cycling voltammetry come to stationary current density/potential (i/E) representations. For the case of highly insoluble metal cations, it can be understood as the formation of a specific stationary solution in-between of the particles of the disrupted external film, grown in thickness. 3. Materials and Methods The electrode was Pb 99.997% in weight, keeping an electroactive area of 0.4 cm2. The working electrode was always polished with 600 emery paper at the beginning of any set of experiments. Then, the electrode was introduced cathodically polarized with the aid of a fourth Pt electrode at a potential of -1.3 Volts, where it was waited up to the stabilization of the current density (in 0.5 mAcm-2). At this moment several cycles of anodic film formation and reduction plus a potential arrest also at -1.3 volts, during 5 minutes, to complete reestablishment of the anodic peak (see the discussion), were performed. Only after this stabilization (about 10 cycles) the electrode was considered to be ready for the measurements and it gives a reproducibility reasonable increased. Between each experiment, the electrode was maintained (as during the stabilizing cycles) at -1.3 Volts, during 5 minutes, for the re-establishment of

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the initial anodic peak. The initial potential Ei, together with the time waited at this potential, ti, were in general -1.3 Volts and 5’. The interval of the swept potentials were between El,c, for the cathodic limit, and El,a, for the anodic one, with possible arrest times tl,c and tl,a, and using equal or different anodic (va) or cathodic (vc) sweep velocities. The average anodic (qa) or cathodic (qc) charge densities were determined through several measurements to reduce their error. The electrolyte solution was always deoxygenated 4.6 M H2SO4 and the reference electrode was always Hg/Hg2SO4/ 4.6 M H2SO4, to which all potentials are referred.

Figure 1. Voltammetric results showing (a) the anodic and cathodic total peaks of PbSO4 film formation and reduction, (b) The initial anodic sweep with the inversion of the anodic voltammetry at the beginning of the anodic peak, showing the nucleation. va = 20 mV s-1, vc = 20 mV s-1, Ei = -1.3 Volts, ti = 5’, El,c = -1.3 Volts, tl,c = 5’; El,a = -0.7 Volts, tl,a= 0’’.

The experiments were cycle voltammetries with va equal to vc or voltammetric measurements with, in different experiments, several va but always the same vc or only one va with different vc (like Fig.1a, with va = vc). For some experiments, the anodic sweep inversion in the initial oxidation peak potential was used to detect if there was or not nucleation (see Fig. 1b with detection of nucleation). 4. Results and Discussions 4.1 Cycles, nucleation and disruption processes In Fig. 2 it is seen first cycle voltammetries in the potential region of the Pb/PbSO4 film formation and reduction [4]. It can be appreciated that qa (the first one of 45 mCcm-2) decreases and qc remains constant (29 mCcm-2), with the number of cycles. This last value is very high to be explained by a dissolution/precipitation model. As the Pb2+ has a very low solubility in H2SO4 (10-8 M [5]), the amount of lost formed film (15 mCcm-2, for the first cycle) can

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not have been dissolved into the solution in the corresponding elapsed time. It must remain attached to the electrode surface. This is in agreement with the fact that, in successive cycles, the qa diminishes up to become equal to that of qc (not shown in the figure). At this stage, the voltammetric i/E plot becomes stationary. A demonstration that this non reduced film remains on the surface of the electrode is that, if in a new anodic sweep of the cycling it is made the inversion of va in the nucleation region, this is not observed. Nevertheless, ending the cycling at -1.3 Volts, waiting there 5’ and then, making in a new anodic sweep with inversion in the nucleation region, this last one is observed (like in Fig. 1b). But, why this remaining part of the PbSO4 film is not reduced at the reduction peak? The explanation is that this part is the disrupted one. Then, as they are particles attached by only superficial tension forces, it can not be reduced by the application of a field directly through the electrode potential. They can be reduced only by a dissolution/precipitation with reduction mechanism, which will take longer time (minimum of 5’) at reasonable cathodic potentials. For example, -1.3 Volts, which presents low hydrogen evolution. The Pb free surface can be then totally recuperated by the use of this methodology. With this procedure the anodic peak and its qa become, in the next anodic sweep, those of the first voltammetry on a free Pb surface, like must be waited from the proposed model.

Figure 2. Cycle voltammetries showing the successive formations and reductions of the PbSO4 films during the first cycles. va = 20 mVs-1; vc = 20 mVs-1, Ei = -1.3 Volts, ti = 5’; El,c = -1.2 Volts, tl,c = 0’’ and El,a = -0.7 Volts, tl,a = 0’’.

4.2 Formation/reduction processes and the disruption process In Fig.3a it is seen several examples of experiments of formation and reduction of PbSO4 films for only the first cycle for different va and always the same vc . In Fig. 3b it is seen the evolution of the observed average qa and qc versus va .

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50 40 30 20 10 0 0

(a)

5

10

15

20

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30

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35

40

45

50

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(b)

Figure 3. Voltammetric first cycles for different va detached in the figure and the same vc. (a) voltammetries; (b) corresponding average qa and qc . Ei = -1.3 Volts, ti = 5’; va variable, in the figure, vc = 20 mVs-1; El,c = -1.3 Volts, tl,c = 5’ ; El,a = -0.7 Volts, tl,a = 5’.

The results show that while the qc is constant, the qa decreases with the increase of va , trending practically to qc, for vc £ va. This shows that the formation rate influences the film passivating thickness, trending to a constant value for high rates, as it is well known [2], and that for slower reduction rates than those of formation, there is practically no disruption. But, why there is the constancy of qc in these experiments when qa is varying? This shows that the disruption, for the same rate of reduction, is higher for lower rates of formation than those of reduction. Then, for va lower than vc, the va influences the disrupted amount of qa. This show that the stability against disruption of the film, it is better for high va versus the same reduction perturbation, for va < vc. This fact will need to be theoretically studied. In Fig.4a it is seen several examples of experiments of formation and reduction of PbSO4 films, for only the first cycle, for always the same va and different vc . In Fig. 4b it is seen the evolution of the observed average qa and qc versus vc . In this case the average qa is, of course, constant, but the qc decreases with the increase of vc . One important point is that, for low vc, the qc trends to the value of qa, in agreement with Fig. 3a and 3b. For vc £ va, the initial formed film is totally recuperated: there is no disruption. This fact was yet previously found for galvanostatic results [6] and need to be theoretically explained. The decrease of qc with the increase of vc must be related with the increase of the injection of point defects, which seems to increase the disruption. This fact must also be theoretically studied. Finally, the tendency for a constant value for high vc of about 25 mCcm-2 (20 mCcm-2 for galvanostatic reductions [6]) is due to the remaining stabilizing field explained in section 2.

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Figure 4. Same conditions as those of Fig. 3 but with va constant (5 mVs-1) and vc variable, detached in the figure. (a) voltammetries; (b) corresponding average qa and qc.

5. Conclusions Through cycle voltammetries; inversion of anodic sweeps, for nucleation detection; holding of potential at cathodic values, for occurrence of dissolution/precipitation mechanism; voltammetries with different va and constant vc, or the inverse, it was possible to show the agreement of experimental results with a proposed model. In this model the passivating films suffer disruption (under some conditions), giving rise to external disrupted films and inner continuous films, glued on the metal. The last ones are the only which can be directly reduced, through solid state reactions. The first ones need to follow the dissolution/precipitation, long time consuming. Some of the qa and qc variations are explained by existent theories but other need to be theoretically studied. References 1. N. Sato and G. Okamoto, in “Electrochemical Passivation of Metals”, Comprehensive Treatise of Electrochemistry. Plenum Press, N. Y., v. 4, Chap. 4, pags. 193-245, 1981. 2. C. V. D’Alkaine, CD-ROOM of the “International Corrosion Congress. Frontiers in Corrosion Science and Technology”. International Corrosion Council. Granada, Spain. Plenary Lecture CD-ROOM (18 pags.), (2002). 3.O. A. Albani, J. O. Zerbino, J. R. Vilche and A. J. Arvia, Electroch. Acta 31 (1986) 14031411. 4. D. Pavlov, Electrochimica Acta 23 (1978) 845-854. 5. H. Bode, Lead Acid Batteries, John Wiley & Sons, N. Y. (1977).

6. C. V. D’Alkaine, L. M. M. de Souza, P. R. Impinnisi and J. de Andrade, submitted paper to the J. of Power Sources.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Morphology, composition and structure of anodic films on binary Al-Cu alloys Jonathan Idrac,a Peter Skeldon,b Yanwen Liu,b Teruo Hashimoto,b Georges Mankowski,a George Thompson,b Christine Blanc,a,* a

CIRIMAT - UMR CNRS 5085 - ENSIACET - 118 Route de Narbonne 31077 Toulouse Cedex 04 - France b Corrosion and Protection Centre, School of Materials, The University of Manchester, PO Box 88, Manchester M60 1QD - United Kingdom *corresponding author: [email protected]

Abstract - Anodizing of Al-Cu alloys in ammonium pentaborate solution was shown to proceed in two stages: initial formation of a Cu-free alumina film followed by the introduction of Cu species into the film. The occurrence of the second step depends on the Cu content in the bulk material and highly influences the electrochemical behavior of the anodized alloys. Keywords: copper; aluminum; anodizing; TEM; RBS

1. Introduction There is significant interest in the corrosion behavior in aerated solutions of Cucontaining Al alloys, which remain of importance for aerospace applications. In such alloys, Cu is present in the matrix regions as well as in various fine strengthening phases and coarse intermetallic particles of relatively high Cu contents. In appropriate media, localized corrosion, such as pitting, can occur, with galvanic coupling between the matrix and particles and alkaline corrosion of matrix regions around local cathodic sites. Pit initiation involves passive film breakdown due to differences between the film on the Al matrix and the intermetallic particles. Thus, to understand pitting mechanisms, knowledge of the behavior of the passive films formed on the constituent phases of the alloy is required.

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It is recognized that commercial Al alloys contain several alloying elements in addition to Cu. However, study is focused here on model binary Al-Cu alloys to separate the individual effects of Cu, over a range of concentrations, on film growth [1-4]. Thus, the morphology, composition and structure of oxide films grown during anodizing of Al-Cu alloys are examined, due to their importance to surface performance [5]. Such aspects of films developed on various metallurgical phases are dependent on the Cu content of the substrates [6]. Attention is paid to the Cu enrichment that occurs during anodizing treatments [1, 7]. The enrichment arises since Cu atoms in solid solution in the Al matrix are not oxidized initially, as a Cu-free alumina film thickens on the alloy. In this study, the oxide films have been observed by transmission electron microscopy (TEM) and analyzed by Rutherford backscattering spectroscopy (RBS) and glow discharge optical emission spectroscopy (GDOES). The electrochemical behaviors of alloys have been studied using open circuit potential (OCP) and potentiodynamic polarization measurements in sulphate solutions. The results are compared to those obtained by Strehblow et al. who performed similar studies but with partially different interpretations [8-9]. 2. Experimental Binary Al-Cu alloys containing up to 7 at.% Cu were deposited using an Atom Tech Ltd magnetron sputtering system with separate high purity Al (99.999%) and Cu (99.99%) targets. The substrates, onto which the alloys were deposited, consisted of electropolished 99.99% Al foils. The deposition chamber was first evacuated to 4x10-7 mbar, with sputtering then carried out at 5x10-3 mbar in 99.998% argon at 300 K. The deposited alloys were subsequently anodized to selected voltages at a constant current density of 5 mA/cm2 in 0.1 M ammonium pentaborate electrolyte at 300 K. Suitable electron transparent sections, prepared by ultramicrotomy, of freshly deposited and anodized alloys were examined by TEM using a Tecnai F30 G2 microscope operating at an accelerating voltage of 300 kV. The composition of the alloys and the anodic films, and the enrichment of Cu, were determined by RBS using 2.0 MeV He2+ ions. Depth profiling analyses of the films were carried out by GDOES using a Jobin-Yvon RF 5000 instrument in an argon atmosphere of 538 Pa by applying RF of 13.56 MHz and power of 40 W. Light emissions were monitored, with a sampling time of 0.01 s, at wavelengths of 396.15, 327.40 and 130.22 nm for aluminum, copper and oxygen respectively. Electrochemical tests on the anodized alloys were performed in a 0.1 M sulphate solution at 293 K. OCP values were determined after 1 h of immersion, when the potential had stabilized. Potentiodynamic polarization was also carried out at a scan rate of 1 V/h. Potentials are quoted relative to the saturated calomel electrode.

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3. Results and discussion Transmission micrographs (Fig. 1(a)) reveal alloys consisting of layers of uniform thickness, about 450 nm, with a nano-crystalline structure. Anodizing in ammonium pentaborate solution led to growth of amorphous alumina films (Figs 1 (b,c,d)). Voltage-time responses during anodizing of alloys containing more than ~ 2 at.% Cu revealed three main regions: (i) an initial linear region; (ii) an intermediate region of increasing slope; (iii) a final region of reduced and variable slope (Fig. 2). Micrographs for the Al-3.5 at.% Cu alloy demonstrate that region (i) corresponds to growth of a uniform amorphous alumina film, with enrichment of Cu beneath the film in a ~ 2 nm thick layer of the alloy (Fig. 1(b)). Region (ii) follows the start of oxidation of Cu, with generation of oxygen gas in sufficient amounts to increase the resistance of the film to ion transport. Hence, the voltage rises at an increasing rate in this stage. In this stage, bubbles of oxygen gas are generated in the film and start to grow (Fig. 1(c)). The bubbles are formed at the alloy/film interface whereas the outer part of the film remains free of bubbles [10]. In this period, the film thickness becomes less uniform, since the current flow through the film is disrupted by the presence of the bubbles. Region (iii) coincides with rupture of the film and release of oxygen gas. Thus, during anodizing, the volume occupied by bubbles in the inner region of the oxide progressively increases, the film thickness becomes increasingly irregular, and the alloy/film and film/electrolyte interfaces are further roughened through rupture of the film, due to high oxygen pressures, and healing following rapid local ingress of electrolyte and re-growth of oxide (Fig. 1(d)). Copper enrichment is maintained in the alloy beneath the film in the various stages of film growth in regions (ii) and (iii). GDOES and RBS confirmed the presence of enrichment of Cu in the alloy. Quantification of the

200 nm

200 nm

a

200 nm

b

c

200 nm

d

Fig. 1: Transmission electron micrographs of the Al-3.5 at.% Cu alloy; (a) as-received; (b) anodized to 70 V; with a mainly uniform oxide and occasional oxygen bubbles near the alloy, (c) anodized to 124 V, with relatively numerous oxygen bubbles; (d) anodized to 150 V, with increased roughening of alloy/film and film/electrolyte interfaces. Anodizing was carried out at 5 mA/cm² in 0.1 M ammonium pentaborate solution at 300 K.

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Time (s) Fig. 2: Voltage-time responses during anodizing of Al-x at % Cu alloys at 5 mA/cm2 in 0.1 M ammonium pentaborate solution at 300 K. x = 0 to 7.

Cu enrichment by RBS revealed that Cu is oxidized and incorporated into the oxide when a sufficient Cu content in the enriched layer is achieved. With achievement of this threshold Cu level, the voltage-time response departs from linear behavior with further anodizing since oxygen bubbles are formed. For other alloys, i.e. containing less than about 2 at.% Cu, the three regions are not evident in the voltage-time responses. Thus, either the Cu has not enriched sufficiently to enter the oxide, i.e. with no generation of oxygen gas in the film, or the Cu species have entered the oxide, but the generation of oxygen is still too low to affect significantly the generally uniform resistance of the film to ionic transport. The RBS and GDOES results for the Al-1.8 at.% Cu alloy indicate the presence of enrichment on anodizing to 150 V, but little or no Cu in the oxide, unlike the behavior for alloys of increased Cu content, for which Cu was clearly located within the film. The corrosion behaviors of Al-Cu alloys, both before and after anodizing, were examined in electrochemical tests in sulphate solution. OCP values are reported on Figure 3. Corrosion potential (Ecorr) values were also obtained from potentiodynamic polarization curves, which revealed a wide passivity range, with anodic current densities of about 2 mA/cm2. Figure 3 shows that OCP values measured on Al-Cu alloys before anodizing progressively increase with Cu content in the bulk alloy, by about 570 mV in the range 0 to 7 at.% Cu. When the alloys are anodized to 50 V, the OCP values are increased relative to those of the non-anodized alloys, associated with the thickening of the oxide film. Further, the potentials of the anodized alloys increase by about 150 mV with increased Cu content of the alloy. The voltage-time responses during anodizing to 50 V are similar for the alloys, corresponding mainly to region (i), with the growth of an alumina film and Cu enrichment in the alloy. The oxide is further thickened with anodizing to 150 V, with OCP values showing a stronger increase with Cu content of the alloy than for anodizing to 50 V. Alloys of high

Morphology, composition and structure of anodic films on binary Al-Cu alloys 0

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Fig. 3: OCP and Ecorr measured in 0.1 M sulphate electrolyte versus x for Al-x at.% Cu alloys anodized at 5 mA/cm2 in 0.1 M ammonium pentaborate solution. ( ) without anodizing, (o) anodizing stopped at E=50 V, (·) anodizing stopped at E=150 V (117 V for Al-7 at.% Cu).

Cu contents reach region (iii) of the voltage-time response, with Cu present in the oxide film and extensive generation of oxygen. For the more dilute alloys, behavior is limited to region (i), with Cu remaining in the alloy. Ecorr measurements reveal similar values for anodizing to 50 and 150 V, with an increasing potential, by up to about 800 mV, for alloys of increasing concentration, to about 3 at.% Cu, followed by a plateau region. The increased positive shift in Ecorr following anodizing, compared with that of OCP, may be due to influence of prior cathodic polarization in the potentiodynamic measurements, resulting in local film thinning at flaws due to increased alkalinity. Such thinning may facilitate cathodic reduction of oxygen, with increase in potential. The observed potential changes of the alloys following anodizing reflect the altered contributions of anodic and cathodic current on the specimen surfaces with thickening of the film from the native oxide, of a few nanometers thickness, to insulating anodic alumina films about 60 and 180 nm thick at 50 and 150 V. Anodizing tends to heal surface defects, reducing anodic currents, although cathodic activity remains from residual flaw sites. Such changes can lead to the observed increases in Ecorr and OCP after anodizing. Further, the anodic and cathodic reactions may also be influenced by the enrichment of Cu in the alloy, that can reach levels of up to 40 at.% for alloys containing 1 at.% Cu in the bulk material. The enrichment can increase OCP values in the absence of Cu incorporation into the oxide, which may also be associated with reduced anodic currents or increased cathodic currents [7]. Strehblow et al. [8-9] proposed a detailed electrochemical and RBS study dealing with similar observations except for the Cu-enriched layer which was 65 nm thick as estimated by these authors while our STEM observations revealed a 2 nm thick uniform enriched layer. They also observed a roughening of the alloy/oxide interface which was attributed to Cu-islands on which the

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alumina film can not grow; its breakdown due to a local high electric field led to direct exposure of Cu to the electrolyte and to the formation of CuO and Cu2O islands. Oxygen evolution can occur on these electron conductive sites. These explanations constitute another hypothesis to explain the anodic film rupture and the generation of oxygen bubbles that we observed in our study. In the same way, these conductive sites could determine the OCP values measured for anodized alloys and explain the shift observed for OCP on fig. 3. 4. Conclusions Anodizing of Al-Cu alloys in ammonium pentaborate solution proceeds in two stages: initial formation of a Cu-free alumina film, followed by the introduction of Cu species into the film when a sufficient enrichment of Cu is achieved at the alloy/film interface. Oxygen gas at high pressure in generated within the film following incorporation of Cu. The open-circuit and corrosion potentials of alloys increase by about 500 mV with increase of Cu content to 7 at.%, and are further increased by about 200 - 500 mV by thickening of the surface oxide by anodizing. The increased potential with film thickening probably reflects a reduced anodic current, with a significant remaining cathodic activity. Acknowledgments - The authors are grateful to the Groupe de Physiques des Solides for provision of time on the Van de Graaff accelerator of the Universités de Paris VI et VII. They also wish to acknowledge assistance to JI through a Marie Curie Training Site funded under the EC Programme “Improving Human Research Potential and the Socio-economic Knowledge Base’.

References 1. H. Habazaki, M.A. Paez, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood and X. Zhou, Corros. Sci., 38 (1996) 1033 2. X. Zhou, G.E. Thompson, P. Skeldon, G.C. Wood, K. Shimizu, H. Habazaki, Corros. Sci., 41 (1999) 1089 3. Y. Liu, E.A. Sultan, E.V. Koroleva, P. Skeldon, G.E. Thompson, X. Zhou, K. Shimizu, H. Habazaki, Corros. Sci., 45 (2003) 789 4. S. Garcia-Vergara, P. Skeldon, G.E. Thompson, P. Bailey, T.C.Q. Noakes, H. Habazaki, K. Shimizu, Appl. Surf. Sci., 205 (2003) 121 5. P. Schmutz and G. S. Frankel, J. Electrochem. Soc., 146 (1999) 4461 6. M. A. Paez, O. Bustos, G.E. Thompson, P. Skeldon, K. Shimizu and G.C. Wood, J. Electrochem. Soc., 147 (2000) 1015 7. S. Garcia-Vergara, F. Colin, P. Skeldon, G.E. Thompson, P. Bailey, T. C. Q. Noakes, H. Habazaki and K. Shimizu, J. Electrochem. Soc., 151 (2004) B16 8. H.H. Strehblow and J.C. Doherty, J. Electrochem. Soc., 125 (1978) 30 9. H.H. Strehblow, C.M. Melliar-Smith and W.M. Augustyniak, J. Electrochem. Soc., 125 (1978) 915 10. P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, H. Habazaki, K. Shimizu, Phil. Mag. A 76, 729 (1997).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Repassivation of aluminium during the ACgraining process by aluminium hydroxide formation Thanasis Dimogerontakisa,*, Herman Terryna, Paola Campestrinib a

Dept. of Metallurgy, Electrochemistry and Materials Science, Vrije Universiteit Brussel, Pleinlaan 2, 1050 Brussels, Belgium *([email protected]) b

Agfa-Gevaert, Septestraat 27, 2640 Mortsel, Belgium

Abstract : In this work the AC-graining in hydrochloric acid of the aluminium alloy 1050 was studied. The influence of the electric charge and of the current density on the smut layer formation and on the graining morphology was investigated. Key words: aluminium, graining, smut layer, repassivation, lithography

1. Introduction The AC-graining is extensively used for the roughening of the aluminium lithographic printing plates. The roughening plays a significant role in obtaining high performance printing plates both in terms of image clarity and plate longevity because it improves the adhesion of the photosensitive coating and enhances the water retentive properties [1, 2]. The AC-graining is carried out by applying an alternative current on the aluminium plates in a suitable electrolyte, which is mainly based on hydrochloric or nitric acid. During the anodic semicycle of the applied ACcurrent, aluminium dissolution takes place according to preferential crystal directions in the case of the HCl based electrolytes. Therefore, a high number of cubic pits are generated each half period. On the other hand, hydrogen evolution takes place during the cathodic semicycle. This hydrogen evolution results in increasing the interfacial pH, which finally provokes the precipitation of aluminium hydroxide and the formation of the smut layer. This layer is mainly composed of amorphous aluminium hydroxide with an enormous quantity of water. The smut layer causes a partial repassivation of the aluminium substrate.

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This repassivation influences significantly the next generation of the cubic pits during the following anodic semicycle [3]. The AC-graining process finally creates a large number of hemispherical pits on the aluminium surface. These pits are broader than deep and grow by lateral expansion until they merge. The hemispherical pits are actually composed of a very high number of small cubic pits created by the crystallographic attack of the aluminium substrate during the anodic semicycle [4, 5]. In this work the influence of the electric charge and of the current density on the smut layer formation as well as on the developed anodic and cathodic potentials during the AC-graining in hydrochloric acid was investigated. The effect of the smut layer formation on the developed graining morphology was also studied. 2. Experimental For the electrodes an aluminium alloy 1050 was used in this work. The electrolytic cell consisted of two parallel aluminium electrodes in a distance of 2 cm. The AC-graining was performed in 0.34 M HCl at 37 oC by applying a constant alternating (50 Hz) current density in the range of 40 to 120 A dm-2. A galvanostat (Spitzenberger GSM 2000) in combination with a wave generator (Wavetek M29) was used for applying the sinusoidal current. The potential was measured versus a reference electrode (Ag/AgCl) and was continuously recorded by a high frequency acquisition system (National Instrument PCI6024E). The weight of the smut layer was gravimetrically estimated by the removal of the layer with phosphoric – chromic acid [6]. The graining morphology was examined by Scanning Electron Microscopy (JEOL 6400) after the removal of the smut layer. 3. Results and discussion The weight of the smut layer vs. the consumed electric charge in the range of 480 to 1440 C dm-2 was determined for samples grained in 0.34 M HCl by 120 A dm-2 (Figure 1). It was observed that the weight of the smut layer increased linearly with the increase of the consumed electric charge in the examining range. The potential signal was recorded from the beginning of the AC-graining until the maximum examined electric charge (Figure 2a). It was observed that the cathodic potential was a bit larger than the anodic one at the first cycle of graining. Afterwards, a significant increase on the cathodic potential took place with the increase of the electric charge. The anodic potential also increased but much less than the cathodic one. These observations indicated that the repassivation induced by the increase of the smut layer formation with the increase of the consumed electric charge, resulted in significantly increasing the

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cathodic potential, while a relatively small increase on the anodic potential was resulted (Figure 2b). Thus, a more intensive hindrance on the cathodic reaction, which is actually the hydrogen evolution, took place with the smut later formation. This more intensive hindrance could be attributed to the fact that according with the literature [7], the hydrogen evolution in the AC-graining is a mass transport controlled reaction. Therefore, the formation of a thick hydrated smut layer might significantly hinder the reduction of the hydrogen ions. Weight of the smut layer / gm-2

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Figure 1: Weight of the smut layer vs. electric charge for samples grained in 0.34 M HCl at120 A dm-2 at 37 oC. (a)

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Figure 2: (a) Signal of the potential after 1.2, 480, 960 and 1440 C dm-2 and (b) anodic and cathodic potential vs. weight of the smut layer.

Aluminium samples were grained up to 1440 C dm-2 by applying current densities in the range of 40 to 120 A dm-2 in order to investigate the influence of the current density on the smut layer formation. The weight of the smut layer vs. the applied current density was plotted (Figure 3a). It was obvious that the smut layer formation increased about linearly with the increase of the current density.

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The anodic and cathodic potentials were also measured after graining to 1440 C dm-2 with different current densities in the range of 40 to 120 A dm-2 (Figure 3b). It was observed that there was a significant increase of the cathodic potential and a much smaller increase of the anodic potential with the increase of the current density. This indicated that the increase of the smut layer formation induced by the increase of the current density influenced both potentials but again much more the cathodic one. 30

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Figure 3: (a) Weight of the smut layer vs. current density and (b) anodic and cathodic potential vs. current density for Al samples grained in 0.34 M HCl by 1440 C dm-2 at 37 oC.

The graining morphology, which was developed by the different current densities, was examined by Scanning Electron Microscopy (SEM). It was obvious even by visual observation of the micrographs (Figure 4) that the increase of the current density and so of the smut layer formation resulted in increasing the size of the hemispherical pits. It also seems that a less uniform graining morphology was developed by increasing the current density.

(a)

(b)

Figure 4: SEM micrographs of samples grained in 12.5 g/l HCl with 960 C dm-2 at 37 oC by (a) 40 and (b) 120 A dm-2.

In order to extract more precise information concerning the influence of the current density and smut layer formation on the graining morphology, image

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analysis was applied to the micrographs of samples grained to a charge of 480 C dm-2 at all the current densities examined (40, 60, 80 and 120 A dm-2). The pit size distribution histograms (Figure 5) indicated that, as the current density increased, the number of hemispherical pits decreased and the dispersion of their size distribution increased. Image analysis indicated also that the pits population density decreased from 3.33 1010 m–2 at the lowest current density to 2.15±0.04 1010 m-2 at the highest current density. In addition, the increase of the current density resulted in increasing the extent of the plateau (non-attacked) areas from 70.3±1.0 % at the lowest current density to 76.9±0.7 % at the highest. The standard deviation of the mean value of the pits size was calculated by the statistic treatment of the image analysis results. It was observed that the increase of the current density induced an increase on the calculating standard deviation from 9.0±0.4 at the lowest current density, and the lowest smut layer formation, to 12.2±0.6 at the highest. This increase on the standard deviation confirmed the increase of the dispersion of the pits size distribution that was also observed by the respective histograms, indicating that a less uniform graining morphology was developed with the increase of the current density. 40 A dm-2

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The quality of the roughening morphology of the aluminium substrate is one of the most important parameters in producing effective aluminium lithographic printing plates since it significantly affects the printability and durability of the plates. The best roughening quality for the printing plates is a uniform graining morphology with relatively small hemispherical pits and as small as possible extent of non-attacked (plateau) areas. The examined graining morphologies indicated that a significant degradation of the quality of the roughening for aluminium printing plates took place with increasing the current density and thereby, increasing the smut layer formation. However, the repassivation induced by the smut layer formation did not depend only on the total weight of

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the layer but also on the uniformity of its thickness. Therefore, further work is in progress for determining the influence of the graining conditions on the thickness of the smut layer. 4. Conclusions The smut layer formation increased linearly with the increase of the consumed electric charge during AC-graining in hydrochloric acid. The repassivation induced by the smut layer caused a significant increase of the cathodic potential that is developed during the graining. The anodic potential also increased by the smut layer formation but to a much smaller extent compared with the cathodic one. This indicated that the smut layer caused a larger hindrance on the cathodic reaction, which is actually the hydrogen evolution, than on the anodic. This significant hindrance on the cathodic reaction can be attributed to the fact that the smut layer is a quite thick hydrated layer, which can significantly affect the hydrogen evolution since it is a mass transport controlled reaction. The smut layer formation also increased about linearly with the increase of the applied current density. This increase of the smut layer, in addition to the influences on the anodic and cathodic potentials, also had significant effects on the developing graining morphology. A less uniform graining morphology with less and larger hemispherical pits and a larger extent of non-attacked (plateau) areas was developed with the increase of the current density. Thus, a significant degradation of the quality of the roughening for aluminium printing plates took place with the increase of the current density and so the increase of the repassivation induced by the smut layer formation. 5. Acknowledgements The authors thank An Dotremont for her help on the experimental part of this work. Agfa-Gevaert N.V. and the Institute for the Promotion of Innovation by Science and Technology in Flanders (IWT) are acknowledged for the financial support of this project. References 1. 2. 3. 4. 5. 6. 7.

J. Vermeersch, E. Verschueren. M. Van Damme, ATB Metallurgie, 43, 130 (2003). R. Hutchinson, Trans. IMF, 2001, 79(4), B57. H. Terryn, J. Vereecken, G.E. Thompson, Trans IMF, 66, 116 (1988). H. Terryn, J. Vereecken, G.E. Thompson, Corros. Sci., 32, 1159 (1991). H. Terryn, J. Vereecken, G.E. Thompson, Corros. Sci., 32, 1173 (1991). ASTM G1-82. P. Laevers, Study of the Mechanism of the AC-Eletcrolytic Graining of Aluminium, Thesis, University of Brussels, 1995.

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Anodization of Ti : Formation of Self-Organized Titanium Oxide Nanotube-Layers P. Schmukia*, H. Tsuchiyaa, L. Taveirab, K. Sirotnac and J. M. Macaka a

University of Erlangen Nuremberg, Dept. of Materials Science, LKO, Martensstr. 7, 91058 Erlangen, Germany b on leave from: UFRGS - Universidade Federal do Rio Grande do Sul Av. Osvaldo Aranha 99, 90035-190-Porto Alegre, Brasil c on leave from: Institute of Chemical Technology Prague, Dept. of Power Engineering, CZ-16628 Prague, Czech Republic

Abstract - The work reports on the electrochemical fabrication of layers of selforganized high aspect ratio titanium oxide nanotubes achieved by an optimized and controlled anodization of Ti in fluoride containing solutions. In general, the morphology and the structure of porous layers are affected strongly by the electrochemical parameters used. Under optimized conditions self-organized relatively highly ordered nanotubes with a length between some 100 nm and several µm are formed consisting of arrays with single tube diameter between 10 and 100 nm and a tube wall of approx. 10 nm thickness. The work discusses factors of the electrochemical pore formation process and key electrochemical factors that lead to self-organization. Keywords : self-organization, porous titanium oxide, tubular structure, electrochemical anodization

1. Introduction Self-organized nanostructures of different metals or semiconductors have received considerable attention due to the anticipated high technological potential of these materials. For a range of metals such as Al [1-4], Ti [5-14], Zr [15-17], Nb [18], W [19-20], Ta [21, 22], Hf [23] it has recently been found that self-organized porous structures can be formed under optimized electrochemical treatments. These nanoarchitectural oxide films can have very specific

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functional properties. For example, titanium oxide (TiO2) is a promising material because of gas sensing capabilities, self-cleaning ability, and its use in solar cells. Nanoporous TiO2 has been formed by electrochemical anodization of Ti in acid electrolytes containing small amounts of hydrofluoric acid (HF) [5-7]. A drawback of acidic HF-based electrolytes is that typically porous TiO2 layers have been grown only up to thicknesses of about 500 nm, unlike Al, Hf or Zr where porous films with thicknesses of several ten micrometers can be obtained. Recently we obtained high aspect ratio self-organized porous TiO2 structures with thickness higher than 2 µm by anodization in buffered neutral electrolytes containing NaF and NH4F instead of HF [8,9]. Thus, by exploiting the decrease of the chemical etching rate of TiO2 in these neutral solutions and establishing a pH gradient by adjusting the electrochemical conditions high aspect ratio TiO2 nanotubes can be grown. The present work discusses various aspects of the pore formation process and several critical factors affecting the pore – or nanotube – morphology. 2. Experimental For electrochemical experiments, Ti foils (from Goodfellow, England) of 0.1 mm thickness and purity of 99,6% were used. Prior to anodization, the samples were degreased in an ultrasonic bath in acetone, isopropanol and methanol, followed by a deionized water rinse and then dried in a nitrogen stream. The samples were contacted by a Cu plate and then pressed against an O-ring in an electrochemical cell, leaving 1 cm2 exposed to the electrolyte. The electrochemical measurements were carried out in a classical three-electrode configuration with a Ag/AgCl (1 M KCl) electrode as a reference electrode and a platinum grid as a counter-electrode. The Ti samples were anodized using a high-voltage potentiostat Jaissle IMP 88 interfaced to a computer. The electrolytes were various acids and salts with additions of small amounts of HF and NH4F (0.05 - 5 wt.%). The experiments were performed at room temperature under aerated non-stirred conditions. All solutions were prepared from chemical reagents of high purity. For anodization, the potential was typically continuously increased from the open-circuit potential (OCP) to an end potential with several scan rates and then held at the end potential for different times. For surface characterization, the specimens were rinsed with deionized water and dried with nitrogen. Scanning electron microscope images were taken for the morphological characterization of the porous titanium oxide using a Hitachi SEM FE 4800. Direct SEM cross-sectional thickness measurements were carried out on mechanically bent samples, where cracking and a partial lift-off of the porous layer occured.

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3. Results and Discussion In general, the morphology (and the thickness) of the anodic oxide layers formed on Ti strongly depends on the electrochemical parameters such as the applied potential, the time of anodization, or the sweep rate of the potential ramp. Additionally, the HF-concentration in the electrolyte must be optimized to achieve self-ordered pore growth. Anodization in solutions with a low HFconcentration, for instance <0.05 wt.% HF in 1 M H2SO4 solution, yields large nonregular pores. On the other hand, anodization in solutions containing > 0.5 wt.% HF leads to uniform etching of the surface (electropolishing). 3.1. Initiation of porous layers A feature common to all porosification experiments is that establishing the final steady state of the self-organized etching process requires an induction time. Figure 1a shows schematic current-time transients for anodization of Ti at the “right” voltage in an electrolyte e.g. H2SO4 in absence and in presence of the “right” amount of fluoride. In pure background electrolyte, the typical exponential current decay is observed due to the growth of a compact oxide layer. For electrolytes containing fluorides, clear deviations from the behavior are observed. In this case, after an initial exponential decay (phase I) the current increases again (phase II) with a time lag that is shorter, the higher the fluoride concentration. Then, the current reaches quasi-steady state (phase III). This steady state current increases with increasing fluoride concentration. This type of current-time curve has been previously reported for self-organized pore formation for other materials, as well [24]. Typically, such a current behavior has been ascribed to different stages in the pore formation process, as schematically illustrated in Fig. 1b. In the first stage, a barrier oxide is formed, leading to a current decay (a). In the next stage, the surface is locally activated and pores start randomly growing (b). Due to pore growth, the active area increases and the current increases. After some time, many pores have initiated and a tree-like growth takes place. Therefore, the individual pores start interfering with each other, and start competing for the available current. This leads under optimized condition to a situation where the pores equally share the available current, and self-ordering under steady-state conditions is established (c). Indeed, direct SEM cross-sectional observation take for each sample at the three phases delineated in Fig. 1 confirms the above description regarding nanotube formation. Fig. 2 shows samples removed from the electrolyte after different times – shown is the evolution of the morphology of the porous titanium oxide in 1 M (NH4)2SO4 + 0.5 wt. % NH4F. Immediately after the sweep from the OCP to 20 V, a thin and rough oxide film (» 50 nm) covers all the surface of the titanium (Fig. 2a). The results from XPS measurements show that the film at this stage is composed of Ti, O and some traces of F. The XPS data suggest that

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the film is composed of an inner layer of TiO2 and a outer layer of Ti(OH)4. This film is essentialy the high-field passive layer – the thickness very well corresponds to a growth factor of 2.5 nm / V. At approximately 3 min, first signs of localized attack become apparent with the FE-SEM; breakdown sites that are randomly distributed over the film surface and round shaped holes in the substrate can be observed. Then, the current increases due to the growth of the pores. At this point, the SEM image reveals that underneath the relatively compact top oxide film some pore structures become visible. The porous oxide film shows in the cross-section a thickness of 200 to 300 nm and a relatively regular but worm-like structure (Fig. 2b). After 30 min from the start of the polarization, there are areas of porous structure visible but still there are patches of the top compact oxide film visible. With increasing anodization time the current drops further while the porous oxide film thicknesses establishes at 500800 nm (Fig. 2c). Anodization for longer times leads to a further establishment of the homogenous self-organized porous structure. 3.2. Thickness and morphology of the porous layers The thickness of the porous layers was determined from SEM cross-sections of mechanically fractioned samples. For anodization in 1 M H2SO4 + 0.15 wt.% HF the thickness of the porous layer was found to first steeply increases from about 47 nm at 20 minutes to 540 nm after 10 hours. This is followed by a steady state with a plateau of thickness values at approximately 580 nm. The fact that, both the current and the porous layer thickness reach a limiting value after a certain polarization time can be explained by a steady-state situation. Both, pore growth at the inner interface and dissolution of the oxide layer at the outer interface take place. Steady state is established when the pore growth rate at the inner interface is identical to the dissolution rate of the oxide film at the outer interface. This limiting porous oxide layer thickness can be overcome by the electrochemical conditions [9]. That is, higher thickness of the porous layer can be obtained, if the anodization treatment is carried out in less aggressive electrolytes. One explains the fact that the hydrolysis reaction during anodization leads to acidification of the pore tips (similar to pitting corrosion). This acidification accelerates the chemical TiO2 dissolution exactly there, where it is desired: at the pore tip (while the rest of the porous structure remains relatively stable due to the more alkaline pH of the background electrolyte) – for more details see Ref. [9]. Fig. 3 and Fig. 4 show top views and crosssectional SEM images of a range of porous TiO2 layers produced in different electrolytes. Clearly the thickness of the oxide layer depends on the aggressivity of the electrolyte and is in line with the idea that the pore growth mechanism is controlled by competing film growth and dissolution. Another crucial factor that was found in various electrolytes to affect the pore morphology is the potential scan rate. This is also well in line with the approach

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that local acidification crucially determines the morphology. The higher the scan rate, the higher the current density and therefore the lower the pH at the pore tip and the higher the H+ flux from the tip to the pore mouth [9]. Hence it is understandable that only optimized scan rates (and therefore pH – profiles within the pore) can produce high aspect ratio tubular structures. 3.3. Electrolytes and pore dimensions The results in Fig. 3 and 4 clearly show that the morphology of the single pores and the pore dimensions that are the pore diameter, the pore spacing, the pore lenght and the occurance of a rippled side-wall strongly depend on the electrolyte used. Fig. 3 and Fig. 4 show pores produced using H2SO4, (NH4)2SO4, CH3COOH and (NH4)H2PO4 electrolytes. The pore diameter varies from a 15 nm in CH3COOH to approx. 100 nm in (NH4)2SO4. The pore lenght can be varied from approx. 200 nm to about 5 mm. 4. Conclusions Self-organized porous TiO2 layers can be electrochemically prepared by anodization in electrolytes containing small amounts of F- species. The self-organized pore growth takes place as a steady-state process, as a competition of activation and passivation reactions. Hence, a limiting thickness of the porous oxide layer is established. By optimizing the anodization parameters, self-organized porous TiO2 layers with a comparably high degree of order and various dimensions and aspect ratios can be achieved. References [1] H. Masuda, K.Fukuda, Science 268 (1995) 1466. [2] H. Masuda, K. Yada, A. Osaka, Jap. J. Appl. Phys. 37 (1998) L1340. [3] H. Masuda, K. Fukuda, Appl. Phys. Lett. 71 (1997) 2770. [4] O. Jessensky, F. Muller, U. Gösele, Appl. Phys. Lett. 72 (1998) 1173. [5] V. Zwilling, E. Darque-Ceretti, A. Boutry-Forveille, Electrochim.Acta 45(1999) 921. [6] D. Gong, C.A. Grimes, O.K. Varghese, W. Hu, R.S. Singh, Z. Chen, E.C. Dickey, J. Mater. Res. 16 (2001) 3331. [7] R. Beranek, H. Hildebrand, P. Schmuki, Electrochem. Solid-State Lett. 6 (2003) B12. [8] J. M. Macak, K. Sirotna, P. Schmuki, Electrochim. Acta 50 (2005) 3679. [9] J. M. Macak, H. Tsuchiya, P. Schmuki, Angew. Chem. 44 (2005) 2100. [10] H. Tsuchiya, J.M. Macak, L. Taveira, E. Balaur, A. Ghicov, K. Sirotna, P. Schmuki, Electrochem. Commun. 7 (2005) 576. [11] L. V. Taveira, J. M. Macak, H. Tsuchiya, L. F. P. Dick, P. Schmuki, J. Electrochem. Soc. 152 (2005) B405. [12] A. Ghicov, H. Tsuchiya, J. M. Macak, P. Schmuki, Electrochem. Commun., 7 (2005) 505.

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[13] J. M. Macak, H. Tsuchiya, L.Taveira. A.Ghicov, P. Schmuki, J.Biomed.Mat.Res.A, in press. [14] P. Schmuki, J. M. Macak, H. Tsuchiya, in Proc. of the Symposium "Pits and Pores: Formation, Properties and Significance for Advanced Materials", Ed. P. Schmuki, D.J. Lockwood, H.S. Isaacs, Y. Ogata, M. Seo The Electrochemical Society, PV of ECS Hawaii meeting 2004. [15] H. Tsuchiya, P. Schmuki, Electrochem. Commun. 6 (2004) 1131. [16] W.J. Lee and W.H. Smyrl, Electrochem. Solid-State Lett. 8 (2005) B7. [17] H. Tsuchiya, J. M. Macak, I. Sieber, P. Schmuki, Small 1 (2005) 722. [18] I. Sieber, H. Hildebrand, A. Friedrich, P. Schmuki, Electrochem. Commun. 7 (2005) 97. [19] H. Tsuchiya, J. M. Macak, I. Sieber, L. Taveira, A. Ghicov, K. Sirotna, P. Schmuki, Electrochem. Commun. 7 (2005) 295. [20] N. Mukherjee, M. Paulose, O. K. Varghese, G. K. Mor and C. A. Grimes, J. Mater. Res. 18 (2003) 2296. [21] I. Sieber, B. Kannan, P. Schmuki, Electrochem. Solid-State Lett. 8 (2005) J10. [22] I. Sieber and P. Schmuki, J. Electrochem Soc. 152 (2005) C639. [23] H. Tsuchiya, P. Schmuki, Electrochem. Commun. 7 (2005) 49. [24] V.P. Parkhutik and V.I. Shershulsky, J. Phys.D 25 (1992) 1258.

II

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Fig 1a) Schematic polarization curve and stages (I,II,III) of pore formation.

Fig.1b) Schematic representation of the stages (I=a, II=b, III=c) of selforganized pore formation.

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a)

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Fig. 2. SEM top-view and cross-section images show the evolution of porous titanium oxide morphology after anodizing for (a) 0 – 1 min; (b) 6 min; (c) 20 min; in 1 M (NH4)2SO4 + 0.5 %wt. NH4F electrolyte at 20 V (after a potential sweep from open-circuit potential to 20 V with weep rate of 1 V s-1).

d = 10 nm

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Fig. 3. Influence of the electrolyte on tube diameter [8-10, 12]. Top-view images of the porous titanium oxides prepared in (a) CH3COOH/NH4F, (b) (NH4)H2PO4/NH4F, (c) H2SO4/HF and (d) (NH4)2SO4/NH4F.

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Self-Organized Nanoporous Valve Metal Oxide Layers Hiroaki Tsuchiya,a Jan Macak,b Irina Sieber,c Luciano Taveira,d Patrik Schmukie Department of Materials Science, LKO, University of Erlangen-Nuremberg, Martensstrasse 7 D-91058 Erlangen Germany a [email protected]; [email protected]; [email protected]; [email protected]; [email protected]

Abstract: The present work examines that formation of self-organized nanoporous oxide layers can be achieved on a whole series of valve metals directly via anodization in fluoride-containing electrolytes. In particular, nanoporous layers on IVB valve metals – Ti, Zr and Hf – are presented. The remarkable difference between the metals is the thickness, which may be essentially ascribed to the difference in chemical dissolution rates of each oxide. Morphology, structure and thickness are strongly affected by the electrolyte pH. Keywords: self-organization; nanoporous; nanotubes; valve metal; fluoride; electrochemical structurization

1. Introduction Nano-scale structurization on metals, alloys and semiconductors by exploiting localized dissolution (controlled localized corrosion) reactions has scientific and technological interests in recent years due to the simplicity of process and high expectations regarding applications. Self-organization or selfstructurization on such materials can be initiated under optimized electrochemical conditions. In particular, one-dimensional porous Al2 O3 with a high aspect ratio has been widely and intensively investigated concerning with formation mechanisms and applications such as sensing or as a photonic crystal[1-4]. In further applications, the one-dimensional materials have been used as a template for secondary material deposition such as metal,

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semiconductor or polymer[5, 6]. For other metals than Al, such as Ti[7, 8], only little information was available on the formation of porous structures until a short time ago. Recently we have achieved the formation of self-organized porous structures on a whole range of valve metals, for example, Ti[9-14], Zr[15, 16], Hf[17], Ta[18], W[19] and Nb[20]. In the present work, we report the formation of self-organized porous oxide layers mainly on IVB group metals. 2. Experimental Ti (99.6 % purity), Zr (99.8 %) and Hf (97.0 %) samples with 0.1 mm thickness were degreased by sonicating in acetone, isopropanol and methanol, followed by rinsing with deoinized (DI) water and drying in a nitrogen stream. The electrochemical setup consisted of a conventional three - electrode configuration with a platinum gauze as a counter electrode and a Haber Luggin capillary with Ag/AgCl (1 M KCl) electrode as a reference electrode. Electrochemical experiments were carried out using a high - voltage potentiostat Jaissle IMP 88 in fluoride-containing electrolytes at room temperature. The electrochemical treatment consisted of a potential ramp from the open - circuit potential (OCP) to a fixed potential, Efix, with a different sweep rate followed by holding the applied potential at Efix for different times. After the electrochemical treatment the samples were rinsed with DI water and dried in a nitrogen stream. Scanning electron microscopy was employed for the structural and morphological characterization of the porous titanium oxide using a Hitachi SEM FE 4800. In order to gain information on thickness of the porous layer direct SEM cross-sectional thickness measurements were carried out on mechanically bent samples – while bending the porous layer cracked into many parts and where a profitable partial lift-off of the porous layer occurred. Phase identification of the porous layers was carried out with X-ray diffractometer (Phillips X’pert-MPD PW3040). 3. Results and discussion Some examples of nanoporous structures formed in 1M H2SO4 containing small amounts of fluoride ions are shown in Fig. 1. It is apparent that selforganized porous structures can be achieved on a whole range of valve metals by anodization in fluoride-containing electrolytes. The average pore diameter strongly depends on the metals, that is, approx. 100 nm for Ti and W, 50 nm for Zr and Hf and 20 nm for Ta and Nb. In fluoride-free electrolytes, a compact oxide layers were present on the surface as expected. Therefore, fluoride ions are considered to act as chemical agents to dissolve electrochemically formed oxides by forming soluble complexes, that is, a key to achieve pore formation on such metals is the dissolution of the anodic oxides as soluble complexes. In the following section, results regarding IVB group valve metals, Ti, Zr and Hf

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will be presented. Figure 2 shows cross-sectional images of porous structures formed on (a) Ti, (b) Zr and (c) Hf. As apparent from this figure, features of cross-sections on Ti are significantly different from those on Zr and Hf. Nanotube structures are formed on Ti, but on Zr and Hf, on the other hand, relatively wavy tubular porous structures are formed in highly parallel configuration. Furthermore, the thickness of TiO2 nanotube layers is thinner than those of ZrO2 and HfO2 porous layers. This can be partially ascribed to the total electric charge generated during anodization, that is, comparably higher electric charge was generated for Zr and Hf than for Ti[21]. However it is impossible to explain from the point of this view as the thickness difference observed in Fig. 2 is much higher than that estimated from the electric charge. Therefore, other factors have to be considered to determine thickness of oxide layers, for example, dissolution rate of the anodic oxide layers by fluoride ions. The electrochemical conditions can influence on morphology, structure and thickness of oxide layers. Figure 3 shows clear examples – cross-sectional SEM images of layers formed on (a) Ti, (b) Zr and (c) Hf in neutral buffer (NH4)2SO4/NH4F electrolytes. Comparing Fig. 3(a) with Fig. 2(a) it is obvious that thickness of the oxide layer formed on Ti in a neutral electrolyte is relatively thicker than that in an acidic electrolyte. This difference is attributed to pH dependence of oxide dissolution rate, that is, the dissolution rate is higher at low pH than that at high pH[10, 11]. Therefore the thickness of nanotube layers on Ti strongly depends on the electrolyte pH though the morphology of nanotubes is same independent of the pH. For Zr and Hf, the features of the topview of the layers formed in neutral buffered electrolytes are similar to that in an acidic electrolyte shown in Fig. 1. However it is noteworthy that in the buffered electrolytes extremely straight-featured nanotubes are formed on Zr and Hf (see Fig. 3(b) and (c)). It is apparent from Figs. 2 and 3 that the crosssectional feature is relatively different depending on the electrolytes although the thicknesses of both layers are almost same independent of the electrolytes. The factor to make the difference in the cross-sectional features could be buffering effect of (NH4)2SO4/NH4 F system. Therefore, the electrolyte can strongly affect morphology and thickness of the resulting layers. Effect of other electrochemical parameters such as sweep rate and applied potential were reported in our previous works[15]. 4. Conclusion Self-organized nanoporous oxide layers can be formed on a whole range of valve metals by electrochemical anodization. A key to achieve self-organized pore formation on such metals is to tune the situation where oxide formation and local dissolution of the oxide take place at the same time. The morphology, structure and thickness are strongly affected by the electrochemical conditions such as electrolyte, sweep rate and applied potential. In the present work the

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effect of electrolyte is stressed. For Ti, pH strongly affect the thickness of the oxide layers – in an acidic electrolyte thickness is approx. 400 nm and on the other hand, in a neutral electrolyte approx. 2 µm. Anodic TiO2 can be dissolved much faster by fluoride ions in the acid than in the neutral electrolyte. This pH dependence of fluoride-induced oxide dissolution rate must result in the achievable thickness depending on the electrolyte pH used. For Zr and Hf, cross-sectional morphology is strongly affected by the electrolyte. In the acid relatively wavy tubular pores are produced. In the neutral electrolyte, on the other hand, extremely straight tubes are formed. This could be attributed to the difference of chemical properties of electrolytes such as a buffering effect. References 1. C. A. Huber, T. E. Huber, M. Sadoqi, J. A. Lubin, S. Manalis and C. B. Prater, Science 263 (1994) 800. 2. H. Masuda and K. Fukuda, Science 268 (1995) 1466. 3. H. Masuda, M. Ohya, H. Asoh, M. Nakao, M. Nohtomi and T. Tamamura, Jpn. J. Appl. Phys. 38 (1999) L1403. 4. H. Masuda, M. Ohya, K. Nishio, H. Asoh, M. Nakao, M. Nohtomi, A. Yokoo and T. Tamamura, Jpn. J. Appl. Phys. 39 (2000) L1039. 5. H. Masuda and M. Satoh, Jpn. J. Appl. Phys. 35 (1996) L126. 6. M. Steinhart, J. H. Wendorff, A. Greiner, R. B. Wehrspohn, K. Nielsch, J. Schilling, J. Choi, U. Gösele, Science 296 (2002) 1997. 7. D. Gong, C. A. Grimes, O. K. Varghese, W. Hu, R. S. Singh, Z. Chen and E. C. Dickey, J. Mater. Res. 16 (2001) 3331. 8. V. Zwilling, E. Darque-Ceretti and A. Boutry-Forveille, Electrochim. Acta 44 (1999) 421. 9. R. Beranek, H. Hildebrand and P. Schmuki, Electrochem. Solid-State Lett. 6 (2003) B12. 10. J. M. Macak, H. Tsuchiya and P. Schmuki, Angew. Chem. Int. Ed. 44 (2005) 2100. 11. J. M. Macak, K. Sirotna and P. Schmuki, Electrochim. Acta (in press). 12. H. Tsuchiya, J. M. Macak, L. Taveira, E. Balaur, A. Ghicov, K. Sirotna and P. Schmuki, Electrochem. Commun.(in press). 13. A. Ghicov, H. Tsuchiya, J. M. Macak and P. Schmuki, Electrochem. Commun. 7 (2005) 505. 14. L. V. Taveira, J. M. Macak, H. Tsuchiya, L. F. P. Dick, P. Schmuki, J. Electrochem. Soc. (in press) 15. H. Tsuchiya and P. Schmuki, Electrochem. Commun. 6 (2004) 1131. 16. H. Tsuchiya, J. M. Macak, I. Sieber and P. Schmuki, Small (in press). 17. H. Tsuchiya and P. Schmuki, Electrochem. Commun. 7 (2005) 49. 18. I. Sieber, B. Kannan and P. Schmuki, Electrochem. Solid-State Lett. 8 (2005) J10. 19. H. Tsuchiya, J. M. Macak, I. Sieber, L. Taveira, A. Ghicov, K. Sirotna and P. Schmuki, Electrochem. Commun. 7 (2005) 295. 20. I. Sieber, H. Hildebrand, A. Friedrich and P. Schmuki, Electrochem. Commun. 7 (2005) 97. 21. H. Tsuchiya, J. M. Macak, L. Taveira and P. Schmuki, Corrros. Sci. (accepted).

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Self-Organized Nanoporous Valve Metal Oxide Layers

(a)

(b)

250 nm

(c)

250 nm

(d)

250 nm

(e)

250 nm

(f)

250 nm

250 nm

Figure 1 Top-view images of nanoporous oxide layers formed on (a) Ti (20 V), (b) Zr (20 V), (c) Hf (50 V), (d) W (40 V), (e) Nb (20 V) and (f) Ta (20 V) in 1M H2SO4 containing small amounts of fluoride ions.

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1 µm

Figure 2 Cross-sectional images of nanoporous oxide layers formed on (a) Ti (20 V), (b) Zr (20 V) and (c) Hf (50 V in 1M H2SO4 containing small amounts of fluoride ions.

Figure 3 Cross-sectional images of nanotube oxide layers formed on (a) Ti (20 V), (b) Zr (20 V) and (c) Hf (50 V) in 1M (NH4)2SO4 containing small amounts of fluoride ions.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Studies on transition of titanium from active into passive state in phosphoric acid solutions Elzbieta Krasicka-Cydzik University of Zielona Gora, Zielona Gora, Poland, [email protected]

Abstract - Studies on active-passive transition of titanium in phosphoric acid solutions of 0.5-4 M revealed the unusual surface chemistry with the electrolyte concentration. The passive film was affected by the applied anodic potential, but under galvanostatic conditions at low current densities (0.1-0.6 Am-2) the slope of dE/dt always showed the minimum at the concentration ~2 M, which resulted in a coating of an oxide film by a gel-like layer. The active–passive transition was a process in which an inhibiting effect of phosphate ions on a dissolution of oxide layer was observed. Capacitance and SEM/EDS examination gave the evidence of two-layered surface film. To explain all the above results, an oxide pre-film model and the mechanism of two-layered surface film formation on titanium in phosphoric acid solutions are proposed. Keywords: titanium, transition region, phosphoric acid

1. Introduction Studies on formation of thin anodic layers on titanium and its implant alloys in phosphoric acid solutions of 0.5-4 M revealed the unusual surface chemistry of titanium as a function of electrolyte concentration [1]. Under galvanostatic conditions at low current densities (0.1-0.6 Am-2) the slope of dE/dt always showed the minimum at the concentration about 2M. The use of that electrolyte resulted in a coating of an oxide film by a gel-like layer [2], similar to the one observed in other media on aluminum [3]. It was found that the active–passive transition was a process in which an inhibiting effect of phosphate ions on a dissolution of oxide layer was observed. The present study aims to present a better understanding of the electrochemical growth behaviour of the

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galvanostatic anodic oxide on commercially pure titanium brought about by changing the electrolyte concentration. 2. Experimental Titanium samples (TIMET Ltd UK) of 0.5 cm2 surface area were abraded (600 paper), polished with alumina powder, degreased with acetone, rinsed with water and then left in air at ambient temperature. Anodizing and all electrochemical tests were carried out with an ATLAS 9831 (Poland) in three electrodes glass cell using 0.5-4M H3PO4 solutions prepared from 85% wt. orthophosphoric acid. Samples were kept in the solution for 30 min to establish the free corrosion potential (Ecor) and then the anodic galvanostatic polarization with current density 0.5 Am-2 was performed for 15 min till the samples reached the potential plateau ~3 V (SCE). Electrodes were removed from the electrolyte with the potentiostat switched on, washed with distilled water, dried and transferred through air to further impedance analysis in 0.9 %NaCl solution using a 10 mV rms perturbation from 105 Hz down to 0.18 Hz.. The potentiodynamic tests were carried out with a scan rate of 3 mV/s from –0.8 V (SCE). The anodically treated titanium surfaces were examined by Scanning Microscopy (JLM 5600), equipped with EDS analysis facilities. 3. Results 3.1. The effect of the phosphoric acid concentration on the oxide growth

Fig. 1. Galvanostatic and potentiodynamic studies, chronopotentiometric E-t curves of Ti anodised at 0.5 A/m2 in 0.5-2 M H3PO4 and active-passive transition region of polarization curve for titanium (scan 3 mV/s) in 0.5-4 M H3PO4

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Typical examples of voltage vs time transients (dE/dt) during the growth of the galvanostatic anodic oxide film on titanium for the current density of 0.5 A/m2 (Fig. 1a) show that the continuous linear growth of potential to the steady state, i.e. dynamic equilibrium between the oxide formation rate and field assisted and/or chemical dissolution rate is dependent on the concentration of H3PO4 [1]. The slope dE/dt, which may be used as a parameter for comparing the rate of oxide formation at initial stages, at potential lower than about 3 V (SCE) [4] when oxygen evolution is excluded, demonstrates (Fig. 1a) the lowest value in 2 M H3PO4.Polarization curves (Fig.1b), show that after an initial range of cathodic depassivation, samples reach the corrosion potential Ecor. An active– passive transition is observed next with passivating currents in the order of a few microamperes, which are typically observed during the passivation of titanium and its alloys [5]. Anodic curves for 0.5 M and 3 M H3PO4 solutions (Fig.1b), illustrate quasi-passive behaviour in active-passive transitions, while curves for 1 and 2 M H3PO4, having the higher corrosion potential Ecor, do not show linear dependence in this potential region. The differences in anodic Tafel slopes are accompanied by a shift of the Ecor value in the positive direction by ~0.15V with the increase of H3PO4 concentration to 2 M. 3.2. The effect of the phosphoric acid electrolyte concentration on the impedance characteristics of anodic layers

Fig. 2. Impedance spectra for titanium anodised in 2 M H3PO4 exposed to 0.9% NaCl solution a) Nyquist spectra, b) Bode diagrams and results their fitting to c) equivalent circuit

The EIS spectra of titanium were recorded in 0.9 % NaCl solution 1 h after anodizing at 0.5 A/m2 in 0.5M and 2M H3PO4 solutions. Examples of the

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impedance spectra, represented in both complex impedance diagram (Nyquist plot), and Bode amplitude and phase angle plots, displayed in Fig.2b, exhibit two time constants: the first in the high-frequency part arises from the ohmic resistance due to the electrolyte and the impedance characteristics resulting from the penetration of the electrolyte through a porous film, and the second in low-frequency part accounting for the processes taking place at the substrate/electrolyte interface. Such a behavior is typical of a metallic material covered by a porous film which is exposed to an electrolytic environment [6].The EIS data are satisfactorily fitted using the equivalent circuit given in Fig.2c, which consists of a solution resistance Rs, the capacitance Cp of the barrier layer, the charge transfer resistance associated with the penetration of the electrolyte through the pores Rp; and the polarization resistance of the barrier Rb as well as the electrical double-layer capacitance at the substrate/electrolyte interface Cb. In the case of surface layer formed in 2 M H3PO4 the specimen is covered by a passive oxide film of higher impedance (Fig.2a). The significant increase of the resistance Rb and Rp values for the 2 M H3PO4 anodized samples, over those determined for the 0.5 M H3PO4 anodized titanium, was also observed. These EIS results were complementary to those obtained by Ecor measurements and potentiodynamic polarization studies [1,2]. 3.3. Scanning microscopy characterisation of thin anodic layers The SEM/EDS examinations reveal that thin films of titanium oxide are covered by gel-like layer with crystalline phosphates nuclei inside. In surface layers formed in 0.5 M H3PO4 phosphorus species were present as randomly

Fig. 3. SEM micrographs of titanium surface anodised at 0.5 A/m2 in a ) 0.5 M, b) 2 M H3PO4

distributed deposits of titanium phosphate (Ti(HPO4)2) (Fig. 3a-1) on a thin layer of amorphous TiO2 (Fig. 3a-2), which is usually formed at the potential no higher than 7.5V [7]. Uniformly dispersed numerous deposits of phosphates are found in a surface oxide of sample anodised in 2 M H3PO4. However, the oxide and phosphates are covered with the additional layer consisting of 76.3±3.6 wt.% of phosphorus and 23.7±1.5wt.% of oxygen (Fig.3a-3 and 3b). Thus, a

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two-layer model of the surface film corresponds to the electrochemical behavior of the system consisting of oxides and phosphates. 4. Discussion Titanium as metal very sensitive to the a pre-treatment [8], due to the polishing, rinsing with water and drying, is usually covered by an air-formed oxide film and on immersion into acid solution shows potentials of the active-transition region [9,10]. However, in solutions of low pH may become active. In sulphate solution the anodic oxide film on titanium dissolves giving Ti3+ ions [11]. Typical activation behaviour and slow decrease in the open-circuit potential (– 0.3 V SCE), is observed on immersing titanium into 1N HCl [10]. Titanium behaves differently in H3PO4 solutions. Although the values of Ecor ranging from -0.1V to -0.6V (SCE) indicate in Pourbaix diagram [12] that the oxide film should dissolve to Ti3+, on immersion to H3PO4 solutions titanium shows the continuous shift of potential towards the anodic direction [2]. Such a tendency indicates that the rate of the anodic reactions is continually decreasing, as a result of the presence of an adsorbed, additional layer on the metal surface [13]. Thermodynamic data [1,8] for the potential E ³ –0.8V (SCE), indicate that the following reactions on titanium are likely: Ti2O3 + 6 H+ + 2e- = 2 Ti2+ +3 H2O Eº = - 0.478 V - 0.177pH – 0.059 log(Ti2+)

(1)

Ti2O3 + H2O = 2 TiO2 + 2 H+ + 2 e- Eº = - 0.556 V – 0.059 pH

(2)

Ti2O3 + 3 H2O = 2 TiO2·H2O + 2 H+ + 2 e- Eº = - 0.139V – 0.059 pH

(3)

2 Ti(OH)3 + H2O=2 TiO2×H2O + 2 H+ + 2e- Eº = -0.091V–0.059 pH

(4)

In acidic solutions within the cathodic region, the only oxide dissolution is reaction (1). This reaction determines the potential-current changes in activepassive region of titanium in 0.5 M and 3-4 M H3PO4, whereas due to the shift of the corrosion potential Ecor towards the anodic direction for titanium immersed in 1-2 M H3PO4 electrolyte, its direct oxidation proceeds according to reactions 3 and 4. These results indicate that, the layer of phosphates (Fig. 3) blocks the oxide dissolution. Insight into adsorption of phosphates to TiO2 surface revealed the hypothesis according to which they form covalent bonding to oxygen [14], or metal ions react with phosphate anions forming a gel of metal hydrophosphates [15]. Both proposed processes would lead to local increase of pH at the oxide surface and in consequence to the increase of concentration of dihydro-phosphate ions (E-pH diagram) [12]. Then, on phosphates covered titanium oxide electrode, the following gel like layer formation could proceed

198

2 H2PO4- + 2 H3O+ + (n-2) H2O® 2 H3PO4 · n H2O

E. Krasicka-Cydzik

(5)

This attribution agrees with Morligde’s et. al. results on aluminum [3]. Both reactions: the phosphates adsorption and gel-like layer formation are nonfaradaic, but are competitive towards the oxide dissolution (reaction 1) with regard to proton consumption. The advantageous effect of these two reactions on the anodizing, may be attributed to an inadequate supply of H+ ions to keeping up with the demand for the reaction (1) of oxide dissolution. The increasing coverage of the anodized titanium surface by phosphates ions with the electrolyte concentration provides the evidence of a direct influence of electrolyte anions in suppressing the formation of dissolved titanium ions. According to potential/pH diagram for P-Ti-H2O system [16], H3PO4×0.5H2O, the product of reaction 5, is stable thermodynamically in solutions of pH ranging to 3. 5. Conclusions Studies on active-passive transition of titanium in H3PO4 solutions revealed a minimum voltage-time growth, dE/dt at concentration 2 M H3PO4 at the relatively low constant current density ~ 0.5 Am-2. It is believed that the observed phenomena can be attributed to the layer of adsorbed phosphate ions which during the anodizing developed into gel-like layer above the anodic film on the electrode surface and played an important role in inhibiting oxide dissolution. References 1. E. Krasicka-Cydzik, Formation of thin anodic layers on titanium and its implant alloys in phosphoric acid solutions, University of Zielona Press, 2003, ISBN 83-89321-80-7 2. E. Krasicka-Cydzik, Corros. Sci., 46 (2004) 2487 3. J.R. Morlidge et. al , Electrochim. Acta, 44 (1999) 2423. 4. G.T. Burstein, R.M. Souto, Electrochim Acta 1995;40: 1881–8 5. R.M. Souto, M.M. Laz, Rui L. Reis, Biomaterials, 24 (2003) 4213. 6. I. Thompson, D. Campbell, Corros Sci 36 91994)187–98. 7. T. Shibata, Y.C. Zhu, Corros. Sci. 37 (1995) 343 8. E.J. Kelly, in: J. O’Bockris, B.E. Conway, R.E. White (Eds.), Modern Aspects of Electrochemistry, Plenum Press, New York, London, 1982, pp. 332–336 9. R.M. Torresi, O.R. Camara, C.P. De Pauli, Electrochim. Acta 32 (9) (1987) 1357. 10. E.J. Kelly, J. Electrochem. Soc. 126 (12) (1979) 2064 11. A.C. Makrides, J. Electrochem. Soc. 111 (4) (1964) 392 12. M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, NASE, 1974 13. I. Felhosi, J. Telegdi, G. Palinkas, E. Kalman, Electrochem. Acta 47 (2003) 2335–2340. 14. K.E. Healy, P. Ducheyne, Biomaterials. 13 (1992) 553. 15. A.S. Wagh, S.Y. Yeong, J. Am. Ceram. Soc. 86 (11) (2003) 1838–1855. 16. A. Roine, HSC Ver. 2. 03, Outkumpu Research Oy, Pori, Finland, 1994

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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The Influence of Temperature on the Passive Film Properties of ASTM Grade-7 Titanium J.J. Noël, L. Yan, D. Ofori, P. Jakupi, and D.W. Shoesmith Department of Chemistry, The University of Western Ontario, London, ON, N6A 5B7 Canada

Abstract - The influence of temperature on the passivity of highly corrosion-resistant alloys is often critical to their industrial performance. We have been studying the effects of temperature (up to ~150°C) on Ti alloys, such as ASTM Grade-7 (Ti-7), in neutral saline solutions, using open circuit potential measurements and electrochemical impedance spectroscopy (EIS). As the temperature increased, up to ~80°C, the oxide film resistance increased and the capacitance decreased, consistent with thickening of, and defect annealing within, the passive film. At higher temperatures, the passive film displayed decreasing resistance and increasing capacitance, and a second time constant developed at low frequencies, indicating the introduction of pores or cracks into the oxide. Thus, this thin oxide film ruptures under certain conditions, leading to a temporary loss of corrosion protection. Such oxide breakdown and repair yields potential transients (“electrochemical noise”), which contain information on the electrochemical behaviour of the material. Keywords ASTM Grade-7 Titanium, Electrochemical Impedance Spectroscopy, Electrochemical Noise, Oxide Film Breakdown, Cathodic Modification

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1. Introduction The thin surface oxide film on Ti (~2-5 nm) provides excellent corrosion resistanceand is highly impermeable to hydrogen [1]. However, the oxide can fracture at T >333K [2], while remaining chemically and thermodynamically stable, and thus temporarily lose its passivity. This fracture may occur at weak points in the oxide -for example, in the oxide covering grain boundaries in the underlying material [3], and local breakdown and repair of the passive film can yield transient signals in the corrosion potential or current. Their analysis provides information concerning the electrochemical behaviour of the alloy under these conditions. Crystallization and film rupture introduce low resistance pathways for ion transport (detectable by EIS), which could lead to hydrogen absorption, the first step to failure of industrial structures by hydrogen-induced cracking (HIC) Figure 1 [4] shows schematically this corrosion process for Ti-7. H2 O

H2

Gauge Block

H+ H2 O2Ti4+

Ti3+/Ti4+

TiO2 Ti (Passive Site)

Electrode Feed-through

Sealing glands

Pd Pd En rich ment Ti

Ti Titaniu m vessel

Habs (Temporary Active Site)

Figure 1. Possible Corrosion processes within a film breakdown site on Ti-7.

Ho memade Ag/AgCl Reference electrode

Teflon liner Titaniu m Planar electrode Solution Titaniu m counter elect rode

Figure 2. Electrochemical cell in a titanium pressure vessel.

In this paper we report on the influence of temperature on the passivity of ASTM Grade-7 titanium (Ti-0.2 wt.% Pd) in neutral pH saline solutions (0.1-3 mol×L-1 NaCl) at temperatures up to ~150°C. Our primary long-term goal is to determine the influence of oxide film breakdown/repair events on the corrosion of and hydrogen absorption by Ti-7 in aggressive environments at high temperatures.

The Influence of Temperature on the Passive Film Properties of Grade-7 Titanium 201

2. Experimental Electrodes were fabricated from ASTM Grade-7 titanium obtained from RMI Titanium. The details of electrode design and preparation have been described elsewhere [5]. The composition of the Ti-7 used is shown in Table 1. Experiments were performed in aerated, NaCl and Na2SO4 solutions (pH 6 to 7.5) of various concentrations prepared with ultra-pure water (Millipore, 18.2 MW×cm). To permit high-temperature aqueous work, experiments were conducted inside a 1 L PTFE-lined ASTM Grade-4 Ti pressure vessel (Parr Instrument Co., model 4621), as shown schematically in Figure 2 and described in detail elsewhere [5]. All potentials were measured versus an internal Ag/AgCl (0.1 M KCl) reference electrode. The counter electrode had the same composition as the working electrode. The temperature was increased gradually from 30°C to 150°C in daily increments of 10°C and the corrosion potential (Ecorr) recorded while it stabilized over a period of about 22 h. Electrochemical impedance spectroscopy (EIS) was then performed using a Solartron 1287/1255B measurement system controlled by Z-Plot software (Scribner Associates). The frequency range for EIS measurements was 105 to 10-3 Hz and the voltage amplitude 10 mV. Subsequently, the temperature was increased and the measurement sequence repeated. Table 1. Chemical composition (wt.%) of Ti-7 used as electrode material Element

Ti

Fe

C

O

N

Pd

Wt.%

Bal.

0.18

0.011

0.011

0.134

0.16

3. Results and Discussion Figure 3 shows an example of the recorded Ecorr as a function of time in 0.1 M NaCl solution at various temperatures. The values of Ecorr after electrode exposure to 0.1 M NaCl or 0.1 M Na2SO4 at each temperature for ~22 h are plotted in Figure 4. In chloride solution, Ecorr increased up to 80oC and then decreased at higher temperatures. In sulphate, Ecorr did not increase to the same extent for T < 80oC and did not show any decrease at high temperatures. In both solutions, potential noise signals due to film fracture events began at 60oC and became quite significant around 80oC. Figure 5 shows an expanded section of the Ecorr vs time plot recorded in 0.1 M NaCl at 110oC. A rapid initial decrease in Ecorr indicates a local film fracture event, and the slower exponential recovery indicates a film re-growth process. The time required for the potential to recover to its original value, the repassivation time (tR), increased with temperature, indicating that repassivation was more difficult at higher

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temperatures. The event frequency increased steadily with increasing temperature up to T = 130oC, consistent with previous results [5], but decreased at 140oC and 150oC, suggesting a limited availability of fracture sites. 0.20 0.1M Na2SO4

0.15

80°C

0.1M NaCl

120°C

Ecorr (Volts)

0.10

100°C 90°C 110°C

0.05 0.00

-0.05

130°C 140°C

-0.10 -0.15

150°C

-0.20

20

40

60

80

100

120

140

160

o

Temperature ( C)

Figure 3. Corrosion Potential of ASTM Grade-7 titanium in 0.1 M NaCl as a function of time at various temperatures.

Figure 4. Corrosion potential (Ecorr) as a function of temperature for Ti-7 in 0.1 M NaCl and 0.1 M Na2SO4.

0 .1 2 4 0

A

CFILM

0 .1 2 3 5

RFILM

E c o r r ( V o lt s )

R

0 .1 2 3 0

CFILM

B

0 .1 2 2 5

R

CDL RPORE 0 .1 2 2 0 1 9 .5

RPOLARIZATION 1 9 .6

1 9 .7

1 9 .8

1 9 .9

2 0 .0

T im e (h o u rs )

Figure 5. Expanded section of the Ecorr - time plot recorded on Ti-7 in 0.1 M NaCl solution at 110°C.

Ti

TiO2

Figure 6. Equivalent circuit models fitted to the EIS data: a) single- and b) two- time constant circuits.

EIS was used to measure film capacitances and resistances. The equivalent circuits used to model the data are sketched in Figure 6, and the parameter values determined are plotted in Figure 7. The single time constant (t) circuit,

The Influence of Temperature on the Passive Film Properties of Grade-7 Titanium 203

representing the EIS response of a continuous oxide film was used at temperatures up to ~60°C; above this temperature, a 2t circuit, representing a fractured surface oxide, was necessary. In chloride solution, the film capacitance (CFilm, Figure 7A) decreased with temperature (T < 80oC) initially, but then increased with temperature up to 150oC. This decrease in CFilm in chloride at T < 80oC could be attributed to either film thickening or the (potential-driven or thermal) annealing of defects, or both. Either possibility would be consistent with the observed increase in Ecorr up to 80oC (Figure 4). In sulphate solutions, CFilm was effectively independent of temperature for T ≤ 60oC, implying that the film properties didn’t change much with temperature. In both solutions CFilm increased for T ≥ 80oC. This could have been caused by either film thinning or an increase in the number and size of defects in the film. Film thinning was ruled out since the solubility of Ti oxide in neutral aqueous solution is extremely low. Instead, the increase in CFilm was most likely due to water incorporation at fracture-repair sites, leading to a substantial increase in the effective dielectric constant of the film. Figures 7B and 7C indicate that at low temperatures the film resistance behaved similarly in both chloride and sulphate solutions. With increasing temperature, the behaviours in the two solutions became increasingly different. Above 80oC the film resistance in chloride was significantly greater than that in sulphate. The film resistance reached a relatively constant value in chloride solution, while it decreased steadily with temperature in sulphate solution. The pore resistance in sulphate solution decreased effectively to zero, whereas that in the chloride solution remained >104 Ω×cm2 at high temperatures. We can surmise that the corrosion and repassivation processes within fracture sites proceed as follows. Upon film fracture, exposed intermetallics (i.e., TixFe) and Fe-containing phase corrode. The Ti4+ ions produced hydrolyze to produce protons, leading to acidification within the site. Proton diffusion to the bulk solution leads to a pH gradient within the fracture site. The Pd in solid solution within the -phase Ti matrix acts as a catalyst [6] for proton reduction coupled to the TixFe and phase anodic dissolution reactions. This leads to an increase in potential within the site, which forces repassivation to occur. At high temperatures, anodic dissolution of TixFe intermetallics and phase leads to a substantial dissolved Ti4+ concentration in the fracture site. During the period of fracture and repassivation, some Ti4+ diffuses out of the low pH region at the alloy surface and is precipitated as TiO2·2H2O under more neutral conditions. The precipitation of TiO2·2H2O in chloride solution seems to occur within the fractures/flaws leading to a high resistance in the blocked pores; however, sulphate appears to complex Ti4+ to form TiOHSO4+ or related species, thereby increasing its solubility. This complex appears to escape the fracture sites before depositing as TiO2·2H2O outside the pores, leading to a negligible pore resistance in sulphate solution.

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120 1

110

?F Res ista 0.1 nce (M Film Resistance (M W .cm )

90

2

Film Capacitance (F/cm2 )

100

A

80 70 60

NaCl Na SO

50

2

B NaCl Na SO 2

4

0.01

4

40 30 20

40

60

80

100

120

140

160

20

40

60

80

100

120

140

160

o

Temperature ( C)

Temperature (

o

C)

100000

Pore Resistance (W.cm2)

10000

1000

C 100

Na SO

10

2

4

NaCl 1

0.1 60

80

100

120

140

160

o

Temperature ( C)

Figure 7. Comparison of A) film capacitance, B) film resistance, and C) pore resistance values for Ti – 7 in 0.1 M NaCl and 0.1 M Na2SO4 solutions determined by the fitting to experimental EIS data.

References 1. D.W. Shoesmith, F. Hua, K. Mon, G. De, P. Pasupathi, and G. Gordon, CORROSION/2005, paper no. 05582: NACE International Houston, TX, (2005). 2. T. Shibata, Y. C. Zhu, Corr. Sci., 36(10) (1994) p. 1735-1749. 3. J.J. Noël, PhD Thesis, The University of Manitoba, Winnipeg, MB, (2000). 4. D.W. Shoesmith and B. M. Ikeda, Atomic Energy of Canada Limited report, AECL-11709, COG-96-557-I (1997). 5. X. He, PhD Thesis, The University of Western Ontario, London, ON, (2003). 6. T. Fukuzu, K. Shimogori, H. Satoh, In Proceedings of the Fourth International Conference on Titanium, Kyoto, Japan, (1980) p. 2695-2703.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Chromatic properties of anodised titanium obtained with two techniques MariaPia Pedeferri,a Barbara Del Curto,a Pietro Pedeferria a

Dipartimento di Chimica, Materiali e Ingegneria Chimica "Giulio Natta" Politecnico di Milano, Via Mancinelli, 7, Milano 20131, Italy [email protected]

Abstract - The aim of the work is to compare the surface properties of titanium oxidised with two different techniques. First technique is the traditional one, using phosphoric acid solution, and the other technique is the one described in [1]. Final potential of the anodisation has been varied from few Volt up to 140 V. Microstructure of the different films show that only the samples anodised with the new technique present an homogeneous microstructure, i.e. all the crystalline grains are coloured, while the traditional technique colours only a part of them. Only having all grains fully coloured it is possible to have at the macroscopic scale saturated and bright colours. Keywords: titanium, anodic oxidation, colour, microstructure, gloss

1. Introduction Until just a few years ago, titanium was only used for its strength, lightness and corrosion-resistance. Nowadays, new potential developments have been found for this metal, because of the properties of the oxides which can be generated on its surface. Titanium is usually covered by a protective film a few nanometers thick. By means of electrochemical techniques of anodisation its thickness can be increased, from a few up to 300 nanometers, giving at the titanium surface a wide range of interference colours. The aim of this work is to study and compare the chromatic properties of titanium oxidised by two different techniques.

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2. Experimental Two techniques of anodic oxidation of titanium grade 2 were considered. The first technique was the traditional one, i.e. using phosphoric acid solution (510% wt) at room temperature and with current density of 10 mA/cm2. The other technique, described in [1], was a two-step anodisation, the first step being anodisation carried out in a hydrochloric acid solution (10-20%) at constant potential (4V), and the second step traditional anodisation in phosphoric acid (5-10% wt) at room temperature and with current density of 10 mA/cm2. Final potential of the anodisation varied from a few volts up to 140 V. Before anodisation, samples were etched in a solution of nitric and hydrofluoric acid. In order to study the effects of surface finishing on colour perception, CP grade 2 titanium samples with various surface finishings were considered: smooth, sandblasted, and chemically polished surfaces. Anodised samples were analysed by means of optical microscopy (Me F3A Reichert Jung), optical spectrophotometry (portable spectrophotometer CM2600d Minolta), laser profilometry (3D laser profilometer Microfocus UBM) and gloss tests (mutligloss 268 Minolta). Gloss and profilometry measurements were carried out according to ISO 7668 and DIN 4768, respectively. 3. Results and discussion Figure 1 shows micrographs of a sample anodized on the left-hand side with the traditional technique and on the right-hand side with the new technique. Final potential in both cases is 80 V. It is worth noting that on both the microscopic and macroscopic scales there is a substantial difference between the two sides of the sample. In particular, the one anodised with the new technique presents all the crystalline grains coloured, while the traditional technique colours only a part of them. Bright, saturated colours on a macroscopic scale can only be achieved in the first case.

Fig.1 Titanium sample anodised with the two techniques: macroscopic and microscopic scale; the right of the sample is anodised with the two-step anodisation.

Chromatic properties of anodised titanium obtained with two techniques

207

The difference between the two techniques lies in the first anodisation of the new technique, which is carried out at constant potential in hydrochloric acid solution. The potential value of this anodisation strongly influences homogeneity of colours grains. In Fig.2 micrographs of titanium samples subjected to double oxidation are illustrated. The sample are first anodised at different values of 1 V, 2 V, 4 V 6 V and 8 V in hydrochloric acid (15%) and then anodised at 80 V in phosphoric acid (5%). In order for all grains to be fully coloured, the first anodisation has to be carried out at the optimal potential value (4 V). As it is shown in Fig.2 the percentage of non-coloured grains increases as one moves away from the optimal potential.

100 mm

100 mm

100 mm

100 mm

100 mm

Fig.2 Effect of the potential value of the first anodisation. The sample were first anodised at different values of 1 V, 2 V, 4 V 6 V and 8 V in hydrochloric acid (15%) and then anodised at 80 V in phosphoric acid (5%).

Five titanium samples have been anodised with the new technique at the optimal potential value of the first anodisation (4 V), and varying the potential of the second oxidation at: 25 V, 50 V, 80 V, 110 V and 140 V. As the applied potential increases, the colour of the five samples changes in the following sequence: blue, yellow, purple, green and pink. Spectral curves of samples with a smooth surface finishing anodised at 25, 50, 80, 110 and 140 V are shown in Figure 3. Also reported in the same figure is the spectral curve of non-anodised titanium (TQ), which presents a uniform behaviour and a silver appearance.

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Fig.3 Spectral curves of samples anodised at 25 V, 50 V, 80 V, 110 V and 140 V.

Maximum intensity, i.e. constructive interference, will take place [3] if: 2nd = ml m = 0, 1, 2, 3...

Eq.1

while destructive interference conditions correspond to the minimum intensity: 2nd = (m+½l) m = 0, 1, 2, 3...

Eq.2

where n, d and l are refractive index, film thickness and wavelength respectively. From spectral curves, assuming an average constant refractive index, titanium oxide thickness can be calculated (Figure 4) using equations 1 and 2. The titanium oxide refractive index used for calculations is the one reported in literature [4].

Fig.4 Oxide thickness as a function of the final potential.

Chromatic properties of anodised titanium obtained with two techniques

209

Increasing the applied potential of the second anodisation, oxide thickness varies linearly from 50 nm at 25 V, up to 340 nm if the final potential is 140 V; the increase in thickness is about 2.2 nm/V. In addition to the thickness of titanium oxide, other parameters also contribute to defining chromatic properties and different shades; in particular, surface finishing. Anodised samples have been analysed by means of profilometry and gloss tests in order to determine their roughness and brightness. Gloss is the term used to describe the perception by an observer of the mirrorlike appearance of a surface, and is defined as the degree to which a surface exhibits specular reflectance. This means that it is a measure of the imageforming ability of a surface. The perception of gloss is influenced by numerous factors, among them being the surface profile (roughness) and the viewing angle. Gloss measurements were taken at angles of 20°, 60° and 85°. Samples with different surface finishings (smooth, sandblasted and chemically polished) were anodised at 60 V. The three samples appear as different shades of yellow. In Figure 5, their spectral curves are reported.

Fig.5 Spectral curves of titanium samples anodised at 60 V with different surface finishing.

It is worth noting that these spectral curves present a similar behaviour and in particular the minimum and maximum of the different curves correspond to the same wavelength, i.e. the same colour, but visual perception of them depends on surface finishing. Actually, the reflectivity value for a given wavelength decreases in the following order: smooth, sandblasted, and chemically polished surface. Roughness values (Ra and related standard deviation) and gloss values (gloss at 20°, 60° e 85°) for the three surface finishing are reported in Table 1.

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MP. Pedeferri et al. Table 1 Roughness and gloss values Sample

Ra (nm)

Std. Dev. Gloss 20°

Gloss 60° Gloss 85°

Smooth

1.72

0.05

1.5

8.7

8.6

Sandblasted

2.14

0.11

1.1

5.4

5.0

Chem. polished

2.21

0.59

6.2

29.1

10.0

Smooth finishing is the one that presents the lowest roughness. Sandblasted and chemically-polished finishings present a comparable average value in terms of Ra, but in particular, the standard deviation of a chemically-polished sample is very high due to its greater heterogeneity. In effect, roughness values in the case of chemically-polished finishing samples are not so significant since they display a macroroughness visible to the naked eye which is much higher with respect to the others. Different surface morphologies are responsible for different gloss values. The lowest gloss values are obtained, for the three surface finishings, in correspondence to an angle of incidence equal to 20°. In correspondence to 60° and 85° angles, the gloss values obtained are comparable for smooth and sandblasted finishing while they are considerably dissimilar for chemicallypolished finishing, due to its surface heterogeneity. 4. Conclusion Only using the two-step anodisation technique described in [1], it is possible to obtain at the microscopic scale all grains fully coloured and at the macroscopic scale saturated and bright colours. Increasing the applied potential, oxide thickness varies linearly up to 340 nm, with an increase of about 2.2 nm/V. A specific potential corresponds to a precise thickness and colour, and colour perception is strongly affected by surface finishing. Bibliography [1] EP 1 199 385 A2 (16.10.2001) Method of colouring titanium and its alloy through anodic oxidation [2] Pedeferri, P., Titaniocromia (e altre cose), Interlinea, Novara 1999. [3] White, M. A., Properties of materials, Oxford University press, 1999. [4] Tilley, R., Colour and optical properties of materials, England Wiley, 2000. [5] Frova, A., Luce colore visione Rizzoli, 2000. [6] Palik, E.D., Handbook of Optical Constants of Solids, Academic Press, Boston 1991.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Amorphous-to-Crystalline Transition of Anodic Niobia H. Habazakia,*, T. Ogasawaraa, H. Konnoa, K. Shimizub, K. Asamic, S. Nagatac, P. Skeldond and G.E. Thompsond a

Graduate School of Engineering, Hokkaido University, Sapporo 060-8628, Japan University Chemical Laboratory, Keio University, Yokohama 223-8521, Japan c Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan d Corrosion and Protection Centre, School of Materials, The University of Manchester, P.O.Box 88, manchester M60 1QD, UK * Corresponding author: [email protected] b

Abstract - An amorphous-to crystalline transition of anodic niobia has been examined using magnetron-sputtered niobium and its alloy substrates, with particular attentions paid to the preexisting surface oxide layer as a nucleation site of crystalline oxide and incorporation of foreign species hindering its nucleation. Keywords: anodic oxide, oxide structure, field crystallization, niobium

1. Introduction Thin anodic films formed on valve metals, particularly on aluminium, niobium and tantalum, are of great importance as dielectrics for electrolytic capacitors. The anodic oxides formed on niobium and tantalum are generally amorphous, but an amorphous-to-crystalline transition occurs during anodizing at increased electrolyte temperatures [1-7]. This phenomenon is known as field crystallization [7], but the mechanism is not yet well understood. The transition and growth of crystalline oxides are detrimental for the capacitor applications, due to increased leakage current. It has been reported that inclusions in metal substrates act as initiation sites of the field crystallization [6]. Recently, Nagahara et al. have reported that the crystalline oxides are preferentially nucleated at the convex surface of the chemically polished niobium sheet [3]. They have proposed that cracking of the

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anodic film was first initiated at the convex regions, associated with strong tensile stress, inducing crystallization at such regions. In the present study, for the fundamental understanding of the field crystallization of anodic niobium oxide, sputter-deposited niobium films without inclusions and precipitates and with flat surface have been used. Factors influencing the field crystallization have been examined using the niobium substrate, which is free from the known preferential sites for field crystallization. Based on the results obtained, the mechanism of the field crystallization has been discussed. 2. Experimental Niobium and Nb-N films were prepared by magnetron sputtering on to glass, silicon and aluminium substrates. Aluminium substrates were electropolished and subsequently anodized in 0.1 mol dm-3 ammonium pentaborate electrolyte to 200 V to provide flat surfaces. The niobium and niobium alloy films prepared were anodized at a constant current density of generally 50 A m-2 to selected voltages with and without current decay in stirred 0.1 mol dm-3 ammonium pentaborate and 0.1 mol dm-3 phosphoric acid electrolytes at 60oC. Platinum sheet was used as a counter electrode. Surfaces and cross-sections of the anodized specimens were observed by a JEOL JSM-6500F field emission gun scanning electron microscope. Ultramicrotomed sections, about 10 nm thick, of the anodic films were also observed by a JEL JEM2010 transmission electron microscope operating at 200 kV. For examining depth distribution of the electrolyte-derived species in anodic films, glow discharge optical emission spectroscopy (GDOES) analysis was carried out using a Jobin-Yvon 5000 RF instrument in an argon atmosphere of 898 Pa by applying RF of 13.56 MHz and power of 40 W. Light emissions of characteristic wavelengths were monitored throughout the analysis with a sampling time of 0.01 s to obtain depth profiles. The wavelengths of the spectral lines used were 313.079, 149.262, 249.678, 177.499 and 130.217 nm for niobium, nitrogen, boron, phosphorus and oxygen respectively. The alloy and anodic film compositions were determined using Rutherford backscattering spectroscopy (RBS) using 2.0 MeV He2+ ions. The scattered particles were detected at 170o to the incident beam direction, which was normal to the specimen surface. The data were analysed using RUMP program. 3. Results and Discussion 3.1. Influence of electrolyte Fig.1 shows the current decay curves of the sputter-deposited niobium

Amorphous-to-crystalline transition of anodic niobia 10

213

2 o

Current Density / A m

-2

Nb anodized at 100V, 60 C 10

10

1

0

-3

0.1 mol dm (NH4 ) 2 B10 O 16 10

-1

-3

0.1 mol dm H3 PO 4 10

-2

0

2000

4000

6000

Anodizing Time / s

Fig. 1 Current decay curves of the sputter-deposited niobium during anodizing at 100 V in 0.1 mol dm-3 ammonium pentaborate and phosphoric acid electrolytes at 60 oC.

during anodizing at 100 V in 0.1 mol dm-3 ammonium pentaborate and phosphoric acid electrolytes at 60oC. Initially, the current decreases rapidly with time in both the electrolytes. In the borate electrolyte, then, the current increases after approximately 3.0 ks, due to formation of crystalline oxide (Fig. 2(a)). Petal-like defects can be seen in the micrograph and the development of the crystalline oxides inside of the defects was confirmed by TEM observations. In contrast, no crystalline oxide is formed in the phosphoric acid electrolyte (Fig. 2(b)), and hence the current continues to decrease at least for 7.2 ks (Fig. 1). The different crystallization behaviours in these two electrolytes should be associated with the depth distribution of the electrolyte-derived species incorporated into the anodic films. The GDOES depth profiles of the anodic films formed to 100 V without current decay in both the electrolyte at 60oC (Fig.

Fig. 2 Scanning electron micrographs of surfaces of the Nb specimens anodized at 100 V in (a) 0.1 mol dm-3 ammonium pentaborate electrolyte and (b) 0.1 mol dm-3 phosphoric acid electrolyte at 60oC for 7.2 ks.

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214

(a)

(b) -3

-3

0.1 mol dm H3 PO4

0.1 mol dm (NH4 )2B 10O16

o

o

60 C intensity / arb. unit

Intensity / arb.unit

60 C

Oxide Film B Nb

Oxide Film P P

Nb

B 0

1

2

3

Sputtering Time / s

4

0

1

2

3

4

Sputtering Time / s

Fig. 3 GDOES depth profiles of the anodic films formed on the niobium films to 100 V in (a) 0.1 mol dm-3 ammonium pentaborate electrolyte and (b) 0.1 mol dm-3 phosphoric acid electrolyte at 60oC.

3) reveal that the phosphorus species distribute approximately in the outer half of the film thickness. RBS analysis revealed that relatively high concentration of phosphorus, i.e., [P]/([P]+[Nb]) = 8 at%, was incorporated in that region. In contrast to the phosphorus species, the incorporation depth of the boron species is only about one tenth of the film thickness. The phosphorus species incorporated up to the deep film region, therefore, may suppress effectively the nucleation of crystalline oxide in the amorphous matrix. This also suggests that the nucleation should occur in the phosphorus-containing region, although Vermilyea reported the nucleation at the metal/film interface [6]. 3.2. Influence of thermal oxide The thermal oxide formed on the niobium films by heat treatment in air accelerated significantly the field crystallization. The current increase, associated with the crystalline oxide formation, during anodizing at 100 V, occurs at shorter time for the specimens heat-treated at 250oC in comparison with the as-deposited specimen. Since the initial current response and the voltage-time response before reaching 100 V are similar for both the specimens, the crystalline oxide should be nucleated during the current decay even for the heat-treated specimen. SEM observations revealed the increased number of nucleation of crystalline oxide with the heat treatment. When the heat treatment was carried out in vacuum, no accelerated crystallization occurs, indicating that the thermal oxide formed on the substrate acts as nucleation sites. The crystallization occurred for the specimen heat-treated in air at 250oC even during anodizing at a low constant current density of 5 A m-2 to more than 100 V. Relatively small crystalline oxides were formed at high

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215

Fig. 4 Scanning electron micrographs of cross-sections of the niobium specimen anodized to 140 V at 5 A m-2 in 0.1 mol dm-3 ammonium pentaborate electrolyte at 60oC.

density by anodizing to 140 V. The cross-sectional observation (Fig. 4) reveals that the crystalline oxide is present between the amorphous top layer and the metal substrate. The niobium substrate has a columnar structure, typical of the sputter-deposited materials. Due to faster growth of the crystalline oxide the residual thickness of the niobium film becomes small below the crystalline oxide region. 3.3. Influencing of substrate composition When the Nb-10 at% N alloy was anodized at 100 V in 0.1 mol dm-3 ammonium pentaborate electrolyte at 60oC, under which condition the niobium film forms crystalline oxide, no field crystallization was observed. The anodic films formed on this alloy contains N2O gas bubbles in the inner 70% of the film thickness [8]. Due to different semi-conductive properties of the anodic films formed on the Nb and Nb-N films, nitrogen species dissolved in Nb2O5 should be present in addition to the gaseous species [9]. The suppressed crystallization of the anodic film on the Nb-10 at% N film is probably due to the incorporated nitrogen species. Similarly, the authors have found the effective suppression of the field crystallization by incorporation of silicon and tungsten species from the substrate. 3.4. Mechanism of field crystallization From the significant influence of the thermal oxide as well as the effective suppression of the field crystallization by foreign species incorporated both from the electrolyte and from the substrate, the crystalline oxide should be

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nucleated at the film region where the original air-formed or thermal oxide is present. The region should be at the approximately outer 30% of the film thickness, since the transport number of cations during growth of amorphous anodic niobia is approximately 0.3. Once the crystalline oxide is nucleated, its growth is quite rapid. The rapid growth is probably associated with the lower field strength in the crystalline niobia compared with in the respective amorphous oxide. 4. Conclusions Amorphous-to-crystalline transition of anodic niobia formed on the sputter-deposited niobium films, which are free from preferable sites, such as inclusions and surface roughness, of the transition, initiated at approximately outer 30% of the film thickness where air-formed oxide is located. The growth of the crystalline oxide to the metal/film interface is rapid. The phosphorus species incorporated from the phosphoric acid electrolyte suppress effectively the amorphous-to-crystalline transition. Similarly, the foreign species incorporated from the substrate, such as nitrogen, silicon and tungsten species suppress the crystalline oxide formation. Acknowledgments The present work was supported in part by the Grant-in-Aid for Scientific Research, No. 16360353 from Japan Society for the Promotion of Science. References [1] [2] [3] [4] [5] [6] [7] [8] [9]

D.M. Lakhiani and L.L. Shreir, Nature, 188 (1960) 49. K. Nagahara, M. Sakairi, H. Takahashi, K. Matsumoto, K. Takayama and Y. Oda, Electrochem., 72 (2004) 624. K. Nagahara, M. Sakairi, H. Takahashi, S. Nagata, K. Matsumoto, K. Takayama and Y. Oda, J. Surf. Finish. Soc. Jpn., 55 (2004) 943. H. Asoh, H. Odate and S. Ono, J. Surf. Finish. Soc. Jpn., 55 (2004) 952. N.F. Jackson, J. Appl. Electrochem., 3 (1973) 91. D.A. Vermilyea, J. Electrochem. Soc., 104 (1957) 542. D.A. Vermilyea, J. Electrochem. Soc., 102 (1955) 207. H. Habazaki, T. Matsuo, H. Konno, K. Shimizu, S. Nagata, K. Takayama, Y. Oda, P. Skeldon and G.E. Thompson, Thin Solid Films, 429 (2003) 159. H. Habazaki, T. Matsuo, H. Konno, K. Shimizu, K. Matsumoto, K. Takayama, Y. Oda, P. Skeldon and G.E. Thompson, Surf. Interface Anal., 35 (2003) 618.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Ellipsometric Analysis of Growth Process of Oxidation Films on Magnesium and its Alloys in Neutral Aqueous Solutions Nobuyoshi Hara,a Daisuke Kagaya,b Noboru Akaoa a

Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan, [email protected] b School of Engineering, Tohoku University, present:Topy Industries, Ltd.

Abstract - The growth process of oxidation films on pure Mg, and AZ31B and AZ91D alloys in pure water, Na2SO4 and NaCl solutions has been analyzed by in-situ ellipsometry. The films with a refractive index of 1.42 – 1.47 grew rapidly in two consecytive stages; stage I where the film thickness increases approximately linearly with time, and stage II where the growth rate decreases with time. In every solution, the film growth rate of 6N-Mg was higher than that of AZ91D. The oxidation films were mainly composed of Mg(OH)2. The films formed on 6N-Mg and AZ91D have protective ability, while they suffer local breakdown at high potentials. Keywords : Magnesium, Mg alloys, Oxidation films, Ellipsometry, Film thickness

1. Introduction Magnesium alloys are being increasingly used in aerospace, automotive and electronics industries because of their excellent physical and mechanical properties such as low density and high strength-to-weight ratio. However, Mg alloys have poor corrosion resistance that hinders the application without the protection by surface coatings [1]. Magnesium is easily oxidized to form a thick oxide/hydroxide film when it comes in contact with humid air or water. This high reactivity of Mg has a detrimental effect on coating qualities including adhesion and uniformity. It is therefore important to get information about the growth of oxidation films on Mg alloys. The information on the surface films is necessary also to understand the corrosion mechanism of the alloys. There have been several studies of the oxidation behavior of pure Mg and Mg alloys in

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vacuum [2-5], humid air [6-8] and water [7-9], but few attempts to elucidate the film growth process in electrolyte solutions. Some previous studies revel that a partially protective film forms on Mg surface in NaCl and Na2SO4 solutions [10-12]. However, the growth process of such a film remains unknown. The purpose of the present study is to analyze the growth process of oxidation films on pure Mg and Mg alloys in neutral solutions by in-situ ellipsometry. 2. Experimental Pure Mg with two different purities, 99.96% (3N-Mg) and 99.9999% (6NMg), and two Mg alloys, AZ31B (3.42%Al, 0.88%Zn) and AZ91D (9.36%Al, 0.75%Zn), were used as specimens. The size of 3N-Mg, AZ31B and AZ91D specimens was 25x15x2 mm, and that of 6N-Mg f15x2 mm. The specimens were polished with SiC paper and diamond paste to a 1mm finish and then decreased ultrasonically in acetone. The test solutions used were pure water, 0.1M Na2SO4 (pH6.3) and 0.1M NaCl (pH6.1) solutions. The temperature of the solutions was kept constant at 20°C. All the solutions were used without deaeration. A rotating-analyzer automatic ellipsometer was used to monitor the film formation in situ. The detail of the ellipsometer is described in a previous paper [13]. Monochromatic light of wavelength 546.1 nm was used as incident light and the angle of incidence was set at 60.00 degrees. Two ellipsometric parameters, the relative phase retardation, D, and the relative amplitude reduction, tanY, were measured at an interval of 1s. The thickness and optical constants of surface films were determined using a theoretical D vs. Y curve which fits the experimental (D, Y ) data with minimal error. The theoretical D vs. Y curves were calculated using Drude’s exact optical equations for a threemedium (ambient-film-substrate) model. The optical constant of film-free surface of a specimen was determined preliminarily by tribo-ellipsometry [14]. The chemical composition of the specimen surface before and after immersion tests was analyzed by X-ray photoelectron spectroscopy, XPS, and Auger electron spectroscopy, AES. 3. Results and Discussion 3.1. Growth process of oxidation films in pure water Figure 1 shows the experimental D vs. Y plot for AZ91D alloy in pure water and a theoretical D-Y curve calculated for the growth of a film having an optical constant, N2=1.473-0.019i. During the initial stage of immersion (0-600s), the experimental plot does not fit a single theoretical curve. This suggests that both the thickness and the optical constant of surface films change with time. After the immersion time exceeds 600s, the experimental plot fits well the theoretical curve for N2=1.473-0.019i, suggesting that a homogeneous film with a fixed

Ellipsometric Analysis of Growth Process of Oxidation Films

219

optical constant grows on the surface of AZ91D alloy. The thickness of the film at each immersion time can be determined from the position of the experimental point on the theoretical curve. The same measurement and analysis were performed on the other three specimens.

Y

1ks B (600s)

42°

100°

20ks

(N2 =1.455–0.013i)

(N2 =1.473–0.019i)

Solution: H2O 0 0

160°

140°

D

AZ31B

50

Film-free surface of AZ91D (N3 = 0.610 – 3.526i)

120°

100 AZ91D

A (55s) 0s

6N-Mg

(N2 =1.458–0.021i)

5ks C (10.8ks) 3ks

50nm

40°

150

150nm

10ks 100nm

44°

3N-Mg (N2 =1.4237–0.019i)

d / nm

46°

200

Calculated (N2 = 1.473 – 0.019i) Experimental

Specimen: AZ91D 48° Solution: H2O

5

10 t / ks

15

20

Fig.1 Experimental D-Y plot during immersion Fig.2 Thickness, d, of surface films on pure of AZ91D in pure water and theoretical D-Y Mg and Mg alloys as a function of immersion curve for growth of film with N2=1.473-0.019i. time, t, in pure water.

Figure 2 exhibits the thickness of surface films as a function of immersion time for pure Mg and Mg alloy specimens in pure water. Although the film thickness could not be analyzed in the initial period of immersion, the film (a) Mg 2s AZ91D/H2O/55s Synthesized Measured

Background Al(OH)3

90

85

Binding energy / eV

Al2O3

Background M-OH

Background C-H

MgCO3

Mg

MgCl2

95

Measured

Synthesized

Synthesized

Measured

Intensity (a.u.)

Background MgO or MgCO3 Mg(OH)2

Intensity (a.u.)

Intensity (a.u.)

Synthesized

Intensity (a.u.)

Measured

(e) C 1s AZ91D/H2O/55s

(d) O1s AZ91D/H2O/55s

(b) Al 2s AZ91D/H2O/55s

Al

125

120

115

Binding energy / eV

H2O/CO2

535

M-O

530

Binding energy / eV

MgCO3

C=O C-O

295

290

C-C

285

280

Binding energy / eV

Fig.3 Measured and deconvoluted Mg 2s, Al 2s, O1s, and C1s XPS spectra on surface of AZ91D alloy after immersion in pure water for 55s.

growth was very fast in this period. The thickness of the films on each specimen exceeds 60 nm within 1000s, after which the film growth rate becomes slow. There is a small difference in the film growth rate among 3N-Mg, AZ31B and AZ91D, while the growth rate for 6N-Mg is higher that that for the other three specimens.

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Figure 3 exhibits XPS spectra of AZ91D alloy immersed in pure water for 55s, which corresponds to the position A of the experimental D vs. Y plot in Fig.1. The Mg 2s spectrum consists of contributions from four Mg species, metallic Mg, MgO/Mg(OH)2, MgCO3, and MgCl2. The Al 2s spectrum is comprised of metallic Al, Al2O3 and Al(OH)3 components. The O1s spectrum is deconvoluted into M-O (M: metal), M-OH, MgCO3 and H2O/CO2 components. These results suggest that the surface films formed on AZ91D in pure water are composed of Mg(OH)2, MgCO3, Al2O3, and Al(OH)3. The same measurement and analysis were done on the surface of AZ91D immersed for 600s and 10.8ks, which correspond to the positions B and C, respectively, in Fig.1. Peak area ratios of each Mg component obtained by the deconvolution of Mg 2s spectra are listed in Table 1. The air-formed film consists mainly of MgCO3 component. The fraction of MgCO3 decreases and that of Mg(OH)2 increases with increasing immersion time. It is therefore presumed that at the initial stage of immersion (0-600s) in water, the air-formed film composed mainly of MgCO3 changes into the film composed mainly of Mg(OH)2, followed by an increase in film thickness. This film growth process may be reflected in a Table 1. Peak area ratios of each component in Mg 2s XPS spectra of AZ91D alloy immersed in pure water and 0.1M Na2SO4 for different time.

Component Mg metal MgO/Mg(OH)2 MgCO3 MgSO4 MgCl2

0s 28 7 61 4

Pure water 55s 600s 23 16 32 60 42 19 3 5

10.8ks 13 71 12 4

0s 28 7 61 4

0.1M Na2SO4 22s 105s 4 3 85 73 4 19 4 4 3 1

complex change in the experimental D vs. Y plot observed in Fig.1. 3.2. Growth process of oxidation films in 0.1M Na2SO4 and 0.1M NaCl Figure 4 illustrates the experimental D vs. Y plot for AZ91D alloy immersed in 0.1M Na2SO4 and a theoretical D - Y curve calculated for the growth of a film with N2=1.418-0.053i. From the beginning of immersion, the experimental plot fits well the theoretical curve, suggesting that the air-formed film on AZ91D alloy dissolves immediately after immersion in Na2SO4 and an alternative oxidation film grows rapidly. This was confirmed by the results of XPS analysis shown in Table 1. From the change in the peak area ratios of MgCO3 and Mg(OH)2 with immersion time, it is evident that the transformation from the air-formed film consisting mainly of MgCO3 to the film of Mg(OH)2 is almost completed after immersion for 22s.

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Figure 5 exhibits the changes in the thickness of surface films on 6N-Mg, AZ31B and AZ91D with immersion time in 0.1M Na2SO4. For each specimen, the d vs. t curve can be divided into two stages; stage I where the film thickness increases linearly with time, and stage II where the growth rate decreases with the lapse of time. Stage I is limited to a short period of time, about 60s. At the end of this stage, the film thickness exceeds 70nm. The thickness of the films formed on 6N-Mg in stage II is larger than that on AZ31B and AZ91D alloys. Dissolution and precipitation may be the predominant mechanism for the film formation in stage I.

2

200

150nm

AZ91D-ThM 48 ° Specimen: Solution: 0.1M Na SO

AZ31B (N2 =1.437–0.0101i)

4

130nm

46 °

150

3ks

100nm 42 ° B (105s)

40 ° 38 ° 110°

d / nm

Y

44 °

50nm A (22s)

120°

100

2ks 1ks 0s

Calculated (N2 = 1.418 – 0.053i) Experimental Film-free surface of AZ91D (N3 = 0.904 – 3.495i)

D

6N-Mg (N2 =1.410–0.013i)

4ks

130°

140°

AZ91D (N2 =1.418–0.053i) 50 Solution: 0.1M Na2SO4 0 0

1

2 t / ks

3

4

Fig.4 Experimental D-Y plot during immersion Fig.5 Thickness, d, of surface films on 6Nof AZ91D in 0.1M Na2SO4 and theoretical D-Y Mg, AZ31B and AZ91D alloys as a function curve for growth of film with N2=1.418-0.053i. of immersion time, t, in 0.1M Na2SO4.

Figure 6 exhibits d vs. t curves for 6N-Mg and AZ91D in 0.1M NaCl. The film growth process in this solution is similar to that observed in 0.1M Na2SO4 (Fig.5). The film growth on AZ91D is suppressed in NaCl solution. The film thickness at the end of stage I is 30 nm, which is about a half of the corresponding thickness in 0.1M Na2SO4. 3.3. Influence of oxidation films on corrosion resistance In order to examine the influence of the formation of oxidation films on the corrosion behavior of underlying metals, anodic polarization curves of 6N-Mg and AZ91D were measured after immersion for 5min and 1h in 0.1M Na2SO4 and 0.1M NaCl solutions. Figure 7 exhibits the results obtained in 0.1M NaCl. The polarization curves of high purity Mg and AZ91D alloy show the passivity region. This indicates that the oxidation films formed during immersion have protective ability. However, passivity breakdown occurs under anodic polarization. The passive current density decreases and the breakdown potential increases with increasing immersion time. The passive current density of 6NMg covered with a thicker film was higher than that of AZ91D alloy with a thinner film. This suggests that the corrosion resistance of Mg and its alloys

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200

10

Solution: 0.1M NaCl

10

0

i / A•m

d / nm

–2

150

Solution: 0.1M NaCl

1

6N-Mg (N2 = 1.4645–0.003i)

100

10

tim = 5min d = 112nm

-1

10

6N-Mg

50

-2

10

AZ91D (N2 = 1.415–0.018i) 0 0

tim = 60min d ~ 250nm

tim = 5min d = 46nm AZ91D

tim = 60min d = 81nm

-3

1

2 t / ks

3

4

10

-1.9

-1.8 -1.7 -1.6 -1.5 -1.4 E / V (vs. Ag/AgCl(3.33M KCl))

-1.3

Fig.6 Thickness, d, of surface films on 6N-Mg Fig.7 Anodic polarization curves for 6N-Mg and and AZ91D alloy as a function of immersion AZ91D measured after immersion in 0.1M NaCl for a given time, tim. Film thickness, d, was time, t, in 0.1M NaCl. determined at the beginning of polarization

should be determined not only by the thickness of surface oxidation films but also by the chemical composition and structure of the films. It is also noted that in the case of conventional Mg-based materials, the intrinsic passivity region of Mg, ranging from -1.85V to -1.55V, should be almost hidden by the overlapping of large cathodic current. 4. Conclusion 1. Surface films with a refractive index of 1.42 - 1.47 grow rapidly on Mg and its alloys in pure water, 0.1M Na2SO4 and 0.1M NaCl solutions. 2. The film growth process was divided into two stages; stage I where the film thickness increases approximately linearly with time, and stage II where the growth rate decreases with the lapse of time. 3. The film growth rate of Mg-based materials depended on their purity and composition and the kind of electrolyte. In every test solution, the film growth rate of high purity Mg (6N-Mg) was higher than that of AZ91D alloy. 4. The results of XPS analysis suggested that the oxidation films formed on AZ91D alloy in air and aqueous solutions are mainly composed of MgCO3 and Mg(OH)2, respectively. 5. Anodic polarization curves measured after immersion in 0.1M NaCl suggested that the films formed on high purity Mg and AZ91D alloy have protective ability, while they suffer local breakdown at high potentials. The intrinsic passivity region of Mg was clearly observed on high purity Mg. References 1. J.E. Gray and B. Luan, J. Alloys Comp., 336 (2002), 88. 2. S.J. Splinter, N.S. McIntyre, W.N. Lennard, K. Griffiths, and G.Palumbo, Surf. Sci., 292 (1993), 130.

Ellipsometric Analysis of Growth Process of Oxidation Films 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.

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S.J. Splinter, N.S. McIntyre, and G. Palumbo, Surf. Sci., 302 (1994), 93. S.J. Splinter and N.S. McIntyre, Surf. Sci., 314 (1994), 157. S.J. Splinter, N.S. McIntyre, P.A.W. van der Heide, and T. Do, Surf. Sci., 317 (1994), 194. T. Do, S.J.Splinter, C. Chen, and N.S. McIntyre, Surf. Sci., 387 (1997), 192. J.H.Nordlien, S.Ono, N.Masuko, and K.Nisancioglu, J.Electrochem.Soc., 142 (1995), 3320. J.H.Nordlien, K.Nisancioglu, S.Ono, and N.Masuko, J.Electrochem.Soc., 143 (1996), 2564. J.H.Nordlien, K.Nisancioglu, S.Ono, and N.Masuko, J.Electrochem.Soc., 144 (1997), 461. G. Song, A. Atrens, D. St John, J. Nairn, and Y. Li, Corros. Sci., 39 (1997), 855. G. Song, A. Atrens, D. St John, X. Wu, and J. Nairn, Corros. Sci., 39 (1997), 1981. P. Schmutz, V. Guillaumin, R.S. Lillard, J.A. Lillard, and G.S. Frankel, J. Electrochem. Soc., 150 (2003), B99. H. Kim, N. Akao, N. Hara, and K. Sugimoto, J. Electrochem. Soc., 145 (1998), 2818. S. Matsuda and K. Sugimoto, J. Jpn. Inst. Met., 45 (1981), 203.

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Corrosion Protection of Magnesium Alloys by Ce, Zr and Nb oxide layers Hélène Ardelean, Isabelle Frateur, and Philippe Marcus Laboratoire de Physico-Chimie des Surfaces, CNRS (UMR 7045), Ecole Nationale Supérieure de Chimie de Paris, 11 rue P. et M. Curie, 75005 Paris (France).

Abstract - Two novel surface treatments (Cr6+-free) for improved corrosion protection of magnesium alloys have been designed. One is based on conversion coatings produced in a bath containing cerium, zirconium and niobium compounds. The second consists in anodic oxidation, which produces Zr and Nb oxides. The corrosion protection was evaluated by polarization curves and electrochemical impedance spectroscopy. X-ray photoelectron spectroscopy (XPS) was used for the characterization of the coating composition and of changes caused by the polarization in the corrosive medium. The electrochemical results show: i) a significant decrease of the corrosion and anodic currents and ii) higher polarization resistances for both coatings compared to other treatments. XPS showed that the novel coatings consist mainly of stable mixed metal oxides. Keywords : Magnesium alloys, Conversion coating, Anodic oxidation

1. Introduction There is an increasing interest in using magnesium and its alloys in a wide variety of structural applications, which include automotive and aerospace equipments, in order to decrease weight and fuel consumption. However, magnesium and its alloys exhibit very poor corrosion resistance. Surface treatments, including anodic oxidation, organic coatings and environmentally undesirable chromate-based conversion coatings, have been used. As part of a research program on the corrosion and corrosion protection of magnesium and its alloys, we have investigated the effects of different elements such as Ce, Al [1], Zr, Nb, on the formation and properties of protective coatings. This paper

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describes the recent advances in the corrosion protection of AZ91 alloys, achieved with new conversion [2] and anodizing [3] Cr6+-free treatments. 2. Experimental AZ91 magnesium alloys provided by Pechiney S.A. were used. In order to study the effects of different elements such as Ce, Zr and Nb, on the formation of a protective coating, magnesium alloy samples were placed in different solutions containing Ce(NO3)3 alone or Ce(NO3)3 mixed with ZrO(NO3)2 or Ce(NO3)3 mixed with ZrO(NO3)2 and NbxOyFz. Finally, a novel conversion coating was obtained in a solution of Ce(NO3)3, ZrO(NO3)2 and NbxOyFz, as described in ref. [2]. A new anodizing treatment based on NbxOyFz, ZrF4, H3PO4, H3BO4, was studied at a current density of 1-2 A/dm2 and a voltage up to 240-340 V [3]. The electrochemical behaviour was investigated using an EG&G Princeton Applied Research 273 model potentiostat/galvanostat and two classical three-electrode cells with a saturated mercurous sulphate electrode (MSE) as reference. Impedance spectra were recorded at the corrosion potential using an AUTOLAB PGSTAT30 FRA2 system (ECO CHEMIE). XPS analyses were carried out with a VG ESCALAB Mark II spectrometer using an Al Kg X-ray source. 3. Results and Discussion 3.1. Evaluation of the conversion coatings 3.1.1. Polarization curves ____

untreated alloy

- - - - DOW 22 treatment - × - Ce and Zr treatment .......

Ce treatment

- - - - - Ce, Zr and Nb [2]

Fig. 1 Polarization curves of AZ91 alloys treated in different conversion baths, recorded in 0.5M Na2SO4 (sweep rate: 0.5 mV/s).

The polarization curves of AZ91 alloys (recorded in 0.5 M Na2SO4 solution) after different conversion treatments are presented in Fig. 1. For all coatings, the corrosion potential was ennobled compared to the untreated alloy. The corrosion currents (obtained from the Tafel extrapolation) and the anodic

Corrosion Protection of Magnesium Alloys by Ce, Zr and Nb oxide layers

227

currents (measured at E=-1.5 V/MSE), were very different depending on the composition of the conversion solution. The polarization curves show the increased protective effects provided by the new Ce, Zr and Nb–based coating, which are: i) a large shift towards positive values of the corrosion potential and ii) a remarkable decrease of the corrosion and anodic currents, in comparison with the untreated alloy and the DOW 22 Cr6+-based conversion film. 3.1.2. Impedance spectra 4000 untreated DOW 22 Ce, Zr and Nb, 3h Ce, Zr and Nb, 24h

-Im Z / W

3000 2000

16.5H z 0.15Hz

1000 0 10.4mHz

-1000

0

1000

2000

3000

4000

5000

Re Z / W Fig. 2 Impedance spectra of AZ91 alloys treated in different conversion baths, plotted at Ecor after 2h of immersion in 0.5M Na2SO4. Minimum frequency: 10.4 mHz. 7 points/decade.

The impedance spectra recorded for AZ91 alloys after different conversion treatments are shown in Fig. 2. The impedance diagrams exhibit two capacitive loops in the high and low frequency ranges (HF and LF, respectively), and one inductive loop. The HF capacitive loop results from both a charge transfer and a film effect. However, in practice, the resistance associated with the HF loop is effectively that of the charge transfer resistance Rt, and can be used as such in determinations of the corrosion current (Stern and Geary relationship into which the value of Rt is introduced). The LF capacitive loop is attributed to relaxation of mass transport in a solid phase (diffusional process). The inductive loop may be due to relaxation of coverage due to an adsorbed intermediate (probably MgOHads) [4]. The spectra show a significant increase of the charge transfer resistance for Ce-Zr-Nb-based coatings. This increase of Rt, more important for longer treatment times, indicates a decrease of the corrosion current.

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3.1.3. XPS analysis Ce4+

Ce3+

Fig 3. XPS surface analysis of the AZ91 alloy with the conversion coating (a) prior to and (b) after electrochemical corrosion testing.

XPS analyses were performed on AZ91 samples after immersion during 30 minutes in a Ce, Zr, Nb-based solution [2]. CeO2, ZrO2 and Nb2O5 were detected as shown in Fig. 3(a), in agreement with literature data for Ce(IV), Zr(IV) and Nb(V) in oxides [5-8]. Fig. 3(b) shows the XPS spectra recorded after polarization from -2.4 to -1.5 V/MSE in 0.5M Na2SO4. The comparison of the Ce 3d peaks, before and after polarization, shows a slight reduction of Ce4+ into Ce3+ after polarization. The regions for Zr and Nb, also present as Zr4+ and Nb5+, respectively, in the coating, are not affected. The above data lead to the conclusion that this new coating provides improved corrosion resistance, due to the presence of stable Ce, Zr and Nb surface oxides. 3.2. Evaluation of the anodizing treatments 3.2.1. Polarization curves The polarization curves recorded in 0.5M Na2SO4 solution for AZ91 alloys after

Corrosion Protection of Magnesium Alloys by Ce, Zr and Nb oxide layers

229

different anodizing treatments are presented in Fig. 4. The polarization curves show: a significant decrease of the corrosion and anodic currents for the new (Nb and Zr-based) anodizing treatments compared to the untreated alloy and the known US-4,978,432 patent treatment. ____

untreated alloy - - - US-4,978,432 patent treatment -× -× Nb/ PCT WO 03/069026 A1 patent treatment ××××××× Nb and Zr / PCT WO 03/069026 A1 patent treatment

Fig. 4 Polarization curves of AZ91 alloys treated in different anodizing baths, recorded in 0.5M Na2SO4 solution (sweep rate: 0.5 mV/s).Impedance spectra

6000

-Im Z / W

5000

untreated US-4,978,432 patent Nb / PCT WO 03/069026 A1 patent Nb and Zr / PCT WO 03/069026 A1 patent

4000 3000

Fig. 5 Impedance spectra of AZ91 alloys treated in different anodizing baths, plotted at Ecor after 4h of immersion in 0.5M Na2SO4. Minimum frequency: 10.4 mHz. 7 points/decade.

11.8Hz

The electrochemical impedance spectra show an 1000 increase of the charge transfer resistance that indicates an 0 10.4mHz improvement of the corrosion 0 1000 2000 3000 4000 5000 6000 7000 protection for the new Re Z / W anodizing treatment compared to the untreated alloy and the known US-4,978,432 patent treatment. 2000

55.6mHz

4. Conclusions A new green conversion coating (Cr6+-free) was produced on AZ91 magnesium alloys, by a treatment in a solution containing cerium, zirconium and niobium compounds. A novel anodizing treatment was designed, based on a zirconium, niobium, phosphate and borate containing solution.

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The electrochemical measurements showed: i) a decrease of the corrosion and anodic currents, and ii) a higher charge transfer resistance for both new conversion and anodizing treatments compared to the untreated alloy and to other existing treatments. Chemical characterization of the new conversion coating by XPS revealed the presence of stable oxides: CeO2, ZrO2, Nb2O5, as main components. The positive effects of the new coatings on the corrosion resistance can be attributed to the high stability of the compounds that constitute the surface oxide layers. The coatings could be employed either as final treatment or as supports for painting. They offer particularly interesting alternatives to the Cr6+based coatings for magnesium alloys. References 1. H. Ardelean, C. Fiaud, P. Marcus, Materials and Corrosion 52 (2001) 889. 2. P. Marcus and H. Ardelean, patent application: FR n° enreg. nat. 0107173, PCT WO 02/097164 A2 (05.12.2002) and Euro-PCT n° 02743345.7-2119 / n° publ. europ. 1390565, sect. I.1, (25/02/2004). 3. H. Ardelean and P. Marcus, patent application: FR n° enreg. nat. 0201772 (2002) and PCT WO 03/0069026 A1 (21.08.2003). 4. N. Pébère, C. Riera, F. Dabosi, Electrochim. Acta 35 (1990) 555. 5. P. Burroughs, A. A. Hammett, A. F. Orchard, G. Thornton, J. C. S. Dalton Trans. (1976) 1686. 6. J. F. Moulder, W. F. Stickle, K. D. Bomben, Handbook of X-ray Photoelectron Spectroscopy, Perkin Elmer, Eden Prairie, 1992. 7. P. C. Wong, Y. S. Li, K. A. R. Mitchell, Surf. Rev. Lett. 2(3) (1995) 297. 8. A. Galtayries, R. Sporken, J. Riga, G. Blanchard, R. Caudano, J. Electron. Spectrosc. Relat. Phenom. 88-91 (1998) 951.

Section B Passivity of Semiconductors

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Passivation of (100) and (111) Silicon in KOH Solution Harold G.G. Philipsen and John J. Kelly Debye Institute, Utrecht University, P.O. Box 80000, 3508 TA, Utrecht, The Netherlands, Tel: +31-30-2532214, Fax: +31-30-2532403, E-mail: [email protected] / [email protected]

Abstract - The anodic passivation of (100) and (111) silicon surfaces in KOH solution was studied by potentiodynamic and current-transient measurements. Striking differences are observed for the two orientations. The slow oxidation kinetics of the (111) surface are attributed to the exceptional stability of this face with respect to the chemical etching reaction. Keywords: Silicon, (100), (111), Anisotropic etching, Anodic oxidation kinetics.

1. Introduction At the previous meeting in this series we reported on the passivation of (100) Si in alkaline solution. In that paper [ 1 ] and in subsequent work [ 2, 3 ] we showed that the anodic oxidation and passivation of silicon in alkaline solution is coupled to the anisotropic chemical etching reaction. If this is the case then one might expect to find considerable differences in the kinetics of oxidation for different silicon surfaces. In this paper we show that this is indeed the case for (100) and (111) surfaces. Anisotropy in anodic oxidation and passivation may be important in MEMS (micro-electromechanical systems).[ 4 ] 2. Experimental For the experiments, 4 inch wafers of p- and n-type (100) Si were masked with a silicon nitride layer and etched anisotropically in such a way that V-grooves defined by (111) facets developed. Figure 1 shows a schematic representation of

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a V-groove during formation. In order to obtain well-defined (111) surfaces, the edges of the mask windows were aligned along the <110> crystallographic direction, with an accuracy of 0.05°.[ 5 ] For the electrochemical measurements, a conventional 3-electrode setup was used. Only the front side of the sample (working electrode WE, see figure 1) was exposed to the etchant and the back side was provided with an Ohmic contact. A platinum sheet and saturated calomel electrode (SCE) were used as counter and reference electrode, respectively. Experiments were performed in a thermostated double-walled vessel, containing 5.0 M KOH (p.a. quality, Merck) solution which was continuously purged with argon to remove dissolved gases.

Figure 1 (a)

(b)

Schematic representation of V-groove formation. At the start of the experiment (t0), only (100) faces are exposed to the etchant. After the grooves are fully formed (t4), the electrode consists of (111) faces. The relevant crystallographic directions are indicated by vectors. SEM micrograph showing the end of the rectangular mask windows. The electrode consists of an array of 30 grooves of 100 mm width.

3. Results and Discussion Figure 2 shows a series of voltammograms measured during the formation of Vgrooves in p- and n-type silicon. The anisotropy in the anodic oxidation and passivation becomes immediately obvious as soon as groove etching starts.[ 6 ] The initial scan prior to anisotropic etching shows a single peak at lower potential, characteristic of the (100) surface.[ 7 ] Subsequently, voltammograms were measured at regular time intervals during anisotropic etching. The (100)

Passivation of (100) and (111) Silicon in KOH Solution

235

peak decreases and a second peak develops at considerably more positive potential as the (111) side walls are exposed (see figure 1). Finally, when the groove is fully formed, only the peak corresponding to the (111) surface remains.

Figure 2 Voltammograms measured during V-groove formation. Scan rate 10 mV/s. (a) p-type Si, 50.0°C. Measured every 30 minutes. (b) n-type Si, 60.0°C. Measured every 20 minutes.

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The voltammograms shown in figure 2 were measured at a constant potential scan rate. For a comparison of the kinetics of anodic oxidation and passivation, the scan rate was varied. The results for n-type and p-type (100) Si did not depend markedly on the scan rate.[ 7 ] On the other hand, a considerable difference was observed between (100) and (111) surfaces, this difference being most pronounced for the n-type electrodes (figure 3). In the (100) case (fig 3(a)) the potential of the current maximum corresponding to the onset of passivation does not depend on the scan rate. At higher scan rates the peak is broadened somewhat. In contrast, the peak potential for the (111) surface shifts strongly to positive potentials with increasing scan rate (fig. 3(b)); a displacement of 3 V is observed for the fastest scan (75 mV/s, not shown in fig. 3(b)). The almost linear dependence of peak potential on scan rate suggests that time, and not potential, is the determining factor in these experiments.

Figure 3 Voltammograms measured at various scan rates at 60°C. (a) n-(100), (b) n-(111)

That kinetic effects are important in the (111) case was confirmed in potentialstep experiments; current transients were measured following an abrupt change in potential at t=0 from the open-circuit value to a value in the passive range. Figure 4 shows that at all temperatures the initial current for the n-type (111) electrode is low. After some time, the current increases and reaches a maximum, after which the electrode passivates. An increase in temperature markedly accelerates the activation process. The inset of the figure shows the anisotropy in the anodic oxidation of both surfaces. For the (100) electrode no induction period is observed after the potential is stepped. An initial high current is followed by a monotonous decrease to a steady-state value. Thermal activation of the (111) surface seen in the current transients of figure 4 is also clear in voltammograms measured at different temperatures.

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Figure 4 Current transients at various temperatures for n-type (111) following a potential step from open-circuit potential to 0.00 V vs. SCE. The inset shows transients for n-(100) and n-(111), measured under identical conditions (55.0°C).

Anodic current in an n-type semiconductor requires electron injection into the conduction band.[ 2, 8 ] Previous work on (100) electrodes in KOH solution showed the injecting species to be activated intermediates of the chemical etching reaction.[ 2 ] This reaction has two main steps: the OH- catalyzed conversion of Si-H surface bonds to give Si-OH (and H2) and the breaking of the polarized Si-Si backbond by reaction with water. Two main factors are responsible for the strong etching anisotropy: the stability of the monohydride on the Si(111) surface as compared to the dihydride on the (100) surface, and the fact that a (111) Si surface atom has three backbonds, whereas a (100) surface atom has only two.[ 8 ] Since the chemical and anodic processes are coupled [ 2 ] the same factors that account for the strong anisotropy in etching can also explain the marked electrochemical differences between (100) and (111) surfaces, in particular the slow activation of anodic oxidation in the case of n-type (111) Si. 4. Conclusions Electrochemical oxidation of Si in KOH solution depends on a chemical activation of the surface. Measurement of current-potential curves and current time transients gives information about the chemical stability of (100) and (111) faces.

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References (1) J.J. Kelly, X.H. Xia, P.M.M.C. Bressers, Electrochem. Soc. Procs., 99-42, 348 (2) X. Xia, C.M.A. Ashruf, P.J. French, J. Rappich, J.J. Kelly, J. Phys. Chem. B, 105, 5722 – 5729 (3) X.H. Xia, J.J. Kelly, Phys. Chem. Chem. Phys., 2001, 3, 5304 – 5310 (4) S.D. Collins, J. Electrochem. Soc., 1997, 144, 2242 – 2262 (5) M. Vangbo, Y. Bäcklund, J. Micromech. Microeng., 1996, 6, 279 – 284 (6) H.-R. Kretschmer, X.H. Xia, J.J. Kelly, A. Steckenborn, J. Electrochem. Soc., 151, C633 – C636 (7) R.L. Smith, B. Kloeck, N. de Rooij, S.D. Collins, J. Electroanal. Chem., 1987, 103 – 113 (8) P. Allongue, V. Costa-Kieling, H. Gerischer, J. Electrochem. Soc., 1993, 140, 1026

336 – 2001,

2004, 238, 1018 –

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On the nature of oscillations of Open Circuit Potential of silicon immersed in CuSO4/HF solution Vitali Parkhutik a), Yuri Makushok a) and Yukio Ogata b) a

b

R&D Centre “Materials and Technologies of Microfabrication, Technical University of Valencia, Cami de Vera s/n, 46022 Valencia, Spain; E-mail: [email protected] Institute of Advanced Energy Sources, Kyoto University, Japan

Abstract - Oscillations of Open Cirquit Potential (OCP) during immersion plating of silicon electrodes with copper has been recently discovered and ascribed to concurrent reactions of copper deposition at the semiconductor surface and silicon oxide growth/dissolution underneath. Here we present the results of studying forced OCP oscillations (in the presence of sinusoidal external excitation) and discuss the physical mechanism of the oscillatory process Keywords: currentless deposition, oscillations, Open Circuit Potential

1. Introduction In a series of recent works a new effect of oscillatory behavior of Open Circuit Potential of silicon immersed in electrolyte consisting of 0.01M CuSO4 and about 0.05M HF has been registered [1,2]. These oscillations vary in their period, shape and regularity depending on the experimental conditions such as temperature of solution, electrolyte stirring, concentration of components, sample history, etc. We have shown that they are due to the deposition of copper layer at the surface of Si sample and simultaneous cyclic growth and dissolution of silicon dioxide films beneath the copper layer. Oscillations can last days without any tendency to hinder. This work summarizes recent experimental findings on the impact of external electrical signals (a.c. and d.c.) on the OCP oscillations and discusses the possibility to observe oscillatory behavior on silicon immersed in solutions of salts of other metals.

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2. Experimental set-up We have used the samples of p- and n-Si oriented in (100) and with different doping levels. The samples were mounted in Teflon sample holder with 1cm2 opening to the electrolyte solution which was a mixture of 0.01M CuSO4 with small variable amount of HF. Pt reference electrode and Cu counter electrode were employed. Electrolyte temperature was controlled with a precision of 0.5ºC by using a Peltier cooler. OCP kinetics was recorded using PAR273A potentiogalvanostat, by maintaining zero current and measuring the voltage drop between working electrode (silicon) and reference electrode External periodic excitation of the system was performed using frequency generator HP 3312A connected to the input of the PAR273A potentiogalvanostat. An example of typical oscillatory behavior of the OCP value is shown in Fig.1. The oscillations emerge after some induction time during which initial copper layer is deposited at the sample surface. They are quite regular and asymmetric as an inset in Fig.1 is showing.

Fig.1. Initiation of OCP oscillations during immersion of p-Si in 0.01CuSO4+0.08HF solution at 26ºC. An insert shows a zoomed fragment of oscillatory kinetics

Fig. 1 shows the very first moments of the initiation of the oscillatory process: it passes through the stage of OCP value transient (first 70 s of immersion) when Cu film is formed at the naked surface of Si wafer. Further on, oscillations emerge, they can be quite irregular (as in the case shown at Fig.1) or very regular [3] with their period and amplitude depending on the experimental variables. For example, their period Tosc depends on the electrolyte viscosity as Tosc ~ exp(h ) on the electrolyte temperature as Tosc ~ exp(-T ) .

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241

Other important parameters of the oscillatory kinetics are: the composition of the electrolyte (there exists a window concentration of both components of the electrolyte where the oscillations are regular and long-lasting), Si doping level (highly doped n-Si do not show oscillatory kinetics), electrolyte stirring and sample history. The summary of the experimental results accumulated till now can be found elsewhere [3,4]. Here we present new results on the influence of the external periodic or steady-state electrical signals onto the oscillatory OCP kinetics. Of course, the application of a.c. and d.c. voltages to the sample using external source results in appearance of some current flow through the system, and in this case we have a superposition of internal effects determining the OCP behavior with external excitation which tends to alter the integral electroneutrality of the electrochemical system. In discussion section of the paper we discuss the conditions necessary to observe the OCP oscillations during currentless deposition of other metals on silicon. 3. Results 3.1. Forced oscillations in the Si/electrolyte system In order to discard existing possibility that the OCP oscillations are due to some artifacts of the electrochemical system or related with experimental setup, we have analyzed the influence of external electrical signals onto the occurrence of oscillations and their ordering.

Fig.2. Result of superposition of external a.c. voltage with amplitude 0.1 V and period 0.26s (left) and 0.68s (right) onto OCP oscillations with period 0.44s.

First, we have studied the application of sinusoidal external voltage (amplitude Aext) of varying period (Text) to the system oscillating with the intrinsic period

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TO. Fig.2 shows the results for the a.c. signal with Aext= 0.1V of both smaller and higher period. The application of fast external oscillations provokes the appearance of some chaos in the system during first 5 seconds, but then OCP oscillation becomes prevailing. Secondary oscillation of double frequency and smaller amplitude emerges as a result of external modulation that is known as bifurcation effect. Slower oscillations (Fig.2,right) are dominating at the beginning (first 8 seconds) but then the oscillatory kinetics becomes chaotic as a result of interference between external and internal oscillatory processes. Both of these results show clearly that the OCP oscillations are the internal property of the Si/electrolyte system and not the artifacts of the experiment. We have also studied the influence of steady state polarization (both positive and negative) onto the OCP oscillations. Applications of small negative or positive currents through the system did not influence qualitatively the OCP kinetics unless a certain critical value is achieved. Then the oscillations become dumped and disappear with time. Fig.3 shows the results.

(a)

(b)

Fig.3 Influence of the superposition of steady- state cathodic (a) and anodic (b) current onto the OCP oscillations.

According to Fig.3, the application of both cathodic and anodic currents (if their value is sufficiently high) causes dumping of the oscillations, but the reason for this is different in both cases. In the case of cathodic current Cu deposition process prevails over the Si oxidation, while in the opposite case Si oxidation reaction is dominating. The results on forced oscillations in Si/electrolyte system demonstrate, on the one hand, that this is a typical case of non-linear system with strange attractors, bifurcations and other features specific for stochastic resonance systems. On the other hand, it does not coincide with the results of previous works on forced

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oscillations in electrochemical systems, where much less stability towards external excitations has been found [5]. 3.2. Morphology of the Cu/Si thin film system Copper film formed at the surface of Si electrode is weakly adherent and can be easily removed and analyzed using electron microscopy. Its general view is shown in Fig.4,a, and in more details both external and internal surfaces of free-standing Cu film are illustrated in Fig.4(b,c).

(a)

(b)

(c)

Fig. 4. SEM imaging of Cu film formed at the surface of Si electrode (a) and more detailed view of its external (b) and internal (c) surfaces

Depending on the experimental conditions Cu film can be very rough and possess a multi-layer structure (Fig.4,a gives an idea of this layering). Very important is the granular structure of the deposit and its easy permeability by an electrolyte. Beneath copper film Si electrode is corroding with a formation of a crater with flat mirror-like polished bottom and depth proportional to the time of the process. 4. Discussion We have shown earlier [3] that the oscillatory behavior of OCP is due to concurrent and mutually conditioned reactions of Cu deposition and Si

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oxidation/dissolution. This mechanism is supported by direct microscopical observations and electrochemical considerations as well. Based on this mechanism It is interesting, to consider the possibility to observe similar oscillations on silicon electrode immersed in salts of metals different from copper. Based on thermodynamic calculations of red-ox potentials for different cationic species, a list of possible candidates for the occurrence of OCP oscillations should include: bismuth, tellurium, rhenium, osmium, ruthenium and mercury. All these metals have their red-ox potentials similar to that of copper and thus could be good candidates for the observation of the OCP oscillations similar to those reported in this paper. In sum, the results reported in this paper allow to conclude that the phenomenon of the OCP oscillations in Si/CuSO4+HF system. on the one hand, fits quite well into the general phenomenology of the complex physical-chemical systems experiencing chaos-order transitions. On the other hand, the present case seems different from the cases of the oscillatory behavior in electrochemical systems. References 1. J.Sasano, Y.Ogata, (private communication) 2002. 2. Yu.Ogata, J.Sasano, T.Itoh, T.Sakka, E.Rayón, E.Pastor, V.Parkhutik, J.Electrochem.Soc., 152(2005)C537. 3. V.Parkhutik, J.Sasano, Y.Ogata, E.Matveeva, In: “Noise in Complex Systems and Stochastic Dynamics” Ed. L.Shimansky-Geller, D.Abbott, C.Van der Broeck, Proceedings SPIE, v.5114, 2003, p.396-405. 4. V. Parkhutik, E. Rayon, E. Pastor, E. Matveeva, J. Sasano, and Y. Ogata, Phys.Stat.Solidi (a), 202(2005)1586. 5. P.Parmananda , M. Rivera, R. Madrigal, Electrochimica Acta, 44 (1999) 4677.

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Passivation and Local Corrosion of p-Silicon in Anhydrous Organic Solutions of Chlorides U. Lelek-Borkowskaa, J. BanaWa a

AGH – University of Science and Technology, Faculty of Foundry Engineering, Department of General and Analytical Chemistry, ul. Reymonta 23, 30-059 Kraków, Poland

Abstract - The electrochemical investigations of p–silicon single crystals have been performed in anhydrous DMF solutions of LiCl and HCl to examine the possibility of etching and passivation of silicon semiconductors in such media. The results obtained by means of cyclic voltammetry (CV) and XPS surface analysis showed that silicon dissolves in anhydrous N,N–dimethylformamide containing chlorides, similarly like in methanol, according to the consecutive two–step mechanism. The presence of Si(II) intermediate was confirmed by means of XPS measurements on potentiostaticallyetched surface. The adsorption of DMF molecules on a silicon surface inhibits the anodic dissolution. Replacement of LiCl for HCl decreases the dissolution process rate. Microscopic observations of surface morphology of Si single crystals after etching show anisotropy of anodic dissolution. Keywords: electrochemical dissolution, silicon, etching, organic solvents

1. Introduction Organic solutions of electrolytes became a subject of high interest during last decade because of a possibility of etching and passivation of silicon [1-11]. Our investigation of metal dissolution and passivation in organic environments shows that these media give a possibility of controlling the passivation process by the reducing or determining the concentration of passivating agent - source of oxygen necessary to oxide formation (water or oxy-acid molecules) [12,13]. The organic solutions of electrolytes enable also the control of the etching

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morphology of metal surface, from polishing to structural etching, dependent on the dielectric permittivity and donor-acceptor properties of the solvent. The earliest papers concerning anodic behavior of silicon in organic solvents [1-3] dealed with formation of oxide film in presence of oxy-anion and small amount of water. The oxy-anion and water were told to be a source of oxygen in oxidation process [2,3]. Rieger, Flake, Kohl et. all investigated silicon etching in acetonitrile containing hydrogen fluoride [4] or complex MFm-1 anions, which were the fluoride donors for anodic reaction [5]. They found, that in these environments silicon dissolves to tetravalent form without hydrogen evolution. Mechanism of silicon dissolution, like in water was initiated by hole consumption, which was followed by a few steps of electron injection [4]. In fluoride-free solutions fluorine containing ions like: BF4-, PF6-, CF3SO3-, AsF6-, SbF6- [5] were the source of F- ions necessary for Si(IV) complexing. There were also investigated the influence of organic solvent addition (Ndimethylformamide, dimethylsulphoxide and ethanol) on morphology of pores growing in silicon surface layer [6-9]. It was found that pore diameters and morphology depends on the type of organic solvent additions [8]. It was found during last years that modification of silicon surface by grafting of organic groups allows to achieve desired properties – chemical stability and electronic passivation (low concentration of electronic surface states in the band-gap) [10,11] The aim of this study was to explain the mechanism of anodic dissolution in anhydrous methanol and N,N-dimethylformamide solutions of chlorides. 2. Experimental Boron doped (p-type) silicon single crystals (111) oriented with the resistivity up to 10 ×cm-1 were prepared by gold sputtering on a back side to obtain metal-semiconductor junction. Anhydrous methanol was prepared according to the procedure described in [16]. Anhydrous N,N-dimethylformamide with water amount less than 0.005% was delivered by POCh Gliwice, Poland. The amount of water in methanol solutions (<0.01%) was evaluated by Karl-Fischer method. CH3OH-LiCl and DMF-LiCl solutions were composed by dissolution of anhydrous lithium chloride solvent. The CH3OH-HCl and DMF-HCl solutions were obtained by saturation of solvent with dried gaseous HCl. The concentration of HCl was established by means of atomic absorption spectroscopy. Three-electrode system: p-Si as working, platinum - auxiliary and Ag/AgCl/Cl-,methanol or Ag/Ag+, methanol as a reference electrode was applied in electrochemical measurements. All the potential values were recalculated to the SHE. The shift was about 600mV (depending on concentration). The investigations were performed at standard temperature in the deaerated solutions (Ar bubbling).

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Cyclic voltammetry (CV) was applied for evaluation of polarization curves with use of PS-6(Meinsberg, Germany) and Voltalab PGZ301(Radiometr Analytical, France) potentiostats. The morphology of etched surface was investigated with the use of X-ray photoelectron spectroscopy (XPS) technique (ESCA-VSV-spectrometer, MgKg X-ray of 1253.6eV was applied) and scanning electron microscopy (SEM) with EDS analyzer (JEOL 5500 LV). 3. Results and discussion 3.1. Anodic behavior of p-Si in methanol-chloride solutions Our research on p-Si dissolution in anhydrous methanol containing chloride showed that this process goes ahead according to two consecutive steps SisSi(II)sSi(IV) [14,15]. At low overvoltage the formation of “porous silicon” was observed during anodic dissolution. At the range of medium anodic overvoltage the silicon surface was covered by relative stable film of Si(II)ad compounds. At high overvoltage the anodic dissolution started at defects of single-crystal`s surface and structural etching of the silicon surface was observed. The presence of oxygen in anodic products indicated the participation of methanol in the early stages of anodic reaction. The reaction order dlogi/dlog[Cl-] equal 0.65 suggested that chloride ions take part in the further stages of reaction. The presence of hydrogen ions accelerates the process rate by participation in the reaction[15]. 3.2. Anodic behavior of p-Si in N,N-dimethylformamide - chloride solutions

3,0

i /mAcm

i /mA*cm

-2

-2

Typical voltammogram for p-Si (111) polarized in N,N-dimethylformamide containing lithium chloride is presented in Fig. 1. p-Si(111) DMF-0,1M LiCl 1V/min

2,5

0,050 0,045 0,040

2,0 0,035

1,5 0,030

1,0

U = 1.9V i = f(v) 2 R =0,948

0,025

0,5

0,020

0,0

0,015

0

500

1000

1500

2000

2500

3000

3500

U /mV vs SHE

Fig.1. Typical anodic polarization curve for p-Si in 0.1MLiCl-DMF solution.

0

1

2

3

4

5

6

7

8

9

10

11 -1

v /Vmin

Fig.2. Potential scan rate vs. current density for p-Si in 0.1MLiCl-DMF solution at 1.9V

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log i

The curve is initially flat and the current density is very low. It seems that the electrode surface is blocked by adsorbed molecules of solvent (DMF). Over the potential of 2V current density increases rapidly, which means the breakdown of the surface film and start of the silicon etching. The current loop appearing on a polarisation curve indicates that silicon surface starts to etch locally. The correlation between the potential scan rate and current density for 0.1M LiCl-DMF solution at activation potential (~2V) is linear (Fig.2.). This means that the adsorption determinates the whole process rate. At higher potential (>2.3V) current density becomes the square root function of potential scan rate, which indicates that diffusion overtakes the control of the process. Fig. 3 presents the influence of Cl- concentration on characteristic potential values: corrosion, activation and repassivation potential. It can be observed that the increase of chloride ions concentration shifts the activation and repassivation potentials to the more noble values. This suggests the inhibiting influence of chloride containing anodic product on the reaction rate. Replacement of LiCl with HCl in N,N-dimethylformamide solutions decreases anodic current density for p-silicon (Fig.4) and thereby the dissolution process rate. This effect is probably connected with a very low dissociation degree of hydrogen chloride in this solvent (a = 10-3). 0,0

p-Si(111) DMF-0,05M HCl DMF-0,01M LiCl 1V/min

-0,5 -1,0 -1,5 -2,0 -2,5 -3,0 -3,5 -4,0 0

500

1000

1500

2000

2500

3000

U /mV vs SHE

Fig.3. The influence of Cl- concentration on characteristic potential values

Fig.4. Influence of H+ presence on current density for p-Si in DMF-Cl- solutions.

Fig.5. presents XPS spectra obtained for silicon surface polarized in 0.5MLiCl-DMF at potentials of 1.9V (breakdown potential). The main component of surface layer is Si(IV) - probably primal silicon dioxide layer stabilized with adsorbed DMF molecules. The presence of adsorbed N,Ndimethyl formamide is confirmed by signals from N-C-O components in O1s band [17] as well as from C bounded with N in C1s band and N-C bounds in N1s band [18]. There is also present a small amount of Si(II) - intermediate surface product [19]. There is no signal in Cl2p band. At high overpotential of 3.8V the layer is much thinner (stronger signal from Si(0)), but the overall content is similar. Analogous results can be observed for silicon polarised in DMF-HCl solution.

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Fig. 6 presents pictures of silicon surface polarized in 0.5MLiCl-DMF solution at high overpotential (3.8V) for 15 min.

Fig.5. XPS spectra for p-Si (111) polarized at ~2V in 0,5MLiCl – DMF solution.

Fig.6. Surface of p-Si (111) etched in DMF0.5M LiCl at 3.8V for 15min. (magn. 250x)

Fig.7. Surface of p-Si (111) etched in DMF0.05MHCl at 3.8V for 15min. (magn.2500x)

There are characteristic figures visible – “crosses” on the etched surface. It seems that locally initiated pits grow in specified directions. Such kind of oriented dissolution of silicon surface can be explained according to [20]. The silicon surface etched in 0.05M HCl-DMF solution at high overpotential (3.8V) looks quite different from that treated in LiCl solution (Fig.7). The figures forming during etching of silicon single crystal’s surface are squareshaped. They are visible only “bottoms” of the tetrahedrons growing into the bulk silicon. This kind of dissolution is similar to that in methanol-HCl [15].

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4. Conclusions Silicon undergoes dissolution and passivation in nonaqueous organic solutions of chlorides. The dissolution process proceeds similarly to aqueous fluoride media – with formation of divalent surface intermediate. Aprotic solvent, such as N,N-dimethyl formamide, stabilizes primary oxide layer. Etching of silicon surface starts only at high anodic overpotentials (>2V) after desorption of solvent molecules. The etching process of p-Si in DMF solutions of chlorides is anisotropic. References 1. P.F. Schmidt, W. Michel, J. Electrochem. Soc. 104, (1957) 230. 2. M. Croset, E. Petreanu, D. Samuel, G. Amsel, J.P. Nadai, J. Electrochem. Soc. 118 (1971) 717. 3. M.J. Madou, W.P. Gomes, F. Fransen, F. Cardon, J. Electrochem. Soc. 129 (1982) 2749. 4. E.K. Propst, P.A. Kohl, J. Electrochem. Soc. 141, (1994) 1006. 5. M.M. Rieger, J.C. Flake, P.A. Kohl, J. Electrochem. Soc. 146, (1999)1960. 6. S. Lust, C. Lèvy–Clèment, Phys. Stat. Sol. (a) 182, (2000) 17. 7. G. Hasse, M. Christophersen, J. Carstensen, H. Föll, Phys. Stat. Sol.(a) 182 (2000). 23, 561 8. M. Christophersen, J. Carstensen, A. Feuerhake, H. Föll, Mat. Sci. Eng. B 69–70 (2000) 194. 9. S. Izuo, H. Ohji, P.J.French, K.Tsutsumi, Sensors Atuators A 97-98 (2002) 720. 10. A. Teyssot, A. Fidelis, S. Fellah, F. Ozanam, J.-N. Chazalviel, Electrochim. Acta 47 (2002) 2565. 11. B. Fabre, D. D.M. Wayner, J. Electroanal. Chem. 567 (2004) 289. 12. J. BanaW, Electrochim. Acta 32, (1987) 871. 13. J. BanaW, Passivity of Metal in Anhydrous Solutions of Oxy–Acids, Mat. Sci. Forum, 185 (1995) 845. 14. U. Lelek-Borkowska, J. BanaW, Electrochim. Acta 47 (2002) 1121. 15. J. BanaW, U. Lelek-Borkowska, M.Starowicz, J. of Solid State Electrochem. 8 (2004) 422. 16. A.I. Vogel, Preparatyka organiczna, ed. by WNT, Warszawa (1984). 17. L.N. Bui, M. Thompson, N.B. Mckeown, A.D. Romaschin, P.G. Kalman, Analyst 118 (1993) 463. 18. P.Y. Jouan, M.C. Peignon, Ch. Cardinaud, G. Lemperiere, Appl. Surf. Sci. 68 (1993) 595. 19. G. Hollinger, Appl. Surf. Sci. 8 (1981) 318. 20. V. Lehmann, H. Föll, J. Electrochem. Soc. 137 (1990) 635.

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Silicon Texturing Under Negative Potential Dissolution (NPD) Conditions Y. Ein-Eli, N. Gordon and D. Starosvetsky Department of Materials Engineering, Technion-Israel Institute of Technology, Haifa 32000, Israel

Abstract - "As – cut" (111) oriented p-type silicon can be textured via NPD process at potentials below -10 V in KOH concentrations between 16 and 32% wt. The surface of textured silicon is characterized with morphology of arrayed pits in the shape of flat bottomed triangles. It was established that the morphology of textured (111) silicon surface, as well as a current-time profile recorded during texturing via NPD are strongly dependent on the applied potential. Increase in KOH concentration and a negative shift in the applied potential significantly reduce texturing time. Key words: silicon, Texturing, NPD, alkaline

1. Introduction Texturing of single and multi crystalline silicon is of great importance in the solar cell industry in order to reduce light reflection [1-6]. Silicon textured surfaces can be obtained by isotropic etching in HF based solutions [1, 4, 5], anisotropic etching in alkaline organic or inorganic solutions [6, 7], reactive ion etching (RIE) [5], and by electrical discharge [3]. Anisotropic etching is considered to be a highly effective silicon texturing technique. It is based on a marked difference in the dissolution rate of silicon surfaces with different crystallographic orientations [7-21]. Anisotropic etching is also widely spread in microelectronics, MEMS fabrication and crystal growth industries. [8, 12, 15, 17-19, 22-26]. Silicon texturing based on anisotropic etching in alkaline solutions is usually performed by the immersion and exposure of the substrate at open circuit potential (OCP). This process is time consuming since the silicon etching rate in alkaline solutions at OCP is relatively low. Previous work conducted by our group with polished and as-cut (100) and (110) silicon surface orientations showed that the etch rate

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of silicon can be significantly accelerated under negative potential dissolution (NPD) conditions, resulting in a major reduction of texturing time [27-29]. In the present work we studied the texturing of as-cut (111) oriented p–type silicon under NPD conditions in the dark. The experimental works with (111) oriented silicon were conducted with "ascut" wafers, allowing NPD of p-type silicon without the need of irradiation [27]. 2. Experimental Texturing of "as – cut" p – type (111) (8 – 12 W-cm) oriented silicon was studied in potassium hydroxide (KOH) solutions at concentrations of 8-50 wt% at room temperature at potential range of 異10V and –50V. A detailed description of the experimental set-up is provided in ref. 27-29. 3. Results and Discussion 3.1. Current-Time Transients and Morphology as a function of NPD Potential Figure 1 presents current-time profiles of as-cut p-type silicon with crystal orientation of (111) recorded during texturing in 24 wt% KOH under different applied cathodic potentials in the dark. At applied potentials of 異10 V and 異20 V the cathodic currents reached values of 異0.013 and ~ 異0.103 A/cm2 respectively, within a short time subsequent to potential application. During further exposure the current slowly increased with time. pS i< 1 11> etch in g 2 4% K O H

-10 V -20 V -30 V -40 V

2

Current (A/cm )

1 .5

1 .0

0 .5

Figure 1: Current-time transient obtained from negatively polarized as-cut p-silicon (111) measured in 24 wt% KOH at different applied potentials of -10, -20, 30 and -40 V vs. SCE.

0 .0 0

1 00 0 20 00 T im e (sec)

3 00 0

At a cathodic voltage of 異30 V the shape of the current-time curves markedly changed: cathodic current soared to a maximum value (1.449 A/cm2) upon voltage application and slowly decreased with time, reaching a minimum after 50 min exposure. A slight increase in cathodic current was detected in the last 10 minutes of the exposure. At an applied potential of -40 V the maximum value of cathodic current was higher (1.67 A/cm2) compared with current values recorded at -30 V polarizations. During further exposure

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the cathodic current rapidly decreased, reaching a minimum (0.5 A/cm2) within ~5 minutes followed by a rapid increase in the current which is observed to be faster than the current increase recorded at a potential of -30 V. Scanning Electron Microscopy (SEM) micrographs obtained from silicon surface subsequent to 60 minutes exposure to potentials of -10, -20, -30 and -40 V in 24 wt% KOH are shown in Figs. 2 a-d. a

b

20mm

20mm

d

c

20mm

Figure 2: SEM micrograph of silicon surface obtained subsequent to 1 hour texturing via NPD: (a)-10 V; (b) -20 V; (c) -30 V and (d) -40 V vs. SCE in 24 wt% of KOH.

200mm

The morphology of silicon surface obtained after 60 minutes etching at – 10 V (Fig. 2a) was practically similar to the morphology of untreated as – cut silicon surface. NPD at a potential of – 20 V for 60 minutes revealed pitting of the surface and a texturing in the shape of coined triangles, which are characteristic of (111) oriented silicon surface (Fig. 2b). The pits in the shape of coined triangles obtained at -20 V are small, with triangle sides of length less than 10 mm. Similar surface morphology of arrayed triangles pits was also obtained while NPD process was performed at -30 V (Fig. 2c). At these conditions the pits triangle side length was slightly longer and was measured to be more than 10 mm. Micrographs of silicon surface subsequent to -40 V NPD process presents arrayed of triangles pits, which are very large and shallow (Fig. 2d). The length of triangle sides was

measured to be in the range of 50-60 µm. The development of a surface morphology during etching time under negative potentials can be described by modification in the current values during etching time. The representative "U" shaped current-time curve (Fig. 1) being a characteristic of NPD texturing process of as-cut silicon is schematically shown in Fig. 3. 4. Conclusions "As – cut" (111) oriented p-type silicon can be textured via NPD process at potentials below -10 V in KOH concentrations between 16 and 32% wt. The surface of textured

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silicon is characterized with morphology of arrayed pits in the shape of flat bottomed triangles. It was established that the morphology of textured (111) silicon surface, as well as a current-time profile recorded during texturing via NPD are strongly dependent on the applied potential. Increase in KOH concentration and a negative shift in the applied potential significantly reduce texturing time.

Current increases due t o surface ar ea increase caused by overt ext uring

Current

Figure 3: Model describing current-time “U” shape profile of ascut p-type silicon NPD in the dark.

Current increases due t o defect s zone

Current decreases due t o defect s zone rem oval

Text ur ing

Tim e

Acknowledgments This research work was funded by EC PF5 under contract No. ENK CT2001 00561 FANTASI, by the German-Israel Foundation (G.I.F) under contract number I-2050. References 1. S. W. Park, D. S. Kim, S. H. Lee, J. Material Science, 12, (2001) pp. 619 – 622 2. M.A. Green, J. Zhao, A. Wang, S.R. Wenham, IEEE, Electron Device Letters, 13, (1992) pp. 317 – 318 3. J. Qian, S. Steegen, E. Vandor Poorten, D. Reynolds, H. Van Brussel, International J. Machine Tools & Manufacture, 42, (2002) pp. 1657 – 1664 4. P. Velinden, O. Evrard, E. Mazy, A. Crahay, Solar Energy and Solar Cells, 26 (1992) pp. 71 – 78 5. J. Szlufcik, F. Duerinckx, J. Horzel, E. Van Kerchaver, H. Dekkers, S. De Wolf, P. Choulat, C. Allebe, J. Nijs, Solar Energy materials and Solar Cells, 74, (2002) pp. 155 – 163 6. D.S. Kim, K.Y. Lee, S.H. Won, M.J. Cho, S.W. Park, S.H. Lee, Current Applied Physics, 1, (2001) pp. 505 – 508 7. E. Vazsonyi, K. De Clercq, R. Einhaus, E. Van Kerschaver, K. Said, J. Poortmans, J. Szluficik, J. Nijs, Solar Energy Materials & Solar Cells, 57, (1999) pp. 179 – 188 8. H. Siedel, L. Csepregi, A. Heuberger, H. Baumgärtel, J. Electrochem. Soc., 137, (1990) pp. 3612 – 3626 9. H. Siedel, L. Csepregi, A. Heuberger, H. Baumgärtel, J. Electrochem. Soc., 137, (1990) pp. 3626 – 3632 10. A. Reisman, M. Berkenblit, S.A. Chan, F.B. Kaufman, D.C. Green, J. Electrochem. Soc., 126, (1979) pp. 1406 – 1415 11. M. Shikida, K. Tokoro, D. Uchikawa, K. Sato, J. Micromech. Microeng., 10, (2000) pp. 522 – 527 12. M. Shikida, K. Sato, K. Tokoro, D. Uchikawa, Sensors and Actuators, 80, (2000) pp. 179 – 188

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13. K. Sato, M. Shikida, Y. Matsushima, T. Yamashiro, K. Asaumi, Y. Iriye, M. Yamamoto, Sensors and Actuators A, 64, (1998) pp. 87 – 93 14. R.M. Finne, D.L. Klein, J. Electrochem. Soc., 114, (1967) p. 965 – 970 15. D.B. Lee, J. App. Phys., 40, (1969), pp. 4569 – 4574 16. I. Zubel, I. Barycka, K. Kotowska, M. Kramkowska, Sensors and Actuators A, 87, (2001) pp. 163 – 171 17. É. Vázsonyi, Z. Vértesy, A. Tóth, J. Szlufcik, J. Micromech. Microeng., 13, (2003) pp. 165 – 169 18. J. Chen, L. Liu, Z. Li, Z. Tan, Q. Jiang, H. Fang, Y. Xu, Y.Liu, Sensors and Actuators A, 96, (2002) pp. 152 – 156 19. E. van Veenendaal, K. Sato, M. Shikida, A.J. Nijdan, J. van Suchtelen, Sensors and Actuators A, 93, (2001) pp. 232 – 242 20. I. Zubel, Sensors and Actuators, 84, (2000) pp. 116 – 125 21. I. Zubel, I. Barycka, Sensors and Actuators A, 70, (1998) pp. 250 – 259 22. N. Gabouze, Surface Science, 507 – 510, (2002) pp. 429 – 433 23. R. E. Oosterbroek, J. W. Berenschot, H. V. Jansen, A. J. Nijdam, G. Pandraud, A. van den Berg, M. C. Elewenspoek, J. Microelec. Micromechan. Systems, 3, (2000) pp. 390 – 398 24. B. C.S. Chou, C. N. Chen, J. S. Shie, Sensors and Actuators, 75, (1999) pp. 271 – 277 25. G. Ensell, J. Micromech. Microeng., 5, (1995) pp. 1 – 4 G. Kuchler, D. Scholten, G. Muller, J. Krinke, R. Auer, R. Brendel, Proceedings of the 16th 26. European Photovoltaic Solar Energy Conference, Glasgow, United Kingdom, May 1-5 (2000) 2 pp. 16951698. 27. Y. Ein–Eli and D. Strarosvetsky, Electrochem. & Solid State Letters, 6, (2003) pp. C47-C50 28. D. Starosvetsky, N. Gordon and Y. Ein-Eli, Electrochem. & Solid State Lett., 7, (2004) pp. G75-G78 29. D. Starosvetsky, M. Kovler and Y. Ein-Eli, Electrochem. & Solid State Lett., 7, (2004) pp. G168G171

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Fermi Level Pinning at n-GaAs(110) Electrodes Valentina Lazarescua, Mariuca Gartnera, Rares Scurtua, Mihai Anastasescua, Adrian Ghitaa, Mihail F. Lazarescub, Wolfgang Schmicklerc a

Institute of Physical Chemistry “I.G. Murgulescu”, Spl. Independentei 202, P.O. Box 12-194, 060041 Bucharest, Romania (e-mail: [email protected]) b National Institute of Material Physics,Str. Atomistilor 105, Magurele, CP MG7, 77125 Bucharest, Romania c University of Ulm, Albert-Einstein-Allee 11, 89069Ulm, Germany

Abstract: Mott-Schottky plots and potential-dependent second harmonic generation measurements performed at n-GaAs(110) electrodes in 1 N H2SO4 revealed the Fermi level pinnig near midgap. The possible nature of the bandgap states is discussed on the basis of the information provided by electrochemical impedance spectroscopy and in situ spectroscopic ellipsometry studies Keywords: GaAs(110), Fermi-level-pinning, SHG, EIS, SE

1. Introduction Vacuum well-cleaved GaAs(110) surfaces are known to exhibit flat bands up to the surface [1, 2], the lack of the intrinsic surface states being attributed to reconstruction of the cleaved surfaces [3]. After exposures to oxygen [4] and hydrogen [5] both surface acceptors and donors were, however, observed to form. Fractional monolayers (~0.1) of oxygen [6], or hydrogen [7] were found sufficient to induce surface states that pin the Fermi level of n-type (110)-orientated samples near mid-gap. Mott-Schottky plots and potentialdependent second harmonic generation (SHG) measurements performed in 1 N H2SO4 revealed that a similar phenomenon occurs as well at n-GaAs(110) electrodes / solution interface [8]. In the present work, the possible nature of the surface states responsible for the Fermi level pinning on the n-GaAs(110) electrodes in an electrochemical environment is discussed on the basis of the

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information provided by electrochemical impedance spectroscopy and in situ spectroscopic ellipsometry studies. 2. Experimental Investigations were performed on Si-doped GaAs(110) electrodes with n=1-2x1017 cm-3 prepared from single crystals grown in a Cambridge Instruments MSR-6 pulling machine by Cz-LEC procedure from highly pure components. All the details regarding the sample preparation could be found in the previous work [8, 9]. The as-polished GaAs(110) electrodes were etched in (3:1:1) H2SO4:H2O2:H2O followed by HCl wash and water rinsing before being placed in the measuring cell. The measurements were performed under potentiostatic conditions in well-deaerated 1N H2SO4 solution in a potential range where no Faradaic processes takes place. Electrochemical impedance spectroscopy (EIS) investigations were carried out with IM-6 Zahner frequency analyzer in the range of 1 Hz – 500 kHz. The EIS data have been analyzed by considering a conventional fivecomponent equivalent circuit which satisfactorily describes the response of the surface/interface states having time constants within this frequency range as Batchelor and Hamnett [10] thoroughly discussed. Second harmonic generation (SHG) investigations were performed by using the fundamental of a Q-switched Nd:YAG laser (1064 nm at 20 Hz, with 12 ns pulse length and 25 mJ in a spot size of 0.2 cm2) as previously described [8, 9]. In-situ spectroscopic ellipsometry (SE) measurements of phase difference (D) and amplitude ratio (Y) were performed at an angle of incidence of 700 in the 800-900 nm range by 10 nm step. 3. Results and discussion The Mott-Schottky plot shown in Fig. 1a exhibits a distinct slope change around – 0.2 V (SCE) as the density of the surface/interface states at Fermi level, NSS(eF), (which can be roughly estimated from the corresponding capacitive contribution [11]) becomes higher than 1012 cm-2. Information on the composition of the interfacial region of the GaAs(110) electrode inferred from in situ UV-VIS-NIR ellipsometric spectra bring some more light in this respect. Thickness and composition of the electrode interfacial region were estimated by fitting the measured spectra of the ellipsometric parameters (D,Y) with Bruggeman Effective Medium Approximation (B-EMA) model [13]. Layer thickness (d), optical constants (n, k) and volume fractions of the components were determined from the best fit obtained with a three-phase model, consisting of GaAs-substrate/GaAstransition-layer/electrolyte. The transition layer having an initial thickness of 100 Å was taken as an effective medium mixture of the following possible

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constituents, whose dielectric functions were found in the literature: GaAs [14], GaAs oxide [15], crystalline Ga [16], amorphous As [17], and H2SO4 solution, whose optical properties were found in literature. The values of the mean square difference (MSD) between the measured and the calculated ellipsometric data, expressing the quality of the fit: (1) MSD={(1/N)S[Reric- Rerim)2 + (Imric-Imrim)2]} (where: N is the total number of measured points; m and c represent the measured and calculated data, respectively; Rer=tany*sinD and Imr=tany*cosD) were in the range of 2-3x10-5.

Fig. 1 Potential dependence of surface states densities at Fermi level NSS(iF) (ﾖ), Mott-Schottky 110 plot (ﾐ) and isotropic field parameter, A pp (ﾒ) for n-GaAs(110) in 1N H2SO4

The applied potential influences both experimental ellipsometric parameters as seen in Fig. 2, where data taken for two incident beam energies relatively close to the optical band gap of the gallium arsenide (1.39 and 1.49 eV) are shown. The influence exerted by the applied field on D and Y is, however, different: D decreases rather monotonously with the applied potential (Fig.1a, c) whereas Y exhibits a pronounced change around -0.2 V (Fig.1b, d). The variation of D reflecting mainly changes in d might be due to the gradual hydroxyl – hydrogen replacement in the electrode surface coverage revealed by in situ IR experiments [18]. The changes in the Y profile occurring in the same potential range where the slope of the Mott-Schottky plot is altered seem to be related to the Fermi level pinning. Such a variation expressing modifications of

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the optical constants, suggests significant changes of the volume fraction of the surface layer components in the potential range where the applied field charges not only the depletion region but also the surface/interface states. The volume fractions of the three main potential suppliers of localized energy levels within the band gap, As-excess, Ga-excess and GaAs-oxide derived from the best fit of the experimental data are presented in Fig. 3.

Fig. 2 Potential dependence of ellipsometric parameters, (left-side) and incident beam energies of 1.49 eV (up) and 1.39 eV (down)

(right-side) at

Fig. 3 Potential dependence of (a) %As (ﾐ), (b) % GaAs oxide ( ) and (c) %Ga (ﾒ)

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It is easy to observe that unlike the oxide fraction (Fig 3b) and the gallium excess (Fig. 3c) the arsenic excess decreases rather monotonously (Fig. 3a) over the entire potential range. According to Pouirbaix diagrams [19], As is insoluble in aqueous electrolytes as long as it is not oxidized. However, the use of HCl as a final cleaning step (which is a standard procedure in preparing GaAs substrates [20]) was found to selectively remove Ga leaving a highly reactive porous As-rich structure consisting in clusters of irregular morphology [21, 22] rather than a compact elemental deposit. The decrease of As content of the transition layer when the electrode potential takes more negative values indicates that As atoms confined to the topmost few substrate layers are unstable. The data in Fig. 3 show that As dissolution is accompanied by an increase of both the oxygen-containing species (Fig. 3b) and the gallium excess (Fig. 3c) within the transition layer. The two processes appear to be concurrent, the first one prevailing at potentials above -0.2 V(SCE) and the latter one to potentials bellow -0.2 V (SCE). The more than double mobility of Ga3+ compared with As3+ [23, 24] should be the reason for which the lattice sites of the As atoms lost by dissolution are taken by gallium atoms. This vacancy exchange mechanism resulting in GaAs antisite defects is strongly supported by the high densities of surface/interface states crossed by the Fermi level at E < 0.2 V, centered at about 0.6 eV below the conduction band (Fig. 1a). It is worth mentioning that although the excess As (AsGa) giving birth to a pair of donor levels is largely considered to be the key defect in the advanced models proposed for Schottky barrier formation [25, 26], there is evidence that compensating acceptor levels should also exist. The shift in the energy distribution of the photoemitted electrons from GaAs surfaces with As excess with respect to those with Ga excess [27] as well as the shift of the Fermi level pinning position entailed by the excess of Ga produced by interfacial chemistry [28] are good arguments in this respect. 4. Conclusions Mott-Schottky plots and potential-dependent second harmonic generation (SHG) measurements revealed the mid-gap Fermi level pinning at nGaAs(110) electrodes / solution interface. Information provided by electrochemical impedance spectroscopy (EIS) and in situ spectroscopic ellipsometry (SE) studies point to the surface/interface states associated to missing arsenic or GaAs-antisite defects as possible responsible for the Fermi level pinning in an electrochemical environment. Nevertheless, other structural defects such as nanometric precipitates of metallic Ga cannot be excluded as potential source of the observed bandgap states, taking into account the relatively high weight of the oxygen containing species in the interfacial layer that might significantly alter its cristalline state.

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Acknowledgement Financial support of the Alexander von Humboldt Foundation and from CNCSIS (Project No. 1190/2005) is gratefully acknowledged. References 1. J. van Laar, A. Huijser, T.L. van Rooy, J.Vac. Sci. Technol. 14 (1977) 894. 2. T.U. Kampen, L. Koenders, K. Smit, M. Rückschloss, W. Mönch, Surf. Sci. 242 (1991) 314. 3. D. J. Chadi, J. Vac. Sci. Technol. 15 (1978) 631. 4. W. Mönch, Appl. Surf. Sci. 22/23 (1985) 705. 5. F. Bartels, L. Surkamp, H.J. Clemens, W. Mönch, J. Vac. Sci. Technol. B1 (1983) 756. 6. W. Mönch, Surf. Sci.132 (1983) 92. 7. W. Mönch, Semiconductor Surfaces and Interfaces, Springer-Verlag, Berlin, 1995. 8. V. Lazarescu, R. Scurtu, M.F. Lazarescu, E. Santos, H. Jones, W. Schmickler, Electrochim. Acta, in press. 9. V. Lazarescu, M.F. Lazarescu, E. Santos, W. Schmickler, Electrochim. Acta 49 (2004) 4231. 10. R. A. Batchelor, A. Hamnett, in Modern Aspects of Electrochemistry 22 (1992) 265. 11. K.W. Frese, S.R. Morrison, J. Electrochem. Soc. 126 (1979) 1235. 12. W.E. Spicer, P.W. Chye, P.R. Skeath, C.Y. Su, I. Lindau, J. Vac. Sci. Technol. 16 (1979) 1422. 13. D.G.A.Bruggeman, Ann.Phys. (Leipzig), 24 (1935) 636. 14. D.E.Aspnes, A.A.Studna, Phys.Rev.B, 27 (1983) 985. 15. D.E.Aspnes, B.Schwartz, A.A.Studna, L.Derick, L.A.Koszi, J.of Applied Physics, 48 (1977) 3510. 16. R.Kofman, P.Chheyssac, J.Richard, Phys.Rev.B, 16 (1977) 5216. 17. G.N.Greaves, E.A.Davis, J.Bordas, J.Philod.Mag.,34 (1976) 263. 18. B.H. Erne, M. Stchakovsky, F. Ozanam, J.N. Chazalviel, J. Chem. Soc. 145 (1998) 447. 19. M. Pourbaix, Atlas d’Equilibres Electrochimiques, Gauthiers-Villards, Paris, 1963. 20. K.W. Frese, Jr., S.R. Morrison, Appl. Surf. Sci. 8 (1981) 266. 21. S. Lingier, W.P. Gomes, Ber. Bunges. Phys. Chem. 95 (1991) 170. 22. C.C. Chang, P.H. Citrin, B. Schwartz, J. Vac. Sci. Technol. 14 (1977) 943. 23. H. Habazaki, P. Skeldon, D. Ghidaoui, S.B. Lyon, K. Shimizu, G.E. Thomson, G.C. Wood, J. Phys.D:Appl.Phys.29(1996) 2545. 24. D. Ghidaoui, S.B. Lyon, G.E. Thomson, J. Walton, Corr. Sci. 44 (2002) 501. 25. D.M. Hofmann, B.K. Meyer, F. Lohse, J.-M. Spaeth, Phys. Rev. Lett. 53 (1984) 1187. 26. W.E. Spicer, Z. Liliental-Weber, E. Weber, N. Newman, T. Kendelwicz, R. Cao, C. McCants, P. Mahowald, K. Miyano, I. Lindau, J. Vac. Sci. Technol. B6 (1988) 1245. 27. R.Z. Bachrach, R.S. Bauer, P. Chiaradia, G.V. Hansson, J. Vac. Sci. Technol. 19 (1981) 335. 28. N. Newman, W.E. Spicer, E.R. Weber, J. Vac. Sci. Technol. B5 (1987) 1020.

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Composition and Growth of Thermal and Anodic Oxides on InAlP S. Kleberb, M. J. Grahama, S. Moisaa, G. I. Sproulea, X. Wua , D. Landheera , A. J. SpringThorpea , P.J. Barriosa , P. Schmukib a

Institute for Microstructural Sciences, National Research Council of Canada, Ottawa, Canada K1A 0R6 E-mail : [email protected]; E-mail : [email protected]; E-mail : [email protected] E-mail : [email protected]; E-mail : [email protected]; E-mail :[email protected]; E-mail : [email protected] b University of Erlangen-Nuremberg, Martensstrasse 7 D-91058 Erlangen, Germany E-mail : [email protected]; E-mail: [email protected]

Abstract - Producing insulating layers on III-V semiconductors is crucial for a number of device applications. Al-containing thermal oxides on AlGaAs, InAlAs, and more recently on InAlP have been found to possess good insulating characteristics. This paper presents data on insulating oxides formed on InAlP layers by thermal oxidation at 500oC in moist nitrogen (95oC) or by anodization in sodium tungstate solution. The oxides (20-300nm thick) have been characterized by Auger electron spectroscopy, Xray photoelectron spectroscopy, Rutherford backscattering spectroscopy, and transmission electron microscopy. Oxides are amorphous and appear to be a mixture of indium phosphates and aluminium oxide. Indium hydroxide is present at the outer surface of anodically grown oxide. Thermal oxide growth follows parabolic kinetics, whereas the thickness of anodic oxides increases linearly with potential up to ~120V. Electrical measurements indicate that the thermal oxides have good electrical properties making the films potentially useful for some device applications. Keywords: III-V semiconductors, InAlP, thermal oxidation, anodic oxidation, surfaceanalytical techniques, Auger, XPS, TEM.

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1. Introduction Passivation layers on III-V semiconductors can be produced by deposition of silicon nitride or oxide or created by a variety of oxidation processes including thermal and anodic oxidation. Thermal oxidation data for AlGaAs and InAlAs in GaAs- and InP-based heterostructure devices have been reported, e.g. [1-4], and the Al-containing oxides have often been found to possess good insulating characteristics. Recently, Al-containing thermal oxides on InAlP have been shown to be even more promising [5-8]. This paper presents data on insulating oxides formed on InAlP layers by thermal oxidation at 500°C in moist nitrogen (95°C), or by anodization in sodium tungstate solution. The composition and nature of the oxide have been determined by Auger electron spectroscopy (AES), X-ray photoelectron spectroscopy (XPS), Rutherford backscattering spectroscopy (RBS) and transmission electron microscopy (TEM). Currentvoltage (IV) measurements on Al-gated capacitors were performed. 2. Experimental InAlP layers, approximately lattice matched to GaAs, were grown by molecular beam epitaxy (MBE). For thermal oxidation experiments, two ~ 1mm thick InAlP layers (In0.525Al0.475P and In0.494Al0.506P) were used. The latter was utilized for electrochemical experiments. For electrical measurements of thermal oxides a 60 nm-thick layer of undoped In0.485Al0.515P was grown by MBE on a doped GaAs layer. The GaAs cap layer was removed chemically [7] before complete oxidation of the InAlP layer and subsequent electrical measurements. Thermal oxidations were performed in a Lindberg/Blue furnace at 500°C in moist nitrogen (N2 bubbled through H2O at 95°C with gas transfer through heated tubes to the oxidation furnace). For electrochemical experiments, contact to the InAlP was established by scratching InGa eutectic on the edge of the frontside of the sample. The electrochemical set-up consisted of a conventional threeelectrode configuration with a Pt counter electrode and a Haber-Luggin capillary with a Ag/AgCl electrode as a reference electrode. Electrochemical experiments were carried out using an EG&G 173 and Jaissle IMP 88PC-200V potentiostats. 0.1M Na2WO4 (pH 8.2) solution was prepared from analytical grade chemicals and de-ionized water. Potential step experiments were carried out from 1V to 120V (Ag/AgCl) from times of 5 minutes to 2 h. After oxidation samples were analyzed by AES (PHI 650 system); XPS (PHI 5500 with a monochromated AlKa source); TEM (JEOL 2100F) operating at 200keV. Metal-insulator-semiconductor (MIS) structures were formed by evaporating Al dots (area 5x10-4cm2) through a shadow mask followed by annealing in forming gas for 5min at 450ºC. IV measurements were performed on these capacitors using a probe station with a HP 5140 pA meter / DC voltage source.

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3. Results and Discussion 3.1. Oxide growth and oxide composition Thermal oxidations were performed at 500ºC in moist nitrogen for periods of time ranging from 6 minutes to 4 h. The oxidation kinetics are found to be parabolic (after a brief incubation period), and the ~1 m thick InAlP layer with the higher Al content oxidizes slightly faster [6]. The thickness of anodic oxides formed in 0.1M Na2WO4 (pH 8.2) solution increases linearly with potential up to ~ 120V as shown in Fig.1. Oxides ranging a thickness from ~20nm to ~300nm have been characterized by Auger, XPS and TEM. A high-angle annular dark field (HAADF), (Z contrast) STEM micrograph of the thermal oxide formed after 1h on In0.494Al0.506P is shown in Fig. 2(a). The oxide is uniform in thickness, and the “bright” particles near the oxide / InAlP interface have been attributed to unoxidized indium [5-7]. The bulk of the oxide is amorphous as deduced from both electron diffraction and x-ray diffraction measurements. A TEM bright field image of the anodic oxide formed on In0.494Al0.506P at 20V for 30 minutes in 0.1M Na2WO4 (pH 8.2) solution is shown in Fig. 2(b). The oxide is uniform in thickness and no particles could be seen near the oxide / InAlP interface by either bright field TEM or dark field (HAADF) STEM imaging.

120

1 min 30 min 120 min Linear fit

Oxide thickness (nm)

100 80 60 40 20 0 0

20

40

60

80

100

120

Potential vs. Ag/AgCl (V)

Figure 1: Growth of anodic oxides on InAlP in 0.1M Na2WO4 (pH 8.2) solution for different oxidation times. Oxide thickness determined from Auger depth profiles calibrated from TEM data. (Auger sputter rate 1.15nm/min at 5 nA).

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(a)

(b)

Figure 2: (a) STEM HAADF micrograph of thermal oxide formed on In0.494Al0.506P after 1 h at 500°C in moist nitrogen. (b) TEM bright field image of the oxide formed on In0.494Al0.506P at 20V for 30min in 0.1 M Na2WO4 (pH 8.2) solution. Cross-sections prepared by ion milling.

Auger profiles of thermal and anodic oxides are shown in Fig. 3(a) and (b). The Auger sensitivity factors for P, Al and In in the oxide are found to be quite different from those in the substrate. Therefore, the sensitivity factors in the oxide have been based on the oxide composition as determined by RBS. RBS analysis indicated that the In, P and Al ratios in the oxide are the same as in the substrate and the oxygen concentration is ~67 at.%. Therefore, as seen in Fig. 3(a), P is the major component in the oxide, the Al level is fairly constant and In appears to be depleted in the outer part of the oxide and increases at the interface. This increase in In at the substrate interface is better seen in the Auger profile obtained by Physical Electronics on a PHI 680 system [7]. An Auger profile of an anodic oxide formed at 20V for 30 minutes in 0.1M Na2WO4 (pH 8.2) solution is shown in Fig. 3(b). In is enriched in the outer part of the film. P is not present in the outer third of the film, but is a major component in the inner two thirds of the film. W from the electrolyte is always found to be incorporated into the outer oxide of anodic films. An indication of the chemical composition of thermal and anodic oxides can be obtained from XPS measurements. O1s, P2p and In3d XPS data were collected. Curve fitting was carried out using binding energies of Hollinger et al. [9,10], Faur et al. [11], and an In2O3 standard for the various relevant species. Curvefitting of the O1s signal from a thermal oxide formed on In0.494Al0.506P after 6 minutes in moist nitrogen [Fig. 4(a)] suggests the possible presence of several species. These include small components of the oxide species In2O3 and P2O5 and the phosphate and polyphosphate species InPO4, In(PO3)3 and In(POy)x [12]. The peak positions for P2O5 and In(POy)x are practically coincidental and thus it is questionable whether both species are present in the layer. Similarly, the peaks for InPO4 and In(PO3)3 coincide. The yield corresponding to the latter two species is significantly greater than the yields from In2O3, P2O5 and In(POy)x indicating that one, or both of these species dominate.

Atomic Concentration (%)

Composition and Growth of Thermal and Anodic Oxides on InAlP 100 90 80 70 60 50 40 30 20 10 0

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O

P Al In 0

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Sputter time (min)

Atomic concentration (%)

(a) 100 90 O 80 70 60 50 40 P(oxide) 30 In 20 Al 10 W 0 0 5 10

P(substrate)

15

20

25

Sputter time (min) (b) Figure 3: Auger electron spectroscopy (AES) profiles of oxide on In0.494Al0.506P. (a) Thermal oxide formed in 1h at 500°C in moist nitrogen [STEM micrograph of the oxide shown in Fig. 2(a)]. (b) Anodic oxide formed at 20V for 30min in 0.1M Na2WO4 (pH 8.2) solution [TEM micrograph of the oxide shown in Fig 2(b)]. The Auger peak shape for P in the oxide is significantly different from that in the bulk, and so the P signal in (b) has been fitted to two peaks. The profiles have been truncated after the oxide /InAlP interface. Sputtering was by 1keV argon ions.

Curve-fitting for the P2p signal is consistent with the data for the O1s peak [6]. In(PO3)3 and InPO4 dominate over P2O5 or In(POy)x. Examination of the Al2p XPS peak confirms the presence of Al2O3 (and

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InPO4 or In(PO3)3

P2O5 or In(POy)x

535

534

In2O3 533

532

531

530

529

528

Binding Energy (eV)

(a) In2O3

In(OH)3 H2Oad 535

534

533

OHad

532

531

530

529

528

Binding Energy (eV)

(b)

In2O3

In(OH)3 OHad H2Oad

535

534

533

532

531

530

529

528

Binding Energy (eV)

(c) Figure 4: Curve-fitted O1s spectra from oxide formed on In0.494Al0.506P. (a). Thermal oxide formed after 6min at 500°C in moist nitrogen. Take-off angle at 45°. (b) and (c). Anodic oxide formed at 20V for 30min in 0.1M Na2WO4 (pH 8.2) solution. Take-off angle at 45° in (b) and 10° in (c).

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not AlPO4) in the oxide. The main components of the thermal oxide formed on InAlP from XPS and Auger data are In(PO3)3, InPO4 and Al2O3. Only small amounts of In2O3 and P2O5 are present. O1s spectra of the anodic oxide formed at 20V for 30 minutes in 0.1M Na2WO4 (pH 8.2) solution are shown in Fig. 4(b) and (c). The XPS take-off angle is 45º in (b), as in (a), and 10º in (c). As found earlier by Auger, In is enriched in the outer part of the film and no P is present in the outer third of the film. The oxygen peaks in the XPS spectra of the anodic film are therefore not attributed to phosphate or other phosphorus-containing species, but to oxide and hydroxide. As indicated in Fig. 4(b) the major deconvoluted peaks are indium oxide and indium hydroxide. The two smaller peaks at higher binding energy may be due to adsorbed hydroxyl and adsorbed water as was concluded in a recent study by Grosvenor et al. [13] for the oxidation of iron in water vapour. The approximate ratios of the amount of oxide: hydroxide: adsorbed OH: adsorbed H2O is 58 : 32 : 6 : 4. Spectrum (c) is obtained with a take-off angle at 10º to the sample surface. The spectral area ratio changes to 44 : 35 : 15 : 5, giving relatively larger spectral ratios from hydroxide, adsorbed hydroxyl and water as would be expected from this more surface-sensitive analysis. 3.2. Electrical measurements Electrical measurements performed on metal-insulator-semiconductor structures indicate that the thermal oxides have good electrical properties [6,7]. A brief oxidation in oxygen at 500ºC following oxidation in moist nitrogen oxidizes the residual indium particles present at the oxide/substrate interface [Fig. 2(a)]. This reduces the current density at a field of 1MVcm1 (8.6V gate potential) by two orders of magnitude to 1.7x10–10Acm-2. The breakdown voltage is increased to 44V corresponding to a breakdown field of 5.1MVcm-1, making the oxide films potentially useful for some device applications. 4. Concluding Remarks Thermal oxidation of InAlP layers at 500ºC in moist nitrogen produces amorphous, insulating oxide which is a mixture of indium phosphates and aluminium oxide. The oxidation kinetics are parabolic. The thickness of anodic oxides increases linearly with potential up to ~120V. Indium hydroxide and adsorbed hydroxyl ions and adsorbed water are present at the outer surface of anodically grown oxide. Electrical measurements on oxidized capacitors indicate that thermal oxides have low leakage currents and high breakdown fields, making them potentially useful for some device applications.

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F.A. Kish, S.J. Caracci, N. Holonyak, Jr., K.C. Hsieh, J.E. Baker, S.A. Maranowski, A.R. Sugg, J.M. Dallesasse, R.M. Fletcher, C.P. Huo, T.D. Osentowski and M.G. Crawford, J. Electron. Mat. 21 (1992) 1133. U.K. Mishra, P. Parikh, P. Chavarkar, J. Yen, J. Champlain, B. Thibeault, H. Reese, S.S. Shi, E. Hu, L. Zhu and J Speck, IEDM’97, 21.1.1. S.J. Caracci, M.R. Kramas, N. Holonyak, Jr., M.J. Ludowise and A. Fischer-Colbrie, J. Appl. Phys. 75 (1994) 2706. R.J. Hussey, R. Driad, G.I. Sproule, S. Moisa, J.W. Fraser, Z.R. Wasilewski, J.P. McCaffrey, D. Landheer and M.J. Graham, J. Electrochem. Soc., 149 (2002) G581. P.J. Barrios, D.C. Hall, U. Chowdhury, R.D. Dupuis, J.B. Jasinski, Z. Liliental-Weber, T.H. Kosel and G.L. Snider, Abstract, 43rd Electronic Materials Conference, Notre Dame, Indiana, June 27-29, 2001. M.J. Graham, S. Moisa, G.I. Sproule, X. Wu, J.W. Fraser, P.J. Barrios and D. Landheer, Proc. 5th Int. Conf. On Microscopy of Oxidation, Limerick, Ireland, Aug. 2002, Materials at High Temperatures (2003) 31. M.J. Graham, S. Moisa, G.I. Sproule, X. Wu, J.W. Fraser, P.J. Barrios, A.J. SpringThope and D. Landheer, Proc. Int. Symp., “Corrosion Science in the 21ST Century”, UMIST, UK, July 2003. Y. Cao, J. Zhang, X. Li, T.H. Kosel, P. Fay, D.C. Hall, X.B. Zhang, R.D. Dupuis, J.B. Jasinski and Z. Liliental-Weber, Appl. Phys. Lett. 86 (2005) 062105. G. Hollinger, E. Bergignat, J. Joseph and Y. Robach, J. Vac. Sci. Tech. A3 (1985) 2082. G. Hollinger, J. Joseph, Y. Robach, E. Bergignat, B. Commere, P. Viktorovitch and M. Froment, J. Vac. Sci. Technol. B5 (1987) 1108. M. Faur, D.T. Jayne and M. Goradia, Surface and Interface Analysis 15 (1990) 641. A. Pakes, P. Skeldon, G.E. Thompson, S. Moisa, G.I. Sproule and M.J. Graham, Corros. Sci. 44 (2002) 2161. A.P. Grosvenor, B.A. Kobe and N.S. McIntyre, Surf. Sci. 572 (2004) 217.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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High-k gate oxide formed by anodic oxidation for organic field effect transistors M. Erouela,b, A. Gagnairea, A.L. Demana, J. Tardya, N. Jaffrezic-Renaultb, Z. Sassib, J.C. Bureaub, M.A. Maarefc a

LEOM UMR CNRS 5512, Ecole Centrale de Lyon - 69134 Ecully CEDEX FRANCE Email : [email protected] b CEGELY UMR CNRS 5005, Ecole Centrale de Lyon - 69134 Ecully cedex France and INSA-LYON 69691Villeurbanne CEDEX France c IPEST LaMarsa Tunis TUNISIE

Abstract - This paper reports on a comparative investigation of anodically grown HfO2 and Ta2O5 films as gate oxides for pentacene field effect transistors (OFETs). Ta and Hf films were deposited by e-beam evaporation onto highly doped Si wafers and anodization is carried out in an AGW (acid-glycol-water) solution. After gate oxide growth, OFETs were completed by vacuum deposition of pentacene and gold drain and source contact through a shadow mask. HfO2 and Ta2O5 exhibit a high dielectric constant (er ~21-25) which leads to devices showing a good field effect mobility at low voltage. Devices with HfO2 show improved characteristics with a higher stability and a lower leakage current at low drain bias as well as a high mobility (µ=2.2 10-2 cm2/V.s). Keywords: anodic gate oxide, Ta2O5, HfO2, OFETs, pentacene

1. Introduction Organic field effect transistors (OFETs) are now widely recognized for their potential applications in all fields of so-called "plastic electronics" (smart labels, chemical and bio-sensors, flexible display driving circuits,…). Nevertheless, because of the low carrier mobility in organic semiconductors, OFETs usually operate at relatively high voltage (say, 20-100V) when conventional SiO2 or polymers are used as gate dielectrics. These voltages are not appropriate for practical use. As the field effect mobility in polycrystalline organic

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semiconductors is largely depending on the density of charge in the channel [1], it is important to get a as high as possible gate capacitance. A way is to use high-k gate oxide dielectrics. Al2O3 [2,3], Ta2O5 [4,5] or TiO2 [6] were proposed as gate oxide in OFETs. It is commonly observed that the operating voltage was noticeably reduced but also that these dielectrics remain somewhat leaky. The observation of gate-source leakage current at low drain voltage is also one of the main drawback of low band gap high-k oxides. Furthermore, the high dielectric constant favors polarization effects and subsequent charge localization. Improvement can be brought out by chemical surface treatments or deposition of a thin polymer film on the surface of the high-k oxide [7,8]. HfO2 has been widely investigated as a potential high-k oxide in replacement of SiO2 in future silicon microelectronics but, to our knowledge, was not used in OFETs so far. The dielectric constant of HfO2 is nominally close to that of Ta2O5 (er~25 for high temperature processed materials) but the energy band gap is higher (Eg=5.7eV for HfO2 vs Eg=4.5eV for Ta2O5). This paper aims at comparing the performances of pentacene FETs with either Ta2O5 or HfO2 gate dielectrics both grown by anodic oxidation of Ta and Hf evaporated films. Anodization was reported to be a particularly well convenient procedure to obtain high quality oxide films at low temperature compatible with plastic electronic requirements [3,9] The anodization procedure and physical properties of obtained films are first reported. Then the electrical characteristics of similar OFETs, besides the use of HfO2 or Ta2O5 are compared. Superior performances of OFETs based on HfO2 are observed, i.e. higher mobility and better stability. 2. Experimental procedure Ta and Hf films were deposited by e-beam evaporation onto highly doped Si-p++ substrates. The evaporation rate and the final thickness were respectively 0.5 Å/s and 100 nm for Ta and 2.2 Å/s and 250 nm for Hf. Anodization was carried out in a so-called AGW solution, i.e. a mixture of acid, glycol and water [10]. A solution of tartaric acid (0.1M) buffered with NH4OH to a pH=6 was mixed with ethylene glycol in a 1:5 ratio. The solution was thoroughly stirred and a flow of nitrogen was bubbling in before and during the anodization process. Ammonium hydroxide increases the density of charge carriers in the solution which in turn enables a sufficient electrical conduction of the electrolyte. Ethylene glycol increases the viscosity of the solution. The decrease of ion mobility improved lateral homogeneity of the oxidation process. The anodization process was carried out at constant current density up to an anodization voltage VA in the 25-100V range. The current density is maintained to 0.1 mA/cm2 and 0.2 mA/cm2 for Ta and Hf respectively. Care was taken not to oxide the entire evaporated metal film. Anodic films were characterized by spectroscopic ellipsometry to determine the thickness of the oxide, the refractive index and the absorption coefficient. Capacitors were processed by evaporation of top Al electrodes of various area and the dielectric constant er

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was deduced from impedance measurements carried out at 1 KHz and 10mV modulation voltage with a HP 4284A impedance analyzer. OFETs were processed by vacuum evaporation of 80 nm thick pentacene on the anodic oxide at a rate of 0.02 nm/s. The process is completed by the deposition of gold source and drain electrodes on the pentacene. The devices were characterized in ambient atmosphere with two Keithley 2400 sourcemeter driven under Labview® environment. 3. Results and discussion 3.1. Ta2O5 gate oxide

Anodization Voltage (V)

Figure 1 shows the potential evolution as a function of time during Ta2O5 growth. The growth rate is about 0.11 nm/s. 133.4 nm

100 80 60

Figure 1: Voltage evolution vs time in the electrolyte cell during the growth of Ta2O5. The thickness obtained for VA=50, 75 and 100V are reported

111 nm 82.6 nm

40 Ta anodic oxidation Tartaric acid (10-1 M) + ethylene glycol constant current density I=0,1 mA/cm2

20 0

0

2

4

6

8

10 12 14 16

Time (min)

Table 1 below gathers optical (thickness d, refractive index n) and electrical (gate capacitance per unit area Cox, dielectric constant er) properties for different anodization voltages VA as deduced from ellipsometry and capacitance spectroscopy respectively. Table 1: Thickness, refractive index, gate capacitance and dielectric constant of Ta2O5 films for different values of anodization voltage VA. VA (Volts)

d (nm)

n (at l= 700 nm)

Cox (nF/cm2)

er

50 75 100

82.6 111 133.4

2.38 2.25

243 173 134

22.7 21.5 20.2

The high refractive index is representative of dense and closely packed films. They are quite similar to those obtained by laser ablation at higher temperature, 250-400°C [11]. The dielectric constant is high and, similarly to the index n, decreases with thickness. This is probably due to some porosity in the growing

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film as the oxide growth proceeds. A way to minimize this effect would be to make the oxidation go on at constant voltage V=VA, once VA is reached, down to a stable minimum current density. Work is in progress on that point. The anodization ratio k=d/VA varies from 1.3 to 1.65 nm/V as the thickness increases. As k-1 can be assessed as the breakdown field, this result bears out the assumption of a more pronounced porosity as VA increases. Figure 2 reports the output (2a) and the transfer characteristics (2b). The output characteristics was recorded with ramping the drain voltage from 0 to –1.5 V and back from –1.5 V to 0 to evidence hysteresis effects. The transfer characteristics in the saturation regime is plotted as: I D ,sat = f (VG ) to determine the threshold voltage VT and the field effect mobility µ from the slope of the curve at high VG according to the equation: I D ,sat = W µ Cox (VG - VT )2 2L where W is the channel width, L is the channel length and Cox is the gate capacitance per unit area. We observe a strong hysteresis in the ID-VD characteristics which could be attributed to charge carrier trapping in the channel due to polarization effect in the oxide. The device reach the saturation regime at very low voltage, ~1 volt due to high gate capacitance. However, because of the thin oxide, a gate to source leakage current is observed at low VD (reverse drain current below – 0.2V). Thicker oxides would lead to smaller leakage. From figure 3b, we can deduce the threshold voltage VT=-0.75V and the mobility µ @ 0.8 10-2 cm2/V.s. A surface treatment (eg OTS) should improve the mobility. -0,015

0,12

Vg=-1.5V

0,10 1/2

Vg=-1V 0,000

Ta2O5 Vg=0 and VA= 25V Pentacene 80nm at r=0.2 Å/s L=100µm W=1000µm

0,005

a

0,010 0,2

0,0

-0.5V

-0,2 -0,4 -0,6 -0,8 -1,0 -1,2 -1,4 -1,6

Vd(V)

1/2

-0,005

IIdI (µA)

Id(µA)

-0,010

0,08 0,06

Vd=-1.5V -3

2

µ = 7.85x10 cm /V.s 0,04 0,02 0,00 0,2

b 0,0

-0,2 -0,4 -0,6 -0,8 -1,0 -1,2 -1,4 -1,6

Vg(V)

Figure 2: Output (3a) and transfer characteristics (3b) for a pentacene OFET with a gate oxide made of Ta2O5 grown by anodic oxidation at VA=25V (gate capacitance: 583 nF/cm2). Open symbol=ramp up, close symbol=ramp down.

3.2. HfO2 gate oxide Figure 3 reports the evolution of the anodization voltage as a function of time for a constant current density (0.2 mA/cm2). The thickness of HfO2 films for VA= 25 and 50V is shown.

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The anodization ratio is found to be 1.8 nm/V between 25 and 50 V inaccordance with results in the literature with H3PO4 and H2SO4 [12]. In the same reference, a decreased anodization ratio is found above 50V attributed to crystallization with oxygen evolution and accumulation of stress in the films (sometimes up to breakdown). Cross section high resolution transmission electron microscopy (X-HRTEM) could reveal the crystallites formation. Table 2 below gives the optical and electrical properties of HfO2 films grown at VA=25 and 50V. Anodization Voltage (V)

120 HfO2 in 0.1M tartaric acid + ethylene glycol I=0,2 mA/cm2

100

Figure 3: evolution of the anodization voltage vs time of oxidation in a 0.1 M tartaric acid - ethylene glycol solution (1:5).

80 60

85.6 nm

40 40.1 nm

20 0

0

2

4

6 8 Time (min)

10

12

Table 2: Thickness, refractive index, gate capacitance and dielectric constant of HfO2 films for two values of anodization voltage VA (25 and 50V). VA (Volts)

d (nm)

n (at l= 700 nm)

Cox (nF/cm2)

er

25 50

40.1 85.6

1.99 1.92

469.4 217.2

21.3 22.2

0,18 -0,03

Vg=-1.5V

0,16 0,14

Id(µA)

1/2

IIdI (µA )

-0,02 -0,01

1/2

Vg=-1V

0,00

Vg=0 and -0.5V 0,01 0,2

a 0,0

-0,2 -0,4 -0,6 -0,8 -1,0 -1,2 -1,4 -1,6

Vd(V)

0,12 0,10 0,08

Vd = -1V

0,06

µ = 0.022 cm /V.s

2

0,04 0,02 0,00 0,2

b 0,0

-0,2 -0,4 -0,6 -0,8 -1,0 -1,2 -1,4 -1,6

Vg(V)

Figure 4: Output (4a) and transfer characteristics (4b) for a pentacene OFET with a gate oxide made of HfO2 grown by anodic oxidation at VA=25V (gate capacitance: 469 nF/cm2). Open symbol=ramp up, close symbol=ramp down.

The refractive index is in the range 1.99-1.92 and seems to decrease with the thickness as for Ta2O5, revealing some increasing porosity. This is however not

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corroborated by the C-V measurements which gives higher er at 50V. This point is still unclear and to be confirmed. The refractive index was previously found to be strongly dependent on the nature of the electrolyte [12]. Lower n (n=1.85 an 1.82 for HNO3 1M) and higher n (n= 2.07 in H3PO4 and H2SO4 1M) were reported in this reference. Figure (4a) shows that OFET with HfO2 gate dielectric operates at very low voltage. We also observe a weak hysteresis and smaller leakage current than with Ta2O5. From the transfer characteristics (fig. 4b), we deduce the threshold voltage VT=-0.75V and the mobility µ @ 2.2 10-2 cm2/V.s. We obeserve that the mobility is improved compared to devices with Ta2O5. 4. Conclusion This paper reported results on the anodic growth of high-k oxides films, Ta2O5 and HfO2, in view to realize gate dielectric for high performances organic field effect transistors (OFETs). The anodization process was carried out in a buffered solution of tartaric acid (0.1 M) mixed with ethylene glycol. Optical and electrical properties of films grown at various anodization voltage were reported which show that high quality films are obtained at room temperature. Pentacene OFETs were processed and characterized. Devices with HfO2 revealed better performance than those with Ta2O5 with an improved mobility, a better stability and low leakage. Work is in progress to improve the anodization process and develop a surface treatment of HfO2 for more efficient devices. References [1] C.D. Dimitrakopoulos, S. Purushothaman, J. Kymisssis, A. Callegari and J.M. Show, Science 283 (1999) 822 [2] J. Lee, J.H. Kim and S. Im, App. Phys. Lett. 83 (2003) 2689 [3] L.A. Majewski, M. Grell, S.D. Ogier and J. Veres, Organic Electronics 4 (2003) 27 [4] C. Bartic, H. jansen, A. Campitelli and S. Borghs, Organic Electronics 3 (2002) 65 [5] Y. Iino, Y. Inoue, Y. Fujisaki, H. Sato, M. Kawakita, S. Tokito and H. Kikuchi, Jap. J. App. Phys. 42, part 1, (203) 299 [6] G. Wang, D. Moses, A.J. Heeger, H-M Zhang, M. Narasimhan, R.E. Demaray, J. Appl. Phys. 95 (2004) 316 [7] L.A. Majewski, R. Schroeder, M. Grell, P.A. Glarvey, M.L. Turner, J. Appl. Phys. 96 (2004) 5781 [8] A.L. Deman and J. Tardy, Organic Electronics 6 (2005) 78 [9] L.A. Majewski, R. Schroeder, M. Grell, Adv. Mater. 17 (2005) 192 [10] E. Jalaguier, PhD Thesis, Ecole Centrale de Lyon (1988) [11] S. Boughaba, M.U. Islam, Mat. Res. Soc. Symp., Vol 617 (2000) J3.7.1 [12] M.J. Esplandu, E.M. Patriti, V.A. Macagno, Electrochimica Acta 42 (1997) 1315

Section C Electronic Properties of Passive Films

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

279

Photoelectrochemical analysis of the passive film formed on Fe-Cr-Ni alloys in pH 8.5 buffer solution HeeJin Jang and HyukSang Kwon Dept. of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1, Guseong-dong, Yuseong-gu, Daejeon, 305-701, Rep. of Korea [email protected]

Abstract - The structure and composition of passive films formed on Fe-20Cr-(10, 15)Ni alloys in pH 8.5 buffer solution were examined with an emphasis on the influence of Ni on the structure and composition of the film by analyzing the semiconducting properties of the film in terms of the photoelectrochemistry. The photocurrent spectrum for the passive films on Fe-20Cr-10Ni alloy was almost same in shape as those for the passive film on Fe and Fe-20Cr with two spectral components responsible, respectively, for the d-d and p-d electron transitions that occur in Cr-substituted g-Fe2O3. The photocurrent spectra for the passive film formed on Fe20Cr-15Ni comprised a spectral component originated from NiO in addition to those from the Cr-substituted g-Fe2O3, demonstrating that the passive film on Fe-20Cr-15Ni consists of Cr-substituted g-Fe2O3 mixed with NiO. It was suggested from the high density of donor or oxygen vacancy (~1020 cm-3) and the dependence of the deep donor (or Cr6+) density on Ni content determined from the Mott-Schottky plots for the passive films on Fe-20Cr-xNi (x = 0, 10, 15) alloys that the passive films consist of a nanocrystalline (Cr, Ni)-substituted g-Fe2O3 mixed with or without NiO depending on the nickel content. Keywords: Stainless steels; Passive film; Photocurrent; Mott-Schottky analysis; Semiconducting property

1. Introduction Many analytical studies have agreed that passive films on stainless steels consist of Cr-enriched (Fe, Cr) oxide/hydroxide, although there is still some

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controversy as to the detailed structure and composition of the passive film. The extreme complexity of the metal/passive film/electrolyte system of stainless steels makes the clarification of the passive film difficult. Recently, photoelectrochemistry based on the semiconducting properties of passive films has proved to be a powerful technique for in-situ analysis of passive films on metals and alloys. The structure and composition of a passive film can be determined by comparing the semiconducting properties (photocurrent spectrum, bandgap energy, flat band potential) for the passive film with those for well characterized oxides as well as ex-situ data of the film. Through photoelectrochemical studies and also surface analytical techniques, many models have been suggested on the structure and composition of the passive film of Fe-Cr-Ni alloys [1-6]; single layered structure of Cr2O3-Fe2O3 solid solution [1-2], a Cr-substituted iron oxide [3], duplex layered structure composed of inner Cr2O3-Fe2O3 mixture and outer Cr(OH)3-Ni(OH)3 mixture [4], and inner Fe2+[Cr3+Cr3+]O42-or Cr2O3 and outer Fe2O3 [5]. Whereas Cr is reported to be enriched in passive films on austenitic stainless alloys, Ni is known as depleted or remained in trace amounts in passive films on the alloys [6]. Though the previous research results mentioned above, it is not yet clear how Ni affects semiconducting properties for a passive film on Fe-Cr-Ni alloys. Research objective of the present work is to examine the structure and composition of the passive films on Fe-20Cr-xNi (x = 10, 15) alloys by analyzing the semiconducting properties of the films using photoelectrochemical technique with the focus on the role of Ni in the passive film of the stainless steel. 2. Experimental High purity Fe-20Cr-xNi (x = 0, 10, 15 wt.%) alloys prepared by a vacuum-arcmelting and pure Ni were used as working electrodes. The working electrodes were mounted in an epoxy resin with an exposed area of 0.2 cm2. A conventional three-electrode cell of 1 L-multi neck flask with a quartz window as a photon inlet was used for the photocurrent measurements. The cell was equipped with a platinum counter electrode and a saturated calomel reference electrode (SCE). All the experiments were carried out in deaerated pH 8.5 buffer solution at room temperature. The working electrode was cathodically cleaned by polarization to –1.0 VSCE for 5 min before potentiostatically forming passive film. Passive films were formed on the alloys by stepping applied potential from –1.0 VSCE to a film formation potential (Ef) in passive region. Photocurrent spectra for the passive film were measured at potentials in passive region under continuous illumination without lock-in amplifier and chopper. A 300 W Xenon arc lamp was used as a light source, and a monochromatic light with wavelength of 200 to 800 nm was provided by a scanning digital monochromator controlled by a stepping motor at

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281

a scan rate of 1 nm/s. Details in the photocurrent measuring system were described elsewhere [7, 8]. For Mott-Schottky analysis, the specimen was passivated at an Ef for 2 h just prior to measuring capacitance with sweeping the applied potential at a rate of 1 mV/s. The excitation voltage was 10 mV (peak-to-peak) and the frequency was 1 kHz. 2.0

3. Experimental results and discussion 3.1. Polarization response

Potential / VSCE

1.5

Ni Fe-20Cr-10Ni Fe-20Cr-15Ni

1.0 0.5 0.0 -0.5 -1.0

Fig. 1 shows potentiodynamic polarization 1E-14 1E-11 1E-8 1E-5 0.01 curves for Fe-20Cr-xNi (x = 10, 15) and Current density / A cm Ni, respectively, which were measured in pH 8.5 buffer solution at a potential scan Fig. 1. Potentiodynamic polarization rate of 0.5 mV/s. The corrosion potential responses of Ni and Fe-20Cr-xNi (x = was raised with Ni content in the alloys. 10, 15) measured at a rate of 0.5 mV/s The slight increase in the current density in deaerated pH 8.5 buffer solution at room temperature. of polarization curve of Fe-20Cr at potentials above 0.6 VSCE appears due primarily to the transpassive dissolution of Cr3+ in Cr2O3 in passive film to Cr6+ in chromate. On the other hand, the abrupt increase in the current density at potentials above 1.1 VSCE, observed in the polarization curves is associated with the oxygen evolution reaction on passive films. -2

3.2. Photocurrent response Fig. 2 shows photocurrent spectra for the passive films formed on Fe-20Cr [7], Fe-20Cr-10Ni, and Fe-20Cr-15-Ni alloys, and for the NiO film grown thermally on Ni in air at 400 如 [8], respectively, in pH 8.5 buffer solution. Photocurrent of the Fe-Cr and Fe-Cr-Ni alloys (Figs. 2(a)~(c)) began to increase from about 2.5 eV, exhibiting a shoulder at 3.8 eV and then a peak current at 4.5 eV, which are almost same to those for the passive film on Fe [9]. These photocurrent spectra demonstrate that the photocurrent of the passive film on Fe-20Cr-xNi (x = 0, 10, 15) is generated by the same electron transition sources occurring in the passive film on Fe. Thus, the semiconducting properties of the passive film on Fe-20Cr-xNi (x = 0, 10, 15) alloys can be governed by the Fe-oxide in spite of the high Cr concentration. In the previous photoelectrochemical study, E. Cho et al. [7] analyzed photocurrent response for the passive film on Fe-20Cr in comparison with that on Fe [9]. It is suggested that the passive film of Fe-20Cr alloy was composed

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Photocurrent / nA

of Cr-substituted Fe2O3 involving two 1 .5 1 .0 electron excitation processes; the d-d 0 .5 transition from Fe3+ band to Fe2+ band and 0 .0 the p-d transition from O-2p band to Fe-3d 3 2 band in spinel g-Fe2O3. Based on these 1 0 previous studies, the passive film on Fe4 20Cr-10Ni alloy is regarded to be composed 2 of Cr-substituted g-Fe2O3. 0 The photocurrent spectra for the passive film 1 .5 1 .0 on Fe-20Cr-15Ni alloy showed the second 0 .5 current peak at 5.7 eV (Fig. 2(c)). By 0 .0 2 3 4 5 6 comparing the photocurrent spectrum for the P ho to n energy / eV passive film on Fe-20Cr-15Ni alloy at photon energy greater than 5.2 eV with that Fig. 2 Photocurrent spectra for the for the thermal oxide (NiO) of Ni shown in passive films formed on (a) Fe-20Cr Fig. 2(d), the second photocurrent peak [7], (b) Fe-20Cr-10Ni, (c) Fe-20Crobserved in Fe-20Cr-15Ni alloy, evidently, 15Ni at 0.2 VSCE in deaerated pH 8.5 buffer solution, and for (d) thermally originated from NiO that may exist in the grown NiO on Ni in air at 400 °C [8]. passive film of the alloy. Therefore, the passive film on Fe-20Cr-15Ni alloy appears to be composed of the Crsubstituted g-Fe2O3 and NiO. The photocurrent spectra for the passive film on Fe-20Cr-10Ni alloy did not exhibit the second photocurrent peak at 5.7 eV, implying that NiO was not formed in the film of the alloy. The results suggest that the NiO is formed in passive film when the Ni content in Fe-20Cr-xNi exceeds a critical concentration that is between 10 wt.% and 15 wt.%. Below the critical concentration of Ni, Ni remained in the film is considered to be present as Ni2+ ion by substituting Fe ions in the Cr-substituted g-Fe2O3. F e -20 C r E f = 0 .2 V S C E

(a)

F e -2 0C r-1 0 N i E f = 0 .2 V S C E

(b)

F e -2 0C r-1 5 N i E f = 0 .2 V S C E

(c)

Ni oxid ize d in a ir a t 4 0 0 'C fo r 2 h E a p p = 0 .4 V S C E

(d )

3.3. Mott-Schottky analysis of passive film Mott-Schottky analysis for a passive film provides valuable information on insitu semiconducting properties of the film such as type of semiconductor, defect density, and flat band potential. Fig. 3 shows Mott-Schottky plots for the passive films formed on Fe-20Cr, Fe-20Cr-10Ni, and Fe-20Cr-15Ni alloys for 2 h at a passive potential (Fig. 3(a)) and a Cr-transpassive potential (Fig. 3(b)) in pH 8.5 buffer solution. The capacitance was measured at frequency of 1 kHz with scanning the applied potential at 1 mV/s. We chose 1 kHz as an optimal frequency for the Mott-Schottky analysis because the capacitance of the space charge region in the passive film of Fe is almost constant [10]. Evidently, the passive films on the three alloys exhibited n-type semiconductivity as confirmed by the positive slope at potentials above –0.3 VSCE. The Efb for the films was determined to be –0.3 VSCE regardless of alloy composition, and well

Photoelectrochemical analysis of the passive film

283

corresponded to that of the Cr-substituted g-Fe2O3 determined from the photocurrent in Fig. 7(d). This confirms again that the passive films on Fe20Cr-(0, 10, 15)Ni alloys have a base structure of Cr-substituted g-Fe2O3. The NiO present in the passive film on Fe-20Cr-15Ni did not affect the type of semiconductivity and Efb due probably to its small amount in the film. The Mott-Schottky plot (Fig. 3(b)) for the passive film formed at a Crtranspassive potential (0.8 VSCE) shows two linear regions indicating two kinds of donors, whereas that (Fig. 3(a)) for the passive film formed at a passive potential (0.6 VSCE) has a single linear region. The shallow donor in the n-type passive film is known to be oxygen vacancy or metal interstitial [11], and the deep donor associated with the second linear region for the passive film formed at a Cr-transpassive potential is reported to be Cr6+ ion [12]. Shallow donor density (Nd1) determined from the slope of linear regions in the Mott-Schottky plots in Fig. 3 was 2.2´1020 cm-3 ~ 3.6´1020 cm-3 for the passive film formed at 0.6 VSCE, and 1.6´1020 cm-3 ~ 1.9´1020 cm-3 for that formed at 0.8 VSCE. Apparently, Nd1 decreased with increasing film formation potential, as E. Cho reported [12], and does not have clear dependence on the concentration of Ni in alloys. However, the density of deep donor (Nd2) linearly increased from 8.4´1020 to 1.5´1021 cm-3 with increasing Ni content in the alloys. Thus, Ni in the passive films can be considered to increase the density of deep donor (Cr6+). The increase in the amount of Cr6+ may be involved by the substituting Fe3+ by Ni2+ in the passive film in the point of charge balance. Therefore, this dependence of Nd2 on the concentration of Ni means that some of Ni2+ is incorporated in g-Fe2O3 structure substituting Fe3+. It is suggested from the high donor (or oxygen vacancy) density and the dependence of the deep donor (or Cr6+) density on Ni content that the films on Fe-20Cr-xNi (x = 10, 15) is a nano-crystalline structure composed of (Cr, Ni)substituted g-Fe2O3 mixed with or without NiO depending on the Ni content. This suggestion corresponds the recent high resolution scanning tunneling E f = 0 .6 V S C E 1.5 x 1 0

10

1.0 x 1 0

10

5 .0 x 1 0

-3

F e -2 0 C r-1 5 N i

9

(a )

0 .0

E f = 0 .8 V S C E

F e -2 0 C r

10

F e -2 0 C r-1 5 N i

-2

-2

C /F cm

4

1.5 x 1 0 1.0 x 1 0

10

5 .0 x 1 0

Ef = 0.6 VSCE Ef = 0.8 VSCE

21

1.8x10

F e -2 0 C r-1 0 N i

Donor density / cm

-2

-2

C /F cm

4

F e -2 0 C r F e -2 0 C r-1 0 N i

Nd1+Nd2

21

1.5x10

Nd2

21

1.2x10

20

9.0x10

20

6.0x10

Nd1

20

3.0x10

9

0 (b )

0 .0 -1 .0

-0 .5

0 .0

0 .5

2

4

6

8

10

12

14

16

Ni content / wt%

1 .0

P o te n tia l / V

Fig. 3 Mott-Schottky plots of the passive films formed on Fe-20Cr-xNi (x = 0, 10, 15) at (a) 0.6 VSCE and (b) 0.8 VSCE for 2 h in deaerated pH 8.5 buffer solution.

Fig. 4 Effect of Ni on donor densities of passive film formed on Fe-20Cr-xNi (x = 0, 10, 15) at 0.6 VSCE and 0.8 VSCE for 2 h in pH 8.5 buffer solution.

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microscopy studies which revealed that passive films formed on Fe-Cr, Fe-CrNi alloys and on Ni in H2SO4 solution are nano-crystalline [4, 13]. 4. Conclusions 1. The photocurrent spectra for the passive films on Fe-20Cr and Fe-20Cr-10Ni were almost same. The photocurrent spectra were resolved into two components each of which originated from d-d and p-d electronic transitions in g-Fe2O3, suggesting that the passive films on Fe-20Cr-10Ni are composed of Cr-substituted g-Fe2O3. 2. Photocurrent spectra for the passive film on Fe-20Cr-15Ni alloy comprise two spectral components; one for the Cr-substituted g-Fe2O3 and the other for NiO, suggesting that the passive film consisted of Cr-substituted g-Fe2O3 and NiO. 3. Passive films on Fe-20Cr-xNi (x = 10, 15) alloys exhibited n-type semiconductivity with Efb of -0.3 VSCE, irrespective of Ni content, confirming that the films have a base structure of Cr-substituted g-Fe2O3. 4. It was suggested from the high shallow donor (or oxygen vacancy) density of ~1020 cm-3 and the dependence of the deep donor (or Cr6+) density on Ni content, that the passive films on Fe-20Cr-xNi (x = 10, 15) alloys appears to be a nano-crystalline (Ni, Cr)-substituted g-Fe2O3 mixed with or without NiO depending on the Ni content. References 1. C. Sunseri, S. Piazza, A. Di Paola, and F. Di Quarto, J. Electrochem. Soc. 134 (1987) 2410. 2. A. Di Paola, F. Di Quarto, and C. Sunseri, Corros. Sci. 26 (1986) 935. 3. A. M. P. Simoes, M. F. S. Ferreira, B. Rondot, and M. Da Cunha Belo, J. Electrochem. Soc. 137 (1990) 82. 4. V. Maurice, W. P. Yang, and P. Marcus, J. Electrochem. Soc. 145 (1998) 909. 5. N. E. Hakiki, S. Boudin, B. Rondot and M. Da Cunha Belo, Corr. Sci. 37 (1998) 1809. 6. C. R Clayton and I. Olefjord, in: P. Marcus et al. (Eds.), Corrosion Mechanisms in Theory and Practice, Marcel Decker, New York, 1995, p.175. 7. E. Cho, H. Kwon, and D. D. Macdonald, Electrochim. Acta 47 (2002) 1661. 8. H. J. Jang, C. J. Park, and H. S. Kwon, Electrochim. Acta, 50, 16-17, 3503. 9. J. Kim, E. Cho, and H. S. Kwon, Corros. Sci. 43 (2000) 1403. 10. S. J. Ahn and H. S. Kwon, Electrochim. Acta 49 (2004) 3347. 11. M. Bojinov, G. Fabricius, T. Laitinen, K. Makela, T. Saario, and G. Sundholm, Electrochim. Acta 45 (2000) 2029. 12. E. Cho, Ph. D. Thesis, KAIST, Korea, 2002. 13. V. Maurice, H. Talah, and P. Marcus, Surf. Sci. 304 (1994) 98.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Published by Elsevier B.V.

285

Photoelectrochemical Response and Corrosion Property of Passive Films on Fe-18Cr Alloy Shinji Fujimotoa a

Department of Materials Science and Processing, Graduate School of Engineering, Osaka University, 2-1 Yamada-oka, Suita Osaka 565-0871 Japan, [email protected]

Abstract - Semiconductor properties of passive films formed on the Fe-18Cr alloy in a borate buffer solution (pH=8.4) and a 0.1 M H2SO4 solution were examined using photoelectrochemical response and electrochemical impedance spectroscopy. Typical an n-type semiconductor behaviour is observed by both photo current and impedance for the passive films formed in the borate buffer solution. However, a negative photocurrent was observed in the 0.1 M H2SO4 solution. This indicates that the passive film formed in acid behaves as a p-type semiconductor. However, Mott-Schottky plot of the capacitance showed typical n-type semiconductor property. It is concluded that the passive films on the Fe-18Cr alloys formed in the borate buffer solution is composed of both n-type outer hydroxide and inner oxide layers. On the other hand, passive film on Fe-18Cr alloy in 0.1 M H2SO4 solution consists of p-type oxide and n-type hydroxide layers with isotype heterojunction. The kinetics of passive film growth and the corrosion behavior are discussed in terms of the electronic structure of the passive film. Keywords: Semiconductor, Capacitance, Electronic property, Stainless steels

1. Introduction Passivity of Fe-Cr alloys is one of the most important topics in the corrosion science. Therefore, passive films on Fe-Cr alloys have been characterized by various ultrahigh vacuum (UHV) analytical techniques such as Auger electron spectroscopy (AES) and X-ray Photoelectron spectroscopy (XPS). Such UHV surface characterization techniques subject specimen to be modified under quite different environment compared with that in which passive films were formed. Therefore, in situ electrochemical surface characterization techniques have been tried. The author of this work has reported the photo electrochemical response and electrochemical impedance spectroscopy of passive films on Fe-18Cr alloys

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to reveal difference in electronic structure of passive films formed in sulphuric acid and borate buffer solution [1-3]. In the present paper, the author summarize the photoelectrochemical response of passive films on Fe-18Cr alloy comparing with the Mott-Schottky relation, which is correlated to corrosion behaviour, and also the influence of additive elements. 2. Semiconductor structure of passive films on Fe-18Cr alloy Assuming that the photo excitation occurs as an indirect transition usually observed for passive films on Fe, Cr and Fe-Cr alloys, the photo current spectra can be described as following;

( i ph × hn/I 0 )1/2 = S刷(hn - E g )

(1)

(Iph hn/I0)

1/2

/A

1/2

eV

1/2

W

-1/2

where I0 is the intensity and hn the photon energy of the incident light, respectively, Eg the band gap energy of the passive film and S constant. Figure 1 shows a typical photo-electrochemical response spectrum for passive film formed on Fe-18Cr alloy in borate buffer solution (pH=8.4), plotted following Eq.1. This spectrum does not show one straight line, but exhibits two regions of different slopes. As usually recognised [4], passive films consist of not only a single phase but of duplex layers; oxide and a covering hydroxide layer. Therefore, the author tried to separate this spectrum into two components. The process of separation was described elsewhere [1]. Eg can be estimated as the photon energy at which the (iph hn/I0)1/2 equals to zero. Egs of approximately 2.4 eV and 3.4-3.5 eV are observed, and are almost constant for any polarization period and film formation potential examined. Referring to Eg reported previously [1,2,5], the components -2 1 10-2 1x10 with Eg of 2.4 eV and 3.4-3.5 eV Passivation time : 86.4 ks are identified to Cr(OH)3 and Em = 0 mV Cr2O3, respectively. XPS revealed Ef = 400 mV Ef = 200 mV that passive films of Fe-18Cr alloy are mainly composed of an inner oxide layer and an outer hydroxide -3 5x10 5 10-3 layer. Therefore, passive films on Fe-18Cr alloy are composed of an inner oxide layer with Eg of 3.4-3.5 eV and a covering hydroxide layer with Eg of 2.4 eV. The intensity of 0 the photo cuurent response was also 0 10 0 2 3 4 5 summarised as follows. Figures 2 Photon energy / eV (a) and (b) show the slope of the spectrum, S, for the larger and Fig.1 Typical photoelectrochemical response smaller Eg components, respectively, spectrum for Fe-18Cr alloy measured in as a function of applied potential. borate buffer solution, pH=8.4.

Passive Films on Fe-18Cr Alloy

2 10

S / A W-1 eV-1

1 10

Ef = Ef = Ef = Ef = Ef =

-5

-5

500 m V 400 m V 300 m V 200 m V 0 mV

E g : 2.4-2.5 eV

(a)

0

6 10

-5

E g : 3.4-3.5 eV Passivation time : 86.4 ks

4 10

2 10

-5

-5

(b)

0 -400

-200

0

200

400

600

Potential / mV Ag/AgCl

Fig.2 Variation of the slopes of the action spectrum, S, for (a) hydroxide layer with Eg of 2.4-2.5 eV and (b) oxide layer with Eg of 3.4-3.5 eV as a function of measuring potential. 5 0

Eg = 2.5 eV

-5

0

Sk / 10-6 A W-1 eV-1

These results were obtained for passive films formed at various film formation potentials, Ef, then the applied potential was shifted in less noble direction sequentially by an interval of 100 mV with measurement of the photo current at each potential. As shown in these figure, the slope of spectrum increases with increasing applied potential, which is a typical ntype semiconductor behaviour. Flat band potentials, that are the potential at which the value S becomes zero, are located around -300 mV. A MottSchottky plot of capacitance, which was obtained by an electro-chemical impedance spectroscopy, also showed the typical n-type semiconductor behaviour with similar flat band potentials that were described in Fig.2. Figure 3 shows the photo current spectrum measured for Cr and Fe-Cr alloys in a 0.1 M sulphuric acid, exhibiting different spectrum compared with that obtained for the passive film formed in borate buffer solution, although two band gap energies are similar to that observed in borate buffer solution, that is, 2.5 and 3.6-3.7 eV. The slope of the spectrum, S, for the component, Eg=3.6 eV, exhibits the typical p-type semi-conductor behaviour with flat band potential located around the noble edge in the passive potential region. Therefore, the oxide layer was in a depleted state of p-type semiconductor. On the other hand, S for the spectrum exhibiting Eg of 2.5 eV, that is the hydroxide layer, increases with potential crossing the x axis in the middle of the passive potential region. Therefore, the conduction type is not clear. Figure 4 shows the

287

-50

E f = 600 mV

-100

Fe- 8Cr Fe-14Cr Fe-18Cr pure Cr

Eg = 3.6 eV -150 -400

-200

0

200

400

600

Potential / mV Ag/AgCl Fig.3 Variation of the slope of the photo current spectrum, Sk, for passive film on pure Cr and Fe-Cr alloys formed on 600 mVAg/AgCl as a function of the measuring potential.

S. Fujimoto

288

C-2 / 10-6 F-2 cm4

Mott-Schottky plot for passive films on Fe-18Cr formed in 0.1 M H2SO4. Fe-18Cr 0 mV There are observed two lines with -3 200 mV 0.1 kmol m H 2 SO 4 400 mV 2 positive slope for potentials more 600 mV noble than 0 mV and with negative slope for potentials less noble than 0 mV. The capacitance should not derive from only one space charge 1 region, but from at least two regions. As mentioned before, the oxide layer clearly showed the characteristic of depleted state of p-type semiconductor. Therefore, the lines with negative 0 -400 -200 0 200 400 600 slope might derive from the oxide layer. On the other hand, the positive Potential / mV Ag/AgCl slope might come from hydroxide Fig.4 Mott-Schottky plots for passive layer. Therefore, the covering films formed on Fe-18Cr at several hydroxide layer is concluded to be an potentials in 0.1 kmol m-3 H2SO4. n-type semiconductor layer, which could not be necessarily concluded from the results obtained from photo electrochemical current response. Figure 5 summarises the electronic structure for passive film on Fe-Cr alloys formed in acid and neutral solution. It is noticeable that passive film formed in borate buffer solution consists of both n-type as depleted state inner oxide and outer hydroxide layers, whereas passive film formed in the acid solution has an inner oxide layer of a p-type semiconductor as a depleted state and outer hydroxide layer of an n-type semiconductor. Such different duplex structures of passive films on Fe-Cr alloys formed in the acid and the neutral or alkaline solutions have already reported using surface analysis techniques [4, 6-8].

n-type semiconductor layer

p-type semiconductor layer

(a) in sulphuric acid

Conduction band

Ef Depleted state

Depleted state

Valence band

Passive film

Substrate

n-type layer

Conduction band

Ef

Electrolyte

Electrolyte

n-type layer Depleted state

Passive film

Depleted state

Substrate

Valence band

(b) in borate buffer solution

Fig.5 Schematic illustration showing electronic structure of duplex passive films formed on Fe-Cr alloys formed in (a) sulphuric acid solution and (b) borate buffer solution.

Passive Films on Fe-18Cr Alloy

289

Strehblow etal. discussed that the enrichment of Cr in the passive film could be explained by the solubility and dissolution kinetics of Fe(II, III) and Cr(III) oxide and hydroxide. In neutral solution, an Fe enriched hydroxide film is formed because of less solubility of Fe(III) oxide. In acid solution, on the other hand, Fe oxide easily dissolved resulting in formation of thin Cr(III) enriched outer hydroxide layer. Cation and anion content and their defects may influence on the electronic properties of passive films. In the present work using both photoelectrochemical and electrochemical impedance techniques, the duplex semiconductor properties of passive film on Fe-18Cr alloy formed in the acid and the neutral solutions are consistently discussed. 3. Corrosion processes correlated with electronic structure of passive films As discussed above, a passive film has various electronic states depending on the applied potential and also environment. Figure 5(b) indicates that for an n-type semiconductor layer, the electronic energy band in the passive film exhibits an ascending slope towards the interface with electrolyte. In this condition, electron transfer across the film / electrolyte interface is blocked by depleted region. However, both cation outwards and anion inwards migration could be assisted by the electric field within the film. Therefore, continuous oxide growth, which is enhanced by the potential slope, is possible. Moreover, the inwards penetration of anions, such as Cl- and SO42-, might be promoted, resulting in the breakdown of the passive film to cause localised corrosion. It is also noted, if an n-type semiconductor layer is cathodically polarized, electron transfer from the passive film to the electrolyte is not blocked, because the surface of semiconductor layer is in the accumulated state. Therefore, cathodic reaction, such as oxygen reduction easily occurs through the passive film. On the other hand, the p-type semiconductor in depleted state effectively blocks cathodic reaction, and also the anodic reaction is suppressed because of an inward positive electronic energy band slope. Moreover, cation outwards and anion inwards migration could be suppressed by the electric field inside the film as well. Therefore, both oxide film growth and anodic and cathodic corrosion reactions are effectively inhibited by a p-type passive film. Kirchheim reported that in a neutral solution a passive film grows continuously for long periods following the Cabrera-Mott mechanism [9], whereas in an acidic solution, the growth of the passive film is terminated in a few minutes [9]. Thus, in an acidic solution the passivity is completed in a very short period. Fujimoto et al. confirmed that the n-type semiconductor properties are observed in a very initial period of passivation for Fe-18Cr in a sulphuric acid solution [10], which means that the passive film grows when a positive potential gradient is established during the initial short period. Semiconductor properties of the passive film significantly influence the corrosion behaviour and passive film growth of stainless steel. These properties may be affected by additive elements. Photoelectrochemical response and

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electrochemical impedance were examined for sputter deposited Fe-18Cr alloys with additive elements of various kinds and amounts up to 5 at.%. The results will be reported elsewhere. It is found that additives did not change the band gap energy of the passive film. However, additives modified the carrier density. The Cr enriched passive films on Fe-Cr alloys may form some sort of threedimensional network structure, like the percolation model proposed by Newman and Sieradzki [11,12]. Additives might not modify a continuously connected domain of Cr oxide and hydroxide, resulting in constant band gap energy independent of additive elements. The correlation among the electronic and the chemical structure of passive films and the corrosion resistance of Fe-18Cr alloys with various additive elements will be reported in the future study. Acknowledgement This work was supported by funds from “Priority Assistance of the Formation of Worldwide Renowned Centres of Research - the 21st Century COE program, (Project: Advanced Structural and Functional Materials Design) from the Ministry of Education, Sports, Culture, Science and Technology, Japan. References 1. S. Fujimoto, O. Chihara and T. Shibata, Materials Science Forum, 289292 (1998) 989. 2. H. Tsuchiya, S. Fujimoto, O Chihara and T. Shibata, Electrochimica Acta, 47 (2002) 4357. 3. H. Tsuchiya, S. Fujimoto and T. Shibata, J. Electrochem. Soc., 151 (2004) B39 4. W. P. Yang, D. Costa, and P. Marcus, J. Electrochem. Soc. 141 (1994) 111. 5. F. Di Quarto, S. Piazza and C. Sunseri, Corros. Sci., 31 (1990) 721. 6. S. Haupt and H.-H. Strehblow, Corros. Sci., 37, (1995) 43. 7. S. Haupt and H.-H. Strehblow, J. Electroanal. Chem., 228, (1987) 365. 8. C. Calinski and H.-H. Strehblow, J. Electrochem. Soc., 136, (1989) 1328. 9. R. Kirchheim, "Modification of Passive Films" Eds, P. Marcus, B. Baroux, and M. Keddam, p.17, (The Institute of Materials, London, 1994). 10. S. Fujimoto, S. Kawachi and T. Shibata, Proc. 8th Int. Sym. Passivity and Metals and Semiconductors, M. B. Ives, J. L. Luo and J. R. Rodda, Editors, PV99-42, p.260, The Electrochemical Society Proceedings Series, Pennington, NJ (2001). 11. K. Sieradzki and R. C. Newman, J. Electrochem. Soc. 133 (1986) 1979. 12. S. Fujimoto, R. C. Newman, G. S. Smith, S. P. Kaye, H. Kheyrandish and J. S. Colligon, Corros. Sci. 35, (1993) 51.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

291

Effects of EDTA on the Electronic Properties of Passive Film on Fe-20Cr in pH 8.5 Buffer Solution EunAe Cho,a SeJin Ahn,b and HyukSang Kwonc a

Fuel Cell Research Center, Korea Institute of Science and Technology, 39-1 Hawolgokdong, Sungbukgu, Seoul 136-791, Republic of Korea b Solar Cell Research Center, Korea Institute of Energy Research, 71-2 Jangdong, Yuseonggu, Daejeon 305-343, Republic of Korea. c Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, GuSongDong 373-1, YuSeongGu, DaeJon 305-701, Republic of Korea; [email protected]

Abstract - Effects of EDTA (ethylene diammine tetraacetic acid) on the electronic properties of passive film formed on Fe-20Cr alloy in pH 8.5 buffer solution were examined by the photo-electrochemical and impedance techniques to elucidate the structure and composition of passive film on the alloy. XPS composition profile in depth for the passive film revealed that Cr content in the outermost layer of the passive film was higher for the passive film formed in the solution containing EDTA than that in solution without EDTA, due presumably to the selective dissolution of Fe oxide by EDTA. In spite of the selective dissolution of Fe oxide in the outer layer of the passive film, the photocurrent spectra and the Mott-Schottky plots for the passive film formed on Fe-20Cr in EDTA containing solution were very close to those in EDTA-free solution, which strongly supports that the passive film on Fe-20Cr in pH 8.5 buffer solution is a single layered Cr-substituted g -Fe2O3 oxide rather than a duplex layered oxide composed of inner Cr-oxide and outer Fe-oxide. Keywords: passive film, electronic property, Fe-20Cr, EDTA, photocurrent, MottSchottky plot

292

E. Cho et al.

1. Introduction It is generally agreed that the excellent corrosion resistance of stainless steels is due to the passive film formed on the surface. Thus, elucidating the nature of the passive film is a prerequisite for understanding such high corrosion resistance of stainless steels. In the previous study using in-situ photocurrent measurement [1], the authors proposed that the passive film formed on Fe-20Cr alloy in pH 8.5 buffer solution was a Cr-substituted g -Fe2O3 oxide film where Cr3+ acted as an effective recombination site for electron-hole pairs, emphasizing on its single layered structure. It was based on the experimental results that photocurrent spectrum for the passive film on Fe-x20Cr (x = 0~20 wt.%) resembled in shape to that on pure Fe, regardless of film formation potentials and Cr content. Many other researchers also reported that passive film on stainless steels was a single layered Fe oxide or a Fe-Cr mixed oxide [2-4]. On the other hand, some other authors claimed that the passive film formed on Fe-Cr alloys are composed of duplex oxide layers with inner Cr oxide and outer Fe oxide [5, 6] or with inner oxide and outer hydroxide [7]. As described above, there are still controversies on the structure of passive film on stainless steels, i.e., single or duplex layered structure? In this subject, EDTA (Ethylene Diammine Tetraacetic Acid) is considered to be a useful chemical additive in that EDTA is known as a chelating agent for the removal of Fe oxide [8]. If a passive film on Fe-Cr alloy is, indeed, composed of duplex layers with inner Cr-oxide and outer Fe-oxide, the passive film on Fe-20Cr alloy should reveal the properties of inner Cr-oxide when immersed in solution with EDTA. In this regard, the objective of the present work is to elucidate the structure and composition of passive film formed on Fe-20Cr ferritic stainless steel in pH 8.5 buffer solution by investigating the effects of EDTA on the electronic properties of passive film on the alloy. 2. Experimental Methods High purity vacuum arc melted Fe-20Cr (wt. %) alloy was used as a working electrode. A conventional three electrode cell of 1 -multi neck flask with a quartz window as a photon inlet was used for the photocurrent and impedance measurement. The cell was equipped with a Pt counter electrode and a saturated calomel reference electrode. All the experiments were carried out at ambient temperature in deaerated pH 8.5 buffer solution made of 0.2 M boric acid, 0.05 M citric acid and 0.1 M tertiary sodium phosphate solutions with and without 0.05 M EDTA. The preparation procedures for sample and solution, and the experimental arrangement were described elsewhere [1]. The working electrode was initially reduced at –1.0 VSCE for 10 min. to remove air-formed oxide on the surface, and passivated at various film formation potentials (-0.1 ~ 0.9 VSCE) for 24 h.

Effects of EDTA on the Electronic Properties of Passive Film on Fe-20Cr

293

The 300 W Xenon arc lamp combined with a scanning digital monochromator was used to impose a monochromatic illumination to the working electrode. The monochromator was controlled at a scanning rate of 1 nm/s by stepping motor, which made it possible to provide monochromatic photons with 200 to 800 nm wavelengths to the working electrode. To increase the photon flux, white light from the Xe lamp was focused to the light inlet using two auxiliary focusing lenses. Details of the photocurrent experimental setup were presented elsewhere [1]. The capacitance measurements for the Mott-Schottky analysis were conducted on the passive films formed on Fe-20Cr at various film formation potentials ranging from –0.1 to 0.9 VSCE for 24 h, by sweeping the applied potential in negative direction from the film formation potential. The Mott-Schottky plots were obtained at a frequency of 1 kHz with excitation voltage of 10 mV (peakto-peak) using a frequency response analyzer. Chemical composition profile in depth for the passive film on Fe-20Cr grown for 24 h at 0.2 VSCE was obtained by XPS analysis using the Al-Kg and Mg Kg x-ray source (15 kV, 20mA, 300W), and a pass energy of 20 eV. Depth profiles of chemical composition for passive film were obtained with sputtering of argon ions (PAr = 5×10-7 torr, base pressure = 5×10-10 torr, energy: 5 kV, current: 3.0 mA/cm2). 3. Results and Discussion

current density / Acm

-2

3.1. Polarization Responses -1

10

Cr

Fe

Fe-20Cr

(b)

(c)

without EDTA with EDTA

-3

10

-5

10

-7

10

(a)

-9

10

-1

0

1

potential / V

-1

0

1

potential / V SCE

SCE

-1

0

1

potential / V SCE

Fig. 1 Effects of EDTA on the potentiodynamic polarization responses of (a) Cr, (b) Fe, and (c) Fe-20Cr in deaerated pH 8.5 buffer solution with or without 0.05 M EDTA measured at a scan rate of 0.5 mV/s. Anodic sweep was employed in the polarization tests.

Fig. 1 shows effects of EDTA on polarization responses of Cr, Fe and Fe-20Cr in pH 8.5 buffer solution at a scan rate of 0.5 mV/s. It is apparent from Fig. 1(a) that polarization behavior of Cr is not markedly affected by the addition of EDTA to the solution. On the contrary, the passive region of Fe was

294

E. Cho et al.

significantly reduced with the passive current density being increased about 5 orders in magnitude (Fig. 1(b)) by the addition of 0.5 M EDTA to the solution, demonstrating clearly that 0.5 M EDTA in pH 8.5 buffer solution effectively acts as a chelating agent for Fe oxide. In the case of Fe-20Cr (Fig. 1(c)), EDTA in the solution increased the passive current density about 2 times due presumably to the selective dissolution of Fe oxide in the passive film, but it was ignorable compared to the significant increase in passive current density of Fe by EDTA as shown in Fig. 1(b). This demonstrates that the outer layer of the passive film on Fe-20Cr is not solely composed of Fe oxide, which is contradicted to the duplex layer model. 3.2. XPS Depth Profile passive film

alloy

0.50 0.45

Fe-20Cr pH 8.5 buffer solution Uf = 0.2 VSCE

Cr/(Fe+Cr)

0.40 0.35

without EDTA with EDTA

0.30 0.25 0.20 0.15

0

5

10

15

20

25

Fig. 2 XPS chemical composition profile in depth for passive film formed on Fe-20Cr at 0.2 V for 24 h in deaerated pH 8.5 buffer solution with or without 0.05 M EDTA

30

sputtering time / min.

In order to investigate the effects of EDTA on the chemical composition of passive film, XPS chemical analysis in depth was conducted for the passive film formed on Fe-20Cr at 0.2 VSCE for 24 h in pH 8.5 buffer solution with or without EDTA. For the passive film formed in solution without EDTA, the ratio of Cr increased with depth of the film at the beginning of sputtering, exhibiting the maximum at the estimated thickness of about 1 nm, and then decreased. On the other hand, Cr content in the outermost part of the passive film formed on Fe-20Cr in the solution containing 0.05 M EDTA increased towards film/solution interface without showing the maximum at the subsurface position. It is due presumably to a selective dissolution of Fe oxide in the passive film by EDTA. In addition, Cr content in the passive film was slightly higher throughout the entire film thickness than that formed in the solution without EDTA. These XPS results confirm that EDTA in the solution selectively dissolves Fe oxide in the passive film on Fe-20Cr, as was already demonstrated in the anodic polarization curve, Fig. 1(c). However, it is significant that the overall shape of XPS depth profile is still far from the

Effects of EDTA on the Electronic Properties of Passive Film on Fe-20Cr

295

duplex layer model suggesting inner Cr oxide and outer Fe oxide, regardless of the presence of EDTA in the solution. Rather, it is close to the single layer model. 3.3. Photoelectrochemical Response 1 .5 1 .0

(a ) F e -2 0 C r p H 8 .5 b u ffe r s o lu tio n U f = 0 .2 V S C E

Photocurrent / nA

0 .5 0 .0 0 .4 0 .2

(b ) F e -2 0 C r p H 8 .5 b u ffe r s o lu tio n w ith 0 .0 5 M E D T A U f = 0 .2 V S C E

Fig. 3 iph vs. hn 伊plot for the passive film formed on Fe-20Cr at 0.2 VSCE in deaerated pH 8.5 buffer solution (a) without or (b) with 0.05 M EDTA

0 .0

2

3 4 5 P h o to n e n e r g y / e V

Fig. 3 shows the photocurrent spectra for the passive film formed on Fe-20Cr at 0.2 VSCE for 24 h in pH 8.5 buffer solution with or without 0.05 M EDTA. Regardless of the presence of EDTA in the solution, photocurrent began to increase at approximately 3.0 eV, exhibiting a shoulder at 3.8 eV, and then the peak current density at 4.5 eV, showing different behavior from the photocurrent spectrum for the passive film of Cr [9]. The only difference in the two spectra with or without EDTA is that current fluctuation or current noise is much greater in the film formed in solution with 0.05 M EDTA than in that without EDTA, which is due presumably to the increase in the amount of point defects in the outer layer of passive film arising from the selective dissolution of Fe oxide by EDTA. It was previously demonstrated [1] that the photocurrent spectrum for the passive film on Fe-20Cr alloy is resemble in shape to that for the passive film on Fe except for the large difference in photocurrent density between them, which led to the conclusion that the passive film on Fe-20Cr is a Cr-substituted g -Fe2O3 in which Cr3+ ion act as a recombination site for electron-hole pair. Further, the shoulder at 3.8 eV and the peak at 4.5 eV appeared in the photocurrent spectrum (Fig. 3(a)) for the passive films on Fe-20Cr were

296

E. Cho et al.

generated by the d-d transition from Fe3+ band to Fe2+ band and the p-d transition from O-2p to Fe2+ band in Cr-substituted g -Fe2O3 oxide film with spinel structure [1], respectively. Thus, it can be concluded, based on the similarities in the photocurrent spectrum and XPS composition profile in depth between the passive film formed in solution with and that without EDTA, that the passive film formed on Fe-20Cr maintains its composition and structure of spinel Cr-substituted g -Fe2O3 even in pH 8.5 buffer solution containing 0.05 M EDTA. 3.4. Mott-Schottky Analysis 10

1.2x10

10

pH 8.5 buffer solution Uf = 0.2 VSCE

1.0x10

without EDTA

-2

CSC / F cm

4

9

8.0x10

9

-2

6.0x10

with EDTA

9

4.0x10

9

Fig. 4 Effects of EDTA on the MottSchottky plots for the passive film formed on Fe-20Cr at 0.2 VSCE in deaerated pH 8.5 buffer solution with or without 0.05 M EDTA.

2.0x10

0.0 -0.6

-0.4

-0.2

0.0

0.2

potential / V

Fig. 4 shows the Mott-Schottky plots for the passive film formed on Fe-20Cr at 0.2 VSCE in pH 8.5 buffer solution with or without 0.05 M EDTA. The MottSchottky plots exhibited positive slopes regardless of the addition of EDTA in the solution, indicating that the passive films are n-type semiconductors in both the solutions. By adding EDTA to pH 8.5 buffer solution, the flat band potential and the donor density was increased from -0.45 to -0.30 VSCE and from 3.6´1020 to 5.6´1020 cm-3, respectively. The increase in the flat band potential is due presumably to the different solution chemistry caused by the addition of EDTA to pH 8.5 buffer solution, and the higher donor density imply that the passive film formed on Fe-20Cr in EDTA containing solution has a more defective structure with higher concentration of probably oxygen vacancy due to the selective dissolution of Fe oxide to the subsurface of the film by EDTA. If the duplex layer model is valid, the Mott-Schottky plot for the passive film on Fe-20Cr should reveal that for inner Cr-oxide layer when the passive film on the alloy has been formed in the solution with EDTA. However, the shape of MottSchottky plot for the film formed on Fe-20Cr alloy in the solution with EDTA is far from that for the passive film on Cr [9], showing p-type semiconducting characteristics, but similar to that for the passive films formed on Fe or Fe-20Cr in the solution without EDTA. This strongly supports again that the passive film on Fe-20Cr in deaerated pH 8.5 buffer solution is close to the single layered Cr-

Effects of EDTA on the Electronic Properties of Passive Film on Fe-20Cr

297

substituted g -Fe 2O3 rather than the duplex layered film of inner Cr oxide and outer Fe oxide. From all of these experimental data, we conclude that the passive film formed on Fe-20Cr alloy in deaerated pH 8.5 buffer solution is a single layered and highly defective n-type semiconductor whose structure and composition are very close to spinel Cr-substituted g -Fe2O3. 4. Conclusions 1. 0.5 M EDTA in pH 8.5 buffer solution acts effectively as a chelating agent for Fe oxide, which was confirmed by the anodic polarization curve of Fe. In contrast, effects of EDTA on the anodic polarization response of Fe-20Cr alloy were negligible. 2. It was shown based on the XPS chemical composition profile in depth that Cr content in the passive film on Fe-20Cr was the maximum at the subsurface, and then decreased gradually with depth of the film in the EDTA-free solution. The addition of EDTA to the solution did not affect the overall Cr content profile in the passive film. 3. Regardless of the presence of EDTA in the solution, photocurrent spectrum for the passive film on Fe-20Cr exhibited that for the Cr-substituted g -Fe2O3. 4. Mott-Schottky plot for the passive film on Fe-20Cr exhibited n-type semiconductivity with similarity to that for the passive film formed on Fe in pH 8.5 buffer solution with or without EDTA. 5. From the findings discribed above, it is suggested that the passive film on Fe20Cr in deaerated pH 8.5 buffer solution is close to the single layered Crsubstituted g -Fe2O3. Reference 1. 2. 3. 4. 5. 6. 7. 8. 9.

E. A. Cho, H. S. Kwon and D. D. Macdonald, Electrochim. Acta, 47 (2002) 1661. C. Sunseri, S. Piazza, A. Di Paola and F. Di Quarto, J. Electrochem. Soc., 134 (1987) 2410. A. Di Paola, F. Di Quarto and C. Sunseri, Corrosion Science, 26 (1986) 935. M J .Klopper, F. Bellucci and R. M. Latanision, Corrosion, 48 (1992) 229. N. E. Hakiki, M. Da Cunha Belo, A. M. P. Simoes and M. G. S. Ferreira, J. Electrochem. Soc., 145 (1998) 3821. V. Mitrovic-Scepanovic, B. MacDougall and M. J. Graham, Corros. Sci., 24 (1984) 479. W . P. Yang, D. Costa and P. Marcus, J. Electrchem. Soc., 141 (1994) 113. D. A. Frey, Mater. Perform., 20 (1981) 49. J. S. Kim, E. A. Cho and H. S. Kwon, Electrochim. Acta, 47 (2001) 415.

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299

Electrochemical and Mott-Schottky Behaviour of the Oxide Films on I625 Exposed to Gamma Radiation James F. Dante, Darrell S. Dunn, Kendra T. Price Southwest Research Institute®, 6220 Culebra Road, San Antonio, Texas 78238, USA [email protected]

Abstract - Polarization and Mott-Schottky experiments were performed on alloy 625 samples previously exposed to 1.2 MeV gamma radiation at dose rates up to 1 Mrad/hour. While the breakdown potential is dependent on dose rate, the repassivation potential is not. These residual effects of the radiation exposure were not observed on samples exposed to the field in an air environment. At the highest dose rate, trapped holes at the surface at potentials above the repassivation potential result in a significant reduction in the crevice corrosion nucleation time. Key Words: oxide, alloy 625, radiation, semiconductor, electrochemistry

1. Introduction Passivity of high chromium nickel alloys has made them attractive for use in nuclear reactors and high level waste storage. Despite the low passive current densities under a wide variety of environmental conditions, however, these alloys can be susceptible to localized corrosion. In the case of radioactive waste disposal containers, the effect of radiation, mainly in the form of gammaray photons, has been considered. The primary effect of radiation has been reduced to the creation of stable radiolysis products (e.g., H2O2), which are oxidizing in nature. This was demonstrated by Glass et al for AISI 304L and 316L stainless steel in the presence of a 3.3 Mrad/hour gamma field 1. The radiolysis of the near-field environment may be considered an “indirect” effect of radiation on corrosion. The “direct” effect of gamma radiation on the alloys has largely been ignored. However, because of the semiconductor nature of the protective oxide film on Ni-Cr alloys 2,3, gamma radiation may influence the surface charge and

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electron transfer within the oxide and, thus, directly influence the protectiveness of the oxide film. Throughout the last several decades, numerous studies have exploited the semiconductor nature of transition metal oxides to explain their behavior as catalysts, their corrosion behavior, and basic physical properties of solid liquid interfaces. The effective band gap of many of these oxides is between 1 eV and 3.5 eV. Thus, most studies have employed the use of IR to UV radiation to explore the oxide properties under illumination. Photo effects have been used to study the passive films on Ti 4,5, Al 6, Zn 7, Fe 8,9, Ni 10 and stainless steel 2,3,11-15. In the case of Zn, Fe, Ni, and stainless steel, UV irradiation resulted in not only a reduction in the nucleation time for pitting but also results in permenant changes in the localized corrosion behavior for Zn and Fe. The residual changes were cited to be a result of physical changes of the oxide film. The effect is observed under UV irradiation where the formation of high concentrations of radical species is less likely (although not specifically addressed). Fujimoto et al point to an enrichment in the Cr content of the oxide during illumination and (similar to Marsh et al) do not observe a change in the repassivation potential. Breslin et al observe that pit nucleation rate decreases with incident UV wavelength. Very few efforts have studied the impact of gamma radiation (~1 MeV) on the semiconductor and photo-electrochemical properties of stainless steel oxides and no work has studied oxides of nickel based alloys under gamma radiation. Capobianco et al 16, reported an enrichment of Cr in the outer oxide film layer of type 446 stainless steel after exposure to gamma radiation but gave no information regarding the localized corrosion behavior. Enhancement of catalytic behavior of NiO was noted upon exposure to gamma radiation in air 17. The main goal of this work is to address both the polarization and Mott-Schottky behavior of alloy 625 (UNS: N26625) as a function of exposure to gamma radiation. In particular, the effect of pre-exposure to the radiation field in a borate solution and in air will be examined. The relationship of these properties to localized corrosion will be examined. 2. Experimental All potentiodynamic and electrochemical impedance spectroscopy (EIS) tests were conducted in a four port, flat-bottom electrochemical cell using a three-electrode system. A saturated calomel reference electrode and platinum counter electrode were used. In all cases, aloy 625 was used as the working electrode. All samples were polished to a 1000 grit finish and cleaned with acetone in an ultrasonic bath for approximately five minutes. To ensure a reproducible surface outside of the crevice region, each sample was pretreated using a –1.5 V (SCE) potential for 5 minutes followed by anodic polarization in a borate buffer solution (0.3 M H3BO3 + 0.075 M Na2B4O7 solution at pH = 8.2) for 1 hour. A Teflon crevice former was attached to the sample during the electrochemical pre-treatment, removed during radiation

Electrochemical and Mott-Schottky Behaviour of the Oxide Films

301

exposure, and re-attached with a torque of 0.14 N.m prior to cyclic polarization and impedance measurements. Following the pretreatment, samples were exposed to either 104 rad/hour 6 or 10 rad/hour of 1.2 MeV gamma radiation from a Co60 source. Samples were either exposed while immersed in the borate electrolyte or in air. Electrochemical measurements were taken within 24 hours after completion of the radiation exposure. Cyclic polarization scans were performed at 0.2 mV/sec in a solution containing 5.0 M NaCl and pH = 7 at 60oC. The repassivation potential was measured as the point where the reverse scan crossed the passive current density of the forward scan. Mott-Schottky curves were developed from EIS measurements made at various potentials. Separate samples were used for the different electrochemical measurements. High frequency capacitance values from each EIS scan were derived using an equivalent circuit model. Details of the EIS analysis will be presented in a future publication. A 10 mV AC signal was applied at each potential. A series of EIS measurements were performed from high potentials to low potentials. In all electrochemical testing, the solutions were deaerated. 3. Results and Discussion 3.1. Cyclic Polarization Cyclic polarization scans for alloy 625 in a pH = 7, 5 M NaCl solution following a prior exposure to gamma radiation in a borate buffer and in air were performed. Both the breakdown potential (Ebr) and the repassivation potential , E : immersed in borate (Erp) are plotted as a function of dose rate br + Ebr: exposed in air in Figure 1. The first data points labeled Erp: immersed in borate at 1 rad/hour actuality has no exposure ' Erp: exposed in air and was plotted in this manner to accommodate the use of a logarithmic time scale. Lines in the Figure are drawn through the average values for each dose rate. Several important features can be noted from Figure 1. 0.4

0.3

E (V SCE)

0.2

0.1

0

-0.1

Figure 1: Ebr and Erp in 5 M NaCl, pH = 7 as a function of dose rate after prior exposure to radiation during immersion in a borate buffer solution or in air.

-0.2

-0.3 1x10 0

1x101

1x10 2

1x103

1x10 4

1x105

1x10 6

As with stainless steel exposed to UV light, alloy 625 displays an increase in the breakdown potential when exposed to a dose rate of 10 Krad/hour while immersed in the borate buffer. At this dose rate, there appears to be no effect of the total dose on the breakdown potential. Dose Rate (rad/hr)

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Examination of Figure 1 reveals that at the higher dose rate (1 Mrad/hr) a significant decrease in the breakdown potential is observed. Interestingly, for these high dose rates, an oxidation peak is observed near –0.160 +/- 0.030 V(SCE) which is slightly anodic of the average repassivation potential. There appears to be a limited effect of gamma radiation on the repassivation potential for alloy 625. This is also consistent with observations of UV exposure of stainless steel. The high values observed at the 10 Krad/hr dose rate are most likely an artifact resulting from the high breakdown potential under these conditions (i.e., the repassivation potential is a function of charge passed during localized corrosion up to a critical value). Also of note is the fact that exposure to the gamma field while in air resulted in no residual effects on the breakdown or repassivation potential for alloy 625. Thus, permanent changes in the oxide properties must be explained as a combined effect of photon/solution/surface film interatctions. It has been shown, for example, that under UV radiation, TiO2 substrates can stabalize reaction intermediates as surface states during the oxidation of water 18. 3.2. Semiconductor Properties Mott-Schottky plots of alloy 625 exposed to gamma radiation are shown in Figure 2. Of particular interest is the region where the 1/C2 curve begins to 1x10

9

2x10

It25 immersed in borate buffer no exposure 1 Mrad/hr for 1 hour 1 Mrad/hr for 1 hour 10 Krad/hr for 100 hours 10 Krad/hr for 1 hour 10 Krad/hr for 1 hour

6x10

8

4x10

8

I625 exposed in air no exposure 1 Mrad/hr for 1 hour 1 Mrad/hr for 1 hour 10 Krad/hr for 100 hours 10 Krad/hr for 1 hour 10 Krad/hr for 1 hour

1x10 9

1/C2 (F/cm2)-2

1/C2 (F/cm2)-2

8x10 8

9

8x10 8

4x10

8

2x10 8

0x10 0

0x10 0 -1.2

-1

-0.8

-0.6

Potential (V SCE)

(a)

-0.4

-0.2

0

-1

-0.8

-0.6

-0.4

-0.2

0

Potential (V SCE)

(b)

Figure 2: Mott-Schottky plots of alloy 625 measured in 5 M NaCl, pH =7 solution after prior exposed to Gamma radiation: (a) in a borate buffer solution and (b) in air.

drop significantly or, in some cases, goes through a maximum. Comparison of Figure 2 with Figure 1 reveals that this region of the M-S plots coincides with the repassivation potential for alloy 625 under the various exposure conditions. It is critical to note that no crevice corrosion was observed on any of the

Electrochemical and Mott-Schottky Behaviour of the Oxide Films

303

samples used to develop the M-S plots. This suggests that the observed changes in capacitance are a result of the oxide properties and NOT a result of localized corrosion initiation. The general appearance of the curves is similar to that observed by others 2,3 in the anodic region of alloy 600 and alloy 690. For the puposes of this paper, it should be noted that the low values of 1/C2 are consitent with the literature 2,3 and most likely result from a high surface area of the porous outer oxide layer19 (see Figure 4.8 in reference 19). Calculations of the flat band potential (corrected for series capacitance 20) yields a value of –0.509 +/- 0.014 V(SCE) for all exposure conditions. Thus, there appears to be no affect of exposure on the flat band potential. Specific adsorption would result in a shift in the flat band potential. It can not be completely ruled out, however, since M-S scans were initiated at high potentials; possibly resulting in the oxidation of any adsorbed species (recall the oxidation peak at –0.160 V(SCE). A slight increase in the charge carrier density from 0.55x1021 to 1.18x1021 cm-3 is observed. This may be a result of an enrichment of Cr (depletion of iron) in the oxide film 16. The Bode phase angle curves at the highest potentials tested for samples exposed to gamma irradiation immersed in the borate solution are shown in Figure 3. A low frequency time constant is clearly apparent for the sample exposed to a dose rate of 1 Mrad/hr. This time constant coinsides with the appearance of an oxidation peak at –160 mV(SCE) in the polarization curves at the same dose rate. -100

Figure 3: Bode phase angle data for alloy 625 exposed to gamma irradiation while immersed in a borate buffer solution: , -0.1 V(SCE), * -0.15 V(SCE), / -0.15 V(SCE), & -0.15 V(SCE)

I625 in 5 M NaCl, pH = 7 no exposure 1 Mrad/hr for 1 hour 10 Krad/hr for 100 hours 10 Krad/hr for 1 hour

Phase Angle

-80

-60

Near the repassivation potential, the space charge layer in the oxide film changes from a depletion layer to an accumulation layer where positive charge carriers can reach the oxide/electrolyte interface. At the 1 Mrad/hour dose rate, it appears that these accumulating positive charges are consumed in the oxidation reaction observed in the polarization scans and as a low frequency time constant in the impedance data. The origin of this oxidation process is not clear but it could be a result of radiation induced structural changes (from induced stress) or compositional changes in the oxide (eg. reduction of species within the film). Another possibility is that radiation induced surface states are re-oxidized. In any event, these processes would lead to trapped holes at the surface and hence a relatively rapid accumulation rate resulting in a decrease in the breakdown potential (dectease in the crevice corrosion nucleation time) of alloy 625. -40

-20

1x10 -2

1x10-1

1x10 0

1x101

Frequency (Hz)

1x10 2

1x103

1x10 4

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4. Conclusions Residual effects of gamma irradiation occur as a result of the combined effect of radiation and solution chemistry. Prior exposure to gamma radiation at 10 Krad/hr increases the breakdown potential by several hundred mV while dose rates of 1 Mrad/hr cause a significant reduction in the breakdown potential compared with samples with no radiation exposure. On the other hand, prior exposure to gamma radiation does not affect the repassivation potential. At the highest dose rates, an oxidation process occurs in the region of the repassivation potential. While the origin of this event is not clear, trapped positive charges (holes) at the surface would lead to a more rapid accumulation of holes resulting in a decrease in crevice nucleation time. 5. References 1. R. S. Glass, G. E. Overturf, R. A. Von Konynenburg, and R. D. McCright, Corrosion Science, Vol. 26 (1986) 577. 2. M. G. S. Ferrreira, M. Da Cunha Belo, N. E. Hakiki, G. Goodlet, M. F. Montemor, and A. M. P. Simoes, J. Braz. Chem. Soc., Vol. 13 (2002) 433. 3. M. F. Montemor, M. G. S. Ferreira, N. E. Hakiki, and M. Da Cunha Belo, Corrosion Science, Vol. 42 (2000) 1635. 4. R. N. Noufi, P. A. Kohl, S. N. Frank, and A. J. Bard, J. Electrochem. Soc., Vol. 125 (1978) 246. 5. A. Michaelis, S. Kudelka, and J. W. Schultze, Electrochim. Acta, Vol. 43 (1998) 119. 6. S. Menezes, R. Haak, G. Hagen, M. and Kendig, J. Electrochem. Soc., Vol. 136 (1989) 1884. 7. A. L. Rudd and C. B. Breslin, J. Electrochem. Soc., Vol. 147 (2000) 1401. 8. P. Schmuki and H. Bohni, Electrochim. Acta, Vol. 40 (1995) 775. 9. D. F. Heaney and D. D. MacDonald, J Electrochem. Soc., Vol. 146 (1999) 1773. 10. S. Lenhart, M. Urquidi-Macdonald, and D. D. Macdonald, Electrochim. Acta, Vol. 32 (1987) 1739. 11. G. P. Marsh, K. J. Taylor, G. Bryan, and S. E. Worthington, Corrosion Science, Vol. 26 (1986) 971. 12. N. E. Hakiki, M. F. Montemor, M. G. S. Ferreira, and M. da Cunha Belo, Corrosion Science, Vol. 42 (2000) 687. 13. S. O. Mousa and M. G. Hocking, Corrosion Science, Vol. 43 (2001) 2037. 14. S. Fujimoto, T. Yamada, and T. Shibata, J. Electrochem. Soc., Vol. 145 (1998) L79. 15. C. B. Breslin, D. D. MacDonald, E. Sikora, and J. Sikora, Electrochim. Acta, Vol. 42 (1997) 137. 16. G. Capobianco, G. Palma, G. Granozzi, and A. Glisenti, Corrosion Science, Vol. 33 (1992) 729. 17. H. G. El-Shobaky, A. El-Mohsen, and M. Turky, Journal of Radioanalytic Chemistry, Vol. 254 (2002) 151. 18. G. Nogami and R Shiratsuchi, J. Electrochem. Soc., Vol 140 (1993), 917. 19. G. Hodes (ed.), Electrochemistry of Nanomaterials, Wiley-VCH, Verlag, (2001) 20. R. De Gryse, O. Gomes, F. Cardon, and J. Vennik, J. Electrochem. Soc., Vol. 122 (1975) 711

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

305

Oxygen Reduction on Passive Steel and Cr Rich Alloys for Concrete Reinforcement Alberto A. Sagüés*, Sannakaisa Virtanen** and Patrik Schmuki**

*Department of Civil and Environmental Engineering, University of South Florida, Tampa FL 33620, U.S.A.** University of Erlangen-Nuremberg, Dept. of Materials Science, WWIV-LKO, Martensstr. 7, D-91058 Erlangen, Germany .

Abstract: Mature passive films (age > 106 s) were grown in aerated saturated calcium hydroxide solution, at the open circuit, on plain steel and alloys with 9% and 22% Cr, as well as austenitic stainless steels. The materials were commercially produced concrete reinforcing steels. Differences in the polarization of the oxygen reduction reaction were evidenced. Oxygen reduction rates at moderate cathodic polarization were found to approximate ideal Tafel behavior, to be greatest for plain steel and to decrease with increasing Cr content. Mott-Schottky analysis indicated n-type semiconductor behavior in all cases, with higher apparent donor density for the plain steel and little differentiation between the other materials. Keywords: Oxygen reduction, chromium, passive, capacitance, semiconductor 1. Introduction Steel reinforcement in concrete normally remains passive due to contact with highly alkaline pore water (pH>~12.5). Chloride ion contamination above a threshold value can induce localized reinforcement corrosion, where much of the cathodic reaction is oxygen reduction on surrounding passive steel surfaces [1]. Stainless and other Cr-containing steels have a higher corrosion initiation threshold and are being increasingly used instead of plain steel in concrete subject to aggressive service conditions [2]. An added benefit of using Cr steels is that, for a given cathodic polarization potential, the rate of the oxygen reduction reaction can be significantly lower than for plain steel [3] with

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consequent lessening of severity of localized corrosion should it eventually start. However, little information is available on the mechanism responsible for this decrease and on its dependence on alloying content. It has been speculated [3] that the decreased cathodic reaction rate reflects lowered conductivity of the passive film but the supporting experimental evidence is limited. To address these issues the electrochemical behavior of the passive films in alkaline solution was examined here for series of commercial reinforcing steel alloys spanning a wide composition range. 2. Experimental Materials, test solution and electrochemical procedures Commercial corrugated reinforcing steel stock was used, turned to 9mm diameter. The alloys were ASTM A-615 plain steel (PS), a proprietary alloy with ~9 wt% Cr (9 Cr), a duplex, low Mo ferritic-austenitic stainless steel (2201), and two 316L stainless steel (316-1 and 316-2) (Table 1). Cross section slices of each alloy were embedded together but mutually isolated in a 40 mm diameter metallographic epoxide holder and ground to an F-4000 grit finish. Immediately after surface grinding the assembly was ultrasonically cleaned in methanol, dried, and immersed in saturated Ca(OH)2 solution (pH~12.6) aerated with decarbonated air. The sensing point of the reference electrode was placed about 6 mm away from the center of the specimen cluster. The sensing tip was either a Luggin probe extended from an Ag/AgCl (1M KCl) reference electrode (RE) -in which scale all potentials are reported in this paper- or alternatively a 1-cm long gold wire periodically calibrated against the RE. The counter electrode was a 25mmx25mm platinum sheet. Temperature was ~22o ± 3o C. The steel surfaces were allowed to evolve at the open circuit potential (OCP) except for electrochemical evaluations conducted after ~1 wk and ~4 wk immersion. Evaluation included OCP electrochemical impedance spectroscopy (EIS), cyclic cathodic (starting at OCP) potentiodynamic polarization at a scan rate of 0.167 mV/sec, and apparent capacitance measurements (simply evaluated from the imaginary component of the impedance) at 10 Hz and 1 kHz also during cathodic excursions from OCP. Cathodic excursions extended no further than -0.5 V, to represent the potential range normally experienced in atmospherically exposed reinforced concrete. Following the tests at ~4wk immersion, the test cell was temporarily deaerated with flowing N2 and the tests repeated in the deaerated condition. All tests were at least in duplicate and results shown are typical of replicate behavior.

Oxygen Reduction on Passive Steel and Cr Rich Alloys for Concrete Reinforcement 307 Table 1 Chemical composition of alloys used (wt%) Alloy

Cr

Ni

Mo

C

S

P

Mn

Si

Cu

PS 9 Cr 2201 316-1 316-2

0.14 9.3 21.57 17.57 16.17

0.09 0.1 1.74 10.22 10.24

0.017 0.03 0.24 2.08 2.15

0.41 0.05 0.03 0.03 0.03

0.054 0.015 0.001 0.002 0.007

0.017 0.012 0.019 0.027 0.029

0.85 0.45 4.73 1.57 1.62

0.18 0.23 0.78 0.57 0.4

0.4 0.12 0.34 0.36 0.75

3. Results and Discussion All alloys showed significant ennoblement during the first few days of exposure, indicating slow maturing of the passive film. The PS potential evolved toward terminal values in the order of -0.1V, typical of those normally observed in concrete and aerated simulated pore solution [1]. The stainless steels evolved toward terminal potentials about 0.1V or 0.2 V lower, also as commonly observed. The 9 Cr potentials were comparable to those of the stainless steels. Terminal potentials of replicate specimens were in the same order but showed some variability ascribed to incipient crevice corrosion at the perimeter in contact with the embedding epoxide. The potentiodynamic tests showed essentially the same behavior (Figure 1) after ~ 1wk or ~ 4wk exposure to the aerated solution. For cathodic polarization beyond ~0.1 V from the OCP all alloys displayed approximately linear (less so for PS) E-log i behavior with apparent Tafel slope ~0.12 V for PS and 9 Cr and ~0.16 V for the stainless steels. There was relatively little hysteresis in all cases in that polarization regime. The greater deviation from Tafel behavior for PS, which had by far the greatest cathodic current densities (~ 20-30 mA/cm2 at -0.5 V, Figure 2) may be an indication of the onset of concentration polarization. The 9 Cr had intermediate cathodic currents densities while the stainless steels had as a group very low and comparable cathodic current density. While Cr content, appears to be the main factor (Figure 2), a somewhat more uniform decreasing trend in cathodic current density is obtained when plotting against Cr+Ni, or total alloying content. Upon deaeration, all current densities decreased by 1-2 orders of magnitude. All currents returned to the former levels upon reaeration, indicating that the cathodic behavior under aeration reflected primarily oxygen reduction and not reduction of species in the passive film. This observation is in keeping with expected behavior based on previous work [4] that showed at pH~13 little reduction of passive films on iron at potentials > ~ -0.7 V.

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EIS measurements at the OCP showed a well-defined real high frequency impedance limit which was consistent with the expected value of the solution resistance Rs (~ 60 W with the cell configuration used). Results were comparable in the 1wk and 4wk tests. At < 100 Hz the impedance could be closely described (after subtracting a nominal constant close to Rs) by a parallel combination of a large polarization resistance and a constant phase angle element (CPE, Z=Yo-1 (jw)-n where w is the angular frequency) that was nearly ideally capacitive (n= 0.95-0.96). The spectrum above 100 Hz could be approximated using the same equivalent circuit, but with n~0.7. The greater frequency dispersion at higher frequencies likely reflects microscopic and macroscopic current distribution artifacts. Thus the Mott-Schottky analyses were reported only using the apparent capacitance C=1/(wIm(Z)) calculated for 10 Hz where capacitive behavior was nearly ideal. It is noted that no systematic difference between the Rs values for PS and for the high Cr alloys was observed. The EIS behavior observed here is at variance with that described for similar systems by Abreu et al [3]. Those authors reported a high frequency loop with a resistive component (interpreted as a film resistance, but that was of the same order as the expected Rs) that was greater for stainless steel than for PS. -0.10 PA1

-0.15 -0.20

E/V

-0.25 -0.30

PS -0.35

316-2

9 Cr

-0.40

316-1

2201

-0.45 -0.50 1.E-09

1.E-08

1.E-07

1.E-06

1.E-05

1.E-04

Current Density / mA cm-2

Fig. 1 Example of cyclic cathodic polarization curves. 1wk in aerated Ca(OH)2.

The Mott-Schottky plots (Figure 3) showed usually linear regions with slope indicative of n-type behavior and a common apparent flatband potential of ~-0.7 V for all alloys. Apparent donor densities Nd were calculated from the slopes assuming a dielectric constant = 10 with results shown in Figure 2. Artifacts may exist from neglecting Helmholtz interfacial capacitance since total

Oxygen Reduction on Passive Steel and Cr Rich Alloys for Concrete Reinforcement 309

1.E-04

1.0E+23

1.E-05

1.0E+22

1.E-06

1.0E+21

Nd / cm -3

i c / A cm-2

capacitance is relatively large (in the order of 20-30 mF/cm2 at the OCP), and from the very narrow (~0.5 nm) depletion zone thickness calculated from the same data. Nevertheless, the resulting Nd values although large are consistent with those observed by other investigators [5]. It should be noted also that actual capacitances and donor densities may be somewhat lower than the nominal values obtained, since the calculations ignored the exaggerating effect of surface roughness.

Ic Nd 1.E-07

1.0E+20 0

5

10

15

20

25

% Cr

Figure 2. Cathodic current density ic at -0.5 V, and apparent donor density Nd from 10 Hz MottSchottky analysis as function of Cr content. 1wk in aerated Ca(OH)2. 4.E+09

4 F-2 C^-2 / cm^4-F^-2 C-2 /cm

3.E+09

PA 1

PS 9 Cr 2201 316-2 316-1

2.E+09

1.E+09

0.E+00 -0.8

-0.6

-0.4

-0.2

E/V

Figure 3. Mott-Schottky analysis. The solid lines indicate slope and extrapolation based on data starting ~ 50 mV below OCP. 1wk in aerated Ca(OH)2.

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In contrast with the large variations in cathodic polarizability, Nd was greater by only about a factor of 2 for PS than for the Cr alloys. There is almost no differentiation in Nd between the Cr alloys themselves, and no dramatic difference in the apparent flatband potential value of any of the alloys. The higher currents densities in PS may reflect the thinner depletion layer (and associated nonlinear increase in electronic current) consistent with the noted differences in donor density. It may also be speculated that in the Cr alloys an outer n-type layer rich in Fe oxides is present, coupled with an inner Cr-oxide rich layer that would be expected to have p-character [6]. Such configuration may possibly require additional polarization (with respect to a plain steel case) to achieve a given current level. Investigation of these issues continues. 4. Summary Oxygen reduction rates at moderate cathodic polarization were found to approximate ideal Tafel behavior, to be greatest for plain steel and to decrease with increasing Cr content. The results indicate that the cathodic current under naturally aerated conditions stemmed from oxygen reduction with little if any contribution from the reduction of oxide formed at potentials typical of steel in atmospherically exposed concrete. Mott-Schottky analysis indicated n-type semiconductor behavior in all cases, with higher apparent donor density for the plain steel and little differentiation between the other materials. Acknowledgements This investigation was made possible by sabbatical leave support for A.A.Sagüés by the University of South Florida, and by the facility and personnel contributions by the University of Erlangen-Nuremberg. References 1. A.A. Sagüés, M.A. Pech-Canul, A.K.M. Al-Mansur, Corrosion Science, Vol. 45, No.1 (2003), pp. 7-32 2. L. Bertolini, F. Bolzoni, T. Pastore and P. Pedeferri, British Corrosion Journal, Vol. 31, No.3 (1996), pp.218-222 3. C.M. Abreu, M.J. Cristóbal, M.F. Montemor, X.R. Nóvoa, G. Pena, M.C. Pérez, Electrochimica Acta, Vol. 47 (2002), pp. 2271-2279 4. P. Schmuki, M. Büchler, S. Virtanen, H.S. Isaacs, M.P. Ryan and H. Bohni, J. Electrochem. Soc., Vol. 146, No. 6 (1999) pp. 2097-2102 5. M. Büchler, "Experimental Modeling of Passive Films on Iron" Doctoral Dissertation Diss. ETH No. 12504, Swiss Federal Institute Of Technology (1998) 6. N. Hakiki, Da Cunha Belo, A. Simoes and M. Ferreira, J. Electrochem. Soc., Vol. 145, No. 11 (1998) pp. 3821-3829

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

311

Diffusivity of point defects in the passive film on Fe SeJin Ahn and HyukSang Kwon Department of Materials Science and Engineering, Korea Advanced Institute of cience and Technology, KuSongDong 373-1, YuSongGu, DaeJon 305-701, Korea [email protected]

Abstract - Diffusivity of point defects (D0) in the passive film formed on Fe in deaerated pH 8.5 buffer solution at ambient temperature was estimated by the Mott-Schottky analysis based on the Point Defect Model and surface charge approach assuming that donors are oxygen vacancies and/or iron interstitials. From the exponential decay of the concentration of donors with film 20 2 -1 formation potential, D0 was calculated to be 1.69 10 cm s . Keywords: Diffusivity; Point defects; Passive film; Fe; Mott-Schottky plot

1. Introduction The growth and breakdown of passive films on metal and alloy surface has been a subject of intense interest because these films are generally considered to be responsible for the corrosion resistance of metals and alloys in contact with reactive environments. Even though there is still some controversy among the models and theories on the phenomena of passivity, it is now generally agreed that passive film is a highly defective oxide film with the point defects being metal and oxygen vacancies and metal interstitial. Especially, the transport behavior of these point defects under a high electric field existed in the passive film is of great interest in describing the growth and breakdown kinetics of the film. The transport properties of point defects can be expressed quantitatively by their diffusivity (D0). However, there are few reports on the defects diffusivity in the passive film. The Point Defect Model (PDM), one of the proposed models on passivity, provides an analytical description of the growth and breakdown of passive films on reactive metal surface [1-3]. The quantitative analysis of the PDM on the concentration and the transport properties of point defects enable us to estimate the diffusivity of point defects (DO) in the passive film, which appears in the Nernst-Planck equation [4]. On the other hand, Bojinov [5] developed a surface

312

S. J. Ahn and H. S. Kwon

charge approach to describe the growth of passive film on metals, which is basically based on the idea of the PDM. However, in his calculation of diffusivity of oxygen vacancy in the passive film on tungsten, he used more general transports equation rather than Nernst-Planck equation, a low field limited diffusion-migration law, used in [4,6]. Because of very high field strength in the passive film (~106 Vcm-1), it is apparent that the approach of Bojinov [5] is more appropriate to calculate DO in the passive film on metals. The objective of this work is to determine DO in the passive film on Fe based on the combination of the PDM and the surface charge approach. 2. Calculation of Do in the passive film on Fe The passive film of Fe is considered to have spinel structure of -Fe2 O3 and/or Fe3 O4, or related structure [7, 8-10], which exhibit n-type semiconductivity with high density (1020 ~ 1021 cm-3) of donor [10-12]. The dominant point defects in the passive film on Fe in this study are considered to be oxygen vacancies and/or iron interstitials acting as electron donors [2, 3]. However, as it is not possible to separate the contribution from oxygen vacancies and iron interstitials on the measured diffusivity value based on the PDM, the diffusivity is considered to be that for the combination of these two point defects. According to the PDM, the donor density, i.e. the concentration of oxygen vacancy and/or iron interstitial, in the passive film depends on the film formation potential expressed by Eq. (1) [4];

ND = w1 exp(-bVff) + w2

(1)

where ND is the donor density in the passive film which can be measured by the Mott-Schottky analysis [4], w1, w2 and b are unknown constants, Vff is the film formation potential. Bojinov [5] proposed that diffusivity of oxygen vacancy in the passive film can be deduced by Eq. (2) with a high field approximation which is appropriate to the case of passive film. DO =

2aJ O w2 exp(2aK )

(2)

where JO is the steady state flux of donors, and K = F /RT, F is the Faraday constant, is the mean electric field strength, R is the gas constant, T is temperature in Kelvin, and a is the half-jump distance. At this point, from Eq. (1) and Eq. (2), we have totally four unknown parameters, i.e. JO, , a and w2, those are to be determined experimentally.

313

Diffusivity of point defects in the passive film on Fe

The steady state flux of point defects (JO) is expressed in terms of steady state current density through the passive film as shown in Eq. (3), regarding that the current flow is mainly due to the flux of oxygen vacancy and/or iron interstitial as discussed above. JO =

iSS 2e

(3)

where iss is the steady state current density through the passive film, e is the charge of an electron. The PDM postulates that the steady state film thickness (LSS) is related to the film formation potential (Vff) and by Eq. (4) [1,2]; LSS

1

(1

(4)

)V ff + B

where is the polarizability of the film/solution interface and B is a constant. To obtain half-jump distance (a), we followed Bojinov’s surface charge approach [5]. In Fig. 1, each circuit elements are defined as follows; Cb is the barrier film capacitance, Rb is the resistance of defect migration, Rsc, Lsc are elements associated with the surface charge at the film/solution interface, Co is the faradaic pseudocapacitance and R is the solution resistance. Among these parameters, Bojinov [5] reported that Rb/Rsc is expressed in terms of as Eq. (4); Rsc

Lsc

Cb

Rb/RSC = /1-

Co

R

(5) Rb

Fig. 1. Equivalent circuit of the

In addition, Rb is dependent on Vff with relationship metal/film/solution interface as Eq (9) containing a; according to the surface charge iSS dRb /dVff = (2Fa /RT)-1

(6)

approach. Circuit elements: Cb – barrier film capacitance, Rb – resistance of defect migration, RSC, LSC – elements associated to the surface charge at the film/solution interface, C0 – faradaic pseudocapacitance, R – solution resistance [10].

In fact, in deriving Eq. (6), it was assumed that iSS is independent on the film formation potential (Vff). This point will be discussed further later. It is apparent that we can determine , and a at a time from Eq. (4), (8) and (9). The value of w2 can be obtained easily by fitting donor density vs. film formation potential curves obtained by Mott-Schottky analysis to Eq. (1). In summary, DO can be expressed in the following final form in which all the terms in the right side are measurable or constants.

314 DO =

S. J. Ahn and H. S. Kwon iSS a w2 e exp(2a F / RT )

(7)

3. Experimental methods A conventional three-electrode cell of 1-L multi neck flask was used to conduct electrochemical tests for the passive film on Fe. A platinum counter electrode and a saturated calomel reference electrode (SCE) were used in the cell. Pure Fe (99.99 %) was used as a working electrode. The working electrode was mounted in an epoxy resin with an exposed area of about 1 cm2. The solution used in this study was a deaerated pH 8.5 buffer solution. The working electrode was cathodically polarized at -1.0 VSCE for 10 min. to remove an airformed oxide on the surface, and then a passive film was grown potentiodynamically by ramping the applied potential from Ecorr to film formation potential (0.1~0.9 VSCE) at a sweep rate of 0.5 mV/s, and then potentiostatically at the film formation potential for 24 h before initiating the measurements of interest. The ac impedance measurements were performed on the passive film formed on Fe for 24 h at various film formation potentials in the frequency range of 100,000 ~ 0.01 Hz at an ac amplitude of 10 mV (r.m.s). Capacitance for the film was measured at a fixed frequency of 1000 Hz using an excitation voltage of 10 mV (peak-to-peak) with a scan rate of 1 mV/s, and then used for the MottSchottky plot. All the experiments were performed at an ambient temperature. 4. Results and discussion Anodic polarization curve of Fe in deaerated pH 8.5 buffer solution at an ambient temperature is shown in Fig. 2 in which applied potentials for the formation of stable passive film were chosen to be ranged from 0.1 to 0.9 VSCE. To get a steady state current density through the passive film on Fe, potentiostatic polarization tests were performed. Evidently, the current reached a steady state value (iSS) after 24 h for all the film formation potentials. The values of iSS were independent of film formation potential, and the average value was estimated to be 0.02607 µAcm-2. To determine the thickness of passive film at a steady state, galvanostatic reduction experiments were performed on the passive films formed for 24 h at various film formation potentials (Fig. 3). The Fig. 2. Anodic polarization curve thickness of passive films was calculated for - of Fe in deaerated pH 8.5 buffer solution at ambient temperature Fe2 O3 according to the following reduction equation with a scan rate of 0.5 mV/s. with a 100 % current efficiency and literature data potential / VSCE

1.5

pure Fe deaerated pH 8.5 buffer solution

1.0 0.5

passive range

0.0

-0.5 -1.0

10-7

10-6

10-5

10-4

current density / Acm-2

10-3

315

Diffusivity of point defects in the passive film on Fe

(MFe2O3=159.69 gmol-1,

-Fe2O3=4.90

gcm-3).

0.2

pure Fe deaerated pH 8.5 buffer solution

+

-

2+

Fe2 O3 + 6H + 2e Ç 2Fe + 3H2 O

potential / VSCE

0.0

(8)

tpassivation = 24 h

-0.2

iapplied = -10 µAcm-2

-0.4 -0.6

The thickness of passive film calculated using the results in Fig. 3 were showed good linear relationship with film formation potentials, and (1 ) / is estimated to be 2.50 nmV-1 by Eq. (4). The next step is to extract circuit elements in Fig. 1, and then calculate , and a using Eq. (5) and (9). Among those circuit elements, CO can be obtained by linear sweep voltammograms of Fe with different scan rates, since one can write [5]:

dipl/dv = CO

0.9 VSCE

0.1 VSCE

-0.8 -1.0 0.000

0.001

0.002

0.003

reduction charge / Ccm-2

Fig. 3. Effects of film formation potential on the galvanostatic reduction charge with -10 µAcm-2 for the passive film formed on Fe for 24 h in deaerated pH 8.5 buffer solution. The dotted line indicates the evaluation of the reduction charge.

(9)

800000

pure Fe deaerated pH 8.5 buffer solution 600000 t = 24 hours passivation

0.01 Hz

Vff = 0.1 VSCE

where ipl is the plateau current in linear sweep voltammogram and v is the scan rate. From the linear sweep voltammograms, CO was calculated to be 6.37 10-3 Fcm-2. Fig. 4 shows the impedance spectra for the passive Fig. 4. Nyquist plots for the films formed on Fe for 24 at 0.1 VSCE. Those for the passive film formed on Fe film formed at 0.2 ~ 0.9 VSCE were similar to it in for 24 h in deaerated pH 8.5 principle. RSC, LSC, Cb, and Rb were extracted from buffer solution at 0.1 VSCE. the impedance spectra. Average Rb/RSC was determined to be 0.63. We obtained =0.39 using Eq. (5), which in turn provided the electric field strength in the film, , of 2.44 106 Vcm-1. Finally, knowledge of and yields half jump distance a = 0.4 nm by Eq. (6). The last parameter to be determined is w2, which appeared in Eq. (1) and (3). Fig. 5(a) shows the Mott-Schottky plots for the passive film formed on Fe for 24 h at various film formation potentials. The donor density (ND) calculated from the linear slopes of the Mott-Schottky plots using =12 [13] is presented in Fig. 5(a) Mott-Schottky plots for the passive Fig. 5(b). It is evident from Fig. 5(b) that film formed on Fe for 24 h at various film ND decreases exponentially with film formation potentials at a fixed frequency of formation potential (Vff). The value of w2 1000 Hz. (b) ND as a function of film was determined to be 2.01 1020 cm-3 formation potential. The donor density of from Fig. 5 (b). With all the parameters passive film was calculated from the linear of iSS, a, w2 and , which have been slope of the Mott-Schottky plot in Fig. 5(a). -Z" /Ωcm

2

400000

200000

experimental simulation

0.1 Hz

0

0

50000 100000 150000 200000 250000 300000

Z' / cm2

2 -4 -2

C / F cm

0.8 VSCE

tpassivation = 24 h

0.7 VSCE

0.6 VSCE

10

3x10

2x1010 1x1010

0 -0.6 -0.4 -0.2 0.0

5x1020

0.9 VSCE

0.5 VSCE 0.4 VSCE 0.3VSCE 0.2 VSCE 0.1 VSCE 0.2

0.4

potential / VSCE

0.6

0.8

-3

4x1010

(a) pure Fe deareated pH 8.5 buffer solution

D / cm N

5x1010

(b)

ND = 4.73x1020exp(-5.74Vff)

+ 2.01x1020

4x1020

3x1020

2x1020

1.0

0.0

0.2

0.4

0.6

film formation potential / VSCE

0.8

1.0

S. J. Ahn and H. S. Kwon

316

found by above procedure, DO is calculated 1.69 10-20 cm2s-1 by Eq. (7). It is apparent that D0 in this study is 2«5 orders of magnitude lower than those reported in the literatures for Fe [14] and carbon steels [6, 15]. The authors think that there are two main reasons for the difference. First, it may be due to that they [6] used low field approximation to calculate diffusivity of point defects. And the other reason is presumably the different film formation time [6, 14, 15]: The short passivation times (0.25 ~ 4 h), evidently, result in the overestimation of diffusivity because the current flow through the passive film at this stage is much higher than that at the steady state. Thus, the authors claim that the relatively low D0 value obtained in this work (~10-20 cm2s-1) is more adequate to represent the steady state D0 in the passive film on Fe than previously reported values. 5. Conclusions Potentiostatic polarization tests revealed that the steady state current density (iSS) through the passive film formed on Fe for 24 h was independent of film formation potential (Vff). The steady state thickness (LSS) of passive film on Fe varied linearly with film formation potential (Vff). The donor density determined from the Mott-Schottky plots for the passive film on Fe was found to decrease exponentially with increasing film formation potential (Vff) according to the following first order exponential decay function. The diffusivity of point defects (DO) in the passive film on Fe was calculated to be 1.69 10-20 cm2s-1. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

D. D. Macdonald, J. Electrochem. Soc., 139 (1992) 3434. D. D. Macdonald, S. R. Biaggio, and H. Song, J. Electrochem. Soc., 139 (1992) 170. J. Liu and D. D. Macdonald, J. Electrochem. Soc., 148 (2001) B425. E. Sikora, J. Sikora and D. D. Macdonald, Electrochim. Acta, 41 (1996) 783. M. Bojinov, Electrochim. Acta, 42 (1997) 3489. Y. F. Cheng, C. Yang and J. L. Luo, Thin Solid Films, 416 (2002) 169. I. Iitaka, S. Miyake, and T. Iimori, Nature, 139 (1937) 156. M. P. Ryan, R. C. Newman, and G. E. Thompson, J. Electrochem. Soc., 142 (1995) L177. A. J. Davenport, L. J. Oblonsky, M. P. Ryan, and M. F. Toney, J. Electrochem. Soc., 147 (2000) 2162. J. S. Kim, E. A. Cho, and H. S. Kwon, Corrosion Science, 43 (2001) 1403. Stimming U and Schultze J. W., Electrochim. Acta., 24 (1979) 859. K. Azumi, T. Ohtsuka, and N. Sato, J. Electrochem. Soc., 134 (1987) 1352. N. E. Hakiki and M. Da Cunha Belo, J. Electrochem. Soc., 143 (1995) 3088. M. Bojinov, T. Laitinen, K. Makela, and T. Saario, J. Electrochem. Soc., 148 (2001) B243. X.P. Guo, Y. Tomoe, H. Imaizumi, and K. Katoh, J. Electroanal. Chem., 445 (1998) 95.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

317

Electrochemical Impedance of Thin Rust Film Fabricated Artificially M. Itagakia,*, H. Arakia, K. Watanabea, H. Katayamab, and K. Nodac a

Department of Pure and Applied Chemistry, Faculty of Science and Technology, Tokyo University of Science, Noda, Chiba 278-8510, Japan. b Corrosion Research Group, National Institute for Materials Science, Sengen, Tsukuba, Ibaraki 305-0047, Japan. c Department of Materials Science and Technology, Faculty of Engineering, Shibaura Institute of Technology, Tokyo 108-8548, Japan.

Abstract - Thin rust films were fabricated by the oxide precipitation on the cellulose membrane set at the interface between NaOH and Fe(NO3)3 solutions. The rust film was identified by X-ray diffraction, and mainly composed of a-FeOOH. The rust film showed anion selective permeability when it was immersed in NaCl solution. The electrochemical impedance of the thin rust film membrane was measured by using a four-electrode method. The electrochemical impedance showed a capacitive loop on the complex plane. The ion permeation resistance decreased with the increase of NaCl concentration in the electrolyte solution where the membrane was immersed. The ion permeation resistance was also varied after the immersion in the solution containing Mo(VI). Keywords: rust film, ion-selective permeability, membrane potential, electrochemical impedance, film resistance

1. Introduction The corrosion resistance of the weathering steel depends on the protective nature of the rust layer formed on the surface. The ion-selective permeability, which is one of significant properties of the protective rust layer, has been investigated by artificial rust film formed by oxide precipitate1, 2) and real rust

318

M. Itagaki et al.

film formed by wet/dry cycles3-5). Sakashita et al.2) reported that hydrous Fe(III) oxide membrane, which was made by the precipitation form the solution, had cation selective permeability and that the adsorption of MoO42- on the surface of hydrous Fe(III) oxide led to the anion selective permeability. On the other hand, Noda et al.3-5) measured the membrane potential of the rust films generated from the steel sheets by wet/dry cycles and estimated their ion-selective permeability. They3-5) stated that the rust film of Fe-Co low alloy steel was effective for decreasing the corrosion rate because of its cation selective permeability, and summarized the relation between the ion-selective permeability and the corrosion rate. Electrochemical impedance spectroscopy (EIS) is useful electrochemical method to investigate the metal corrosion because the time constants for charge transfer and the surface film can be determined. Kihira et al.6) measured impedance spectra of the weathering steel covered with the rust film. They6) discriminated the film resistance from the whole impedance spectrum, and proposed new corrosion monitoring method on the basis of the rust film resistance. Nishimura et al.7-9) investigated the rust film formation on the lowalloy steels under wet/dry cycles by EIS, and reported that the rust impedance decreased and its capacitance increased with the cycle. Recently Itagaki et al.10) determined the impedance spectra of rust film membrane formed by wet/dry cycles. In this paper10), it was found that the rust film membrane showed an apparent capacitive behavior and that the resistance of the rust film of Fe-0.6 %Mo was quite large. In the present paper, the electrochemical impedance of artificial rust film is determined in order to investigate the electrochemical nature of rust film. Furthermore, the effect of MoO42- adsorption on the film resistance is discussed since it can be expected that the reciprocal of the film resistance corresponds to the permeation rate of the ions in the rust film. 2. Experimental 2.1. Fabrication of the rust film The artificial rust film was fabricated by referring the report by Sakashita et al.2). The cellulose membrane was fixed between two vessels that have a circular mouth. One vessel contained 0.1 mol/dm3 Fe(NO3)3 solution and another one contained 0.1 mol/dm3 NaOH. The diameter of the circular mouth were 3 cm, thus the both sides of the cellulose membrane were contacted with the solutions at the interfacial area of 7 cm2. The cellulose membrane was immersed between two solutions for 3 days, and the rust film was deposited on the side contacted with Fe(NO3)3 solution. Consequently, the rust film was deposited on another side for 5 days after replacing the vessels. The uniform

Electrochemical Impedance of Thin Rust Film Fabricated Artificially

319

rust was deposited on the cellulose membrane. The thickness of the rust film involving the cellulose membrane was from 0.5 mm to 0.9 mm. The rust film membrane was immersed in 0.1 mol/dm3 Na2MoO4 solution for one hour. MoO42- was adsorbed on both sides of rust film membrane. The film was rinsed by distilled water after the immersion. 2.2. Measurement of membrane potential The two-electrode method to measure the membrane potential is described in the reference 3. The rust film was sandwiched by two acrylic sheets, which have hole (diameter: 11 mm ) at the center. These sheets were fixed between two cells. The reference electrodes were saturated KCl/AgCl/Ag electrode (SSE). The electrolyte in one cell was 0.01 mol/dm3 NaCl aqueous solution, and the NaCl concentration of another electrolyte solution was from 0.01 mol/dm3 to 0.1 mol/dm3. The electrolyte solution used in the present paper was prepared by doubly distilled water and analytical grade reagents. The electrometer equipped with the potentiostat (Hokuto, HA1010mM1A) was used to measure the potential difference between two reference electrodes. All measurements in the present paper were performed at room temperature (298 K). 2.3. Electrochemical impedance of thin rust film membrane The measurements of electrochemical impedance of rust film membrane were performed by four-electrode method 10). The rust film was fixed between two cells by the same arrangement as the measurement of membrane potential. The electrolyte solutions contained 0.01 mol/dm3 or 0.1 mol/dm3 NaCl. The working and reference electrodes were platinum wire and SSE, respectively. The LCR meter (Hioki, 3522-50) was used to measure the impedance spectra. The current flowed between two working electrode was measured by controlling the potential between two reference electrodes. The amplitude of the potential modulation of the working electrode was 10 mV and the frequency range was from 10 mHz to 10 kHz. 3. Results and discussion 3.1. Membrane potential of the rust film The composition of the rust film was analyzed by X-ray diffraction. The diffraction pattern was compared with ASTM cards for a-FeOOH, b-FeOOH, g-FeOOH and Fe3O4. It was found that the rust film was mainly composed of aFeOOH.

M. Itagaki et al.

320

Membrane potential / mV

Figure 1 shows the plots of membrane potential and the ratio of concentrations of NaCl in two cells. In Fig. 1, the dotted lines that have positive (60 mV/decade) and negative (-60 mV/decade) slopes mean the perfect cation and anion-selective permeability, respectively. The plots of the rust film that are denoted by solid circle show negative slope, indicating that the rust film has anion-selective permeability. On the other hand, the plots represented by open circle show positive slope in the case of MoO42- adsorption. These results indicate that the adsorption of MoO42- changes the property of the rust film into cation-selective permeability, and agree with the report by Sakashita et al.2). Since the fixed charge inside the a-FeOOH has positive sign, the anions can permeate though the film. MoO42- adsorbs on the positive charge site and varies the fixed charge to negative sign. 100 80 60 40 20 0 -20 -40 -60 -80 -100 -1.5

Fig. 1 Plots of membrane potential and the ratio of concentrations of NaCl in two cells. Open and solid circles denote the plots for rust film with and without MoO42adsorption, respectively. -1

-0.5

0

0.5

1

1.5

log (c2 / c1) cation i

cation

anion Fe2珽

i

anion Fe2珽

Fe

Fe

anion–selective permeability

cation–selective permeability

Fig. 2 Schemes of steel covered with two kinds of rust films

Figure 2 shows the schemes of steel covered with two kinds of rust films. If the steel covered with rust film becomes anodic site in the localized cell, the current flows from steel surface to film surface. When the steel is surrounded by the corrosive atmosphere containing NaCl, the migration of chloride ions maintains the current through the anion-selective film. As the result, the chloride ions may concentrate at the interface between the rust film and steel, and accelerate the iron dissolution. On the other hand, sodium ions can't contribute to the current flow through the cation-selective film at the anodic site. Therefore, the migration of ferrous or ferric ions must maintain the anodic

Electrochemical Impedance of Thin Rust Film Fabricated Artificially

321

current. The above discussion indicates that the cation-selective rust film has more protective property than the anion-selective film. 3.2. Electrochemical impedance of the rust film membrane Many researchers measured the impedance spectra of steel covered with rust film6-9). Since the metal/solution interface has a complicated structure, the measurement of the electrochemical impedance of the rust film itself should be effective for the general discussion on the properties of the rust film. In the present paper, the electrochemical impedance of the artificial rust film was measured by four-electrode method. Figure 3 shows Nyquist plots of the electrochemical impedance of the rust film membrane in 0.01 mol/dm3 and 0.1 mol/dm3 NaCl solutions. The thickness of the film was 0.53 mm. Both results show the capacitive behavior. The time constant of the capacitive semicircle is composed of the film resistance Rf and the film capacitance Cf. It is clearly shown that the impedance in 0.01 mol/dm3 is larger than that in 0.1 mol/dm3. The plots at low frequency range diverge from a true semicircle due to the low S/N ratio when the electrolyte is 0.01 mol/dm3 NaCl solution. The simple equivalent circuit to represent the experimental results is shown at the rightupper in Fig.3, where Rsol means the solution resistance. The curve-fitting for the electrochemical impedance was performed by this equivalent circuit. The following parameters were obtained by the curve-fitting using Zview software (Solartron). Rsol: 171 ohm, Rf: 1275 ohm, and Cf: 5.0x10-3 F for 0.1 mol/dm3 NaCl solution. Rsol: 908 ohm, Rf: 7900 ohm, and Cf: 4.7x10-4 F for 0.01 mol/dm3 NaCl solution. The Rsol decreases with the increase of NaCl concentration since the cell arrangement is identical. Moreover, the Rf also decreases with the increase of NaCl concentration. This result means that the water is absorbed into the rust film and ions involved in the water take a charge as electric carrier. The Cf takes the order of mF approximately. This result agrees with the huge capacitance reported in the literatures9, 10). Figure 4 shows the Nyquist plots of electrochemical impedance of the rust film membranes with and without MoO42- adsorption. The electrolyte was 0.1 mol/dm3 NaCl solution. The thicknesses of the rust film with and without MoO42- adsorption were 0.78 mm and 0.88 mm, respectively. It was found that the MoO42- adsorption increased the diameter of the capacitance loop on the complex plane. The parameters obtained by curve-fitting are as follows. Rsol: 186 ohm, Rf: 2014 ohm, and Cf: 2.0x10-3 F without MoO42- adsorption. Rsol: 207 ohm, Rf: 2619 ohm, and Cf: 1.4x10-3 with MoO42- adsorption. The Rf increases remarkably by the adsorption of MoO42-, indicating that the presence of MoO42- gives the low migration rate of ions in the rust film. Itagaki et al.10) reported that the Rf of the rust film from Fe-Mo low-alloy steel is larger than Rf of pure Fe. These facts suggest that the cation selective permeability and low permeation rate of chloride ions decrease the corrosion rate of the steel in

M. Itagaki et al.

322

the presence of Mo. The Rf without MoO42- adsorption in Fig. 4 is larger than the Rf in 0.1 mol/dm3 NaCl solution in Fig. 3, where the experimental conditions except for the film thickness are identical. The value of Rf is in proportion to the film thickness. This result supports the above-mentioned discussion that the ions involved in the water take a charge as electric carrier inside the rust film. -8000 -7000

Rsol

Rf

-6000 Im{Z}/

-5000

Cf

-4000 -3000 -2000 -1000

Fig. 3 Nyquist plots of electrochemical impedance of rust film membrane in 0.01 mol dm-3 (open circle) and 0.1 mol dm-3 (solid circle) NaCl solutions.

0 0

2000 4000 6000 8000 10000 12000 Re{Z}/

-2500 -2000

Fig. 4 Nyquist plots of electrochemical impedance of rust film membrane in 0.1 mol dm-3 NaCl solution. Open and solid circles denote the plots for rust film with and without MoO42adsorption, respectively..

Im{Z}/

-1500 -1000 -500 0 0

500 1000 1500 2000 2500 3000 3500 Re{Z}/

4. Summary The impedance spectra of the rust film fabricated artificially were measured, and the following results were obtained. The impedance of the rust film shows capacitive behavior. The film resistance increases by the decrease of NaCl concentration and the increase of film thickness, since the migration of ions incorporated in the rust film with water controls the film resistance. The immersion of the rust film in Na2MoO4 solution changes the ion-selective permeability and increases the film resistance. Noda et al.5) reported that the addition of a few percents of Co, Ni, Al and W as alloy element decreased the

Electrochemical Impedance of Thin Rust Film Fabricated Artificially

323

corrosion rate of steel, and that the cation selective permeability of Fe-Co and Fe-Al low-alloy steels contributed to the corrosion protection. Though they5) discussed the variation of ion-selective permeability by the alloy element addition, the reason for the corrosion protection by alloy elements has not been explained sufficiently. The present method must be useful to investigate the corrosion protection mechanisms in the presence of various metallic ions because the resistance of rust film itself can be determined. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10.

I. Suzuki, N. Masuko and Y. Hisamatsu, Boushoku Gijutsu, 20, 319 (1971). M. Sakashita, Y. Yomura and N. Sato, Denki Kagaku, 45, 165 (1977). K. Noda, T. Nishimura, H. Masuda and T. Kodama, J. JIM, 63, 1133 (1999). K. Noda, T. Nishimura, H. Masuda and T. Kodama, J. JIM, 64, 767 (2000). K. Noda, T. Nishimura, H. Masuda and T. Kodama, Proc. APCCC, (2001) p.1369. H. Kihira, S. Ito and T. Murata, Corrosion, 45, 347 (1989). T. Nishimura, H. Katayama, K. Noda and T. Kodama, Zairyo-to-Kankyo, 49, 45 (2000). T. Nishimura, K. Noda and T. Kodama, Zairyo-to-Kankyo, 49, 734 (2000). T. Nishimura, H. Katayama, K. Noda and T. Kodama, Corrosion, 56, 935 (2000). M. Itagaki, R. Nozue, K. Watanabe, H. Katayama, K. Noda, Corros. Sci., 46, 1301 (2004).

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

325

A new approach to describe the passivity of nickel and titanium oxides Román Cabrera-Sierraa,b, Ignacio Gonzálezb,*, Jorge Ávalos-Martínezb, Gerardo Vázquezb, Máximo A. Pech-Canulc a

Escuela Superior de Ingeniería Química e Industrias Extractivas (ESIQIE-IPN) Academia de Química Analítica, Edificio Z5. A.P:75-874, C.P. 07338, México, D.F. (MEXICO). b Universidad Autónoma Metropolitana-Iztapalapa. Departamento de Química. Apartado Postal 55-534. 09340 México, D.F. (MEXICO). [email protected] c Departamento de Física Aplicada, Centro de Investigación y de Estudios Avanzados del IPN, A.P. 73 Cordemex, C.P. 97310, Mérida, Yucatán, (MEXICO).

Abstract - Nickel and titanium oxide films were grown by means of potentiostastic pulse in aqueous borate solution at different pH. The passive properties of these films were evaluated by EIS technique. These results were used to validate a new model, used in this work, to describe the conduction through the different semiconductors films. This model considers the initial assumptions of point defect model, and also, the presence of molecular hydrogen into the films. Keywords : Point Defect Model, EIS, Passive films, nickel, titanium

1. Introduction Passivity of metals and alloys is a phenomenon of great technological importance since it helps reducing the rapid deterioration of many construction materials. Thus, there are still continuous efforts to improve understanding of the formation of passive films on metal surfaces. Starting with the work of Cabrera and Mott in the late 1940s [1], several models have been proposed in the literature to describe the mechanism of growth and dissolution of passive films on metal substrates in contact with different aqueous media. The point

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defect model (PDM) developed by Macdonald et. al.[2] is perhaps the most well-known oxide film model nowadays. It was originally developed to explain the growth, breakdown and impedance characteristics of passive films on Ni, Fe and Fe-based alloys. It was later extended to consider the properties of barrier passive films under steady state conditions and finite-rate interfacial kinetics. Furthermore, it has been used recently by different authors as a framework to model different oxide [3] and sulfide [4,5] passive films. In particular, new models can be envisioned extending the applications of the PDM to account for phenomena relevant to the growth and dissolution of passive films and which has not been considered so far in a mechanistic model. 2. The role of hydrogen in the electrochemical impedance response of passive films The effect of hydrogen incorporation in passive films is described in this work. Pyun and Oriani [6] reported that hydrogen in passive film is present as protons. Thus the reaction of protons with O2- or OH- increases the H2O/O2- and OH-/O2ratios in the passive film. Z. Grubač and M. Metikoš-Huković [4] considered the incorporation of OH- ions into anionic vacancies of a bismuth sulfide film as an essential step in the solid-state transformation of sulfide to oxide. For hydration of passive oxide films on aluminum, B.C.Bunker et al. [7] suggested that OH- ions can readily migrate between oxygen vacancies leading to the formation of oxyhydroxide or hydroxide phases such as AlOOH and Al(OH)3. In the case of nickel it is generally agreed that the passive film consists of an inner oxide layer and an outer layer of hydroxyls or hydroxide [8,9]. On the other hand in situ STM observations by Scherer et al [9] have shown for passivation of nickel in sulfuric acid solutions that passive layer formation proceeds via the nucleation and growth of small 3D oxide and/or hydroxide islands followed by a slow restructuring to a structural well defined NiO layer. This suggests the possible presence of hydroxyl ions and vacancies in the passive layer. In the case of titanium STM measurements and density functional theory (DFT) calculations have shown that oxygen vacancies are active sites for water dissociation on TiO2 with one OH group filling the vacancy and H being adsorbed on the nearby bridging oxygen atom [10,11]. Thus it is likely that in this case hydroxyl ion vacancies may be formed. In this work a new model is presented (see Figure 1), based on the PDM but incorporating reactions at the interfaces which consider that diffusion of hydroxide ions (OH-) may take place through a mechanism similar to that proposed for pitting. In this case, a condensation region may form in substrate-selective regions, where atomic hydrogen (Ho) accumulates due to OH- reduction at the metal – film interface. This reduction is a result of reaction 8, and fosters the formation of H2 due to atomic hydrogen recombination. The subsequent release of this gas at high pressure would originate film rupturing in certain regions of the metal – film interface (a phenomenon often related to blistering type of corrosion). Analysis

A new approach to describe the passivity of nickel and titanium oxides

327

of the impedance response of passive films exhibiting n-type and p-type electronic conduction character was performed following an approach described in a previous publication [5]. The overall impedance function ZT is given by:

(ZT - Rs )-1 = jw C p + Z -f 1

(1)

where Rs is the solution resistance, Cp is the interfacial capacitance, Zf the overall faradaic impedance. We assumed that for a passive film with p type electronic character, Z f = Z fm + Z fh , where Zfm and Zfh are the faradaic impedances for cation and hydroxil ion vacancies at the metal-film interface, respectively.

Figure 1 Schematics of processes that occur inside a passive film according to the model presented in this work, within the framework of the PDM.

⎡ eg m L (g m + K ) + e -g m L (g m - K )⎤ Z fm = R1, M x+ ⎢1 + k1 ⎥ jw e g m L - e - g m L ⎣ ⎦ 1

R1, M x+

(

⎡ e Z fh = R2, OH - ⎢1 + k - 3b ⎣ 1 R 2,OH -

)

= 2 F k1 b1 Cm0 + k -1 b-1 a1

(

(

-g h L / 2

)

(K + g h ) - e

(

2 jw e

)

0 a1 = F b3b k 3b + b-3b k -3b COH

g hL / 2

g hL / 2

-e

(2)

(K - g h ) ⎤

-g h L / 2

)

⎥ ⎦

(3) (4) (5)

Similarly, for a passive film with n type electronic character, Z f = Z fo + Z fh , where Zfo and Zfh are respectively, the faradaic impedances for anion (oxygen) and hydroxyl ion vacancies at the metal-film interface.

328

R. Cabrera-Sierra et al.

⎡ e-gOL (K +g O )- egOL (K -g O )⎤ Z fo = R1,O2- ⎢1+ k-3a ⎥ jw egOL - e-gOL ⎣ ⎦

(

(

)

1 = 2 F b3a k3a + b-3a k -3a CO0 a1 R1,O 2-

)

(6) (7)

with: g m = K 2 + jw DM , g h = K 2 + 4 jw / DOH , g O = K 2 + jw / DO In order to test the applicability of this model, an analysis of the experimental impedance response of passive films formed on nickel and on titanium was carried out. Oxide passive films on these two metals are representative of films with p-type and n-type semiconductor behavior, respectively. These metallic oxide films were electrochemically formed by a potentiostatic treatment (at three different potentials) in a borate media. 3. Results and Discussion The diagnostic criteria proposed in the PDM [2] for the steady state current dependence on potential corroborated that nickel and titanium oxide films exhibited p-type and n-type semiconductor behavior, respectively. Figures 2 and 3 show the experimental impedance spectra for these films obtained in the frequency range from 40 kHz to 0.03 Hz and also their corresponding simulated diagrams (using equations 1-7 and parameters in Table 1). For NiO the diagrams exhibit at least two time constants with the low frequency one resembling a diffusion tail. In the case of TiO2 the diagrams resemble a capacitive semicircle.

Fig. 2 Impedance diagrams for Ni in 0.642M H3BO3+ 0.145 M NaOH solution, pH 8.4 obtained at: a) -0.4 V, b) 0.0 V and c) 0.4 V vs SSE.

The very good agreement between experimental and calculated impedance diagrams, and the reasonable values of parameters used in the simulations (see

A new approach to describe the passivity of nickel and titanium oxides

329

Table 1) support the validity of the model presented in this work. In consistency with the observation that the steady state current increases with potential, the values of the kinetic constant (k1) associated to the annihilation of cation vacancies also increases. An additional observation is that the values of the kinetic constant (k3b), corresponding to the formation of hydroxyl ion vacancies were two orders of magnitude higher than k1. On the other hand, the diffusion coefficient for hydroxyl ion vacancies for the three potentials is about one order of magnitude smaller than that for cation vacancies. This would suggest that the diffusion time constant of hydroxyl ion vacancies is larger than that for cation vacancies. Thus, the effect of hydroxyl ion vacancies can be ascribed to the impedance response in the low frequency range.

Fig. 3 Impedance diagrams for Ti in 0.642M H3BO3+ 0.2 M NaOH solution, pH 9 obtained at: a) -0.3 V, b) -0.6 V and c) -0.9 V vs SSE.

In the case of TiO2 passive films, the kinetic constant k3a shows a very poor dependence with the potential, an observation that is in agreement with the lack of dependence of steady state current on the applied potential. With respect to the diffusion coefficient for hydroxyl ion vacancies it can be observed that for the three potentials it is much smaller than for oxygen vacancies. This would suggest that the diffusion time constant of hydroxyl ion vacancies is larger than that for oxygen vacancies. 4. Conclusions The impedance model, derived within the framework of the PDM allowed the simulations of impedance diagrams typical for passive films having n-type and p-type semiconductor behavior. Contribution of hydroxyl ion vacancies to the impedance response was shown to consist of a diffusional impedance with a

R. Cabrera-Sierra et al.

330

time constant higher than that for cation vacancies (p-type electronic character) and higher than that for anion vacancies (n-type electronic character). A very good agreement was obtained between calculated impedance diagrams and typical impedance spectra for passive films on Nickel and Titanium. Table 1. Values of parameters used in the simulations. e = 1 x 106 V.cm-1 NiO film

-0.4 V vs SSE -1

0 V vs SSE

0.4 V vs SSE

2.287 x 10

-6

k3b (cm seg )

4.833 x 10

-4

DM (cm2 seg-1)

2.152 x 10-18

1.00 x 10-18

0.91 x 10-18

DOH (cm2 seg-1)

6.121 x 10-19

3.52 x 10-19

1.639 x 10-19

L (cm)

2.7 x 10-7

3.0 x 10-7

3.2 x 10-7

10

22

18

0.082 and 0.357

0.125 and 0.20

0.161 and 0.222

-0.9 V vs SSE

-0.6 V vs SSE

-0.3 V vs SSE

k1(cm seg ) -1

Cp (mF) R1,M

c+

(W) and

R2,OH-

(W)

TiO2 film -1

-5

k3a (cm seg )

6.83 x 10

k3b (cm seg-1)

2.758 x 10-4

2

-1

-18

6.049 x 10

-5

2.08x 10-4

8.633 x 10

-4

7.766 x 10-4

-4

1.148 x 10-3

2.003 x 10-3

2.348 x 10-3

1.983 x 10

6.724 x 10-19

DO (cm seg )

7.75 x 10

DOH (cm2 seg-1)

8.92 x 10-19

5.41 x 10-19

1.99 x 10-19

L (cm)

2.5 x 10-7

2.5 x 10-7

3 x 10-7

25.5

22.5

21.8

1.25 and 0.625

0.435 and 0.086

0.075 and 0.074

Cp (mF) R1,O2-

(W) and

R2,OH-

(W)

9.739 x 10

-19

The authors are grateful for the financial support of CONACyT project 47162. References [1] N. Cabrera and N.F. Mott. Rep. Progr. Phys. 12 (1948), 163. [2] D.D. Macdonald. J. Electrochem. Soc. 139 (1992), 3434. [3] A.C. Lloyd, J.J. Noël, S. McIntyre, D. Shoesmith, Electrochim. Acta. 49 (2004), 3015. [4] Z. Grubač, M.Metikoš-Huković. J. Electroanal. Chemistry. 565, (2004), 85. [5] R. Cabrera-Sierra, E. Sosa, M.A. Pech-Canul and I. González. Electrochim. Acta. In press. [6] Su-Il Pyun, Ch. Lim and R.A. Oriani, Corros. Sci., 33 (1992) 437. [7] B.C. Bunker, G.C. Nelson, K.R. Zavadil, J.C. Barbour, F.F. Wall, J.D. Sullivan, C.F. Windisch Jr. M.H. Engelhardt and D.R. Baer. J. Phys. Chem. B. 106 (2002), 4705. [8] N. Pineau, C. Minot, V. Maurice and P. Marcus. Electrochem. Solid-State Lett. 6 (2003), B47. [9] J. Scherer, B. M. Ocko, O. M. Magnussen. Electrochim. Acta. 48 (2003), 1169. [10] R. Schaub, P. Thostrup, N. Lopez, E. Lægsgaard, I. Stensgaard, J. K. Nørskov and F. Basenbacher. Phys. Rev. Lett. 87 (2001), 266104-1. [11] M. Menetrey, A. Markovits, C. Minot. Surf. Sci. 524 (2003), 49.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

331

Relation of the Photocurrents to the Corrosion Rates of Pure Aluminum having Various Oxides Maia C. Romanes, Kathleen A. Donovan and T. David Burleigh Materials and Metallurgical Engineering Dept., New Mexico Tech, 801 Leroy Place, Socorro, NM 87801 USA [email protected]

Abstract - Photoelectrochemical measurements of aluminum in saltwater revealed an inverse relationship between the magnitude of the photocurrent and the polarization resistance of the oxide layer. It was postulated that the photocurrent originated from the mid-bandgap traps. Pitting appeared to be preceded by photocurrents from the high energy defects (trap depth E=3.8 eV). Keywords: aluminum, pitting, photocurrents, oxygen vacancies, EIS

1. Introduction Al2O3 is a wide-band gap semiconductor with a bandgap around 8-9 eV [1]. However, photocurrents have been observed at much lower photon energies. Burleigh [2] used photoelectrochemistry and reported that a 2.4–3.6 eV envelope defines the sub-bandgap energies for passive films formed on aluminum, ranging from the native wet-polished oxide to the boiled oxide. He corroborated the observation of Goodman [2] that “significant differences in electrical properties of the layers may result from relatively subtle variations in preparation techniques.” Menezes et. al. [3] extended this observation, noting that “corrosion resistance in Al and Al alloys is intimately related to the characteristics of the oxide film, and changes in the corrosion behavior are often connected with the subtle chemical/electronic changes in the film.” In this study, photoelectrochemistry (PEC) was used to characterize different oxide films on aluminum. The advantage of PEC is that it is a non-destructive, in-situ method that can provide information on semiconductor properties of aluminum oxide-hydroxide films. The secondary objective was to use

332

M.C. Romanes and H.S. Kwon

photoelectrochemistry to understand how the oxides’ defect density controls the corrosion resistance. 2. Experimental In photoelectrochemical research, the metal sample is immersed in an electrolyte with an electrical lead connected to it. The weak electrical current coming from the sample is measured (dark current), and the sample is then illuminated with monochromatic light. The resulting change in the electric current during illumination is also measured and this change is the photocurrent. The magnitude of the photocurrent depends on the applied voltage and the wavelength (color) of the illuminating light. This experimental procedure has been described in previous publications [2]. This current study investigated two different oxides, “thermal” and “plasma,” grown on the surface of pure aluminum. The oxide samples were provided by Sandia National Laboratories, located in Albuquerque, NM. The samples were prepared by first sputter-depositing pure aluminum on an oxidized silicon wafer. The nascent aluminum was next exposed either to pure oxygen at room temperature (the thermal oxide) or exposed to oxygen in an RF plasma (the plasma oxide). The wafer was then broken into 1x1 cm squares, a copper wire was attached using conducting silver paint, then the sample and wire were coated with Devcon 2-ton clear epoxy so that only a small window of the aluminum was exposed to the electrolyte. They were tested in either 50mM NaCl or 3.5% NaCl deaerated solutions. The reference electrode was the saturated calomel electrode and the counter electrode was a platinum wire. Four treatments, with four replicates each, were used and were coded as follows: AT Thermal Oxide in 50 mM NaCl AP Plasma Oxide in 50 mM NaCl BT Thermal Oxide in 3.5% NaCl BP Plasma Oxide in 3.5% NaCl Each replicate was exposed to the electrolyte for 5 hours unless it pitted whereby no photocurrent could be measured because the dark current was large and noisy. Initial PEC measurement was made upon immersion and every hour thereafter. Electrochemical Impedance Spectroscopy (EIS) was used after every PEC measurement to measure thickness and polarization resistance (Rp) of the films. After the final EIS measurement, the flatband potential was measured. The flatband test used the same experimental set-up but the wavelength was held at the peak photocurrent while the applied potential was increased by increments until the photocurrent was unmeasureable. All measurements were made at room temperature.

Relation of Photocurrents to the Corrosion Rates of Pure Aluminum

333

3. Results and Discussion Figure 1 shows the photocurrent for a representative thermal oxide. Noteworthy the photocurrent increases continually over the 5 hour span. 2 at2-a i(photo) nA/cm2, 0 hr at2-b i(photo) nA/cm2, 1 hr at2-c i(photo) nA/cm2, 2 hr at2-d i(photo) nA/cm2, 3 hr at2-e i(photo) nA/cm2, 4 hr at2-f i(photo) nA/cm2, 5 hr

i(photo) nA/cm2

1.5

1

0.5

0 150

200

250

300

350

400

450

500

550

wavelength (nm)

Figure 1: The growth in photocurrent spectra for the thermal oxide AT2 as measured during 5 hours of immersion in saltwater. The photocurrent was measured at OCP; -1.1 V at 0 hr, and -1.5 V at 1-5 hr.

The PEC and EIS was repeated hourly for five hours for all sixteen samples unless pitting occurred. The following observations were made: 1. There was a wide variation in the photocurrent behavior for the supposedly identical samples. We propose that the sputtered deposited Al varied across the wafer, and produced different qualities of oxide. 2. The photocurrent shifted to shorter wavelengths (higher energy) over time for those samples that did not pit. 3. The aluminum samples only pitted in 3.5% NaCl solution. 4. The photocurrent grew and then decayed for some samples. 5. The plasma oxides had both a higher initial photocurrent and a larger growth in the photocurrent in the 3.5% NaCl solution than did the thermal oxides. EIS measurements showed that the capacitance increased for all of the samples, indicating that the insulating oxide thickness decreased. When the EIS polarization resistance, Rp, is plotted against the PEC photocurrent at 300 nm (approximately the peak wavelength), the results are

M.C. Romanes and H.S. Kwon

334

shown in Figure 2. There is an inverse relationship between the Rp and the photocurrent. The larger photocurrent corresponds to a smaller polarization resistance and vice versa. We propose that the magnitude of the photocurrent corresponds to the number of defects in the oxide film and thus, its corrosion resistance. 3.5% 3

0 .4

2.5

0 .3

2

0 .2

1.5

0 .1

1

0

Figure 2: At 5 hr, there is an inverse relationship between the photocurrent and the polarization resistance.

BP4

BP3

BP2

BT4

BP1

BT3

BT2

BT1

AP4

AP3

AP2

AT4

AP1

AT3

AT2

0.5

AT1

-0 .1

iphoto (nA/cm2) at 300nm at 5hr

Rp (MOhm.cm2)

50 mM 0 .5

0

S a m ple

The photocurrent measurements are plotted in Fowler plots in which the squareroot of the quantum efficiency, F (the number of electrons emitted per photon striking the surface), is plotted against the energy of the incident light. For samples that did not pit, a representative Fowler plot is shown in Figure 3a whereas Figure 3b is a graph for those that pitted. After the first hour, a change in slope occurred and divided the plot into a low energy side and a high energy side. There is little growth in the low energy side but large growth in the high energy side. This allows us to infer that the photocurrent is related to two types of traps (defects) in the oxide. Each type has a corresponding energy level or trap depth, “Et.” By extrapolating to F1/2 = 0, the trap depth can be estimated. Et = 2.4 eV for the low energy defect while Et = 3.8 eV for the high energy defect (after subtracting off the low energy signal). Pitting appeared to be preceded by photocurrents coming from the high energy defects/traps. The low energy source (shallow traps) does not seem to participate in photocurrent generation prior to pitting. Electron photoemission from the aluminum metal does not seem to be a significant contributor to the photocurrent because the energies and voltages are higher than reported for photoemission [1,4], and the low energy photocurrent should grow substantially if the oxide is thinning and there is photoemission.

335

Relation of Photocurrents to the Corrosion Rates of Pure Aluminum 0.014

0.0035

50 mM NaCl

0.003

0.01

0.0025

0.008 0.006

0.002 0.0015

0.004

0.001

0.002

0.0005

0

b

F

a

Squareroot

Squareroot

F

0.012

3.5% NaCl

0

2

3

4

5

2

6

3

4

5

6

Photon Energy (eV)

Photon Energy (eV)

Figure 3: The graph on the left (a) shows a sample that did not pit during the 5-hour test period with a high energy and a low energy. The graph on the right (b) is a sample that pitted with only the high energy region.

Flat band potential in Figure 4 also indicated the presence of these two sources of photocurrent. On the left is a sample that generated a large photocurrent and a small Rp (AT2) and the right produced a very small photocurrent but a very high Rp, (AT3). AT2

3

AT3

0.3 at2-fba nA/cm2 at2-fbb nA/cm2

2

Photocurrent (nA/cm )

2

Photocurrent (nA/cm )

at3-fb nA/cm2

2

1

0

-1.6

-1.4

-1.2

-1

-0.8

-0.6

-0.4

Potential (V vs. SCE)

-0.2

0

0.2

0.1

0

-2

-1.5

-1

-0.5

0

Potential (V vs. SCE)

Figure 4: The flatband potentials demonstrated two types of defects; the left one with a flat band at –1.2 V and the right one with a flatband of –0.7 V.

The low energy source (shallow traps) corresponds to more positive flatband potentials (FBP) and higher polarization resistance (Rp) while the deeper traps correspond to more negative FBPs and lower Rp. Some oxide samples exhibited both sources by having two slopes in the plot.

336

M.C. Romanes and H.S. Kwon

4. Summary In this study, we were able to shed light on the following: 1. There was a significant difference in corrosion resistance between plasma and thermal-grown oxides. 2. The polarization resistance (Rp) was inversely related to the magnitude of the photocurrent. 3. Fowler and Flatband plots indicated two sources of photocurrents: a low energy (~2.4 eV) source and a high energy (~3.8 eV) source. We postulate that these sources are trap/defects (possibly oxygen vacancies). 4. Unlike the low energy defect, the high energy defect increases significantly with time in deaerated NaCl. 5. Pitting was only observed in the higher 3.5% NaCl solution. Both the tendency to pit and the growth of the photocurrent appear to be connected to the number of high energy defects. Acknowledgements The authors thank Dr. Kevin Zavadil of Sandia National Laboratories, Albuquerque, NM, for providing the samples and the funding through the Sandia-University Research Program. References 1. Goodman, A. M., Photoemission of holes and electrons from aluminum into aluminum oxide, J. of App. Physics, 41, 5,(July 1970). 2. Burleigh, T.D., Photoelectrochemical analysis of the hydroxide surface films on aluminum and its alloys, Materials Science Forum, 185-187, (1995). 3. Menezes, S., Haak, R., Hagen, G. and Kendig, M., Photoelectrochemical characterization of corrosion-inhibiting oxide films on aluminum and its alloys. J. Electrochem.Soc. 136,7 (1989). 4. Di Quarto, F., Piazza, S., Santamaria, M., Suseri, C., Handbook of Thin Films, editor H.S. Nalwa, Vol. 2, p. 373-414, 1, Academic Press, NY ((2002).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

337

Growth and Characterization of Anodic Films on Al-Nb Alloys M. Santamariaa, F. Di Quartoa, M. Gentilea, P. Skeldonb and G. E. Thompsonb a

Dipartimento di Ingegneria Chimica dei Processi e dei Materiali, Università di Palermo, Viale delle Scienze, 90128 Palermo, Italy, e-mail: [email protected] b Corrosion and Protection Centre, School of Materials, The University of Manchester, PO Box 88, Manchester, M60 1QD, UK

Abstract - The anodizing behaviours of sputtering-deposited aluminium, niobium and Al-Nb alloys, containing 0.4, 7.5, 21, 40 and 55 at.% niobium, have been examined in 0.1 M ammonium pentaborate electrolyte with interest in the morphology, structure and electronic properties of the anodic oxides. Transmission electron microscopy revealed amorphous oxides, containing units of Nb2O5 and Al2O3, with formation ratios intermediate between those of anodic alumina and anodic niobia. Photocurrent spectroscopy provided increased understanding of the electronic properties of the anodic films, suggesting the formation of "mixed oxides" with insulating behaviours. The estimated band gap values are correlated with film compositions through the electronegativites of the constituent species. Keywords: Al-Nb alloys, Photocurrent spectroscopy, mixed oxides.

1. Introduction Anodic films on valve metals alloys are of interest for investigating mechanisms of oxide growth, with previous work focused mainly on morphology, structure and composition, leading to insights into the mechanism of anodic oxidation and ionic transport in "mixed oxide". Since these films are also of interest for possible application in electronics, for instance in metal-oxide semiconductor junctions, the possibility of controlled modification of the solid state properties, such as band gap, flat band potential and dielectric constant, by use of "mixed oxides", is appealing from practical and theoretical viewpoints. Of relevance, it

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M. Santamaria et al.

has been proposed recently that the band gaps of crystalline binary oxides correlate with the electronegativities of their constituents. Such correlation suggests the possibility of predicting the band gaps of ternary oxides (1), using an average electronegativity parameter for the cationic group and hence, tailoring of oxide properties. Moreover, the proposed correlation has been extended to amorphous oxides, taking into account the effect of the amorphous structure on the band gap of the corresponding crystalline oxide (1). Thus, in the present work, the morphology, structure and electronic properties of anodic films grown on sputter-deposited Al-Nb alloys are examined by means of both ex-situ transmission electron microscopy (TEM) and in-situ photocurrent spectroscopy (PCS). 2. Experimental Aluminium, niobium and Al-Nb alloys containing 0.4, 7.5, 21, 40 and 55 at.% niobium were deposited by magnetron sputtering, using an Atom Tech Ltd system, with targets of aluminium (99.999%) and niobium (99.9%). Sputtering was performed in 5´10-3 mbar argon after previous evacuation to 5´10-7 mbar. The substrates consisted of high purity aluminium sheet that had been electropolished for 180 s at 20 V in a solution of ethanol and perchloric acid (4:1 by vol.) at 283 K. Alloy compositions were determined by Rutherford back scattering spectroscopy (RBS) using 1.83 MeV He+ ions supplied by the Van de Graaff accelerator of the University of Paris. The scattered ions were detected at 165° to the direction of the incident beam. Data were interpreted by the RUMP program. The deposited layers were anodized either potentiodynamically to 9 V (SCE) at 100 mV s-1 or galvanostatically to increased voltages at 5 mA cm-2 in ammonium pentaborate electrolyte (ABE) at 293 K. The experimental set-up for the photoelectrochemical measurements has been described elsewhere (1). A 450 W UV-vis xenon lamp coupled with a monochromator, allows irradiation of the specimen through a quartz window. A two-phase, lock-in amplifier, with a mechanical chopper, enables separation of the photocurrent from the total current in the cell. The photocurrent spectra are corrected for the relative photon efficiency of the light source at each wavelength, so that the photocurrent yield in arbitrary current units is represented on the y-axis. Ultramicrotomed sections, of 10 nm nominal thickness, of anodized specimens were prepared by ultramicrotomy and examined in a JEOL FX 2000II transmission electron microscope. 3. Results and Discussion Aluminium, niobium and Al-Nb alloys were anodized potentiodynamically to 9 V (SCE), with polarization curves revealing typical of growth of a barrier type anodic films as shown for the example of the Al-7.5%at%Nb alloy (Fig.1). Notably, the plateau current densities of the alloys were between those of

Growth and Characterization of Anodic Films on Al-Nb Alloys

339

aluminium (i @ 0.25 mA cm-2) and niobium (i @ 0.35 mA cm-2). Transmission electron micrographs of ultramicrotomed sections of the anodized alloys disclosed amorphous anodic films of uniform thickness, attached to relatively flat alloy substrates, as shown in Fig. 2 for the Al-7.5 at.%Nb alloy anodized to 100 V galvanistically. The formation ratios for the films increase with increase of niobium content and are intermediate between those for films on niobium, 2.3 nm V-1 (2) and aluminium, 1.2 nm V-1 (3). From the photocurrent spectrum for an anodic film grown potentiodynamically on the Al-55at.%Nb alloy and polarized in the same ABE electrolyte at UE = 3 V(SCE), a band gap, Eg, of 3.75 eV can be estimated by extrapolating to zero the (Iphhn)0.5 vs hn plot (see inset) and assuming non direct optical transitions. Using the same procedure for other anodic films, grown in the same conditions, a decrease in Eg as the niobium content in the base alloy increases is revealed (Table 1). It is important to stress that no photocurrent was detected for anodic films on alloys with a niobium content < 21at.%. Table 1 - Band gap values of anodic films grown on Al-Nb sputtered alloys at 100 mV s-1 in 0.1 M ABE at 293K

Alloy Composition Eg / eV

Al-21at.%Nb 4.11

Al-40at.%Nb 3.90

Al-55at.%Nb 3.75

Nb 3.35

Photocurrent vs potential curves were recorded at 10 mV s-1 in 0.1 M ABE by irradiating the films of Table 1 at different wavelengths. With shift of the polarizing voltage toward the cathodic direction, an inversion of the photocurrent sign occurs for all the investigated alloys (see Fig. 4), as sharp changes in the phase angle suggest, thus indicating behaviour typical of insulating materials (1). The presence of steady-state cathodic photocurrent has been confirmed by recording the total current in the dark and under irradiation at potential more cathodic than the inversion potential. In previous works (1) a correlation between the optical band gap of crystalline oxides, MxOy, and the square of electronegativity difference of their constituents, (cM - cO)2, was proposed. Such a correlation was derived by assuming a direct relation between the optical band gap and the single M-O bond energy, the latter determined using the Pauling equation. The following two fitting equations were found for s,p metal and d metal oxides respectively, with the noticeable exception of nickel that followed the s-p metal oxide correlation. s,p)

Eg - DEam (eV)= 2.17 (cM - cO)2 – 2.71 (1a)

d)

Eg - DEam (eV)=1.35(cM - cO)2 – 1.49

(1b)

in which DEam increases as the degree of crystallinity decreases and DEam = 0 for crystalline oxides. DEam values ≤ 0.35 eV are expected according to the

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theory of amorphous semiconductors in the presence of lattice disorder only (4). In the case of "mixed" oxides, the average single bond energy is estimated taking into account the contributions of both cations involved in the oxide network formation. Thus, the same correlation was extended to the case of "mixed" amorphous oxides, considering an average cationic electronegativity, defined as follows:

cM = xAcA+xBcB

(2)

where A and B are the two metals in the "mixed" oxide, and x represents their cationic fractions. Study of the photoelectrochemical behaviour of a large number of "mixed" oxides (5-8) has enabled the validity of eqs. (1) and (2) to be dismantled. The photoelectrochemical results of this work allowed further assessment of validity. As already found for anodic films on Al-W alloys, the comparison between the experimental data and the theoretical predictions according to eqs. (1) suggests that the d-d metal correlation holds also for sp-d metal mixed oxides, provided that the percentage of d-metal in the film is higher than or equal to a minimum value, which is 21at.Nb%, very close to the value already found for Al-W mixed oxides (8). The difference between the experimental and theoretical values provides an estimate of DEam (see eq. 1b); in agreement with the theoretical expectation (4), the largest value was measured for oxide richest in aluminium. DEam depends on the relative amounts of aluminium and niobium incorporated into the film. In this frame the amorphizing nature of the aluminium oxide reduces but does not cancel the tendency of oxides grown on d-valve-metals to crystallize during the anodizing process. 4. Conclusions The results reported in this study show the effects of alloying aluminium with niobium on the growth of mixed oxides with morphologies and electronic properties between those of the pure partner oxides. Thus, transmission electron microscopy revealed the amorphous films, with formation ratios intermediate between those of anodic alumina and anodic niobia. Further photocurrent spectroscopy disclosed insulating films with the estimated band gaps consistent with an expectant theoretical correlation with the difference of electronegativities of the film constituents. References 1. F. Di Quarto, S. Piazza, M. Santamaria and C. Sunseri, in: H. S. Nalwa (Ed), Handbook of Thin Film Materials, vol. 2, Academic Press, S. Diego, 2002, Ch.8. p. 373. 2. J.P.S. Pringle, Electrochim. Acta, 25 (1978) 1423. 3. A.C. Harkness and L. Young, Can. J. Chem., 44 (1966) 2409. 4. N.F. Mott and E.A. Davis, Electronic Processes in Non-crystalline Materials, 2nd Ed., Clarendon Press, Oxford (1979).

Growth and Characterization of Anodic Films on Al-Nb Alloys

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5. M. Santamaria, D. Huerta, S. Piazza, C. Sunseri, F. Di Quarto, J. Electrochem. Soc., 147, 1366 (2000). 6. F. Di Quarto, M. Santamaria, P. Skeldon, G. E. Thompson, Electrochim. Acta, 48, 1143 (2003). 7. M. Santamaria, F. Di Quarto, G. E. Thompson, P. Skeldon, ATB Metallurgie, 40-41, 431 (2000-2001). 8. S. Piazza, M. Santamaria, C. Sunseri, F. Di Quarto, Electrochim. Acta, 48, 1105 (2003).

i / mA cm-2

0.5 0.4 0.3 0.2 0.1 0 -1

0

1

2

3

4

5

6

7

8

9

Formation voltage / V(SCE) Figure 1 – Polarization curve for Al-7.5at.%Nb anodized at 100 mV s-1 in 0.1 M ABE at 293 K

135 nm

Mixed oxide Alloy

Figure 2 - Transmission electron micrograph of an ultramicrotomed section of Al-7.5 at.%Nb anodized to 100 V at 5 mA cm-2 in 0.1 M ABE at 293 K.

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Photocurrent yield / a.u.

10

200

(Iph hn)0.5 / a.u.

8 6 4

100 0

2

3

3.5 4

4.5

5

hn / eV

0 200

250

300

350

400

450

wavelength / nm

Figure 3 - Photocurrent spectrum for an anodic film grown on Al-55at.%Nb at 100 mV s-1 to 9 V(SCE)in 0.1 M ABE at 293K and polarized in the same electrolyte at UE = 3 V(SCE). Inset: estimation of the optical band gap value by assuming non direct optical transitions.

0

Iph / nA

-60 100

-120

0

Phase angle / °

200

-180 -1

0

1

2

3

4

5

U E/ V(SCE) Figure 4 - Photocurrent and phase angle vs polarizing voltage recorded at 240 nm for an anodic film grown on to Al-55at.%Nb to 9 V(SCE) at 100 mV s-1 in 0.1 M ABE at 293K. Sol: 0.1 M ABE and vscan = 10 mV s-1.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

343

Amorphous semiconducting passive film-electrolyte junctions revisited. The influence of a non homogeneous density of state on the differential admittance behaviour of anodic a-Nb2O5 F. La Mantia, M. Santamaria and F. Di Quarto Dipartimento di Ingegneria Chimica dei Processi e dei Materiali, Università di Palermo, Viale delle Scienze, 90128 Palermo, Italy, [email protected]

Abstract - An analysis of the electronic properties of amorphous semiconductor-electrolyte junction is reported for passive films grown on Nb in alkaline solution and in a large range of thickness (~20nm ÷ ~250nm). A modelling of electronic density of state (DOS) has been carried out by fitting EIS spectra, at different potentials and in a range of frequencies (0.1 Hz ≤ f ≤100 kHz), and differential admittance (DA) data of a-Nb2O5/El interface. The fitting of EIS and DA curves was performed by using the theory of amorphous semiconductor Schottky barrier and a non-homogeneous DOS distribution. Keywords: Amorphous Semiconductor, Passive film, EIS, Differential Admittance

1. Introduction The study of electronic properties of passive films is an important issue in the physico chemical characterization of passive film-electrolyte junction. Semiconducting properties are frequently investigated by measuring the differential capacitance of the junction as a function of electrode potential, UE, and frequency of the superimposed ac signal. Interpretation of impedance data is performed by using a simple Mott-Schottky analysis, for getting out the flat band potential, Ufb, and the donor or acceptor concentration. The limits of such an approach have been reported in the literature (1). In an attempt to overcome these limits we proposed years ago a different approach based on the theory of amorphous semiconductor (a-SC) Schottky barrier in the low band bending regime (2-4). More recently we tried to extend such an approach to the high

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band bending regime by including also possible non-homogeneity in the distribution of electronic density of states (DOS) (5-6). In this paper a test of the proposed model is performed by studying the electronic properties of a-Nb2O5 anodic films grown in NaOH solution with different thickness (20 – 250 nm). The growth in NaOH solution has been preferred in order avoid any influence of electrolyte species incorporation on DOS. The EIS and admittance measurements were carried out in 0.5 H2SO4 testing solution so allowing a direct comparison of with Ufb values of thin films grown in the same testing solution (6). It will be shown that the theory of aSC/electrolyte junction is able to explain the admittance data of a-Nb2O5/El interface in a large range (10 Hz – 10 kHz) of frequency and electrode potential (UE- Ufb ≤ 8 V). A modelling of experimental EIS at different potentials and frequencies (0.1 Hz < f < 100 kHz) will be presented and shortly discussed. 2. Experimental Niobium (purity 99.9) rod from Goodfellow Metals (Cambridge) was used for the experiments. It was sealed into a teflon cylinder with an epoxy resin leaving a flat circular surface (0.5 cm diameter) in contact with the electrolyte. The metal surface was treated mechanically with abrasive paper followed by polishing with alumina powder (0.5 mm), then electropolished in HF/H2SO4 solution (DV = 15 V for 45 minutes at 1500 rpm) and final cleaned in sonicated distilled water. The metal was then anodized at constant growth rate (10 mV s-1) in 0.1 M NaOH to formation voltages, UF, < 10 V(SCE) using a potentiostat (EG&G PAR 173) equipped with a signal generator (EG&G PAR 175), and kept for 2 h to the formation voltage (stabilization). Thicker films were obtained by anodizing at constant current density (500 mA cm-2) until the desired potential was reached (UF ≤ 80 V) by using a Keytley Mod. 227 current source. The impedance spectra and differential admittance curves were recorded by using a Parstat 2263 (PAR), connected to a computer for the data acquisition. The data were then processed according to the theory of a-SC (5-6). For all the experiments a Pt net with very high surface area was used as counter electrode. 3. Results and discussion According to a previous study (6) an anodising ratio ranging between 26.8 and 26.5 Å V-1 was derived from capacitance measurements during the growth of Nb2O5 anodic film up to 9 V(SCE) at 10 mV s-1 by using a dielectric constant of 42 for anodic niobia and an efficiency of film formation equal to one. A further growth during the stabilisation process was observed and a new estimate of the final thickness was performed on the basis of circulated charge and/or of the interference colours (refractive index 2.3). Amorphous niobia anodic films, 20 nm and 250 nm thick, were prepared and investigated by EIS and differential admittance technique.

Amorphous semiconducting passive film-electrolyte junctions revisited

345

Optical band gap values around 3.4 eV were determined as well as zero photocurrent potential values in good agreement with flat band potential, derived from the fitting of admittance curves, were obtained from photocurrent spectroscopy studies. These results will be omitted now for brevity. 3.1. Electrochemical Impedance Spectroscopy (EIS) data analysis After stabilisation in the forming electrolyte the anodic films were investigated in 0.5 M H2SO4 solution by means of EIS from the open circuit potential, UOC, up to 5 V (SCE) for thin (20 nm) film and up to 8 V (SCE) for thicker films. EIS spectra were acquired in step of 1 V: further details on the choice of the initial and final potentials are reported in ref. (6). In separate experiments the curves of differential capacitance and conductance of a-SC/El junction, as a function of UE at different frequencies (10 Hz - 10 kHz), were obtained in a range of electrode potential starting from 8 V or 5 V(SCE) down to - 0.35 V(SCE). In Fig. 1 we report the impedance spectra of a-Nb2O5/H2SO4 junction, in Bode representation, of two films grown in NaOH solution up to 5 V(SCE) and 80 V. At different investigated electrode potentials the Bode diagram displays a similar behaviour with a characteristic inflection point in the module of the impedance and a corresponding minimum in the phase vs frequency plots. It can be explained with the presence of surface states at the aSC/electrolyte interface as suggested by the presence of an arm in parallel with the equivalent admittance of a-SC in the equivalent circuit (6). At higher thickness (UF ≥ 40 V) a slightly different equivalent circuit (see Fig. 2) was adopted including an arm still accounting for a kinetic of charge transfer trough surface states. Now we like to stress that from EIS data analysis we derived also the DOS distribution as well as the values of both components of admittance in parallel with the equivalent circuit of a-SC Schottky barrier (5). 3.2. Differential Admittance Curves In Figs. 3 and 4 we report the (CP)-1 and GP vs. UE plots of a-Nb2O5 Schottky barrier for two different thicknesses. The fitting of experimental data was performed by using the following theoretical expression derived previously (6 and refs. therein) in the hypothesis of a DOS spatially non-homogeneous but still constant (or slightly changing) in energy in an interval equal to KBT/e. According to such expressions we can write for 10 Hz ≤ f ≤ 10 kHz: 1 1 (1) = f (x(y )) C(x,w, y S ) C(w,y S )

w

S

G(x,w, y S ) = G (w,y S ) gw (x(yS ))

(2)

where fw(x(yS)) and gw(x(yS)) are two different trial functions depending only on electrode potential but changing with employed frequency. The term

346

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multiplying the two trial functions can be considered as coincident with the expression of (CP)-1 and GPC for homogeneous film, but averaged in energy ( N(E) = N ). For constant DOS (homogeneous films) these expressions were previously derived (5) as (for yS ≥ yg): ⎞ ⎛ y (3) 2 1 1 g ⎟ ⎜ C HBB (w,y S )

G HBB(w,y S ) = p2f

=

(

)

yS - yg ⎟ + 1+ ⎜ ln yg ⎠ ee 0e2 N ⎝ yC

⎛ y 2 kT g ee 0e2 N⎜⎜ ln + 1+ yS - y g y y e yC C g ⎝

(

⎞ -2 ⎟ ⎟ ⎠

)

(4)

for capacitance and conductance respectively. In previous equations N is the DOS in eV-1cm-3, yS = (UE –Ufb) is the potential drop within the semiconductor, yg is the potential at which the Fermi level cross the half gap of semiconductor (1.3 Volt in our case), yC is a characteristic potential showing a dependence from ac frequency of - 59 mV/decade in the theory of a-SC Schottky barrier (14). The other symbols have their usual meaning. A noticeable influence of the investigated frequency is foreseen on the values of both components of admittance. The experimental results from EIS and differential admittance experiments were fitted by using a density of states increasing as the space charge regions widens toward the the metal/oxide interface. The differential capacitance curves were fitted in a large range of frequencies (10 Hz < f < 10 kHz) in the case of thinner films (thickness < 100 nm). In the case of thicker films (thickness > 100 nm) and higher frequency region, where a growing flattening of the differential capacitance curves was observed, the fitting of capacitance curves failed owing to the appearance in the capacitance vs potential plots of a behaviour typical of insulating materials. Such a behaviour, unexplainable within the limits of the crystalline SC, can be rationalized (1-6) in the frame of the theory of a-SC Schottky barrier. 4. Conclusions We have shown that by using the theory of a-SC Schottky barrier it is possible to explain the admittance behavior of a-Nb2O5 oxide film-electrolyte interface in a large range of frequencies and electrode potential. By fitting the admittance curves in a large range of electrode potential values, it has been possible to derive some information on the spatial distribution of DOS in the bulk of the passive film. Further studies are now in progress aimed to correlate the electronic properties of passive films with the kinetics of growth of such films. References 1. F. Di Quarto and M. Santamaria, Corrosion Eng. Sci. Tech., 39 (2004) 71. 2. F. Di Quarto, C. Sunseri and S. Piazza, Ber. Bunsenges. Phys. Chem., 90 (1986) 549. 3. F. Di Quarto, S. Piazza and C. Sunseri, Electrochim. Acta, 35 (1990) 97.

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Amorphous semiconducting passive film-electrolyte junctions revisited

4. F. Di Quarto, V.O. Aimiuwu, S. Piazza and C. Sunseri, Electrochim. Acta, 36 (1991) 1817. 5. F. Di Quarto and M. Santamaria, ECS PV 2003-25, Ed. V. Birss, L. Burke, A. R. Hillman and R. S. Lillard, p. 116 (2004). 6. F. Di Quarto, F. La Mantia and M. Santamaria, El. Acta in press.

1.E+06

90

1.E+04

60 f/°

½Z½ / W cm-2

1.E+05

1.E+03 1.E+02 1.E+01

30

|Z|exp 5 V |Z|exp 80 V phase exp 5 V phase exp 80 V

1.E+00 0 1.E-01 1.E+00 1.E+01 1.E+02 1.E+03 1.E+04 1.E+05 Frequency / Hz Figure 1 - Experimental and simulated ( ----- ) impedance spectra, in the Bode representation of a-Nb2O5/H2SO4 junction for films grown in 0.1 M NaOH to 5 V(SCE) and 80 V, polarized at UE = 3 V(SCE).

CPE Rel

RSS Rct RSC

CH

CSC

Figure 2 - Equivalent circuit for a-SC/El interface in presence of electron charge transfer and/or recombination of electrons and holes trough surface states

348

F. La Mantia et al. 0.6 0.5 0.06 0.4 0.04

0.3

Gp-exp Cp-exp

0.2

(Cp)-1 / mF-1cm2

Gp / ohm-1cm-2

0.08

0.02 0.1 0

0 -1

0

1

2

3

4

5

Voltage / V(SCE)

3.E-04

2

3.E-04

1.6

2.E-04

1.2

2.E-04

Gp-exp Cp-exp

1.E-04

0.8

(Cp)-1 / mF-1cm2

Gp / ohm-1cm-2

Figure 3 - Fitting of the experimental admittance curves of a-Nb2O5 (Vf = 5 V in 0.1 M NaOH, f = 10 kHz; scan rate = 10 mV s-1, sol: 0.5 M H2SO4). a) GP vs UE: eyC = 0.20 eV and Ufb = - 0.196 V(SCE); b) (CP)-1 vs UE: eyC = 0.20 eVand Ufb = - 0.190 V(SCE).

0.4

5.E-05

0

0.E+00 -1

0

1

2 3 4 5 Voltage / V(SCE)

6

7

8

Figure 4 - Fitting of the experimental admittance curves of a-Nb2O5 (Vf = 80 V in 0.1 M NaOH, f = 20 Hz; scan rate = 10 mV s-1, sol: 0.5 M H2SO4). a) GP vs UE: eyC = 0.20 eV and Ufb = - 0.29 V(SCE); b) (CP)-1 vs UE: eyC = 0.14 eV and Ufb = - 0.29 V(SCE).

Section D Passivity Issues in Biological Systems

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

351

Influence of Protein Adsorption on the Passivation of Dental Amalgams Christopher M.A. Bretta,b, Elsa Jorgea, Carla Gouveia-Caridadea, Humberto Diasb a

Departamento de Química, Universidade de Coimbra, 3004-535 Coimbra, Portugal Instituto Pedro Nunes, Rua Pedro Nunes, 3030-199 Coimbra, Portugal [email protected]

b

Abstract The influence of protein adsorption on the surface properties and corrosion resistance of the high copper dental amalgam Tytin® has been investigated, using the model protein bovine serum albumin, in chloride-containing and artificial saliva solutions by employing voltammetric techniques and electrochemical impedance. It is demonstrated that the effect of this protein can be to inhibit oxide formation and increase the amalgam corrosion rate. The results are discussed and compared with others obtained without protein in solutions of artificial saliva with and without the addition of the organic components. Keywords: bovine serum albumin, artificial saliva, Tytin dental amalgam, passive oxide film, electrochemical impedance

1. Introduction Dental amalgams have been used for over 150 years. Modern high-copper amalgams with sufficient plasticity, high strength and low creep for use in tooth restoration remain the most popular materials in many countries, although they undergo a small amount of corrosion, which must be minimised. A critical factor is to ensure that the system resembles the oral cavity: the variety of different conditions of acidity, temperature, organic compounds etc. makes a study of corrosion and surface passivation of the utmost importance. Proteins can adsorb on metal or naturally-oxidised metal surfaces and self-assemble in a nano-structured environment [1], positively or negatively influencing the oxide film stability, as corrosion inhibitor or accelerator. This is of great importance

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for dental amalgams, which have a complex phase microstructure. A recent study on Ti-6Al-4V alloy in artificial saliva evaluated the influence of added bovine serum albumin (BSA) and fluoride – BSA avoided fluoride ion attack on the passive oxide film [2]. However, to our knowledge, no studies regarding BSA adsorption have been carried out on dental amalgam. Dental amalgam is formed by rapid reaction of liquid mercury with particles (irregular, spherical or spheroidal) of a powder alloy containing principally Ag (40-70%), Sn (15-30%) and Cu (10-30%) in the form of Ag3Sn and Ag-Cu. Mercury diffuses into the alloy particles and reacts, forming various phases of the systems Sn-Hg, Ag-Hg; Ag-Cu and Ag-Sn remain from the reactants. For the currently used, high copper amalgams the main reaction is [3]: g-Ag3Sn + Ag-Cu + Hg ® g1-Ag2Hg3 + g2-Sn7Hg + g-Ag3Sn + Ag-Cu

(1)

The Sn-Hg phase, which corrodes relatively easily, undergoes further reaction: g2-Sn7Hg + Ag-Cu ® h -Cu6Sn5 + g1-Ag2Hg3

(2)

Electrochemical studies of corrosion, beginning in the 1990s, concentrated on factors such as pH [4], and mercury release [5,6] or involved direct comparison between different dental amalgams [7-9]. It was shown that different corrosion mechanisms can occur and that the corrosion current varies over several orders of magnitude according to which amalgam phases are exposed and their surface condition - freshly abraded or with a naturally-formed passivating oxide film. Individual phases of the systems Ag-Hg, Ag-Sn, Ag-Cu and Sn-Hg have been investigated electrochemically and by surface analysis in 0.9% NaCl [10-12]. The behaviour of each phase was significantly different; the g1-Ag2Hg3 phase was the most corroded. Comparison was made with commercial amalgams. A large number of aqueous solutions has been proposed to simulate biological fluids for testing [14,15]; the choice can have a significant effect [14-17]. In [18], using Tytin FC amalgam, the importance of adsorption of the organic component, lactate or lactic acid, and its relation with oxide formation was demonstrated. Similar conclusions were reached in [19] with citric acid. In [20], the concentration of lactic acid in the artificial saliva used in [18] was varied, and compared with results in a standard solution AFNOR S90-701 [21], which has a higher ionic concentration, mainly of sodium and chloride ions, but no organic component. The influence of dissolved oxygen was also probed. Corrosion was greater in the absence of oxygen, demonstrating the importance of passive oxide formation on the amalgam surface. Corrosion in AFNOR was also greater, where no organic component was present. In this work, the influence of protein adsorption on the surface properties and corrosion resistance of Tytin® amalgam has been investigated, using the model protein BSA in chloride-containing solutions and in artificial saliva, by voltammetric techniques and electrochemical impedance. The results are

Influence of protein adsorption

353

discussed and compared with others obtained in solutions of artificial salivas with and without the addition of organic components. 2. Experimental 2.1. Electrodes and instrumentation Tytin® FC dental amalgam (KerrDental, USA) was obtained in two-part plastic capsules separated by a membrane, one with mercury and the other with powder alloy. They were mixed in a mechanical vibrator. The powder alloy composition of Tytin® is 59 % Ag, 28 % Sn, 13 % Cu, and a mercury-to-alloy ratio of ~43%. It was cast into cylinder-shaped pieces of diameter ~6 mm and depth ~3 mm. Amalgam pieces were made into electrodes by fixing a copper wire to one face with flash-dry silver paint (SPI). When dry, this face, together with the edges, was covered with Araldite® (Ciba-Geigy, Switzerland) epoxy-resin and dried for at least 24 hours, leaving just one face exposed. Surface preparation consisted of polishing immediately before use with SiC paper down to 1200 grit and then rinsing with acetone. The exposed electrode area was 0.28 cm2. The three electrode system used for the electrochemical experiments included an Ag/AgCl (3 M KCl) electrode as reference and a Pt foil auxiliary electrode. Experiments were carried out using an Autolab PGSTAT 30 with FRA2 module (Ecochemie, Netherlands) controlled by GPES 4.9 software for open circuit potential (OCP) and polarisation curve measurements and by FRA 4.9 software for impedance measurements. FRA and GPES software was also used for data analysis. Impedance spectra were recorded at the OCP from 65 kHz down to 0.1 Hz, 5 steps per frequency decade with a sinusoidal perturbation of 5 mV rms. 2.2. Bathing solutions The bathing solution was NaCl in different concentrations or artificial saliva as used previously [19] to which concentrations of bovine serum albumin (BSA) up to 10 g L-1 were added. One litre of artificial saliva had 1.5 g KCl, 1.5 g NaHCO3, 0.5 g NaH2PO4, 0.5 g KSCN plus 0.9 g lactic acid, present as lactate. Solutions were prepared using Milli-Q ultrapure water of resistivity > 18 MΩ cm and analytical grade reagents. The temperature was 25±1ºC. 3. Results and Discussion Examples of results obtained in the different solutions will be presented and comparison made with results previously obtained.

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3.1. Sodium chloride solutions Three different concentrations of sodium chloride solution were tested – 0.01 M, 0.015 M (0.9%), 0.5 M and 1.5 M NaCl, the concentration of BSA being varied up to 10 g L-1. Fig.1 shows results regarding OCP values obtained in 0.5 M NaCl solution at two different pH values. -0.09

-0.12

pH 3.0

pH 5.0 -0.15

E / V vs Ag|AgCl

E / V vs Ag|AgCl

-0.12

-0.15

-0.18 -1

0 g L BSA -1 2 g L BSA -1 10 g L BSA

-0.21

0

2

4

6

8

-0.18

-0.21 -1

0 g L BSA -1 2 g L BSA -1 10 g L BSA

-0.24

10

0

2

4

6

8

10

t / ks

t / ks

Fig.1. Time variation of open circuit potential in 0.5 M NaCl at pH 5.0 and pH 3.0.

The influence of BSA is greater at pH 3.0. However, at both values of pH, the corrosion current, Icor, remains low, at ~0.1 µA cm-2, independent of the amount of BSA added. This suggests that adsorption of BSA, shown by the change in the OCP, is not occurring over the whole surface and that corrosion predominantly occurs at points on the surface uncovered by BSA.

9

- Z " / kW cm

2

0.1 Hz

6

0.1 Hz 3

-1

0 g L BSA -1 2 g L BSA -1 10 g L BSA

0 0

3

6

Z ' / kW cm

9

12

2

Fig.2. Complex plane impedance spectra after 4 h immersion in 0.5 M NaCl, pH 3.0.

Further information can be obtained through electrochemical impedance, Fig.2, illustrated for three different concentrations of BSA at pH 3. Successful equivalent circuit fitting consisted in a constant phase element (CPE) in parallel with a series combination of a resistance with a second parallel RCPE, a circuit previously used for adsorbed species [1]. CPEs were modelled as non-ideal

Influence of protein adsorption

355

capacitors and led to exponents of the order of 0.75 and capacitances of 50-100 µA cm-2, as expected. These spectra demonstrate that a large amount of BSA is needed in order to influence the corrosion rate. In agreement with the values of OCP, at pH 5 changes in the spectra on increasing BSA concentration are less. Increasing the NaCl concentration up to 1.5 M, at natural pH, also shows only a small effect of BSA. Nevertheless, absolute values of the impedance remain high showing that chloride ion does not succeed in destroying the oxide layer. 3.2. Artificial saliva Experiments in artificial saliva, AS, (composition in experimental section), were done with and without lactate. The variation of OCP was similar to that in [19], BSA making the value a little more positive after several hours immersion, both without and with lactate. Polarisation curves showed a similar variation: in the absence of lactate, Icor was essentially independent of the amount of added BSA at ~0.2 µA cm-2. With lactate, it was lower in the absence of BSA at 50 nA cm-2 and increased to 0.25 µA cm-2 with increasing BSA concentration, Fig.3. 0.2

E / V vs Ag|AgCl

0.1 0.0 -0.1 -0.2 -1

0 g L BSA -1 2 g L BSA -1 10 g L BSA

-0.3 -0.4 -9

-8

-7

-6

-5

-2

lg (|I| / A cm )

Fig.3. Polarisation curves after 4 h immersion in AS with lactate.

12

- Z " / kW cm

2

0.1 Hz 9

6

-1

0 g L BSA -1 2 g L BSA -1 10 g L BSA

3

0 0

3

6

9

Z ' / kW cm

12

15

2

Fig.4 Complex plane impedance spectra after 4 h immersion in AS without lactate.

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This suggests that lactate acts as an inhibitor but that it is displaced on the surface by BSA, which allows corrosion to occur. This result can be compared with that obtained in acidified artificial saliva, where lactic acid was also found to inhibit oxide formation [19]. Complex plane impedance spectra for artificial saliva without lactate are shown in Fig.4. In fact, the spectra are similar in the presence of lactate and show only a small effect of its presence, only becoming evident at high concentrations. 4. Conclusions The study has shown the importance of studying protein adsorption and its influence on the corrosion of dental amalgams, which have a complex phase microstructure. The competitive effects of protein adsorption, adsorption of other organic components and passive oxide formation have been brought into evidence. Acknowledgements Financial support from Fundação para a Ciência e Tecnologia (FCT), ICEMS (Research Unit 103), Portugal, is gratefully acknowledged. CGC thanks FCT for a PhD grant (SFRH/BD/18659/2004). References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

N.P. Cosman, K. Faith, S.G. Roscoe, J. Electroanal. Chem., 574 (2005) 261. H.-H. Huang, Biomaterials, 24 (2003) 275. R.G. Craig, “Restorative Dental Materials”, Mosby, St. Louis, 1985. M. Marek, J. Dent. Res., 76 (1997) 1308. D.B. Mahler, J.D. Adey, L.E. Simms, M. Marek, Dent. Mater., 18 (2002) 407. T. Okabe, B. Elvebak, L. Carrasco, J.L. Ferracane, R.G. Keanini, H. Nakajima, Dent. Mater., 19 (2002) 38. B. Westerhoff, M. Darwish, R. Holze, J. Appl. Electrochem., 22 (1992) 1142. B. Westerhoff, M. Darwish, R. Holze, J. Oral Rehabil., 22 (1995) 121. H.A. Acciari, E.N. Codaro, A.C. Guastaldi, Mater. Lett., 36 (1998) 148. H.A. Acciari, A.C. Guastaldi, C.M.A. Brett, Electrochim. Acta, 46 (2001) 3887. H.A. Acciari, A.C. Guastaldi, C.M.A. Brett, Corros. Sci., 47 (2005) 635. C.M.A. Brett, H.A. Acciari, A.C. Guastaldi, Key Eng. Mat., 230-232 (2002) 463. V.W.H. Leung, B.W. Darvell, J. Dentistry, 25 (1997) 475. J.Y. Gal, Y. Fovet, M. Adib-Yadzi, Talanta, 53 (2001) 1103. R.I. Holland, Dent. Mater., 8 (1992) 241. C.M.A. Brett, I. Muresan, Key Eng. Mat., 230-232 (2002) 459. C.M.A. Brett, I. Ioanitescu, F. Trandafir, Corros. Sci., 46 (2004) 2803. A.U.J. Yap, B.L. Ng, D.J. Blackwood, J. Oral Rehab., 31 (2004) 595. C.M.A. Brett, F. Trandafir, J. Electroanal. Chem., 572 (2004) 347. K. Elagli, M. Traisnel, H.F. Hildebrand, Electrochim. Acta, 38 (1993) 1769.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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BSA adsorption on Fe-17Cr in acid solution: electrochemical behaviour and surface composition Luis Lartundo-Rojas, Isabelle Frateur, Anouk Galtayries, Philippe Marcus Laboratoire de Physico-Chimie des Surfaces, CNRS (UMR 7045), Ecole Nationale Supérieure de Chimie de Paris, 11 rue Pierre et Marie Curie, 75005 Paris, France

Abstract - The adsorption of BSA and its effect on the electrochemical behaviour and on the surface composition of a Fe-17Cr stainless steel was studied in deaerated 0.05M H2SO4 solution as a function of potential and protein concentration. EIS measurements reveal a corrosion inhibition by the BSA at the corrosion potential Ecor and show that the presence of the protein has no effect on the electrochemical behaviour of the alloy polarized in the passive domain. XPS analyses demonstrate that chemically intact BSA molecules are adsorbed on the steel surface both at Ecor and in the passive region, and that the thickness of the BSA layer is about 4 nm. XPS analyses also show that BSA has no effect on the chemical composition of the mixed Fe2O3-Cr2O3 oxide layer whatever the potential. Thus, the corrosion inhibition effect evidenced at Ecor is due to the protein and not to a change in the chemical composition of the oxide layer. Keywords : stainless steel, BSA, sulphuric acid, XPS, EIS

1. Introduction Stainless steels are often used in protein-containing environments, in particular for marine, food, and biomedical applications. The initial stage of biofilm formation on metallic surfaces in aqueous solution is generally the adsorption of biomolecules, in particular proteins, present in the medium. The adsorption on stainless steel surfaces of a model protein, the Bovine Serum Albumin (BSA), has been studied by analytical chemistry methods (potentiometric titration, intrinsic viscosity measurements, zeta potential measurements, Lowry-Folin method) [1] and surface analysis techniques (FT-IRRAS, XPS, ToF-SIMS) [2-

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4]. Electrochemical methods, such as Electrochemical Impedance Spectroscopy (EIS), have been scarcely used until now [5]. The aim of this work was to study, as a function of the electrochemical potential and the protein concentration in the electrolyte, the adsorption of BSA and its effect on the electrochemical behaviour and on the surface composition of a ferritic stainless steel Fe-17Cr (wt. %) in deaerated 0.05M H2SO4 solution. For this purpose, EIS was coupled to X-ray Photoelectron Spectroscopy (XPS). 2. Experimental The Fe-17Cr alloy was provided by UGINE (France). Its chemical analysis is given in Ref. [6]. The samples were covered with cataphoretic paint, and, before each experiment, were mechanically polished with diamond paste down to 1mm. All the experiments were carried out at room temperature. For electrochemical measurements, a classical three-electrode cell was used. The working electrode was a Fe-17Cr rotating disk electrode (RDE) of 0.16 cm2 surface area, embedded in epoxy resin. The rotation speed was 600 rpm. The counter electrode was a large platinum grid, and the potentials were measured with a saturated sulfate electrode (SSE). The electrolyte was 0.05M H2SO4 (pH 1.3) deaerated by N2 bubbling. The BSA (from Sigma-Aldrich) concentration was comprised between 1 and 20 mg l-1. Ac impedance diagrams were collected with an AUTOLAB PGSTAT30/FRA2 system from ECO CHEMIE (frequency domain ranging from 100 kHz to a few mHz and amplitude of 10 mV). For the X-ray Photoelectron Spectroscopy (XPS) analyses, the sample was a Fe17Cr RDE of 0.61 cm2 surface area, electrochemically pretreated using the same three-electrode cell as for the electrochemical measurements, then rinsed in deionized water for 30 s, and finally dried in air. XPS analyses were performed with a VG ESCALAB Mk II spectrometer, using an Al Ka X-ray source (hn = 1486.6 eV). In addition, the binding energy component of the resolved C 1s peak corresponding to carbon in a hydrocarbon environment was set at 285.0 eV, as surface charging effects may affect this value. All the spectra shown in this paper were obtained at a take-off angle of the photoelectrons of 90° with respect to the sample surface. The data processing (background subtraction and peak fitting) was done with the ECLIPSE software from VG, using a Shirley type background and gaussian/lorentzian peak shapes.

BSA adsorption on Fe-17Cr in acid solution

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3. Results and discussion 3.1. Electrochemical impedance measurements The impedance diagrams plotted at the corrosion potential Ecor without protein and with 10 and 20 mg l-1 of BSA are presented in Fig. 1, after 30 min of immersion in deaerated H2SO4 solution. The diameter of the high frequency capacitive loop, corresponding to the charge transfer resistance [7], markedly increases with the BSA concentration. This result shows a corrosion inhibition by the BSA at Ecor. 400

-Im Z / W cm

2

300 3.1 Hz

200 100

3.1 Hz 11.8 Hz

0

10.4 mHz

-100

0

100

200

10.4 mHz

300

400

10.4 mHz

500

2

Re Z / W cm

Fig. 1. Ac impedance diagrams of a Fe-17Cr RDE, plotted at Ecor after 30 min of immersion in deaerated 0.05M H2SO4, ( ) without BSA, and with (̌) 10 and (r) 20 mg l-1 of BSA.

In the passive domain (-0.6≤E≤+0.25 V/SSE), the impedance diagrams exhibit one single capacitive loop of very high diameter (in the MW cm2 range), which shows that the alloy is covered by a protective barrier layer. Impedance measurements also show that the presence of BSA has no effect on the electrochemical behaviour of the passivated alloy surface [7]. 3.2. XPS analysis 3.2.1. Adsorbed BSA The N 1s and C 1s XPS spectra of the Fe-17Cr sample after 1h of immersion at Ecor in deaerated 0.05M H2SO4 with 20 mg l-1 of BSA are presented in Fig. 2. Similar spectra were obtained after 1h of polarization at 0V/SSE in the same solution. The nitrogen and carbon signals (except carbon contamination

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contributing to the C 1s signal at 285.0 eV) come from the protein and are therefore a fingerprint of the adsorbed biomolecule. 34000

40000

C–C; C=C; C–H 285.0 eV

32000

35000

Intensity [cps]

Intensity [cps]

30000 28000 26000 24000 22000

C=N; N–C=C 287.5 eV

30000

25000

COOH; CONH 288.3 eV

C–N; C–OH 286.3 eV

20000

20000 18000 410

15000 405

400

Binding Energy [eV]

(a)

395

390

294

292

290

288

286

284

282

280

Binding Energy [eV]

(b)

Fig. 2. (a) N 1s and (b) C 1s XPS spectra for Fe-17Cr, after 1h of immersion at Ecor in deaerated 0.05M H2SO4, with 20 mg l-1 of BSA. (___) Experimental spectrum and (＿ ＿ ＿) peak synthesis.

The N 1s peak is symmetric, centered at 400.4 (± 0.1) eV, as expected for the amine or amide groups of the BSA [8]. From the N 1s and C 1s XPS spectra, the atomic ratio C/N can be estimated (the total intensity for carbon and nitrogen were considered). After 1 h of immersion at Ecor, C/N = 3.7. This experimental value is in very good agreement with the theoretical one, calculated from the amino acid composition of BSA molecules, which is equal to 3.7 [9]. This result confirms the presence of the protein on the steel surface. The C 1s spectrum was fitted with four contributions corresponding to well identified carbon bonds associated with chemical groups present in the BSA molecule. The XPS intensity of each contribution (in percent) was compared to the fraction of each type of carbon in the BSA molecules [9] (see Table 5 in Ref. [7]). A good agreement between experimental and theoretical values is reached, providing evidence that chemically intact BSA molecules are adsorbed on the surface at Ecor. A 3-layer model was assumed: an inner mixed Fe2O3-Cr2O3 oxide layer and an outer Cr(OH)3 hydroxide layer covered by an adsorbed BSA layer. The amount of BSA adsorbed on the stainless steel surface can be assessed by calculating the atomic ratio N/(Fe3++Cr3+) (the intensities of iron and chromium in the oxide were considered). The variation of this ratio as a function of the bulk BSA concentration is shown in Fig. 3. A plateau corresponding to 1.6 (± 0.3) is reached above 2 mg l-1 of BSA.

BSA adsorption on Fe-17Cr in acid solution

361

2.5

3+

N/(Fe +Cr )

2.0

3+

1.5 1.0 0.5 0.0

0

5

10

15

20

-1

[BSA] / mg l

Fig. 3. Atomic ratio N/(Fe3++Cr3+) versus bulk BSA concentration in deaerated 0.05M H2SO4 solution, for the Fe-17Cr sample after 1h of immersion at Ecor. alloy The thickness of the oxide layer (e) was calculated from the I oxide Fe + Cr / I Fe + Cr

hydroxide / I alloy ratio, the thickness of the hydroxide layer (x) from the I Cr Fe+ Cr ratio,

M and the thickness of the BSA layer (d) from the I BSA / I alloy N Fe+ Cr ratio, where I X is the X intensity in the M matrix. The values of the different inelastic electron mean free paths were assessed by the TPP2M formula [10]; in particular, BSA lBSA Cr = l N =3.0 nm.

Table 1. Thicknesses of the oxide, hydroxide and BSA layers, and atomic percentage of Cr3+ in the oxide layer, as a function of potential and BSA concentration.

Potential

[BSA] / mg l-1

e / nm

x / nm

d / nm

%Cr3+at

Ecor Ecor Ecor Ecor Ecor Ecor Epassive Epassive

0 1 2 5 10 20 0 20

2.1 2.8 2.1 2.4 2.2 2.5 2.0 2.2

0.5 0.4 0.5 0.4 0.4 0.5 0.9 1.3

3.7 3.9 3.9 4.0 4.0 4.6

34 24 31 24 29 26 64 65

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L. Lartundo-Rojas et al.

For the Fe-17Cr sample immersed 1 h at Ecor with BSA (from 1 to 20 mg l-1), the average value of d deduced from the 3-layer model is 3.9 (± 0.2) nm (Table 1). As the dimensions of BSA molecules in aqueous solution are between 2´25 and 4´13 nm in acidic solution, depending on the pH and the existing electric charges [11], 1 or 2 monolayers of BSA are adsorbed on the steel surface. Similar results are obtained in the passive domain. For the Fe-17Cr sample immersed 1 h at 0 V/SSE with 20 mg l-1 of BSA, C/N = 4.7, N/(Fe3++Cr3+) = 3.3 and d = 4.6 nm (Table 1). 3.2.2. Oxide and hydroxide layers The Cr 2p and Fe 2p XPS spectra of the Fe-17Cr sample after 1h of immersion at Ecor in 0.05M H2SO4 with 20 mg l-1 of BSA are presented in Fig. 4. The Cr 2p and Fe 2p spectra were also recorded after 1h of polarization at 0 V/SSE in the same solution. 90000

60000

Fe (ox)

3+

Cr (ox)

50000 45000

Cr(met)

40000

80000

Intensity [cps]

Intensity [cps]

3+

3+

Cr (hy)

55000

Fe(met)

3+

Fe (sat)

70000

60000

35000 50000

30000 595

590

585

580

Binding Energy [eV]

(a)

575

740

735

730

725

720

715

710

705

700

Binding Energy [eV]

(b)

Fig. 4. (a) Cr 2p and (b) Fe 2p XPS spectra for Fe-17Cr, after 1h of immersion at Ecor in deaerated 0.05M H2SO4 with 20 mg l-1 of BSA. (___) Experimental spectrum and (＿ ＿ ＿) peak synthesis.

From the decomposition of the Cr 2p3/2 and Fe 2p3/2 spectra and using the 3layer model, the thicknesses of the oxide and hydroxide layers were estimated. The corresponding values are reported in Table 1, as a function of potential and BSA concentration. At Ecor and in the passive domain, the thickness of the oxide layer is about 2.2 (± 0.3) nm, independently of the presence/absence of BSA. The thickness of the hydroxide layer is similar with and without protein but depends on the potential: about 0.5 (± 0.1) nm at Ecor and about 1.1 (± 0.2) nm in the passive domain. The atomic percentage of Cr3+ in the mixed Fe2O3-Cr2O3 oxide layer, %Cr3+at=Cr3+/(Fe3++Cr3+)´100, where the intensities of iron and chromium in

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the oxide were considered, is about 28 (± 6) % at Ecor and about 65 (± 7) % in the passive domain, independently of the presence/absence of BSA (Table 1). The chromium enrichment observed in the passive domain is in agreement with the previously reported data [6,12]. The chemical composition of the oxide layer is not modified by the protein whatever the potential. Thus, the corrosion inhibition effect evidenced at Ecor by EIS is attributed to the protein and not to a change in the chemical composition of the oxide layer. 4. Conclusions

The electrochemical behaviour and the surface composition of a Fe-17Cr stainless steel in deaerated 0.05M H2SO4 solution was studied as a function of potential and BSA concentration. The electrochemical measurements show a corrosion inhibition by the BSA at the corrosion potential Ecor and no effect of the protein on the electrochemical behaviour of the passivated alloy surface. XPS analyses demonstrate that the adsorption of BSA takes place on the Fe17Cr surface both at Ecor and in the passive region, and that the adsorbed BSA molecules are chemically intact. From a 3-layer model, the thickness of the protein layer is estimated to about 4 nm whatever the BSA concentration and the potential, which would correspond to 1 or 2 monolayers. XPS analyses also show that BSA has no effect on the chemical composition of the mixed Fe2O3Cr2O3 oxide layer both at Ecor and in the passive region. Therefore, the corrosion inhibition effect evidenced at Ecor by EIS is due to the protein and not to a change in the chemical composition of the oxide layer. References 1. S. Fukuzaki, H. Urano, K. Nagata, J. Ferment. Bioeng., 80 (1995) 6. 2. C. Rubio, D. Costa, M.N. Bellon-Fontaine, P. Relkin, C.M. Pradier, P. Marcus, Colloids and Surf. B : Biointerfaces, 24 (2002) 193. 3. C.M. Pradier, D. Costa, C. Rubio, C. Compère, P. Marcus, Surf. Interface Anal., 34 (2002) 50. 4. C. Poleunis, C. Rubio, C. Compère, P. Bertrand, Surf. Interface Anal., 34 (2002) 55. 5. S. Omanovic, S.G. Roscoe, Langmuir, 15 (1999) 8315. 6. W.P. Yang, D. Costa, P. Marcus, J. Electrochem. Soc., 141 (1994) 111. 7. I. Frateur, L. Lartundo-Rojas, C. Méthivier, A. Galtayries, P. Marcus, Electrochim. Acta (2005) in press. 8. G. Beamson, D. Briggs, in: High Resolution XPS of Organic Polymers, Eds. J. Wiley & sons, London, 1992. 9. T. Peters, Adv. Protein Chem., 37 (1985) 161. 10. S. Tanuma, C.J. Powell, D.R. Penn, Surf. Interf. Anal., 21 (1993) 165. 11. D.C. Carter, J.X. Ho, Advances in protein chemistry, 45 (1994) 153. 12. W.P. Yang, D. Costa, P. Marcus, J. Electrochem. Soc., 141 (1994) 2669.

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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XPS characterisation of BSA adsorption on stainless steel Sandrine Zanna,a Chantal Compère,b Philippe Marcusa a

Laboratoire de Physico-Chimie des Surfaces, CNRS-ENSCP (UMR7045), Ecole Nationale Supérieure de Chimie de Paris, 11 rue Pierre et Marie Curie, 75005 Paris, France b IFREMER, Centre de Brest, Service Matériaux et Structures, BP70, 29280 Plouzané, France

Abstract - In order to better understand the biofilm formation mechanisms, the adsorption of a model protein (BSA) on a 316L stainless steel surface in aqueous solution and the effect of different salts (NaCl, NaCl+MgCl2, NaCl+CaCl2) have been investigated by X-ray photoelectron spectroscopy (XPS). The N1s XPS spectrum exhibits a signal at 400.4 eV which is characteristic of the amine or amide groups of the protein. The XPS data reveal that the protein is adsorbed on the stainless steel surface and that the amount of adsorbed protein depends on the nature and the concentration of the dissolved salts. NaCl has no effect on the BSA adsorption whereas MgCl2 and CaCl2 promote the adsorption of the protein on the stainless steel surface. The C1s signal was decomposed in four components corresponding to the different chemical groups in the BSA (C-C, C-O and C-N single bonds, C=N or N-C=N, CONH or COOH groups), providing evidence that the adsorbed BSA is chemically intact. The stainless steel surface composition was investigated by XPS. The enrichment of Mo and Ni under the oxide layer generally observed in the absence of protein is also found after adsorption of protein, as well as the enrichment of Cr3+ and Mo oxides in the oxide layer. Angleresolved XPS measurements indicate that the Mo6+ is partly converted in the presence of adsorbed BSA. This finding suggests the existence of a specific interaction between the protein and the molybdenum present on the surface. Keywords: XPS, stainless steel, biomolecule, protein

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1. Introduction The formation of biofilms on surfaces of materials immersed in seawater is often detrimental. On stainless steel it has been shown that the adsorption of biomolecules, in particular proteins, constitutes the first step of the biofouling process.1,2 Several parameters can influence the interactions between proteins and surfaces, such as substrate nature, properties of proteins, pH, concentration, ionic strength and temperature.3,4,5 Adsorption of proteins and its influence on the interfacial properties of solid surfaces have been studied by contact angle measurements, Fourier transform infrared reflection-absorption spectroscopy and electrochemical techniques.3,6,7 An effect of the presence of dissolved salts on protein adsorption in seawater has been observed 8,9. The aim of this work was to study by XPS (X-ray Photoelectron Spectroscopy) the adsorption of a model protein bovine serum albumin (BSA) on a 316L stainless steel surface, to determine the influence of Na+, Mg2+ and Ca2+ in solution on the interaction of proteins with the stainless steel surface, and to investigate the possible effect of adsorbed protein on the stainless steel surface composition. BSA is not a marine protein, but its choice for this work was driven by the fact that it is a wellcharacterized protein10 that is often chosen to study the interaction of proteins with solid surfaces11,12,13 and its effect on bacterial adhesion2. 2. Experimental Prior to each measurement, the 1 cm2 stainless-steel samples (AISI 316L) were mechanically polished down to 3 mm, rinsed in hexane and three times in demineralized water at 50°C. The XPS spectra were obtained on a VG Escalab Mk II spectrometer equipped with an Al Ka X-ray source (hn=1486.6 eV) under ultrahigh vacuum conditions. High-resolution spectra were obtained by applying a pass energy of 20 eV. The binding energies were calibrated against the binding energy of Au 4f7/2 and Cu 2p3/2; with this calibration, the low-energy carbon peak, attributed to hydrocarbon contamination, was measured at 285.4 eV.The transmission factor was checked as T=a(Ec)-0.54; the attenuation lengths were calculated from the Tanuma, Pen, Powel formula.14. The spectra were curve fitted using the VG software ECLIPSE. All adsorption experiments were performed in aqueous solutions prepared with high purity water and salts: NaCl (0.62 mole l-1), NaCl+CaCl2 (0.60 mole l1 +0.01 mole l-1), NaCl+MgCl2 (0.59 mole l-1+0.01 mole l-1). The total concentration of salt was identical to that of the natural seawater, except for NaCl+MgCl2, where the MgCl2 concentration corresponds to the CaCl2 concentration in natural seawater. The concentration of BSA (from Sigma) was 20 mg l-1. The stainless steel samples were immersed in the BSA solutions for 24 hours under agitation. They were rinsed in water containing the same amount of salt but no BSA.

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3. Results and discussion Figure 1 shows the N1s and the C1s XPS spectra of stainless steel samples after immersion in aqueous solutions containing NaCl, NaCl+CaCl2 and NaCl+MgCl2. (a)

(b)

Figure 1 (a) : N1s spectra of stainless steel after immersion in aqueous solutions with BSA and different salts, (b) C1s spectrum after immersion in a BSA+NaCl+MgCl2 solution.

A N1s peak at 400.4±0.1 eV is observed in all spectra (figure 1a) except the reference (sample non exposed to BSA). This peak is associated to the amine or amide groups of the protein,15 and is characteristic of the amount of BSA adsorbed on the stainless steel surface. The intensity of the N1s peak is not modified when NaCl is added to the solution, whereas it is markedly increased when Ca2+ or Mg2+ are added. This clearly shows that Na+ has no effect on the adsorption of the protein, whereas Ca2+ and Mg2+ enhance adsorption of BSA on stainless steel. This indicates that monovalent ions (Na+) do not play any role, whereas divalent ions (Ca2+ or Mg2+) promote the BSA adsorption. The C1s spectrum was fitted with four contributions (figure 1b): the first peak, at the lowest binding energy of 285.4 eV is assigned to carbon bound only to C or H; the second peak at 286.5 eV is attributed to carbon in C-N and C-O single bonds; the third peak, at 287.8 eV is attributed to carbon in C=N or N-C=C; the fourth peak, at the highest binding energy of 288.8 eV is thought to include signals from carbon in CONH and COOH groups. The identified bonds correspond to the different chemical groups present in the BSA molecule.8 The relative XPS intensities of the N1s signal and of the various C1s contributions provide evidence that the adsorbed BSA is chemically intact. The thickness of the protein layer was calculated using the measured XPS intensities, and was found to be 3.5 nm.

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Figure 2: Cr2p3/2, Ni2p3/2, Mo3d and Fe2p3/2 core level spectra for the reference sample.

The Cr2p3/2 spectrum (fig.2) was well fitted with three components at 574.5 eV, 576.6 eV and 577.9 eV, corresponding to the metal, the oxide (Cr2O3) and the hydroxide Cr(OH)3, respectively. The Fe2p3/2 spectrum was decomposed into two main components at 707.7 eV and 711.3 eV, attributed to the metal and the oxide/hydroxide (Fe3+) respectively. The Ni2p3/2 spectrum was fitted with a single metallic component at 853.0 eV. The Mo3d spectrum was fitted with three doublets: 227.8 eV (230.9 eV) for the metal, 231.0 eV ( 234.1 eV) for the hydroxide, and 232.6 eV (235.7 eV) for Mo6+. A component at 229.4 eV attributed to Mo4+ was observed after immersion in the aqueous solution. Using the measured XPS intensities corresponding to the metallic and the oxidized chromium, and assuming that the oxide layer is a continuous and homogeneous Cr2O3 layer, the thickness of the oxide layer was estimated at 17Å. Considering that the stainless steel is covered by a homogeneous and continuous oxide layer of 17Å and a BSA layer of 35Å, the molar ratios of the different components on and under the oxide layer were calculated. The oxide layer is enriched in Cr and Mo, and an enrichment in Ni and Mo under the oxide layer is observed. These results are consistent with those of Clayton and

XPS characterisation of BSA adsorption on stainless steel

369

Olefjord.16 It is interesting to note that the 24H immersion in a saline BSA solution does not change the concentrations of Cr, Fe, Mo and Ni on the the stainless steel surface.

Figure 3: Mo6+/ Moox as function of the N1s peak intensity at 90° ( Moox = Mo6++Mo4++Mohydroxide).

Mo3d XPS spectra were acquired at 90° and 45° take-off angles, before and after a 24H immersion in BSA+NaCl+MgCl2 solution. The molar ratio Mo6+/(Mo6++Mo4++Mohydroxide) was calculated from the intensities of the different components. The values of this ratio are 0.72 and 0.74 at 90° and 45° , respectively, for the reference sample and 0.39 at 90° and 0.21 at 45° for the sample immersed in the BSA+NaCl+MgCl2 solution. It can be concluded that after adsorption of BSA, the surface is depleted in Mo6+. This suggests a preferential interaction of the protein with the Mo6+ at the surface. In order to check this point, the molar ratio Mo6+/(Mo6++Mo4++Mohydroxyde) was measured as a function of the N1s peak intensity. The data (figure 3) indicate that the higher the N1s peak intensity (i.e. the higher the amount of adsorbed BSA), the lower the Mo6+ proportion in the passive film. These results confirm a prefetential interaction of the protein with the Mo6+ at the surface. 4. Conclusion The XPS analysis of stainless steel samples immersed in saline solutions containing proteins (BSA) shows that the BSA is adsorbed on the stainless steel surface and that there is an effect of the salt on the protein adsorption. The data

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support the conclusion that the divalent cations (Ca2+ and Mg2+) enhance the BSA adsorption, whereas monovalent cations (Na+) have no effect on the BSA adsorption. The C1s and N1s XPS spectra are characteristic of the BSA, and it is concluded that the BSA remains chemically intact after adsorption on stainless steel. The stainless steel surface composition was analysed after exposure to the aqueous solution of BSA with the different salts. In all cases the enrichment of Mo and Ni under the oxide layer as well as the enrichment of Cr3+ and Mo oxides in the oxide layer are observed, which shows that the adsorption of the protein does not modify the composition of the passive film. Angle-resolved XPS analyses indicate that the Mo6+ is partly converted in the presence of adsorbed BSA. This suggests the existence of a direct interaction between the protein and the Mo6+ present on the surface. References 1. Pradier CM, Bertrand P, Bellon-Fontaine MN, Compère C, Costa D, Marcus P, Poleunis C, Rondot B, Walls MG, Surf. Interface Anal. 2000 ;30 :45. 2. Compère C, Bellon-Fontaine MN, Bertrand P, Costa D, Marcus P, Poleunis C, Pradier CM, Rondot B, Walls MG, Biofouling 2001 ;17 :129. 3. Omanovic S, Roscoe SG, Langmuir 1999; 15:8315 4. Davies J, Nunnerley CS, Paul AJ, Colloids Surf. B 1996;6:181. 5. Su Tj, Lu JR, Thomas RK, Cui ZF, J. Phys. Chem. B 1999;103:3727. 6. Taylor GT, Troy PJ, Sharma SK, Mar. Chem. 1994;45:15. 7. Schakenraad JM, Busscher HJ, Colloids Surf. 1989; 43:331. 8. Pradier CM, Costa D, Rubio C, Compère C and Marcus P, Surf. Interface Anal. 2002;34:50. 9. Kirchman DL, Henry DL, Dexter SC, Mar. Chem. 1989;27:201. 10. He XM, Carter DC, Nature 1992;358:209. 11. Zeng H, Chittur KK, Lacefield WR, Biomaterials 1999;20:337. 12. Ishida KP, Griffiths PR, Appli. Spectrosc. 1993;47:584. 13. Servagent-Noinville S, Revault M, Quiquampoix H, Baron M-H, J. Colloid Interface Sci. 2000 ;221 :273. 14. Tanuma S, Powell CJ, Penn DR, Surf. Interf. Anal.1993;21:165. 15. Dufrene Y, Marchal TG, Rouxhet PG, Appl. Surf. Sci., 1999, 30 :45. 16. Clayton C, Olefjord I, “Passivity of Austenitic Stainless Steels”, chapter 7 in Corrosion Mechanisms in Theory and Practice, ed. P. Marcus, M. Dekker, NY(2002)p217 and references therein.

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Study of corrosion behaviour of Ti-coated AISI 316L stainless steel in a simulated body fluid solution F. Hellala, F. Atmani a, B. Malkib, H. Sedjalc, M. Kerkarc, F. Dalard b a

Ecole Nationale Polytechnique, Département de Métallurgie, LSGM, Algiers, 16200, Algeria, e-mail : [email protected] b Institut National Polytechnique de Grenoble, LTPCM/LEPMI/ENSEEG, Grenoble, 38402, France c Université Abderahmane Mira, Faculté des Sciences et des Sciences de l’Ingénieur, LTMGP, Bejaia, 06000, Algeria

Abstract - In the field of implantology, biomaterials must satisfy appropriate physicochemical and mechanical characteristics on the one hand and be biocompatible on the other hand. In this work, we have studied an AISI 316L stainless steel/titanium system. The stainless steel substrate, onto which titanium film was deposited by a PVD (physical vapour deposition) process, was cold-rolled at different rates (20%, 40%, 60% and 80%). Linear sweep voltammetry was used, in complement to microstructural and mechanical characterisations, to study the deposit and the substrate. The electrochemical results indicate that the best electrochemical behaviour in a simulated physiological environment is obtained for the titanium deposit on a substrate cold-rolled at a rate of between 20% and 60%. Even though the deposit generates internal stress at the film/substrate interface, titanium deposits present good fretting features and wear resistance. Keywords: biomaterials, AISI 316L, titanium, PVD, cold rolling, hardening, anodic polarisation.

1. Introduction In order to exploit the surface properties of titanium (corrosion resistance, wear resistance and biocompatibility) while conserving the mass properties of steel and

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reducing the cost of the implant, we have chosen to use Ti/AISI 316L stainless steel biomaterial. Titanium is deposited on the stainless steel by cathodic sputtering. The steel is cold-rolled in order to determine the effect of the deformation on the titanium deposition. Despite the availability of studies on titanium, there is still much to be learned [1-5]. The aim of this work is therefore to investigate the corrosion kinetics, in a physiological serum, of AISI 316L stainless steel subjected to different cold-rolling rates and then coated with a thin titanium film. Different characterization techniques were used in this study, including linear voltammetry, scanning electron microscopy (SEM), atomic force microscopy (AFM) and energy depressive X-ray analysis (EDX). The mechanical properties of the Ti-deposit were also studied by micro-hardness and wrenching tests.

2. Materials and methods The material used in this work was AISI 316L stainless steel. Flat 160 x 10 x 1mm3 samples were deformed by cold-rolling at different rates (20%, 40%, 60% and 80%) and then cut into 8mm-diameter disks, which were then mechanically polished with SiC paper to grade 4000. Titanium was deposited on the stainless steel disks by physical vapour deposition (PVD) at room temperature, at a power of 300W and in an argon atmosphere with a pressure of 10-2 mbar. Prior to the deposition, the substrate was placed in a high vacuum chamber (10-6 mbar). SEM observation indicates that the deposit takes the exact shape of the polished surface structure. Preliminary EDX spectra seem to indicate that the Ti-deposit on the deformed substrate is thicker than that on an undeformed material. After 3 hours of deposition, the Ti film thickness was estimated to about 1µm. Further examination of the surface was made by atomic force microscopy (Figure 1). This shows the deposit homogeneity in more detail and, in particular, indicates that some of the substrate lands are not completely coated.

Figure 1: 3D AFM image of titanium deposit on AISI 316L steel substrate (50µm x 50µm), showing open porosities.

The film grip was first tested by wrenching and by measuring its micro-hardness. The aim of this was to relate the structural stressed state of the substrate to the film behaviour. In a second step, in order to simulate natural biological conditions, physiological serum was used as an electrolyte: a NaCl 9g/l solution at pH6.3, thermostatically controlled at 37°C. This electrolyte was used to immerse a platinum

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counter electrode and a saturated calomel electrode (SCE) alongside the working electrode in an electrochemical cell controlled by a Solartron SI 1286 potentiostat. The working electrode rotated at 1000rpm and all experiments were performed after an open circuit time of 8 hours with a linear weep of 5mV/s.

3. Results 3.1. Mechanical behaviour Figure 2 shows the effect of the cold rolling ratio on the coating and substrate hardness [6]. The substrate hardens as the rolling rate increases; this is obviously due to the associated deformation [7].

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200 100 0 0

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Cold-rolling ratio (%) Figure 2: Effect of cold rolling ratio on the coating and substrate hardness [6].

The deposited titanium shows a higher hardness, even though it decreases above a rolling rate of 40%. The deposit therefore provides it with good fretting features and wear resistance. Finally, it should be noted that there may be a hardening at the film/steel interface, which can be explained by the formation of internal stress during the deposition.

3.2. Electrochemical behaviour Voltammetry (Figure 3) shows that the electrochemical behaviour of both the stainless steel and the titanium coating is different. As far as the results of linear sweep voltammetry for Ti-coated AISI 316L stainless steel are concerned, there is a small diminution in its potential as the rolling rate increases to 40%. The associated deformation is due to the increase in the dislocation density [1, 8] which represents the anodic dissolution sites. Furthermore, a partial transformation of austenite into martensite [5, 9] may lead to the formation of local micro-cells between the martensite and austenite phases.

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AISI 316L cp Ti

-9

-1.0

-0.5

0

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E (V/ECS)

Corrosion current density (A/cm²)

Figure 3: Corrosion current curves versus potential, for AISI 316L (solid line) and titanium coatings (dotted line), in a serum medium.

0

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1,00E-04 1,00E-05

Ti-coating Ti-coating

1,00E-06

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1,00E-07 1,00E-08 1,00E-09

Cold-rolling ratio (%)

Figure 4: Effect of cold rolling on corrosion potential of AISI 316L substrate (white markers) and titanium coating (black markers).

The evolution of corrosion potential (Figure 4) during the anodic polarisation indicates that at rolling rates of above 20%, there is a small rise in the potential value of the substrate, even though no changes in the chemical composition of the passive film were revealed by XPS analysis [10]. This is probably due to the mechano-chemical effect [4], attributed to the structural reorganization of the dislocations under the effect of the applied stress. Indeed, their multiplication and their arrangement in pile-ups increase stress concentrations. This effect may most likely change the local equilibrium potential, which is well described by the basics of the chemical potential of dislocations

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[4]. These results match well with those obtained in the micro-hardness studies [6], in which the micro-hardness of the titanium deposit was found to present a similar evolution. In the case of the AISI 316L coated with titanium deposit, if the deposited film had a similar microstructure to that of the massive titanium, its potential should be very close. We have recorded that the potential of a coated undeformed substrate is inferior to that of the non-covered substrate; a reasonable explanation for this is that galvanic couples occur between the titanium and iron in the steel. The titanium deposit may be accompanied by a local breakdown of the spontaneous thin passive film (2nm 5nm) formed on the stainless steel. As a consequence, both the titanium and the iron are brought into contact with the solution. 0

20

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corrosion potential (V/ECS)

-0,1 -0,2

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-0,3 -0,4 -0,5 -0,6 -0,7

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-0,8 cold rolling ratio (%)

Figure 5: Effect of cold-rolling on corrosion current density of AISI 316L substrate (white markers) and titanium coating (black markers).

As far as the samples cold-rolled to 20%, 40% and 60% then coated with titanium are concerned, lower corrosion potentials can be observed than on the non-covered steel. On the other hand, AFM investigations reveal that the deposit does not grip well as can be observed in the film morphology (Figure 1). This can be explained by the substantial increase in the corrosion current at higher cold-rolling rates (Figure 5) due to galvanic coupling between the anodic areas (deposit defects) and the rest of the surface. The poor deposit quality may be attributed to the stress field inherent to the cold-rolling process that occurs on the substrate surface. Clearly, if the interfacial forces are high and not reduced by some relaxation mechanism, they may overcome the resistance generated by gripping, particularly if the latter is reduced by the influence of environmental factors such as water and temperature.

4. Conclusion This work has shown the effect of the cold-rolling deformation of the austenitic AISI 316L stainless steel on the physical vapour deposition of titanium film and on its electrochemical behaviour in a simulated physiological environment.

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The micro-structural characterisation shows that the titanium deposit takes the exact shape of the polished surface structure, regardless of the deformation rate of the substrate, but it does not cover all of the surface lands. The film thickness was estimated to about 1µm and the EDX analysis shows that the thickness is more important for the deposit on the deformed substrate. By contrast, in the case of the titanium deposit on AISI 316L stainless steel deformed at different rates, the evolution of the corrosion potential was not found to be homogeneous. The results show that the deformation potential increases from 0% to 20%, then it stabilises until 60%. After that, the corrosion current value for the noncovered substrate increases too much. These results have been correlated to those for micro-hardness, and it appears that the micro-hardness of the titanium deposit varies in a similar way to the evolution of the potential for low deformation rates. The microhardness results obtained for both covered and non-covered substrates lead us to conclude that the deposit generates residual stress at the film/substrate interface. For high deformation rates, although the titanium film remains hard, it seems to have less influence on the substrate. This is possibly due to an absence of cohesion at the interface.

Acknowledgements The authors would like to acknowledge the financial support of the French-Algerian CMEP Project 03MDU586.

References 1. A. Cigada, B. Mazza, P. Pedeferri, D. Sinigaglia, Journal of Biomedical Material Research 11 (1977) 503. 2. A. Barbucci, G. Cesisola, P.L. Cabot, Journal of Electrochemical Society, Vol. (1968) 1. 3. B. Mazza, P.Pedeferri, D. Sinigaglia, A. Cigada, G.A. Mondora, Journal of Electrochemical Society 126 (1979) 2075. 4. E.M. Gutman, G. Solovioff, Corrosion Science 38 (1996) 1141. 5. F. Navaï, O. Debbouz, Journal of Material Science, Vol. (1999) 1073. 6. H. Sedjal, Mémoire de Magister, Génie Mécanique, Université de Béjaia, Algeria, (2004). 7. F. Székely, I. Groma, J. Lendvai, Scripta Materialia 45 (2001) 55. 8. A. Bouzina, C. Brahim, J. Ledion, La Revue de Métallurgie, CIT/Science et Génie des Matériaux, December 1998. 9. U. Kamachi Mudali, P. Shankar, S. Ningshen, R.K. Dayal, Corrosion Science 44 (2002) 2183. 10. L. Peguet, PhD Thesis, INP Grenoble, France (2005).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Effect of Al on the passivity of Ti base implant alloys Sannakaisa Virtanen, Helga Hildebrand, Mariana Ruczickova

University of Erlangen-Nuremberg, Dept. of Materials Science, WWIV-LKO, Martensstr. 7, D-91058 Erlangen, Germany. Abstract - The influence of Al as an alloying element in Ti on the passive behavior was studied. Fast open-circuit activation of native passive films on Ti-6Al-4V, Ti-6Al-7Nb and Ti-50Al - in contrast to commercially-pure Ti - takes place in 0.5 M H2SO4. Passivation in an acidic solution, however, leads to a higher protectiveness of the passive films, as well as to a higher stability against open-circuit activation. Surface analytical experiments were carried out, with the aim to correlate the composition of the passive film with its stability (resistance against activation in acidic solutions) and protective quality. Keywords: Ti, Ti-Al-V. Ti-Al-Nb, implant alloys, passivity, passive films

1. Introduction Titanium alloys such as Ti-6Al-4V and Ti-6Al-7Nb are widely used for biomedical applications. High mechanical stability and corrosion resistance are required, in order to insure negligible ion release from the implant into the human body. Pure Ti shows a better corrosion resistance than the alloys, while the alloying elements (Al, V, Nb) are added to improve the mechanical properties. A release of Al and V from the implant into the tissue is highly undesirable due to the toxic effect of these elements for the human body. Therefore, a detailed understanding of the stability of the passive state is required. For titanium it is known that open-circuit depassivation can take place under deaerated acidic conditions. Most of the more recent studies on the corrosion resistance of Ti and the Ti-base implant alloys have been carried out in

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relatively mild conditions representing a typical chemistry in a "bulk" biological environment - under these conditions all implant materials show a stable passivity. However, under certain conditions, for instance in crevices or other types of occluded cells, much more aggressive environments can be present [1,2], due to pH changes resulting from chemical reactions (hydrolysis of Ti cations [3]) or from inflammation [4,5]. Under such conditions, the relative stability of the different materials could strongly vary; this behavior, however, has not been studied in detail. Nevertheless, it has been reported that activation of Ti and Ti alloys can take place in HCl solutions, and that this activation behavior depends on the alloy composition as well as on the temperature [6]. In the present work, electrochemical and surface analytical investigations of commercially pure titanium (cp-Ti), two implant materials Ti–6Al-4V and Ti6Al-7Nb as well as a Ti-50Al intermetallic compound were carried out. The alloys were investigated in their passive, as well as active states. The main goal was to further clarify the role of the alloying elements in the passive/activetransition of the materials. 2. Experimental Specimens of cp-Ti, Ti-6Al-7Nb and Ti–6Al-4V alloys were used for the investigations. The Al-containing implant alloys show an a/b-structure, Al being enriched in the a-phase and V or Nb being enriched in the b-phase. A Ti50Al intermetallic compound was also studied, to obtain some information on the passive behavior of a high-Al containing material. The microstructure of this materials is very different from the microstructure of the a/b-alloys and, hence, only a rough comparison of the behavior of the different materials can be obtained. Prior to the electrochemical measurements, the samples were ground with SiC up to #4000, then ultrasonically cleaned in ethanol, rinsed with distilled water and dried. Samples were afterwards mounted in an O-ring cell with an exposed area of 1 cm2, and immersed into 0.5M H2SO4. For surface analysis, the samples were polished with diamond paste to 0.25 mm finish. The XPS spectra were obtained on a PHI 5600 spectrometer using monochromated Al x-ray excitation source run at 300 W. Auger spectra were obtained on a PHI 670 spectrometer (acceleration voltage: 10 kV, acceleration current: 10 nA, beam size ~50 nm). 3. Results and Discussion The open circuit potential (OCP) of the different materials was recorded in 0.5 M H2SO4 during a period of 2 hours (Fig. 1). Large differences can be observed between the alloys and cp-Ti: Titanium shows a stable passivity over the 2 h exposure time, whereas the OCP of the Ti-Al materials shifts after a certain period of time from the passive to the active region, i.e. the native oxide layer on the specimens dissolves in the acidic medium causing a dramatic drop in

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potential. The least stable behavior (the shortest activation time) was found for the Ti-50Al intermetallics. -100

Potential (mV Ag/AgCl)

-200

cp-Ti Ti-Al-V Ti-Al-Nb Ti-50Al

-300 -400 -500 -600 -700 -800 -900 0

2000

4000 Time (s)

6000

Figure 1. Open-circuit potential as a function of time for the different materials in 0.5 M H2SO4.

A drastically different behavior can be observed, if the alloys are polarized at a potential in the passive region in the acidic solution prior to the open-circuit exposure. In this case, the alloys TiAlNb and TiAlV remain in the passive state – clearly the passive film formed in the acidic solution in more stable than the native passive film. Contrarily to these two alloys, the Ti-50Al loses its stable passivity, even it was previously passivated under the same conditions as the two implant alloys. A high Al-content in the material hence prevents a formation of a highly-resistant passive film. Impedance spectra measured at the OCP in the acidic solution indicate that the native passive film on cp-Ti shows a somewhat higher lower-frequency resistance than the native film on the Ti-Al-V alloy, demonstrating a better protective quality for the passive film on cp-Ti. Further, the protective quality of the passive film is increased by passivation in the acidic solution, as compared with the film formed in air. The protective quality of the passive films formed in the acidic solution for the three Al-containing materials was found to decrease in the order of Ti-Al-V > Ti-Al-Nb > Ti-50Al, which is in the same order as the relative resistance against open-circuit activation. As the stability of the passive films formed in the acidic solution is strongly enhanced as compared with the native air-formed passive films, preliminary studies on the composition of the passive films by X-ray photoelectron spectroscopy (XPS) were carried out. The results are qualitatively in line with

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data reported in the literature. Typically, the passive film on Ti-base implant alloys has been found to consist mostly of TiO2, with incorporation of oxides of the alloying elements (see e.g. [7]). The XPS results indicate that the native passive film formed in air shows a somewhat higher concentration of Al than the passive film formed in the acidic solution. Concerning the low stability of the passive film formed on the Ti-50Al, the surface analytical results suggest that this could be due to a very high concentration of Al oxides in the film (§50% Alox).

Potential (mV Ag/AgCl)

0 -200

Ti-Al-V -400

Ti-Al-Nb

Ti-50Al

-600 -800 -1000 0

2000

4000 Time (s)

6000

Figure 2. Open-circuit potential as a function of time for the different materials in 0.5 M H2SO4 after the samples were passivated in 0.5 M H2SO4 at E=+0.25 V for 30 min.

As the materials show a multiphase microstructure, the distribution of the alloying elements is expected to be inhomogenous laterally. It has been indeed shown in the literature that the concentration of the alloying elements in the passive films shows laterals variations which reflect the composition of the underlying phase [8,9]. The surface morphology of the activated samples shows that a non-uniform attack has taken place – this could be an indication that the stability of different phases with a different composition may be different. Therefore, preliminary laterally-resolved Auger measurements were carried out on passive and active samples, to obtain first information on the distribution of the alloying elements on the surface. For the passive samples in all cases, clear differences in the ratio of Ti, Al (and Nb/V) in the passive film at different surface locations (different phases) could be detected. For activated samples, the measurements at different sites of the surface show much less variation in the Al and V/Nb concentrations. Most significantly, in the air-formed oxide films on activated samples, no sites of high Al-concentrations could be observed. This

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finding is an indirect indication that selective dissolution of Al out of the highAl containing phases takes place during active dissolution. Even though further experiments are required, to correlate the sites of dissolution with the composition of the substrate and the passive film, the present results nevertheless suggest that the presence of aluminium is detrimental for the stability and protectiveness of passive films on Ti-Al alloys. 4. Summary Fast open-circuit activation of native passive films on Ti-Al-V, Ti-Al-Nb and Ti-50Al takes place in 0.5 M H2SO4. Passivation in an acidic solution leads to a higher protectiveness of the passive films, as well as to a higher stability against open-circuit activation. Preliminary surface analytical experiments suggest that the stability of the passive films is related to the amount of oxidized Al in the film. A more detailed surface analytical study on the distribution of the alloying elements (in depth and laterally) in differently passivated samples is required, to obtain a good correlation between the nature of the passive film and its stability and protectiveness. References 1. 2. 3. 4. 5. 6. 7. 8.

9.

J. L. Gilbert, C. Buckley, J.J. Jacobs, J. Bone and Joint Surgery 76-A (1994) 110 J.R. Goldberg, C. Buckley, J.L. Gilbert, Corrosion Testing of Modular Hip Implants, ASTM STP 1301 (1997). D.G. Kolman, J. Scully, J. Electrochem. Soc. 141 (1994) 10 P.G. Laing, Biomaterials Science and Engineering, Plenum Press, New York (1984). H.-G. Willert, L.-G. Brobäck, G.H. Buchhorn, P.H. Jensen, G. Köster, P. Ochsner, R. Schenk: Clinical Orthopedics and Related Research No. 333 (1996) 51-75 Y. Yu, C.W. Brodwick, M.P. Ryan, J.R. Scully, J. Electrochem. Soc. 146 (1999) 44294438 I. Milosev, M. Metikos-Hukovic, H.-H. Strehblow, Biomaterials 21 (2000) 2103-2113 J. Lausmaa: Journal of Electron Spectroscopy and Related Phenomena 81 (1996) 343361 C. Sittig, G. Hähner, A. Marti, M. Textor, N.D. Spencer, J. Materials Science: Materials in Medicine 10 (1999) 191-198

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Improved adhesion of titanium oxide film to titanium-base alloy by Ti/O compositional gradient using reactive sputter-deposition Tsutomu Sonoda*, Akira Watazu, Kiyotaka Katou, Tadashi Asahina

National Institute of Advanced Industrial Science and Technology(AIST) 2266-98 Anagahora, Shimoshidami, Moriyama-ku,Nagoya 463-8560, Japan

Abstract - Deposition of titanium oxide film having Ti/O compositionally gradient layer onto titanium-based alloy such as Ti-6Al-4V was carried out by reactive DC sputtering, in order to improve not only the biocompatibility of the alloy but also the adhesion at the interface between the deposited titanium oxide film and the alloy substrate with preserving the high hardness and the passivity of the titanium oxide film. The effects of Ti/O compositional gradient on adhesion of the film at the interface to the alloy substrate were investigated by comparing the adhesion of Ti-O compositionally gradient films at the interface boundary with that of Ti-O compositionally constant films at the interface. The compositional gradient was realized by varying continuously the oxygen content in Ar-O2 sputter gas mixture during the sputter-deposition. According to AES in-depth profiles, the oxygen(O) concentration in the deposited film decreased gradually in-depth direction from the surface toward the substrate, confirming that a film with Ti/O compositional gradient, i.e., a Ti-O compositionally gradient film had formed on the alloy substrate, and thereby expecting that the stress concentrated at the interface between the deposited film and the alloy substrate could be relaxed. On the basis of indentation-fracture tests, it was found the compositionally gradient film had more adhesion at the interface than the compositionally constant film, concluding that the Ti/O compositional gradient layer improved the adhesion of the deposited titanium oxide film. Keywords: titanium oxide film, adhesion, reactive sputtering, passivity, biocompatibility

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1. Introduction Functional titanium-base alloy such as Ti-6Al-4V attracting attention as a biomaterial features excellent mechanical properties, corrosion resistance and super plasticity, so enables the forming of denture bases of complicated shapes. However, this alloy contains aluminium and vanadium liable to do serious harm to human bodies1,2, so actual use will require prevention of direct contact with biological tissues. Hence it can be noticed that the hardness of titanium dissolving oxygen atoms increases with increasing the oxygen concentration, and that titanium oxide has high hardness and its coatings enhance the bonding of titanium-base alloy implants to living bone3,4. Thus deposition of titanium oxide film having Ti/O compositionally gradient layer onto titanium-based alloy such as Ti-6Al-4V was carried out by reactive DC sputtering, in order to improve not only the biocompatibility of the alloy5 but also the adhesion at the interface between the deposited titanium oxide film and the alloy substrate with preserving the high hardness and the passivity of the titanium oxide film6. The effects of Ti/O compositional gradient on adhesion of the film at the interface to the alloy substrate were investigated by comparing the adhesion of Ti-O compositionally gradient films at the interface boundary with that of Ti-O compositionally constant films at the interface. The compositional gradient was realized by varying continuously the oxygen content in Ar-O2 sputter gas mixture during the reactive sputterdeposition. 2. Experimental A planar magnetron sputtering system(ANELVA Corp. type SPF-210H) with a 200mm-diameter, 130mm-height stainless steel chamber was used. The planar target used for this study was a 100mm-diameter 99.99mass% pure titanium disk. Ti-6Al-4V alloy substrates(13×9mm , thickness 0.55mm) cleaned with organic solvent were mounted on the water-cooled substrate holder. Ti/O compositional gradient films were deposited onto the substrate by reactive DC sputtering under the Ar-O atmosphere in which the oxygen content was continuously varied with depositing time. Discharge voltage and current for the sputtering were 350V and 1A, respectively. The procedure of the deposition was as follows. First, pure titanium was deposited for 1minute by sputtering in the atmosphere of pure argon. Then, oxygen was gradually introduced into the chamber on condition that the oxygen flow rate was constantly increased for 18.5minutes up to 3.0mL/min, aiming at the formation of Ti/O compositional gradient film by reactive sputtering in the Ar-O gas mixture. Finally, the reactive sputtering under the constant argon flow rate (1.4mL/min) mixed with the constant oxygen flow rate (3.0mL/min) was carried out for 0.5 minute.

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The thickness of deposited film was measured by tracing the substrate-film step using a surface roughness tester. The surface morphology of the film was studied on SEM images. In-depth profiles of the chemical composition (Ti/O ratio) of the film were analyzed by AES with an ion sputter etching method using argon ion beam. Indentation-fracture tests were carried out using a Vickers hardness tester under the load of 50g, to estimate the adhesion of the obtained film to the alloy substrate. 3. Results and discussion The thickness of the obtained film was approximately 3 m. The film appeared to be uniform and adhesive. In regard to the mechanical durability, the film was hard and durable without being damaged by a tip of pincette. The surface of the film was observed to consist of smooth accumulated deposit and fine particles dispersed on the deposit. Under the more detail observation on the topography of the accumulated deposit, some pits and bumps in the order of several microns, which might reflect those existing originally on the substrate surface, were found on the deposit. Thus the obtained film exhibited to be adhesive to the substrate surface. AES in-depth profiles of titanium, oxygen and nitrogen for the deposited film and the substrate are shown in Fig.1. The AES signals for the analysis were Ti-LMM(480eV), O-KLL(510eV) and Al-KLL(1396eV). To convert the signal intensity of each element into its relative concentration, we used RSF (Relative Sensitivity Factor), that is, Ti-LMM:2.836, O-KLL:5.135 and Al-KLL:11.875. And each relative concentration(Ci) was calculated as follows: Ci = ( Ii / Si ) / ( Ij / Sj ), where Ii or Ij is respectively a measured signal intensity of element i or j, and Si or Sj is respectively the relative sensitivity factor of element i or j. In the figure, the Al Auger signal in the deposited layer was assumed to be due to the greater enlargement of noise, because the RSF for Al is much larger, compared with that for Ti or O. And the O Auger signal in the Ti-6Al-4V alloy was also assumed to be mostly due to noise because the signal for real oxygen content in the alloy must be much less. A specific increase of aluminium concentration was detected along the depth direction of the sample. This implies the position of the interface between the film and the substrate. It is shown that the oxygen concentration in the film decreases gradually in depth direction from the surface toward the substrate while the titanium concentration increases gradually in contrast with the oxygen. Therefore it was confirmed that a Ti/O compositional gradient film had formed on the alloy substrate. Then, a specific increase of oxygen concentration was also detected in the vicinity of the interface between the film and the substrate. The amount of the oxygen concentration integrated with depth ranged in the vicinity corresponded to that

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of oxygen concentration integrated with depth ranged in the surface layer of the alloy substrate where oxygen atoms concentrate, which was detected by AES analysis for the surface of the alloy substrate. Therefore it was concluded that the oxygen atoms which had concentrated in the surface layer of the alloy substrate diffused into the titanium layer in the compositional gradient film.

Fig. 1. AES in-depth profiles of titanium, oxygen and nitrogen for the deposited film and the Ti-6Al-4V alloy substrate

An indentation method for determining the adhesion of interfaces between thin films and substrates has been developed by S.S.Chiang5 et al. in 1980. This is achieved by comparing lateral crack lengths for the same load and film thickness. The typical SEM image of 50g indentation-fracture tests on the surface of the Ti/O compositional gradient film is shown in Fig.2(b), compared with that of the Ti-O compositional constant film [Fig.2(a)], which was deposited by reactive DC sputtering under the same sputtering conditions as in this study except the gas mixing condition, i.e., in 1.4mL/min constant Ar+3.0mL/min constant O 琬 mixture. Not only lateral cracks but also radial cracks were observed on the both films. The lateral cracks on the Ti/O compositional gradient film exhibited to be segmental such as circular arcs, while those on the Ti-O compositional constant film exhibited to be continuous such as a circle. Therefore it was found that the compositional gradient film was more adhesive to the alloy substrate than the compositional constant film. Then,

Improved adhesion of titanium oxide film to titanium-base alloy

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the radial cracks on the Ti/O compositional gradient film exhibited to be short propagated and few in number, while those on the Ti-O compositional constant film exhibited to be long propagated and many in number. Therefore it was found that the compositional gradient film was tougher than the compositional constant film.

------------ 10 m Fig. 2. SEM image of 50g indentations on the surface of Ti-O compositional constant film deposited under the same sputtering conditions as in this study except the gas mixing condition, i.e., in 1.4mL/min constant Ar+3.0mL/min constant O琬 mixture(a) and the Ti/O compositional gradient film obtained in this study(b).

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4. Conclusions Coating of Ti-6Al-4V alloy substrates with Ti/O compositional gradient films was carried out by reactive DC sputtering, not only to improve the biocompatibility and the surface hardness of the alloy but also to relax the stress concentrated at the interface between the film and the alloy substrate. The compositional gradient was realized by varying continuously the oxygen content in Ar-O琬 sputter gas during depositing. Under SEM, the surface of the deposited film was found to have fine particles dispersed on a smooth accumulated deposit. Under AES, the oxygen concentration in the film decreased gradually in depth direction from the surface toward the substrate, confirming that a Ti/O compositional gradient film had formed on the alloy, and a specific increase of the oxygen concentration was detected in the vicinity of the interface, concluding that the oxygen atoms which had concentrated in the surface layer of the alloy substrate diffused into the titanium layer in the compositional gradient film. Based on indentation-fracture tests, it was concluded that this depositing method improved the adhesion of the coated film to the alloy substrate and the toughness of the film References [1] S.G.Steinemann and S.M.Perren, Titanium Sci. Technol. 2(1985)1327. [2] P.Galle, Compt. rend. 299(1984)536. [3] A.Dubertret and P.Lehr, Compt. rend. 263(1966)591. [4] T.Kitsugi et al., J. Biomed. Mater. Res. 32(1996)149. [5] S.S.Chiang et al., in: Surface and Interface in Ceramics and Cermics-Metal Systems, Proc. 17th Univ. Conf. on Ceramics, 1980, p. 603.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Liquid crystal behavior in solutions, electrode passivation, and impedance loci in four quadrants C. V. Krishnana,b, Merrill Garnetta a

Garnett McKeen Lab, Inc., 7 Shirley Street, Bohemia, NY 11716-1735 Department of Chemistry, University at Stony Brook, NY 11794-3400, [email protected] b

Abstract - We report here that the phenomenon of impedance with negative resistance and the impedance loci passing through two, three or four quadrants in the complex plane are common in many biological systems such as prothrombin, collagen, lysine, DNA-lysine, DNA-H2O2, EDTA, nitrilotriacetic acid, iminodiacetic acid, methyliminodiacetic acid, and molybdate with FADH2, NADH, lysine, Ni-acetyl-L-lysine, histidine, arginine, 2,3-diaminopropionic acid, 2,4diaminobutyric acid, spermidine, hexanediamine, ethylene diamine, 1,3-diaminopropane, diethylenetriamine, triethylenetetramine, and N,N,N',N'-tetraethylethylenediamine. Impedance loci occurring in two, three or four quadrants depend on the nature of the substrate, counter ions, pH, concentration, hydrogen-bonding characteristics, and the potential of the mercury working electrode. The behavior is limited to a narrow band potential. Many of these systems exhibit liquid crystal characteristics or self-assembly in solution and seem to facilitate global coupling. Keywords: impedance, self-assembly, nonlinear phenomena, bifurcations, global coupling.

1. Introduction The importance of self organized nonequilibrium spatial pattern formations of the interfacial potentials and concentrations of chemicals near the electrochemical interface is now well recognized in electrochemical processes [1-4]. The dynamical spatiotemporal periodicities are attributed to the negative Faradaic impedance (hidden or not) characteristics of the electrode. Impedance spectroscopy is used for gaining insights into the instabilities or bifurcations and to distinguish between saddle-node and Hopf bifurcations. Electrochemical

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oscillators seem to exhibit impedance loci in four quadrants [2-4] as in corrosion or passivation studies [5-8]. We had reported that for palladium lipoic acid complex [9], an experimental chemotherapy agent developed in our laboratory, the impedance loci pass through all four quadrants in the complex plane. We had also observed impedance loci in the first two quadrants for the DNA-H2O2 system [10]. All living systems exhibit spatiotemporal periodicities. To understand the chemical origin of these, electrochemical passivation of metals has been used extensively as a model. We are not aware of any experimental data where the impedance technique has been used to obtain information regarding selforganized dynamical states in biochemical systems. We report here that the phenomenon of impedance loci passing through two, three or four quadrants in the complex plane is common in biological systems. While the details of most of these systems will be discussed in separate manuscripts, here we wish to demonstrate that many of these systems exhibit liquid crystal characteristics or self-assembly in solution that seem to facilitate global electronic coupling. 2. Experimental An EG & G PARC Model 303A SMDE tri-electrode system (platinum counter electrode and Ag/AgCl (saturated KCl, reference electrode) along with Autolab eco chemie was used for electrochemical impedance measurements at 298 K. Palladium lipoic acid complex (1:1) was prepared according to the literature [11]. All the chemicals used were from Sigma and the experimental solutions were prepared in distilled water. No other background electrolyte was used. The solutions were purged with N2 for about 10 minutes before the experiment. Impedance measurements were carried out using about 7 mL solutions in the frequency range 1000 Hz to 5 mHz. The amplitude of the sinusoidal perturbation signal was 10 mV. 3. Results and Discussion The Nyquist and Bode plots for palladium lipoic acid complex are shown in Figure 1. The impedance locus occurring in four quadrants is possibly the first example of a system involving a metal-biomolecule. When the solution is diluted from 0.2 M to 0.04 M, the impedance locus is seen in only the first two quadrants. This kind of impedance plot observed in passivation studies of copper [5] has been attributed to the transition from the nonminimum phasetype to the minimum phase-type and corresponds to the Hopf bifurcation under current control. We also observed the resonance-like maximum of the amplitude characteristic or impedance modulus indicative of harmonic relaxation along with the phase change behavior for the first order (discontinuous) nonminimum phase to the second order (continuous) minimum phase.

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Figure 1: Impedance plots for 0.20 M palladium lipoic acid complex (1:1) at pH 11.45 and 0.04 V. (a) Complex impedance coordinates; (b) Bode coordinates.

The Nyquist and Bode plots, shown in Figure 2 for EDTA is part of our series of studies of molybdate-small molecule interactions [12]. EDTA, involved in oscillations in mitochondria [13], is a versatile ligand and the pH here is much lower than that of palladium lipoic acid. For

Figure 2: Impedance plots for 0.048 M disodium salt of ethylenediaminetetraacetic acid at pH 4.67 and 0.5 V. (a) Complex impedance coordinates; (b) Bode coordinates.

the same EDTA concentration, when the pH is raised to 7.82, we have increased the number of charges on EDTA. Now the impedance locus is confined to the first and second quadrants and the potential is shifted to 0.46 V. For molybdateEDTA complex in the ratio 2:1 at pH 7.95 with the same EDTA concentration, the impedance locus is again limited to the first and second quadrants and the potential is further shifted to 0.08V. These results indicate that the negative impedance or the bifurcation points are sensitive to the charge and nature of the substrate. While the impedance loci in four quadrants are also observed for nitrilotriacetic acid, they are observed in only the first two quadrants for iminodiacetic acid (pH 8.46) and methyl iminodiacetic acid (pH 8.83). The data shown in Figure 3a are part of our series of studies of DNA-small molecule interactions. We have shown two sets of data in these plots to demonstrate the robustness of the results. We did not observe

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Figure 3: Complex impedance plots for (a) 4.9 mg/mL calf thymus DNA, 0.10 M L-lysine, and 0.02M HCl at pH 9.00 and 0.30 V, two sets of data to demonstrate reproducibility and (b) 1.25 mg/mL DNA, 0.05 M sodium acetate, and 0.38 % H2O2 at pH 4.79 and 0.4 V.

this impedance behavior at pH 9.88 with DNA-lysine only, thus showing the importance of protonation and probably counter ion. Another role of the ubiquitous hydrogen peroxide in body fluids is evident from our observation of impedance loci in four quadrants shown in Figure 3b for the DNA-peroxide system. Figure 4a is for the spermidine-molybdate system. Spermidine is known to condense DNA. It can also interact with the molybdate charges with selfassembly in solution. The data in Figure 4b is for L-lysine with and without molybdate. We have also observed impedance loci in four quadrants for molybdate with Ni-acetyl-L-lysine, histidine, ethylenediamine, N,N,N',N'tetraethylethylenediamine, triethylene tetramine and impedance loci in the first two quadrants for prothrombin, collagen, and molybdate with arginine, 2,3diaminopropionic acid, 2,4-diamino butyric acid, diethylene triamine, flavine adenine dinucleotide, and nicotinamide adenine dinucleotide. But with molybdate- N,N'-diethylethylene diamine the impedance locus was seen only in one quadrant. These data demonstrate not only a variety of biological molecules but also the importance of steric hindrance, pH, and hydrogen

Figure 4. Complex impedance plots for (a) 0.06 M spermidine hydrochloride, 0.09 M sodium molybdate, and 0.08 M HCl at pH 10.03 and 0.3 V and (b) (1) 0.19 M L-lysine, 0.095M sodium molybdate, and 0.12M HCl at pH 8.87 and 0.1V; (2) 0.095 M L-lysine and 0.021 M HCl at pH 9.24 and 0.2 V.

Liquid crystal behavior

393

bonding effects. The data with L-lysine and Nİ-acetyl-L-lysine demonstrate the probable role of these involved in bifurcations or oscillations during the histone deacetylation process and in the area of drug development, such as HDAC inhibitors for cancer treatment. 4. Liquid crystal behavior or self-assembly in solutions The phase microscopy pictures (300X) of palladium lipoic acid (1:1) and molybdate-spermidine species are shown in Figure 5. As far as we know the influence of such self- assembled molecules near the double layer has not been investigated on the formation of spatiotemporal

a

b

Figure 5. Phase microscopy 300X. (a) Palladium lipoic acid (1:1) complex, 0.01 M at pH 11.3 (b) 0.06 M spermidine hydrochloride, 0.09 M sodium molybdate, and 0.08 M HCl at pH 10.03.

periodicities. In homogeneous systems, autocatalytic reactions and diffusion resulting from chemical instabilities lead to the formation of spiral waves and other concentration patterns of spatiotemporal phenomena. Our data suggest the electrical phenomena in the electrode/electrolyte interface in the presence of self-assembled systems propagate among the packing units and extend into the bulk. Thus the self-assembly of molecules, especially biological molecules, facilitates local disturbances to be felt at long distances by global coupling. 5. Conclusions Non-stationary regimes far from equilibrium are very prevalent in living systems. We provide here, for the first time, several examples of molecules of biological importance that exhibit impedance loci in four quadrants, a characteristic of systems that exhibit bifurcations and spatiotemporal periodicities. Coherent liquid crystal or self-assembly behavior represents a long range signaling system for biological regulation. Multiquadrant impedance with negative differential resistance quantitates this global coupling.

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References: [1] K. Krischer in: B.E. Conway, J.O’M. Bockris and R. E. White (Editors), in: Modern Aspects of Electrochemistry, No. 32, Plenum Publishers, New York 1999, Chapter 1. [2] M.T.M. Koper, Adv. Chem. Phys., 92 (1996) 161 [3] P. Strasser, The Electrochemical Society Interface, Winter (2000) 46 [4] M.T.M. Koper, J. Chem. Soc. Faraday Trans., 94 (1998) 1369 [5] A. Sadkowski, M. Dolata, and J.P. Diard, J. Electrochem. Soc., 151 (2004) E-20 [6] M. Keddam, H. Takenouti and N. Yu, J. Electrochem. Soc., 132 (1985) 2561 [7] D. D. Macdonald, Electrochim. Acta, 35 (1990) 1509 [8] B.A. Boukamp, J. Electrochem. Soc., 142 (1995) 1885 [9] C.V. Krishnan, M. Garnett, Proc. 1st Spring Meeting of ISE, Abs. P06, Spain 2003 [10] M.Garnett, and C.V.Krishnan, 204th Meeting of Electrochemical Society, Orlando, Florida, Abstract No. 1379(2003) [11] M. Garnett, Palladium Complexes and Methods for Using Same in the Treatment of Tumors and Psoriasis, U.S.Patent, No. 5,463,093, Oct.31(1995) [12] C.V. Krishnan and M. Garnett, 228th American Chemical Society National Meeting, San Diego, Inor. 914, 2005 [13] B. Hess, and A. Boiteux, Annu. Rev. Biochem., 40 (1971) 237

Section E Passivity in High-Temperature Water

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Passive film growth and oxide layer restructuring on stainless steel in a high-temperature borate electrolyte Martin Bojinova, Petri Kinnunenb, Klas Lundgrenc and Gunnar Wikmarkd a

Department of Physical Chemistry, University of Chemical Technology and Metallurgy, 1756 Sofia, Bulgaria, e-mail [email protected] b VTT Technical Research Centre of Finland, FIN-02044 VTT, Espoo, Finland c ALARA Engineering AB, P.O.Box 26, SE-730 50 Skultuna, Sweden d Advanced Nuclear Technology AB, Uppsala Science Park, SE-751 83 Uppsala, Sweden Abstract - The oxide films formed on AISI 316L(NG) in the temperature range 150 300°C in a simulated light water reactor coolant have been characterised by impedance spectroscopy and ex-situ analysis using Auger electron spectroscopy. The results show that the nature of the barrier layer does not change drastically with temperature, although the growth mechanism of the oxide film is different at 150…300°C than at room temperature. A procedure for the calculation of the kinetic constants of the interfacial reactions, as well as the diffusion coefficients of ionic / electronic defects in the oxide has been developed on the basis of the Mixed-Conduction Model for passive films. The effect of temperature on the parameters has been quantified, and their relevance for the corrosion behaviour of stainless steel in a high-temperature electrolyte is discussed. Keywords: stainless steel, passive film, light water reactor coolant, high-temperature borate electrolyte, impedance spectroscopy, ex-situ analysis, conduction mechanism, kinetic model

1. Introduction Well-controlled growth of the passivating film on stainless steels in light water reactors (LWRs) is expected to limit the impact of the coolant on such materials and minimise the concentration of impurities that may reach the nuclear fuel surfaces and become radioactive. The oxide film forming on stainless steel in reactor conditions has a duplex structure [1-8]. The outer layer of the oxide grows via a dissolution/precipitation mechanism, while the inner barrier-like layer grows via a solid-state mechanism [1,5,9-11]. The inner layer contains a significant number of ionic defects which offer routes for ionic species to be transported through the film, making the dissolution of the metal through the film an important process in addition to film growth. The Mixed-Conduction

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Model (MCM) [10,11] treats the barrier layer as a homogeneous medium created and maintained at a steady state thickness via generation, transport and consumption of point defects (interstitial cations, cation and anion vacancies). The MCM includes both interfacial kinetics and solid-state transport as rate determining steps for oxide growth and metal dissolution through the film. The specific aim of the present paper is to elaborate a procedure for the estimation of the interfacial rate constants and diffusion coefficients of the solid-state transport by using in-situ impedance and ex-situ analytical data for hydrothermal oxide films on stainless steel. 2. Experimental AISI 316 L(NG) stainless steel has been studied in a Ti-cladded autoclave at temperatures 150..300°C using a Pt sheet as a counter electrode and an external pressure-balanced 0.1 M KCl/AgCl/Ag electrode as reference. 0.1 M Na2B4O7 deaerated with N2 (99.999%) was used as electrolyte. Impedance spectra in the alloy/oxide/electrolyte configuration were obtained with a Solartron 1287/1260 system in a frequency range 0.02 - 30,000 Hz at an ac amplitude of 20 mV (rms). Impedance spectra in the alloy/oxide/inert metal configuration were performed using an Ir probe as both counter and reference at zero dc current, ac current amplitude of 10 µA (rms). Auger Electron spectra and depth profiles were measured using a Perkin Elmer PHI model 610 scanning Auger microprobe at a base pressure below 2.67 ×10-7 Pa. For depth profiling, sputtering by an Ar+ ion beam (energy 3 keV, current 2.5 mA) was performed. Sputtering rate was 4.5 nm min-1 vs. a Ta2O5/Ta standard. 3. Results Figure 1 shows the impedance spectra at 150°C after 1 h of oxidation at 0.22 V (at the foot of the transpassive region) and subsequent polarisation with 0.1 V/30 min. steps down to the hydrogen evolution region. Qualitatively similar spectra have been measured at the other temperatures. The impedance magnitude at low frequencies does not depend significantly on potential in the passive region, decreasing both at more negative and more positive potentials. The spectra comprise a high-frequency time constant at 1-100 Hz and a low frequency part (0.01 - 1 Hz) with a phase angle around 45°. The high-frequency time constant in the spectra can be ascribed to a relaxation process in the space charge layer, whereas the low-frequency part is best described with a transport impedance in a finite-length layer. shows the impedance spectra measured at 150°C in the alloy/oxide/inert metal configuration after a similar oxidation procedure. These spectra comprise a high-frequency time constant (100-5000 Hz) that can also be related to electronic processes in the space-charge layer, and a low-frequency time constant (1-50 Hz) described once more by a finite-

399

Passive film growth and oxide layer restructuring…

length transport impedance. The characteristic frequencies of the transport-type time constant in the symmetrical configuration are ca. two orders of magnitude higher than those in the assymetrical configuration (cf. Figure 1-). 104

-75

AISI 316L(NG) / 0.1 M Na2B4O7 , 150 °C

AISI 316L(NG) / 0.1 M Na2B4O7, 150 °C

-50

phase / deg

|Z| / Ohm cm2

103

102 0.22 V 0.02 V -0.18 V -0.38 V -0.58 V -0.78 V

101

100 10-2

10-1

100

0.22 V 0.02 V -0.18 V -0.38 V -0.58 V -0.78 V

-25

10 1

10 2

10 3

10 4

0 10-2

10 5

10-1

Frequency / Hz

100

101

10 2

10 3

10 4

10 5

frequency / Hz

Figure 1. Electrochemical impedance spectra of AISI 316L(NG) in 0.1 M Na2B4O7 at 150 °C measured in the alloy/oxide/electrolyte configuration. Points - experimental values, solid lines best-fit calculation according to the procedure outlined in the text. -75

AISI 316L(NG) / 0.1 M Na2B4 O7, 150 °C

-0.14 V -0.04 V 0.06 V 0.16 V 0.26 V 0.36 V

100

-55

phase / deg

|Z| / Ohm cm2

101

10 -1

-35

-0.14 V -0.04 V 0.06 V 0.16 V 0.26 V 0.36 V

-15

AISI 316L(NG) / 0.1 M Na2B 4O7, 150 °C 10-2

10-1

100

10 1

10 2

frequency / Hz

10 3

10 4

10 5

10-2

10-1

100

10 1

10 2

10 3

10 4

10 5

frequency / Hz

Figure 2. Electrochemical impedance spectra of AISI 316L(NG) in 0.1 M Na2B4O7 at 150 °C measured in the alloy/oxide/inert metal configuration. Points - experimental values, solid lines best-fit calculation according to the procedure outlined in the text.

This means that the transport process associated with this time constant is related to current carriers that attain their steady state transport rate faster than those in the assymetrical configuration. This observation implies that in the alloy/ oxide / electrolyte configuration the ionic point defects are the unblocked current carriers, whereas in the alloy/oxide/inert metal configuration this role is played by the electronic defects [11]. The fact that the impedance of this film does not change significantly with temperature suggests that its electrical and electrochemical properties remain qualitatively unaltered.

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Figure 3 summarises the Auger depth profiles of the atomic concentration of oxygen, as well as the relative concentrations of Fe,Cr and Ni normalised to the total amount of metallic elements for the films formed at 150 °C. Estimates of film thickness obtained by computing the depth at which the oxygen concentration drops to half of its maximum value show that the thickness in general increases with potential (from -0.24 to 0.16 V) and temperature (from 150 to 300 °C). The films in the passive region are relatively thin and Crenriched. On the other hand, the transpassive films are depleted in Cr and enriched in Fe and especially in Ni at temperatures above 200 °C. Accordingly, a transformation from a Cr-like to a Ni-like inner layer seems to take place at the passive-to-transpassive transition.

0

70

-0.44 V 50 -0.24 V -0.04 V 40 0.16 V

40 30

30 20

20

10

10 10

20

30

40

depth / nm

50

60

0

AISI 316L(NG) / 0.1 M Na2 B4O7 , 150 °C

40

-0.44 V -0.24 V -0.04 V 0.16 V

35

60 50 40

30 25 20

30

15

20

10

10 0

5

0

10

20

30

40

50

normalised Ni content,%

60

50

70

normalised Fe content,%

AISI 316L(NG) / 0.1 M Na2B4 O7, 150 °C

normalised Cr content, %

oxygen concentration, %

60

0 60

depth / nm

Figure 3. Auger depth profiles of the atomic concentration of oxygen and the normalised content of iron (left), chromium and nickel ( right) for oxides formed on AISI 316L(NG) in 0.1 M Na2B4O7 at 150 °C for 72 h. Estimated positions of the alloy/oxide interface are indicated by arrows.

4. Discussion A picture of the oxide formed on stainless steel in LWRs is presented in Figure 4 [10,11]. The barrier layer is considered to be a normal spinel of the FeCr2O4 type [2,7,8]. Its growth proceeds via the reaction sequence k2-k4 coupled by the oxygen vacancy flux. As the growth of the barrier layer is completed in a short time scale, it is neglected as a slow reaction in parallel with metal dissolution through the layer. The outer layer is considered to be an inverse spinel of the NiFe2O4 type [8]. It grows via dissolution-precipitation mechanism, limited by the reaction sequence k1i-k3i-kr coupled via the interstitial cation flux. As the precipitation reaction is governed by the solubilities of Fe and Ni, it is near equilibrium enabling the formation of large crystallites [1,2,8]. The transpassive dissolution of chromium (reaction sequence k1-k3 coupled via the cation vacancy flux) is also included in the model.

Passive film growth and oxide layer restructuring…

V ··O

k2

CrIIICr

Crm k1

Vm

k1i

eFe Steel

V ··O OO

k4

CrIIICr

V ´´´ Cr

V ´´´ Cr

Fe ··i

Fe ··i

Ni ··i

Ni ··i

k3

H+

H2O 6+

Craq

k3i kr

Barrier inner layer

401

2+ Fe y+ aq ,Niaq

Precipitated outer layer

Electrolyte

Figure 4. A simplified picture of the main processes taking place during oxidation of stainless steel in high-temperature electrolyte according to the Mixed-Conduction Model. Table 1. Kinetic parameters for the oxide growth and restructuring determined by the proposed calculation procedure. Parameter

150 °C

2 -1

-13

200 °C

250 °C

1.5·10

2.0·10

4.0·10

6.0·10-13

Di / cm2s-1

2.5·10-15

6.5·10-15

0.9·10-14

2.5·10-14

k1i / mol cm-2s-1

3.2·10-12

1.0·10-11

3.5·10-11

6.7·10-11

k3i0 / cm s-1

4.0·10-10

6.0·10-9

3.0·10-8

4·10-7

a

0.81

0.83

0.85

0.88

5

E / V cm b3i / V 0

-2 -1

2.2·10

2.0·10

1.4·10

0.8·105

7.3

7.7

9

11

-12

5

-13

300 °C

De / cm s

-1

-13

-12

5

-11

k3 / mol cm s

4.0·10

6.0·10

3·10

7.0·10-11

k1 / cm s-1

7.0·10-10

1.0·10-9

1.5·10-9

4·10-9

b3 / V

4.3

4.0

3.9

3.7

A fit of all the impedance data at a given temperature to the transfer functions of the model [11] gives the possibility to compute the rate constants k1i, k1, k3i and k3 (the last two potential dependent with exponential factors b3i and b3), the diffusion coefficients De and Di, the field strength E, and the polarisability of the barrier layer /electrolyte interface, a. The rate constant of the precipitation reaction, kr, was estimated from the solubility of NiFe2O4 [12]. Calculated spectra are shown in Figure 1- with solid lines and demonstrate the good correspondence between theory and experiment. Using the obtained set of estimates, the steady state currents and the film thicknesses at all studied temperatures have also been predicted. Calculations point to a bigger influence of the reactions involving interstitial cations (k1i-k3i) on the behaviour of stainless

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steel when compared to the reactions involving cation vacancies (k1-k3). The field strength is an order of magnitude smaller at high temperatures than at room temperature, which can be explained by an increased defectiveness of the film at higher temperatures [10]. The relatively large values of Di suggest a grainboundary diffusion mechanism in agreement with the presumed nanocrystalline structure of the barrier layer. The relatively low values of De can be tentatively explained by assuming that the mobile ionic defects act as effective traps for electronic carriers. 5. Conclusions The MCM treats the inner layer of the oxide as a finite homogeneous medium created and maintained at a certain steady state thickness via the generation, transport and consumption of ionic point defects (interstitial cations, cation vacancies and anion vacancies). The porous overlayer formation is modelled as a first-order reaction of reprecipitation. The kinetic parameters of the rate limiting steps were determined in the temperature range 150-300°C by comparison of the model predictions to experimental impedance spectra in two measurement configurations. The obtained estimates can successfully account for both the quasi-steady state current and the film thickness at all temperatures. As a further step in the present investigation, on the basis of the obtained database of kinetic constants and transport parameters, an assessment of the rates of transport and incorporation of minor species in the compact oxide layer is foreseen. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12.

B.Stellwag, Corros. Sci., 40 (1998) 337. Y.-J. Kim, Corrosion, 51 (1995) 849. Y. Hemmi, Y. Uruma and N. Ichikawa, J. Nucl. Sci. Technol., 31 (1994) 443. H. Inagaki, A. Nishikawa, Y. Sugita and T. Tsuji, J. Nucl. Sci. Technol., 40 (2003) 143. D. Lister, E. McAlpine and R. Davidson, Corrosion Science 27 (1987 113. M. Montemor, M. Ferreira and M. Da Cunha Belo, Corrosion Science 42 (2000) 1635. M. Montemor, M. Ferreira, M. Walls, B. Rondot and M. Da Cunha Belo, Corrosion 59 (2003) 11. S. Ziemniak and M. Hanson, Corrosion Science 44 (2002) 2209. J. Robertson, Corros. Sci., 32 (1991) 443. B. Beverskog, M. Bojinov, P. Kinnunen, T. Laitinen, K. Mäkelä and T. Saario, Corros. Sci., 44 (2002) 1923. I. Betova, M. Bojinov, P. Kinnunen, K. Mäkelä and T. Saario, J. Electroanal. Chem. 572 (2004) 211. Y. Hanzawa, D. Hiroishi, C. Matsuura and K. Ishigure, Nucl. Sci. Eng. 124, (1996) 211.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Kinetics of passivation of nickel-base alloys (Alloy 600 and Alloy 690) in high temperature water A. Galtayries,a,* A. Machet,a,b P. Jolivet,b P. Scott,b M. Foucault,c P. Combrade,c P. Marcus a,* a

Laboratoire de Physico-Chimie des Surfaces, CNRS (UMR 7045), Ecole Nationale Supérieure de Chimie de Paris, 11 rue Pierre et Marie Curie, 75005 Paris, France * [email protected], [email protected] b Framatome ANP, Tour Framatome, 92084 Paris La Défense, France c Framatome, Centre Technique, 71205 Le Creusot, France

Abstract - The kinetics of passivation, in high temperature and high pressure water, of nickel-chromium-iron alloys (Alloy 600, Alloy 690) have been investigated by X-ray Photoelectron Spectroscopy (XPS). The samples have been exposed for short (0– 8 min) and longer (0 – 400 hours) time periods to high temperature (325°C) and high pressure water, with boron and lithium, under controlled hydrogen pressure. The fit of the data for short passivation times by three classical kinetics models (parabolic, logarithmic and inverse logarithmic laws) and the extrapolation to longer times reveal that the kinetics of passivation (formation of the internal Cr2O3 layer) of Alloy 600 and Alloy 690 in high temperature and high pressure water are in good agreement with a logarithmic model. Keywords: Alloys 600 and 690; passivation ; X-ray Photoelectron Spectroscopy (XPS); high temperature, high pressure water ; PWR ; kinetics

1. Introduction The oxidation behaviour of nickel-base alloys in high temperature and high pressure water, simulating the primary circuit of steam generators (SG) of pressurised water reactors (PWR), has been studied by several authors [1-5]. It is generally recognised that [6] the passivity of this alloy in such conditions is due to the formation of a chromium-rich oxide layer which provides a diffusion barrier and reduces the corrosion rate. The passive layer reduces the release of corrosion products, such as nickel cations, in primary side water. It is of crucial

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importance because of the activation of 58Ni into 58Co in the primary circuit, which increases the global radioactivity of the primary circuit of PWRs. Another important application of a better understanding of the formation of the passive film is the understanding of stress corrosion cracking mechanisms [6]. Very few published papers deal with the initial stages of passivation in high temperature and high pressure water. The aim of this work was to determine the kinetic law for the growth of the barrier oxide layer for both short and longer times. Prior identification of the nature of the oxide layer on Alloy 600 and Alloy 690 was sought by using X-ray Photoelectron Spectroscopy (XPS). 2. Experimental Polycrystalline samples of commercial Alloy 600 (Ni-16Cr-9Fe (wt %)) and 690 (Ni-29Cr-11Fe (wt %)) were mechanically polished to a 1 mm diamond finish. Prior to the short passivation time periods, further cleaning was performed, in the XPS spectrometer, with argon ion sputtering (4 kV, 0.9 mA). The short passivation times in high temperature, high pressure water (325°C, ~155x105 Pa) were performed in a titanium microautoclave, allowing sample transfers without air exposure. The longer passivation times were performed in a static autoclave, and the samples were rinsed, dried and transferred in air before XPS analysis. In both systems, the aqueous solution simulating unsaturated PWR primary water conditions contained 2 mg.l-1 Li and 1200 mg.l-1 B. A hydrogen overpressure of 0.3x105 Pa was maintained in the autoclave. XPS analyses of the Ni 2p, Cr 2p, Fe 2p, O 1s, C 1s, B 1s and Li 1s core levels were recorded with a VG ESCALAB Mk II X-ray photoelectron spectrometer, using an AlKa or MgKa radiation. For longer passivation times, the ion sputtering depth profile mode was used, with 3 keV argon ions, and a target current of 2.5µA.cm- 2. 3. Results and Discussion 3.1. Alloy 600 From the examination of the Cr 2p, Ni 2p, O 1s XPS core level spectra, obtained after short oxidation times, three compounds have been identified in the passive layer: Cr2O3, Cr(OH)3, Ni(OH)2 [7]. In agreement with the angledependent XPS measurements, a layer model, based on the stratification of the passive layer, has been adopted to describe the oxide layer. It consists of an outermost layer of Ni(OH)2 and/or Cr(OH)3, and an inner Cr2O3 oxide layer, in contact with the alloy. Iron oxide was not included in this model, due to its low

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Kinetics of passivation of nickel-base alloys in high temperature water

concentration in the alloy (<5 at.%), and to the difficulty to locate the small amount of detected Fe2O3. 3

Thickness of the chromium oxide layer (nm)

1

3

2

2

plateau

1 Experimental data for Alloy 600 Inverse Logarithmic curve fitting Parabolic curve fitting Logarithmic curve fitting

0 0

1

2

3

4

5

6

7

8

9

10

Passivation time (min) Figure 1: Measured thickness of the Cr2O3 internal layer formed on a Ni-16Cr-9Fe (wt. %) alloy as a function of passivation time in the microautoclave, and the parabolic, logarithmic and inverse logarithmic fitting of the growth of the Cr2O3 layer in Step 3.

The thickness of the Cr2O3 internal layer has been systematically calculated from the XPS intensities of the metallic and oxide features, using the layer model. Figure 1 displays the results corresponding to the Cr2O3 inner layer, plotted as a function of passivation time. Three domains have been identified: the formation of chromium oxide (1), a plateau corresponding to a non-growing 1 nm thick Cr2O3 oxide layer (2), and a re-oxidation step (3), with the further growth of the Cr2O3 layer. After 8 minutes of passivation, the oxide layer is still growing, exhibiting a duplex layer with islands of Ni(OH)2 on top of the continuous inner layer of Cr2O3 (2 nm thick). Figure 1 also displays the curve fitting of the kinetics of the Cr2O3 growth by three classical kinetic models: parabolic, logarithmic and inverse logarithmic. At this point, it is not possible to discriminate between the three growth laws. However, the discrimination becomes possible by extrapolating to longer oxidation times and comparing with the experimental data.

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To determine the thickness of the chromium oxide inner layer at longer oxidation times, from the XPS profiles, we have used the sputtering time between the maximum of the O2- signal intensity (when the chromium oxide is

Thickness of the chromium oxide layer (nm)

160 Experimental data for Alloy 600 Inverse Logarithmic curve fitting

140

Parabolic curve fitting Logarithmic curve fitting

120 100 80 60 40 20 0 0

100

200

300

400

500

Passivation time (h) Figure 2: Kinetics of passivation of the Ni-16Cr-9Fe (wt. %) alloy in high temperature (325°C) water: thickness of the inner Cr2O3 layer vs. passivation time. The three curves extrapolated from the data obtained for short oxidation times are shown: parabolic, logarithmic and inverse logarithmic models.

reached), and the intensity at half maximum (assigned to the oxide/alloy interface). The sputtering time is then converted into a thickness, using the calibration obtained from Nuclear Reaction Analysis (NRA) [8]. To relate the data of the kinetics of passivation of the alloy for short and long oxidation times (up to 400 hours), the three kinetic laws fitting the data for the short oxidation times have been extrapolated, using the initial value of the Cr2O3 oxide layer measured by XPS (blank test). The three calculated curves are plotted on Figure 2. It appears clearly, from the comparison of the fitted and experimental data, that the parabolic and inverse logarithmic laws are out of range, while a satisfactory agreement is obtained with the logarithmic law: d Cr2O3 = 6 + 0.39 ´ ln(4.06 × t + 1) [9].

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Kinetics of passivation of nickel-base alloys in high temperature water

3.2. Alloy 690 The same experimental approach has been adopted for Alloy 690. The XPS spectra, obtained after short oxidation times, have shown that the passive layer is composed of Cr2O3 (inner layer) and Ni(OH)2 (outer layer). Therefore, the layer model of the passive film has been used to determine the equivalent thickness of chromium oxide (and of nickel hydroxide). Contrary to Alloy 600, the changes of the thickness as a function of oxidation time do not reveal the reoxidation (step 3) with the growth of the Cr2O3 layer. Thus it was not possible to determine the growth law. We decided to use the data obtained, in the same conditions, for short oxidation times of pure Cr (not shown here). They revealed the continuous growth of a Cr2O3 layer (with an outer layer of Cr(OH)3), that could be fitted by the same three classical growth laws. As with Alloy 600, the discrimination becomes possible by extrapolating to longer oxidation times and comparing with the experimental data. A satisfactory agreement with the logarithmic model ( d Cr2O3 = 4 + 0.77 ´ ln(4.06 × t + 1) ) was obtained. The parabolic model could clearly be excluded.

Thickness of the Cr2O3 layer (nm)

180

parabolic m odel (param eters obtained from pureCr)

Alloy 690

160 140 120 100 80

logarithm ic m odel (param eters obtained from pureCr)

60 40 20 0 0

50

100

150

200

250

300

350

400

Passivation time (h)

Figure 3: Kinetics of passivation of Ni-29Cr-11Fe (wt. %) alloy in high temperature (325°C) water: thickness of the inner Cr2O3 layer vs. passivation time. The extrapolations from the data obtained for short oxidation times of pure Cr are used to calculate the parabolic, and logarithmic models.

For Alloy 690, the experimental data of the Cr2O3 equivalent thickness, obtained from the XPS depth profiles (long oxidation times), were fitted by both

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the parabolic and logarithmic models, obtained from short oxidation times of pure chromium. Figure 3 shows the calculated curves as well as the experimental data. The parabolic model does not fit the data but the logarithmic model gives a satisfactory agreement ( d Cr2O3 = 7 + 0.77 ´ ln(4.06 × t + 1) . For Alloy 600, Alloy 690 and pure chromium, the comparison of the extrapolated growth law to longer times shows that in the range of passivation times up to 400 hours, it is possible to unambiguously rule out the parabolic and inverse logarithmic laws (not shown here for Alloy 690). Only the logarithmic law fits well the experimental data. The mechanism associated to a logarithmic law involves the mobility of ions in a high field (with the tunneling of electrons through the oxide). The thicknesses of 10 nm (Alloy 600) and of 30 nm (Alloy 690), obtained here for the Cr2O3 oxide layer, seem too large for the electron tunneling effect. However, one has to take into account the fact that the chromium oxide in Alloy 600 is crystalline and grain boundaries are present in the oxide layer [7]. Such defects, and possibly other defects, may allow the transfer of electrons. 4. Conclusion The kinetics of passivation in high temperature and high pressure water of a Ni16Cr-9Fe and a Ni-29Cr-11Fe (wt. %) alloys (and pure chromium as a reference sample) have been investigated. The composition and the thickness of the surface oxide layers have been measured for very short oxidation times (minutes) and longer oxidation times (up to 400 hours). The thickness of the inner chromium oxide barrier layer has been measured as a function of oxidation time. From the three kinetic laws that fit well the experimental data for the short oxidation times, only the logarithmic law can be retained from the comparison of the extrapolated growth laws with the experimental data for long oxidation times. The logarithmic law is characteristic of an apparent mechanism of high charge field effect for the growth of the internal Cr2O3 layer. This approach has revealed that it is possible to follow the kinetics of the growth of the inner Cr2O3 layer in the initial stages of passivation in high pressure, high temperature water. References 1. J. E. Castle, C. R. Clayton, Passivity of Metals, R. P. Frankenthal and J. Kruger Eds., The Electrochemical Society, Princeton, N. J., USA, (1978), 714. 2. N. S. McIntyre, D. G. Zetaruk, D. Owen, J. Electrochem. Soc., 126 (1979) 750. 3. C. Y. Chao, L. F. Lin, D. D. Macdonald, J. Electrochem. Soc., 128 (1981) 1187. 4. R. L. Tapping, D. Davidson, E. McAlpine, D. H. Lister, Corros. Sci., 26 (1987) 563.

Kinetics of passivation of nickel-base alloys in high temperature water

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5. P. Combrade, M. Foucault, D. Vançon, P. Marcus, J. M. Grimal, A. Gelpi, Proc. 4th Intern. Symp. Environ. Degradation of Materials in Nuclear Power Systems-Water Reactors, D. Cubicciotti Ed., NACE, 1989, 79-95. 6. F. P. Ford and P. L. Andersen, in Corrosion Mechanisms in Theory and Practice - Second Edition, Revised and Expanded, Ed. P. Marcus, Marcel Dekker, Inc., New York (USA), 2002, and reference therein. 7. A. Machet, A. Galtayries, S. Zanna, L. Klein, V. Maurice, P. Jolivet, M. Foucault, P. Combrade, P. Scott, P. Marcus, Electrochimica Acta, 49 (2004) 3957. 8. A. Machet, A. Galtayries, P. Marcus, A. Gelpi, C. Brun, P. Combrade, J. Phys. IV, 11 (2001) 79. 9. A. Machet, A. Galtayries, P. Jolivet, M. Foucault, P. Combrade, P. Scott, P. Marcus, EFC Series, European Federation of Corrosion, "Corrosion Issues in Light Water Reactors: Focus on Stress Corrosion Cracking and Field Experience", in press.

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Behaviour of oxide films in boric acid – lithium hydroxide solutions V. Ignatova*, M. Vankeerberghen, S. Gavrilov, R.-W. Bosch, S. Van Dyck, S. Van den Berghe SCK•CEN, Boeretang 200, B-2400 Mol, Belgium, *[email protected]

Abstract - In this paper we try to relate Stress Corrosion Cracking (SCC) initiation in stainless steels (SS) to the performance and nature of the anodic oxide film that develops during exposure to 300°C water containing boric acid and lithium hydroxide in an autoclave. In-situ electrochemical impedance spectra (EIS) have been obtained while exposing the samples. EIS reveals the behaviour of the oxide layer and the electrochemical processes that take place at the material-oxide-solution interface. After these exposures the oxide films were analyzed by X-ray photoelectron spectroscopy (XPS) to obtain the composition of the oxide film. The results showed that the oxide layers consisted of a mixed oxihydrate (iron-nickel based), dominating at the subsurface region and in larger depth, mainly consisted of chromium-iron oxides. The behaviour of the oxide film is being modelled using an extension to the point defect (PDM) / mixed conduction (MCM) series of models describing interface chemistry and vacancy and electron transport in the film. Keywords : Oxide films, XPS, EIS, Modelling

1. Introduction Many engineering alloys, such as austenitic stainless steels and nickel-based alloys protect themselves from their environment by forming a passive oxide layer [1]. This passive layer makes the material resistant to uniform corrosion but, potentially, makes it susceptible to localised corrosion such as crevice corrosion or stress corrosion cracking [1]. Initiation of localised corrosion is related to the stability of the surface film. Film stability is influenced by the mass and charge transport processes through the film and the interfacial elechtochemical process as described in the point defect model [3]. These processes lead to accumulation of vacancies which might lead to en eventual breakdown of the film. The key issues for describing the film in terms of the point defect model are the nature of the charge carriers, their concentration and mobility in the film. In order to characterise these film properties, complementary experimental techniques are required. In situ measurements, such as electrochemical impedance spectroscopy, can give information about the net exchange rates at the interface between the film and the environment, as

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well as on the electric properties of the film itself. Ex-situ characterisation of the film by surface analysis techniques, such as X-Ray photo-electron Spectroscopy (XPS), yield information about the type and valence of the ionic species in the film, as well as the composition of the film itself. This work is a first attempt to obtain a qualitative picture of the metal-oxide-solution interface. In the future the model will be extended to include more charge carriers, mass transport through the solution, the porous nature of the oxide layer, etc.

2. Experimental approach of characterisation of the oxide layers 2.1. In-situ measurements The stainless steel 316 samples were polished and then oxidised in water with a temperature of 300°C for one week. The test solution contained 2000 or 5000 ppm boric acid and was deaireated with nitrogen in a flow through autoclave. EIS has been used to get a picture of the electrochemical response of the metal-oxide-solution interface. The use of frequencies over a large range of values (for example from 100 kHz to 10 mHz) allows to separate the different processes that take place at this interface. For instance the high frequency response can be related to thickness and composition of the oxide layer which has a much smaller capacitance than the electrochemical double layer. The intermediate frequency response can be related to the charge transfer kinetics and the capacitance of the electrochemical double layer. The low frequency response generally reflects processes that are related to mass transfer. As these processes can interfere with each other a complete physical description of the whole impedance spectra can be quite complicated [3,6]. In this first attempt we try to get good quality EIS data for stainless steels exposed to high temperature water in an autoclave. An example of EIS data is shown in Figure 1. This Nyquist diagram was taken from a stainless steel 304 electrode in PWR water at 300°C. A difference in water chemistry has been obtained by purging either with nitrogen gas (oxygen free) or hydrogen gas. Two different concentrations of hydrogen have been used to see their influence on the impedance response. A first qualitative interpretation shows that there is a clear difference in response between the nitrogen and hydrogen condition. At high frequencies the results are a bit noisy, but still the small semi-circle that can be related to the oxide layers can be noticed. At intermediate and low frequencies this difference (Fig.1) can be related to a difference in electrochemical behaviour. The latter can be explained by the dominance of the hydrogen oxidation/reduction reactions in hydrogenated water at elevated temperatures. The difference between the two hydrogen concentrations is also remarkable. The impedance decreases significantly with hydrogen concentration, which makes sense as the hydrogen reaction goes faster with increasing concentration. Unfortunately the data of the high frequency part is quite scattered and so can not be used to determine a difference in oxide layer thickness and/or composition. Furthermore the aim of the impedance measurements is to try to correlate the electrochemical response with the XPS results. For instance a compact oxide layer will give another result in the Nyquist diagram then a thick porous oxide layer [7]. Generally, the low frequency responce is thought to be related to interfacial reactions [8] or the electric properties of the compact inner layer [9], whereas the high frequency response is thought to be related to the total oxide layer [9]. Therefore continuous impedance measurements are underway to investigate the formation of the oxide layer in time.

413

Behaviour of oxide films in boric acid – lithium hydroxide solutions

20

6000 nitrogen

[H2] 17 cc/kg

[H2] 26 cc/kg

[H2] 31.5 cc/kg

5000

0

-20 3000 -40 2000 -60

1000

0 1.E-03

1.E-02

1.E-01

1.E+00

1.E+01

1.E+02

-80 1.E+03

Frequency, Hz

Figure 1. Bode plot of stainless steel 316 in PWR water at 300°C with nytrogen and three different hydrogen concentrations.

2.2. Ex-situ measurements The depth profiles shown below were carried out on an upgraded ESCALAB – 220i-XL ThermoVG Scientific instrument, using monochromated Al Kg X-rays (energy of 1486.6 eV), with an energy spread of about 1 eV FWHM. The ion beam for layer-by-layer erosion is operating with 3 keV Ar+ ions, which is high enough to cause ion-beam mixing, visible in the long oxygen tails in the profiles. The depth calibration is obtained by direct measurements of the crater depth with surface profilometer. A typical result of the XPS study of the chemical composition of the oxide layer is shown on Fig. 2. The signals, which are displayed against the sputtering time, correspond to the deconvoluted lines of the raw spectra of the metal lines Cr 2p3/2, Ni 2p3/2, and Fe 2p3/2. The identification of the peaks is based on the chemical shift and the overall distribution in depth. For example, in the OH group, most probably Ni(OH)2 or NiOH contribute, because it can be identified together with the Ni metal peak. Also some contribution of the iron group is present, but the shape and the position of the iron compounds do not allow separation of the hydroxides from the oxides. In contrast, Cr mainly forms oxide in the form of Cr2O3. The latter conclusion is confirmed by the coinciding shape of the lines Cr(III) and O1s(II). The surface is enriched with highly oxidised watercontaining Ni-Fe hydroxides, which with the depth are reducing to lower oxidation states, but this is partly due to the oxygen preferential sputtering. The maximum in the distribution of Cr2O3 is observed at about ~ 120-150 nm below the surface, i.e. closer to the interface – oxide – metal substrate. The surface is also relatively enriched with Ni, also present in a metal form, while Cr and Fe in pure metal form appear only after sputtering

Phase shift, °

Zmod, Ohm

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V. Ignatova et al.

414

of the first 120 nm. This result suggests that the film on this sample consists of two layers, as is often reported in literature [10].

depth, nm 0

50

100

150

200

250

300

70

O1sOH O1sII Fe2p

60 50 40 30

at. concrt.%

20 10 0 30

Cr2p Cr2pIII Ni2p Ni2pOH2 Fe2pOOH

25 20 15 10 5 0 0

500

1000

1500

2000

sputtering time, s

Fig. 2 AISI-316 XPS depth profile with total sputtering time of 2000 sec. and total sputtered depth of 280 nm

3. Computer simulation approach The compact layer, or barrier layer of an anodic oxide film, is usually thought of as a semiconductor [11, 11, 12]. In this layer there exists transport of charged carriers such as anion

415

Behaviour of oxide films in boric acid – lithium hydroxide solutions

vacancies, cation vacancies, cation interstitials, electrons and holes. The layer interacts with the substrate material and the solution at its interfaces. Recently we developed a 1D steady-state model for a bicarrier monolayer oxide film [14]. The modelling of the porous outer layer was not taken into account. The use of Poisson's equation allowed us to distribute the potential drop over the metal-film interface, across the barrier layer and over the film-solution interface without the assumption of constant field strength across the film or of a film-solution potential drop which is linearly dependent on the applied electrode

240 elements across film 3.E+06

1.5E+07

1.E+06

1.0E+07

0.E+00

-1.E+06

5.0E+06

Field strength (V/m)

Charge density (C/m³)

2.E+06

-2.E+06

-3.E+06 0.0E+00 0.E+00 1.E-09 2.E-09 3.E-09 4.E-09 5.E-09 6.E-09 7.E-09 8.E-09 9.E-09 1.E-08

Figure 3 Typical charge density and field strength profile across a 1D bicarrier monolayer oxide film. potential. The model containing two mass transport equations, Poisson's equation and appropriate boundary conditions was solved using finite-elements. Figure 3 shows typical charge density and field strength profiles across the film. The finiteelement discretisation and solution technique makes the model readily expandable to include multiple carriers, multiple layers and carrier transport in 2D/3D geometries and transient calculations. Transient calculations can be compared to EIS and film growth measurements. Calcations with multiple layers can be compared to XPS measurements. The number of calculation layers (intra-layer the processes are fixed and inter-layer they can vary) is to be selected on the basis of literature, structural and XPS information.

4. Conclusion We propose a combined approach for studying oxide scales formed on stainless steels surface in exposure to 300°C water containing boric acid and lithium hydroxide in an autoclave. The EIS measurements are obtained in-situ in order to relate the behaviour of the oxide layer to the electrochemical processes taking place during the exposure. Furthermore, ex-situ analysis by XPS

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has been performed in order to determine a depth profile of the composition of the oxide layer. Additionally a point defect computer model is used to simulate the oxide layer behaviour in respect to the charge transfer through the interfaces. The first EIS results showed that the high frequency part of the Nyquist diagram confirms formation of an oxide layer; however the data are too spread to be interpreted in terms of thickness and composition. Complementary, XPS profiles provides well defined composition depth profiles, possible to be quantified both in depth and composition (molecular speciation) within some inherent limits. Preliminary results of the PDM

simulation reveal the typical charge distribution in 1D case for a bicarrier monoxide. The next step will be to extend the model for multilayer and multicarrier case and to fit the calculations to the experimental observations.

References 1. 2. 3. 4. 5. 6. 7. 8.

9. 10. 11.

12.

13. 14.

M.G. Fontana, Corrosion Engineering, McGraw-Hill, ISBN 0-07-100360-6. A.J. Sedriks, Corrosion of Stainless Steels, John Wiley & Sons, ISBN 0-471-00792-7. D. D. Macdonald, J. Electrochem. Soc., vol. 139, n°12, December 1992, p.3434-3449. J.R. Macdonald, Impedance Spectroscopy, Emphasizing Solid Materials and Systems, John Wiley & Sons, 1987. D.D.Macdonald, Corrosion 46 (3), 1990, p. 229. K. Juttner, Electrochimica Acta 35 (10), 1990, p. 1501. D.D. Macdonald, The point defect model for the passive state, J. Electrochem. Soc., Vol. 139, No. 12, pp. 3434-3449, 1992. B. Sala, I. Bobin-Vastra, A. Frichet, Proceedings 6th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems - Water Reactors, Vol. 2, 1997, p 1005.

J. Schefold, D. Lincot, A. Ambard, O. Kerrec, Journal of the Electrochemical Society, Vol. 150, No. 10, 2003, p. B451-B461. S. E. Ziemniak, M. Hanson, Corr. Sci., 44, p. 2209, (2002). Beverskog, B., Bojinov, M., Englund, A., Kinnunen, P., Laitinen, T., Mäkelä, K., Saario, T., Sirkiä, P., A mixed-conduction model for oxide films on Fe, Cr and Fe-Cr alloys in hightemperature aqueous electrolytes – I. Comparison of the electrochemical behaviour at room temperature and at 200°C, Corrosion Science, Vol. 44, pp. 1901-1921, 2002. Beverskog, B., Bojinov, M., Kinnunen, P., Laitinen, T., Mäkelä, K., Saario, T., A mixedconduction model for oxide films on Fe, Cr and Fe-Cr alloys in high-temperature aqueous electrolytes – II. Adaptation and justification of the model, Corrosion Science, Vol. 44, pp. 1923-1940, 2002. Kirchheim, R., Growth kinetics of passive films, Electrochemica Acta, Vol. 32, No. 11, pp. 1619-1629, 1987. M. Vankeerberghen, 1D steady-state finite-element modelling of a bicarrier monolayer oxide film, submitted to Corrosion Science.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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The Effect of CO2 and H2S on the Passivation of Chromium and Stainless Steels in Aqueous Solutions at Elevated Temperature and under High Pressure J. BanaWa , B. Mazurkiewicza, U. Lelek-Borkowskaa , H. Krawieca, W. Solarskia, K. Kowalskib a

AGH – University of Science and Technology, Faculty of Foundry Eng., Dept. of General and Analyt. Chemistry, bAGH-UST, Faculty of Metallurgy and Material Eng., Dept. of Eng. and Material Analysis, Reymonta 23 St., 30-059 Kraków, Poland,

Abstract - The effect of temperature, CO2 and H2S pressure on passivation of chromium and on Fe-Cr and Fe-Cr-Ni-alloys was investigated in H2O-Na2SO4 and H2O-NaCl solutions by means of linear sweep voltammetry, electrochemical impedance spectroscopy (EIS) and X-ray photoelectron spectroscopy (XPS). The investigations have been performed in the temperature range 20 – 1500C under the pressure of 20 bars. The growth of the temperature from 200C to 1500C increases the (Fe2++Fe3+)/Cr3+ and O2-/OH- ratio in the passive film formed on Fe-Cr alloys in aqueous solutions of 0.1M Na2SO4. It seems that the presence of dissolved CO2 also leads to dehydration of the passive film. The stability of passive layer formed on the surface of Fe-Cr and Fe-Cr-Ni alloys decreases in the solutions saturated with CO2 and CO2 + H2S mixture. The effect is more distinct for austenitic Fe-Cr-Ni alloys than for ferritic Fe-Cr ones. The presence of nickel deteriorates the formation of the passive film in solutions saturated with H2S and CO2 + H2S mixture because of the high susceptibility of nickel to sulfide formation. Keywords: chromium, Fe-Cr-Ni alloys, passivity, corrosion in CO2

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1. Introduction Corrosion of metals in geothermal installation is very often determined by the presence of CO2 and H2S dissolved in the thermal water. Metals used in such systems work usually at elevated temperatures and under high pressure. Under such conditions chromium or chromium-nickel steels are applied. The effect of temperature on the passive properties of Fe-Cr alloys in aqueous media can be attributed to the dehydration of the oxide layer [1,2] and to the change of the transport processes through the defected oxide film [3]. Detrimental effect of dissolved CO2 and H2S on the corrosion of carbon steel is relatively well recognized [4,5]. Little attention was paid to the problems of influence of those gases on the stability of passive film on Fe-Cr alloys. The aim of the presented work was to explain the above-mentioned problem. 2. Experimental The investigations were carried out on chromium (99,95% Cr) and Fe-Cr-Ni cast alloys (Table 1). All investigated steels were solution annealed. Investigations were performed in aqueous solutions of Na2SO4 and NaCl, saturated with argon, CO2, Ar+3%H2S, and Ar+48%CO2+3%H2S mixtures. Table 1. Chemical composition and thermal treatment of investigated Fe-Cr-Ni alloys No

C

Si

3

0,040 0,148

7

0,03

Mn 0,11

0,3659 0,10

P

S

0,0004 0,011

Cr

Ni

Mo

22,79 0,015 0,015

Cu

Solution annealing

0,0005 0,128 0,06

1050/2h/air

V

Al

0,0152 0,0639 21,54 25,66 0,1744 0,0497 0,049 0,1386 1150/2h/air

Special high temperature and high pressure cell (autoclave) equipped with the three-electrode system: working electrode, Pt-counter electrode and Ag/AgCl reference electrode was used in the research. Experiments were performed under total pressure of 20bar at temperature of 20, 50, 100, 150, 180OC. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were applied to obtain anodic behavior of investigated metals. Compositions of surface film formed on the metals were analyzed using X-ray photoelectron spectroscopy (XPS) with AlKa excitation of 1486,6eV. In-depth analysis was done by removing the layer by sputtering with Ar+ ions (2kV).

The effect of CO2 and H2 S on the passivation of chromium and stainless steel…

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3. Results and discussion 3.1.

Passivation of chromium

The oxide film formed on chromium in sulfate solutions (Fig.1a) undergoes local etching at high anodic potentials over 0.7VSHE, in the temperature range of 500–1000C. There also appears a current loop on polarization curves. At high temperatures over the 1500C the stability of passive layer increases again and anodic current loop on the CV curve disappears. Contrary to sulfate medium chromium undergoes pitting corrosion in 0.1MNaCl (Fig.1b) in the temperature range over 1000C. Fig.2 compares the effect of temperature on activation and repassivation potentials of Cr in 0.1MNa2SO4 and 0.1MNaCl. In solution saturated with CO2 the stability of passive layer on chromium decreases (Fig.2,3). The shift of activation and repassivation potentials in negative direction was found in sulphate and chloride media (Fig.2). Impedance spectra confirm this negative effect of dissolved CO2 on the polarization resistance of passive chromium (Fig.3).

Fig.1. Effect of temperature on the polarization curves of chromium in: (a) H2O-0.1M Na2SO4-Ar, (b) H2O-0.1M NaCl-Ar system.

The XPS analysis of the passive film, formed on Cr in sulfate solution saturated with Ar or CO2 at the temperature of 500C and 150oC shows that the film consists mainly of CrOOH independent on etching potential and temperature (Fig.4). The growth of the temperature causes only increase of thickness of the passive film. It was confirmed by the intensity of signals from metallic chromium higher at 500C than at 1500C.

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Fig.2. Effect of temperature on stationary ES , activation Ea and repassivation potential Er of Cr in: (a) 0.1MNa2SO4-Ar and 0.1MNaCl-Ar, (b) 0.1MNa2SO4-CO2 and 0.1MNaCl-CO2 system.

Fig. 3. Effect of CO2 on Nyquist diagrams of chromium in 0.1M Na2SO4 (a)

and 0.1M NaCl (b) at ESHE = 0.750 V.

Fig.4. In-depth profile of molar ratio of x

O2-

/ x O 2- + x OH - in passive film obtained on

chromium in 0.1M Na2SO4 at 500C (a) and in 1500C (b).

The effect of CO2 and H2 S on the passivation of chromium and stainless steel…

421

Saturation of the solution with CO2 does not change the composition of the film, but it seems that at low overvoltage, the layer is thicker at the presence of CO2. At high temperature the effect of CO2 on the thickness of passive film is negligible, but dehydration of the passive film is stronger for solution saturated by carbon dioxide than for the medium saturated by argon. 3.2. Passivation of stainless steel Saturation of sulfate solution with CO2 makes the active-passive transition more difficult as well for ferritic (alloy 3) as for autenitic steel (alloy 7)(Fig.5, 6). At the presence of carbon dioxide it was observed that dissolution of metals proceeds faster than for argon atmosphere. The effect of temperature was also distinct, effecting in the increase of anodic current density in all examined cases. This negative effect of CO2 is more pronounced for austenitic steel (Fig.5b, 6b).

Fig.5. Effect of CO2 on CV diagrams of steel 3 (a) and 7 (b) in 0.1M Na2SO4 in 500C.

Fig.6. Effect of CO2 and H2S on CV diagrams of steel 3 (a) and 7 (b) in 0.1M Na2SO4 in 1500C.

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XPS analysis of the passive films obtained on the surface of both investigated steels, polarized in 0.1MNa2SO4, shows that they consist of hydroxylated forms of Cr2O3, Fe2O3 and Fe3O4. Carbon dioxide dissolved in the neutral aqueous solutions stimulates dehydration of passive layer. Fig.7 presents this effect for ferritic steel. Gaseous H2S present in sulphate solutions (fig.7b) behaves in a similar way. Hydrogen sulphide, in neutral solutions saturated with CO2+3%H2S, decreases the stability of passive film formed on the surface of both kinds of steel (Fig.6). This effect is the result of presence of sulphides in the passive layer, which was confirmed by XPS analysis. It was found that the layer contains sulphides of Fe, Cr, Ni (in amount of 2-5%). More distinct effect observed for austenitic Fe-Cr-Ni steel follows from the high affinity of nickel to formation of sulphide and its accumulation in the inner part of the layer. Therefore the austenitic stainless steel is more susceptible to pitting corrosion in NaCl-CO2-H2S system than ferritic one. 0,7

1,0 0

Steel 3 / 0.1M Na2SO4, 20 C, ESHE = 1.20V

2-

0,6

2-

0,4

0,5

2-

2-

0,6

-

xO / (xO + xOH )

-

xO / (xO + xOH )

0,8

0

Steel 3 / 0.1M Na2SO4, 150 C, ESHE = 1,20V

CO2 Ar

0,2

Ar CO2 H2S

0,4

0,3

0,0 0

10

20

30

40

0

50

10

20

30

40

50

sputtering time / min.

sputtering time / min.

Fig.7. In-depth profile of molar ratio x 0

O2-

/x

O 2-

+ x OH - in passive film obtained on the steel 3 0

in 0.1M Na2SO4 at 20 C (a) and at 150 C (b). Influence of CO2 and H2S.

4. Conclusions The passive film formed on Cr in solutions of Na2SO4, at elevated temperatures, consists of hydroxylated Cr2O3 oxide. It seems that in the presence of argon the increase of temperature up to 1500C does not change substantially the composition of the film (the ratio of O2-/OH-). In the case of the carbon presence at elevated temperatures the dehydration of the film takes place but the passive layer still consists of significant amount of hydroxylated chromium oxide. In both cases - in argon and CO2- saturated solutions the temperature rise causes decrease of resistance and capacity of the passive layer. At high temperature (>1500C) the anodic dissolution is limited by the diffusion across the film. The maximum of the susceptibility to local corrosion is observed in the temperature range of 50-1000C. The transformation from the high field conductive film to

The effect of CO2 and H2 S on the passivation of chromium and stainless steel…

423

the diffusion controlled thick layer also occurs in this range of temperature. The growth of the temperature over 1500C increases protective properties of passive film on chromium surface and diminishes susceptibility to the local attack. Stability of the passive film on Fe-Cr-Ni alloys decreases at the presence of dissolved CO2. This effect can be explained by the dehydration of outer part of oxide film by the carbonic acid. The similar effect can be expected for molecules of weak acid - H2S. In this case, the incorporation of S2- ions in oxide film occurs, what additionally makes the mechanism of passivation more complicated, especially in the active-passive transition range. In the neutral solutions saturated with the mixture of CO2+H2S, synergetic effect of dehydration of the outer part and incorporation of S2- ions in the inner part of oxide layer, decreases stability of passive film and therefore the susceptibility to local corrosion increases. This effect is more pronounced for austenitic stainless steels due to formation of nickel sulphides and accumulation of S2- in the passive layer. Acknowledgments: The work was supported by Polish Committee of Scientific Research, grant No7 T08B 001 21 References [1] H. Saito, T. Shibata, G. Okamoto: Corr. Sci. 19 (1979) 693 [2] B. Beverskog, M. Bojinov, A. Englund, P. Kinnunen, T. Laitinen, R. Mäkelä, T. Saario, P. Sirkiä, Corr. Sci. 44 (2002)1901. [3] D.D. Macdonald, A. Sun, N. Pryantha, P. Jayaweera, J. Electroanalyt. Chem. 572 (2004) 421 [4] M.G.S. Fereira, N.E. Hakiki, G. Goodlet, S. Faty, A.M.P. Simões, M. Da. Cunha Belo, Electrochim. Acta 46 (2001) 3767. [5] S.J. Ahn, H.S. Kwon, Electrochim. Acta 49 (2004) 3347.

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) Published by Elsevier B.V.

425

Effect of Lead on Passivation of Alloy 600 Surface Zhongquan Zhou*,a, Jangyul Parkb, J. Ernesto. Indacocheaa, Roger. W. Staehlec, Seong S. Hwangd, Nancy Finnegane, and Rick Haasche a

University of Illinois at Chicago, 842 W. Taylor St., M/C 246, Chicago, IL 60607. Argonne National Laboratory, 9700 S. Cass Ave., Bldg. 212, Argonne, IL 60439 c University of Minnesota, 22 Red Fox Rd., North Oaks, MN 55127 d Korea Atomic Energy Research Institute, P.O. Box 105, Yuseong, Daejeon, Korea e University of Illinois at Urbana-Champaign, 104 S. Goodwin Ave., Urbana, IL61801 * Corresponding author. Email: [email protected] b

Abstract- Effects of Pb on passivity of Alloy 600 surface were investigated in mild acidic aqueous solutions at 90°C using polarization, electrochemical impedance spectroscopy (EIS), Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS). The results indicate that Alloy 600 surface consists of an inner oxide layer and outer hydroxide layer, containing Cr3+ and Ni2+. Pb is incorporated into the surface film and increases the electronic conductivity of Cr oxide in the surface film. Oxidation of Ni is inhibited by the presence of Pb. Keywords: Alloy 600, steam generator, AES, XPS, EIS, polarization, lead, stress corrosion cracking

1. Introduction Lead, dissolved in aqueous solutions, has been found to produce aggressive stress corrosion cracking (SCC) in a broad range of iron-chromium-nickel alloys [1-5]; in some cases, only 0.1 ppm of dissolved Pb is necessary to produce SCC at 300°C [6]. It is hypothesized that PbSCC of various alloys is caused by the incorporation of Pb into the surface film, which reduces its passivity [7-9]. The passive film on Alloy 600 has been characterized as consisting of an inner oxide layer and an outer hydroxide layer, containing Cr3+ and Ni2+ [10-14]. The surface layer with Cr or Ni oxide, which consists of excess oxygen ions or metal ions vacancies, is a cation-selective layer. Pb2+ in the solution may interact with these oxides reducing the passivity of the protective surface film. However, no study has been done on the interaction of Pb with the Ni and Cr oxide film formed in aqueous solution. It is thus still not clear how Pb affects the passivation behavior of Alloy 600. The objective of the present investigation is to characterize the effects of Pb on the passivation of Alloy 600 in order to provide a basis for extending such work to higher temperatures.

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2. Experimental Work Sample preparation.¾The coupons were fabricated from Alloy 600 tubing (Heat M81) with the chemical composition given in Table I. They were wet ground with SiC paper up to 600 grit. The coupons for surface analyses were further polished to 0.05 m. Table I. Chemical composition (wt.%) of Alloy 600 Heat M81 C

Mn

S

Si

Cr

Ni

Fe

Cu

0.037

0.27

0.001

0.3

15.85

74.81

8.16

0.01

Test solution.¾The test solution consisted of an aqueous solution with 110 ppm hydrochloric acid. For Pb-containing solutions, PbO (99.999% purity) was added. The solutions were prepared with A.C.S. grade reagents and high purity water (the resistivity is greater than 18´106Ohm.cm). The pH of the test solution was adjusted to 4.5 at room temperature by adding sodium hydroxide. Experimental set-up.¾The tests were performed in a three electrode cell. The cell was purged with high-purity argon gas during the test and the temperature was controlled at 90±2ºC. A Gamry PC4/750 potentiostat, Solatron 1287 potentiostat, and a Solatron 1252A frequency response analyzer were used for electrochemical testing. Polarization scan.¾Polarization curves were performed at a scan rate of 0.1 mV/sec, from open circuit potential (OCP) to +0.4 V/SHE in the anodic direction and from OCP to -0.6 V/SHE in the cathodic direction. EIS.¾The test specimen was polarized at each test potential for 1 hr before the EIS measurement was performed in the frequency range of 100 kHz to 1 mHz with a perturbation amplitude of 10 mV. AES and XPS.¾Auger depth profiles were performed with a Physical Electronics PHI660 Scanning Auger Microprobe. The specimen surface was sputtered with argon ion beam at a rate of 7 nm/min. XPS spectra were performed with Physical Electronics PHI5400 equipped with a Mg-Ka X-ray source (1253.6 eV). The pass energy was 35.75 eV with a step size of 0.1 eV. Angle-resolved spectra were collected at 30º and 90º takeoff angles. The surfaces were also sputtered with argon ion beam at a rate of 3 nm/min. 3. Results and Discussion Polarization curves.¾ Polarization curves of Alloy 600MA, which were measured in the solutions with 0, 3, 30, and 300 ppm Pb2+, are shown in Figure 1. The OCP is shifted to the anodic direction in the solutions containing Pb2+. In the anodic branch, at potentials between OCP and -0.1 V/SHE, the anodic current is inhibited in all Pb2+containing solutions. The superposition of the cathodic current from underpotential reduction of Pb2+ might be attributed to the reduced anodic current. Another possibility could be the underpotential deposited metallic Pb, or the adsorbed Pb2+-containing species on the electrode inhibit the oxidation of Ni. At potentials between -0.1 V/SHE and +0.1 V/SHE, the inhibiting effect of Pb is not observed in 300 ppm Pb2+-containing solutions. We inferred that there are at least two effects of Pb on the corrosion of Alloy

427

Effect of Lead on Passivation of Alloy 600 Surface

600, depending on the electrode potential and Pb2+ concentration. In the cathodic branch from OCP to -0.5 V/SHE, the cathodic current density increases with increasing Pb2+ concentration due to the reduction of Pb2+ to metallic Pb. 0.3

1: E Pb2+(3ppm)/Pb0 2: E H+(10-4.5)/H 2(1 atm) 3: E Pb2+(30ppm)/Pb0 4: E Pb2+(300ppm)/Pb0

0.2 Potential (V/SHE)

0.1 0

3ppm Pb

2+

0 Pb

2+

30ppm Pb

2+

300ppm Pb

2+

-0.1 4

-0.2 -0.3 -0.4

3

-0.5 10

-10

1,2

10

-9

10

-8

-7

-6

-5

10 10 10 2 Current density (A/cm )

10

-4

10

-3

10

-2

Figure 1. Polarization curves (0.1 mV/sec scan rate) of Alloy 600MA at 90ºC in solutions containing 110 ppm Cl- plus 0, 3, 30, and 300ppm Pb2+, with pH 4.5 (25° C). -7

-0.19 V

-3 30ppm Pb

2+

3ppm Pb

2+

-2

2+

0 Pb

-80

3ppm Pb2+

2

-Imag(KOhm.cm )

-4

-100

-50

-1 0

2+

2

-5

0.05 V

30ppm Pb2+

300ppm Pb -Imag(KOhm.cm )

2

-Imag(KOhm.cm )

-100

-150 -0.35V

-6

-60 0 Pb2+ -40

300ppm Pb2+

-20

3ppm Pb2+

30ppm Pb

2+

0 Pb2+ 0

1

2 3 4 5 2 Real(KOhm.cm )

6

7

0 0

50 100 2 Real(KOhm.cm )

150

0

0

20

40

60

80

100

2

Real(KOhm.cm )

a: -0.35 V/SHE b: -0.19 V/SHE c: 0.05 V/SHE Figure 2. Nyquist plots of Alloy 600MA at 90ºC in solutions containing 110 ppm Clplus 0, 3, 30, and 300 ppm Pb2+ with pH 4.5 (25 º C).

EIS.¾Figure 2 shows the Nyquist plots for Alloy 600 measured in 0, 3, 30, and 300 ppm Pb2+-containing solutions at -0.35, -0.19, and 0.05 V/SHE. At -0.35 V/SHE, the observed capacitance loops, which correspond to the time constants at intermediate frequency (~1Hz), reflect the reduction of Pb2+ and H+ resulting from electron transfer through the Cr oxide layer [15]. As shown in Figure 2a, the capacitance loop size decreases with increasing Pb2+ concentration, indicating that Pb2+ enhances the electronic conductivity of Cr oxide in the surface film. At -0.19 V/SHE in Figure 2b, the observed capacitance loops, which correspond to the time constants at low frequency (~0.01Hz), reflect the ionic transfer process for Ni oxidation [15]. The larger size of the low-frequency capacitance loop in Pb2+-containning solutions indicates that either adsorbed Pb2+ or the underpotential deposition of metallic Pb increase the ionic transfer resistance for Ni oxidation. At 0.05 V/SHE in Figure 2c, in the solution without Pb2+, the capacitance loop bends toward the second quadrant, indicating that an activepassive transition is taking place [15]. A possible factor which produces this activepassive transition is the growth of a Ni compound layer due to the enhanced oxidation

428

Z. Zhou et al.

of Ni and the formation of NiO, which was observed by AES and XPS [15]. However, the active-passive transition is not observed in all Pb2+-containing solutions. This suggests that Pb2+ in the solution retards the growth of the Ni compound layer on Alloy 600. This is confirmed by surface analysis, which shows that Pb2+ inhibits the oxidation of Ni at +0.05 V/SHE, and the formation of NiO is retarded in the surface film exposed to a 300ppm Pb2+ containing solution. Moreover, as seen in Figure 2c, the size of the capacitance loop declines with increasing Pb2+ concentration from 3 ppm to 300 ppm, indicating that the resistance of the surface film is also reduced due to Pb2+. AES and XPS. ¾ Figure 3 shows the normalized Ni and Cr depth profiles of Alloy 600 surface film exposed to air, as well as the film electrochemically produced at 0.05 V/SHE. The surface layer is depleted of Ni and relatively enriched in Cr within the first 4 nm for both air-exposed and electrochemically produced films. The enrichment of Cr can be interpreted as a preferential oxidation of Cr due to the larger oxygen affinity for Cr compared with Ni and Fe. As a result, Cr segregates towards the alloy surface. For the film electrochemically produced at 0.05 V/SHE, a larger depletion of Ni and a larger enrichment of Cr occurs in the first 4 nm. This finding can be attributed to the selective dissolution of Ni combined with preferential oxidation of Cr. However, the surface film electrochemically produced at 0.05 V/SHE in the solution containing Pb2+ shows the similar Ni and Cr profiles as the air-exposed film, suggesting that the selective dissolution of Ni might be inhibited in this case. 0.8

0.3 2+

300ppm Pb

=0.74

bulk

air formed film

0.75

0.25

0.7

0 Pb

Cr/(Ni+Cr+Fe)

Ni/(Ni+Cr+Fe)

Ni

2+

3.5nm 0.65

Ni depletion

Cr

2+

1.7nm 1.9nm 3nm Cr enrichment

0.15

1.6nm 0.6

0.1 0

2

4

6

Depth (nm)

8

10

=0.18

bulk

0 Pb

300ppm Pb

0.2

0

2

2+

air formed film

4

6

8

10

Depth (nm)

Figure 3. Normalized Auger depth profile of Ni and Cr in the surface of Alloy 600MA polarized at 0.05 V/SHE for 12 hr at 90ºC in solutions containing 110 ppm Cl- plus 0 and 300 ppm Pb2+, with pH 4.5 (25° C). The Ni 2p3/2 spectra is the superposition of a sharp peak centered at 852.7 eV±0.2 eV and a broader peak at higher binding energy. These peaks are assigned to metallic Ni (Ni0) and Ni2+ (hydroxide and oxide), respectively [16]. Figure 4 shows the concentration of Ni2+ and Ni oxide in the film, which are displayed by the ratio of corresponding peak intensity to the total Ni intensity. It can be seen that the oxidation of Ni is inhibited in Pb2+-containing solution, and the formation of NiO is retarded, resulting in the lower NiO concentration at a 90º take-off angle. The Cr 2p3/2 spectra is a broad peak centered at 577.1±0.1 eV before ion etching, which is assigned to Cr3+ compounds [10,11,13,17]. An additional peak centered at 574.4 eV±0.1 eV, which is attributed to metallic Cr [16], is superimposed on the Cr3+ peak after ion etching. Figure 5 shows the concentration of Cr3+ in the film displayed by the ratio of its peak intensity to the total Cr 2p3/2 intensity at each sputtering step in the

429

Effect of Lead on Passivation of Alloy 600 Surface

films formed in the solutions with and without Pb2+. The oxidation of Cr extends deeper into the surface layer, and the thicker Cr compound layer is formed in Pb2+-containing solution. 0.7

Intensity ratio

2+

o

30

90

0.6

Ni =Ni(OH) +NiO

o

Ni

0.5

=Ni +Ni(OH) +NiO 2

30

o

90

2

0

Tot

o o

90

0.4 0.3

90

0.2 30

o

30

o

o

0.1 0

2+

Ni /Ni

Tot

NiO/Ni

2+

Ni /Ni

Tot

2+

NiO/Ni

Tot

Tot

2+

300 ppm Pb

0 ppm Pb

Figure 4. The intensity ratio of Ni2+ and NiO in the surface of Alloy 600MA polarized at 0.05V/SHE for 12 hr at 90ºC in solutions containing 110ppm Cl- plus 0 and 300ppm Pb2+, with pH 4.5 (25 ºC). 2+

300ppm Pb

3+

Intensity ratio (Cr /Cr

Tot

)

1 0.8 0.6 2+

0 Pb

0.4

3+

Cr =Cr(OH) +Cr O 3

0.2

0

30 take off angle

Cr

Tot

2

3

0

=Cr +Cr(OH) +Cr O 3

2

3

0 0

2

4

6

8

10

12

Sputter depth (nm)

Figure 5. The intensity ratio of Cr3+ at each sputtering step in the surface of Alloy 600MA polarized at 0.05V/SHE for 12 hr at 90ºC in solutions containing 110 ppm Clplus 0 and 300 ppm Pb2+, with pH 4.5 (25 º C). 0.5

Pb

2000

0

Total

=Pb +Pb

2+

0.4 0

Pb /Pb

0.3

Tot

0

1500

500 0

Tot

0.1

0

30 take off angle 0

2

4

Total

0.2

1000

Pb

Pb /Pb

Pb 4f Intensity (Counts.eV/sec)

2500

6

8

10

12

0

Sputter Depth (nm)

Figure 6. The intensity ratio of Pb0 at each sputtering step in the surface of Alloy 600MA polarized at 0.05V/SHE for 12 hr at 90ºC in solutions containing 110 ppm Cland 300 ppm Pb2+, with pH 4.5 (25 ºC). The Pb 4f spectra is a sharp peak centered at 138.7±0.2 eV attributed to PbO before sputtering. The metallic Pb peak is resolved after sputtering located at 136.9 eV. The Pb 4f peak intensity and the ratio of metallic Pb to total Pb intensity at each sputtering step are shown in Figure 6. It appears that Pb incorporats into the film to ~10 nm, and more than ~40% of Pb is Pb2+. Since the growth of Ni oxide is retarded in the passive film of

430

Z. Zhou et al.

Alloy 600 formed in Pb2+-containing solutions, Cr2O3 which is a p-conductor is predominant in the film and the electronic current is assured by electron holes [18-20]. The substitution of Pb2+ into a Cr2O3 lattice contributes one negative effective charge, which is balanced by creation of one electron hole. Therefore, the concentration of electron holes is increased and the electron transport resistance is reduced by incorporating Pb2+ into the film. 4. Conclusions Pb2+ inhibits the oxidation of Ni, and consequently, the growth of a Ni oxide layer is retarded. In addition, the incorporation of Pb2+ into Cr oxide enhances its electronic conductivity by increasing the concentration of the electron holes. Thus, the passive film on Alloy 600 is less protective in an environment containing Pb2+, and corrosion of Alloy 600 is intensified as a result. Acknowledgements This work was performed at Argonne National Laboratory, with support by the Division of Engineering Technology, Office of Nuclear Regulatory Research, U.S. Nuclear Regulator Commission. Argonne is operated by the University of Chicago for the Department of Energy under contract W-31-109-ENG-38. The test materials were provided by Electric Power Research Institute. The AES and XPS were performed at the Center for Microanalysis of Materials, University of Illinois, which is partially supported by the U.S. Department of Energy under grant DEFG02-91-ER45439. The authors thank J. M. Sarver of Babcock & Wilcox, for his helpful comments on this work. Reference 1. M. Helie, I. Lambert, and G. Santarini, Proc. 7th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, p. 247, Breckenridge, Colorado, 1995. 2. S. S. Hwang, H. P. Kim, D. H. Lee, U. C. Kim, and J. S. Kim, J. of Nuclear Materials, 275 (1999) 28. 3. T. Sakai, N. Nakagomi, T. Kikuchi et al., Corrosion, 54 (1998) 515. 4. R.W. Staehle, Proc. 11th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, p. 381, Stevenson, WA, 2003. 5. J. M. Sarver and B. P. Miglin, EPRI NP-7367-S, 1991. 6. H.Takamatsu, T. Matsunaga, B. P. Miglin et al., Proc. 8th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, p. 216, Amelia Island, 1997. 7. T. Sakai, K. Aoki, and T. Shigemitsu, Corrosion, 48 (1992) 745. 8. S. S. Hwang, U. C. Kim, and Y. S. Park, J. of Nuclear Materials, 246 (1997) 77. 9. A. McIlree, ISG-TIP-3 Meeting, Argonne National Laboratory, April 5-8, 2004. 10. S. Boudin, J. L.Vignes, G. Lorang et al., Surf. Interface Anal., 22 (1994) 462. 11. T. Jabs, P. Borthen, and H. H. Strehblow, J. Electrochem. Soc., 144 (1997)1231. 12. A. C. Lloyd, D. W. Shoesmith, N. S. McIntyre, and J. J. Noël, J. Electrochem. Soc., 150 (2003) B120. 13. A. Machet, A. Galtayries, S. Zanna et al., Electrochimica Acta, 49 (2004) 3957. 14. A. Machet, A.Galtayries, P. Marcus et al., Surf. Interface Anal., 34 (2002) 197. 15. Z. Zhou, Ph.D. thesis, Lead Effect on the Corrosion and Passivation Behavior of Alloy 600, University of Illinois at Chicago, 2005. 16. J. F. Moulder, Handbook of X-ray Photoelectron Spectroscopy: A Reference Book of Standard Spectra for Identification and Interpretation of XPS Data, Eden Prairie, Perkin-Elmer Corp., 1995. 17. V. Maurice, W.P. Yang, and P. Marcus, J. Electrochem. Soc., 145 (1998) 909. 18. P. Kofstad, Nonstoichiometry, Diffusion and Electrical Conductivity in Binary Metal Oxides, Robert E. Krieger Publishing Company, Malabar, FL, 1983. 19. A. Holt and P. Kofstad, Solid State Ionics, 69 (1994) 137. 20. M. F. Montemor, M. G. S. Ferreira, M. Walls, B. Rondot, and M. C. Belo, Corrosion, 59 (2003)11. The submitted manuscript has been created by the University of Chicago as Operator of Argonne National Laboratory (“Argonne”) under Contract No. W-31-109-ENG-38 with the U.S. Department of Energy. The U.S. Government retains for itself, and others acting on its behalf, a paid-up, nonexclusive, irrevocable worldwide license in said article to reproduce, prepare derivative works, distribute copies to the public, and perform publicly and display publicly, by or on behalf of the Government

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

431

Interaction of Oxide Layers on Structural Materials with Light Water Reactor Coolants - its influence on the mechanism of oxide growth and restructuring Martin Bojinov a and Björn Beverskog b a

Department of Physical Chemistry, University of Chemical Techology and Metallurgy, P.O. Box 92, 1164 Sofia, Bulgaria, e-mail [email protected] b OECD Halden Reactor Project, Institutt for Energiteknikk, Box 173, N-1772 Halden, Norway, e-mail [email protected]

Abstract - A four-layer model of the oxide formed on stainless steels in nuclear power plant coolants features a barrier layer similar to room temperature passive films, which grows via solid state defect transport; an inner layer, that is more defective, its mechanism of growth including both solid state transport and dissolution/redeposition reactions; an outer layer formed by dissolution/precipitation, and a deposit layer which growth is dictated by precipitation of material onto the outer layer. Support for such a view is sought in compositional profiles of oxides on stainless steel in nuclear power plant coolant conditions with or without the intentional addition of foreign species. It is suggested that data on the interaction of oxides with such species could provide an additional way to estimate the kinetic and transport parameters of oxide growth on construction materials in light water reactor primary circuits. Keywords: stainless steel; high-temperature water; oxide layer model; barrier layer; point defect; solid-state transport; cation incorporation; dissolution-precipitation

1. Introduction The important role of the oxide layers formed during exposure of stainless steel to Light Water Reactor (LWR) coolants in controlling the general corrosion behavior, acting as a barrier against localized corrosion and a reservoir for radioactivity build-up in the primary circuit, has been recognized for quite some time. The investigations of the mechanism(s) of influence of the oxide layers on such processes have been directed in two ways. The first has been related to the study of the role of surface oxides in material corrosion and based on ex-situ

432

M. Bojinov and B. Beverskog

characterization of the oxides formed during nuclear power plant operation. Electrochemical, photoelectrochemical and mechano-electrochemical methods have started to contribute only recently, their importance being widely recognized [1-4]. The second type of studies has been promoted by the need to understand the mechanism of activity incorporation and has thus been focused on the interaction of oxides with coolant-originating species [5-8]. The aim of the present paper is to propose a cross-link between these approaches based on recent data on the composition of in-reactor oxides and studies on the effect of Zn on the incorporation of radioactive Co in such layers [5-8]. The data are rationalized on the basis of the Mixed-Conduction Model (MCM) for oxide films on stainless steels [2] and a recently presented view of the interaction of cationic species with LWR oxides [8]. 2. Theoretical background (Fe1-xNix)(Cr1-yFey)2O4

Vm

VM´´´

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Figure 1. A simplified picture of the hydrothermally formed oxide layer on stainless steel

A scheme of the oxide on stainless steel in LWRs is presented in Fig. 1. It is assumed that such oxides are spinels, which is consistent with both thermodynamic calculations [9] and experimental observations [10-1]. The barrier layer (BL) grows by ingress of oxygen transported via oxygen vacancies and is considered to be a normal spinel of the chromite (FeCr2O4) type with cation vacancies as main ionic current carriers, ensuring a p-type behavior [8]. The inner layer (IL) is supposed to grow both at its interface with the barrier layer via injection of cation vacancies and by dissolution-redeposition mechanism limited by the transport of interstitial cations. As a result, it contains more Fe and Ni than the BL [8]. The outer layer (OL) is considered to be an inverse spinel of the trevorite (NiFe2O4) type [1] with interstitial cations as main ionic carriers, ensuring an n-type behavior. A deposited layer (DL) formed by precipitation of material from the coolant is also included (Fig. 1). In the MCM, the transport of defects is governed by both concentration and potential gradients [2]. As for hydrothermal oxides, the electric field strength is low in comparison to room-temperature passive films [2], the main driving force for

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transport is the concentration gradient, which explains the occurrence of inward transport of cations from the electrolyte [8]. The transport of matter through the OL and DL is taking place by diffusion-migration through pores and cracks [1]. However, as incorporation of cations in these layers is expected to include solid-state transport in the crystallite interior, the approach outlined above remains of importance for both of them. The possible steps of the interaction reaction of a cation from the coolant, Me2+, with the oxide are [8]: adsorption and surface complexation at the outer layer/coolant interface, incorporation in the outer layer via surface precipitation, incorporation in the inner layer via filling of interstitial lattice positions, Vi : Vi + Me2+ ® Me i··

(1)

1. incorporation in the barrier layer via filling of cation vacancies, VM´´ : VM´´ + Me2+ ® Me M

(2)

As according to the MCM, interstitial cations are generated at the BL/alloy interface and consumed at the IL/BL boundary, whereas cation vacancies are generated at the coolant/IL interface and consumed at the BL/alloy interface [2], it can be concluded that the incorporation of foreign cations will slow down the solid state transport. 3. Experimental AISI 316L(NG) samples (0.015%C, 16.5%Cr, 0.26%Cu, 1.73%Mn, 10.5%Ni,0.54%Si, 2.55%Mo, 0.056%N, balance Fe) have been in-plant exposed for 8000 h to either a VVER (1200 ppm B, 15 ppm K, 15 ppm NH3, dissolved H2 40 cm3 kg-1 STP, ECORR = -0.8..-0.75 V, 250 °C) or a BWR coolant (RT conductivity <0.1 µS cm-1, dissolved O2 300 ppb, ECORR= 0.05..0.1 VSHE, 273 °C). The in-depth composition of the oxide has been determined by Glow Discharge Optical Emission Spectrometry (GDOES) using a GDA 750 instrument with a polychromator of 750 mm focal length and 2400 grooves per mm grating, at an excitation voltage 750 V, discharge current 20 mA and discharge pressure 3.2 hPa. 4. Supporting Results and their Discussion The depth profiles of the oxides on AISI 316 after exposure to VVER and BWR coolants are shown in Figs 2-3. The IL in the VVER coolant is strongly enriched in Cr (Fig. 2). Part of the it is probably restructured via interaction with the electrolyte, as evidenced by the incorporation of B. The thickness of the BL (located between the maximum in B and the interface) seems to be comparable

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Figure 2 : GDOES depth profiles of the oxide layer on AISI 316L(NG) after exposure to a VVER coolant. Estimated oxide thickness is indicated with a vertical double arrow. (a) Fe, Cr and Ni content normalized to the total content of metallic constituents; (b) B content, wt.%, Mo and Co contents normalized to the total content of metallic constituents,%.

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Figure 3 : GDOES depth profiles of the oxide on AISI 316L(NG) after exposure to a BWR coolant. Estimated oxide thickness is indicated with a vertical double arrow. (a) Fe, Cr and Ni contents normalized to the total content of metallic constituents; (b) Mo and Co contents normalized to the total content of metallic constituents, %.

The oxide in a BWR coolant is thicker due to the environment’s higher redox potential. The enrichment of Co in the IL is due to its lower transport rate via vacancies compared to that in the OL via interstitials. That is also the reason why 58Co and 60Co from the coolant are enriched at the OL/IL interface [8]. This boundary can be related to the peak in the Ni content (Fig. 3), which also results from a different transport rate due to change from p- to n-type behavior. The plateau in the Cr profile corresponds to the non-porous part of the OL and is due to a higher transport rate compared to the BL and IL, whereas the sloping part - to Cr depletion in contact with water in the pores. The shallow maximum of Co is due to absorption in the non-porous part of the OL [8]. The final position of Co can be correlated with the Cr rich oxide. In order to further quantify the electrolyte-oxide interaction, samples preoxidized in a BWR

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coolant were exposed to a VVER coolant. The depth profiles after pre-oxidation were similar to those obtained in the real BWR coolant. The significant B incorporation (Fig. 4) indicates that the layer has been restructured. The IL/BL boundary could be identified with the maximum in Cr content, the thickness of the BL being of the order of few tens of nanometers. Although Co is depleted in the oxide, the two-maxima structure of the Co profile evident in the layers formed in BWR conditions (Fig. 3) seems to be unaltered (Fig. 4). The above considerations can be extended to the interaction of oxides with cations added into the coolant, as follows from Fig. 5 which depicts the profiles of Co and Zn in oxides on AISI 304L after a 1000 h exposure to simulated BWR coolant containing Co or Co+Zn (data taken from [7], modified). In the absence of Zn, a Co profile closely similar to that depicted in Fig. 3 is observed.

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Figure 5 : GDOES depth profiles of Co (wt.%) and normalized Zn (%) in the oxide layers formed on AISI 304L after 1000 h of exposure to simulated BWR water containing 1 ppb of Co and different amounts of Zn: 0, 10 and 50 ppb [7, modified]. The position of the oxide/alloy interface as estimated from the normalized depth profile of Fe is shown with vertical lines.

The plateau of the Co curve represents incorporation in the non-porous OL via interstitial lattice positions. The peak is due to the Co accumulation at the OL/IL

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interface, because of the slower diffusion rate in the IL via cation vacancies. Addition of 10 ppb of Zn causes the total oxide thickness to decrease twice, i.e. oxide reactivity and growth rate are reduced. Besides, only a single Co peak representing the accumulation at the OL/IL interface is detected at ca. 1/3 of the depth of the layer. One way to interpret such a behavior is to propose the formation of zincochromite (ZnCr2O4), a p-type semiconductor with a large stability area [11], or a FeCr2O4/ZnCr2O4 solid solution. The rate of inward transport of Co in such a layer will be substantially reduced. For 50 ppb of added Zn, the Co curve has an even smaller and narrower peak, which is due to a further decrease in oxide reactivity and growth rate. At the higher Zn concentration, Zn is enriched at the outer interface due to the possible formation of ZnO [8]. The shoulder of the Zn curve could represent the solid solution of chromite and zincochromite, in agreement with the change in the Co behavior. 5. Conclusions On the basis of the above results, it can be concluded that the growth and restructuring of the barrier and inner oxide sublayers on stainless steel in hightemperature water is determined mostly by kinetics and solid-state transport. The interaction of coolant-originating species with these two layers can be treated also as a solid-state process. The outer layer growth is governed by the stability of the oxide phases formed, and the incorporation of foreign species in them involves surface complexation and reprecipitation reactions. The authors believe that the proposed concept demonstrates the usefulness of the data on the interaction of coolant-originating species for the refinement of the models for the growth and restructuring of oxide layers in LWRs. References 1. B. Stellwag, Corros. Sci., 40 (1998) 337. 2. B. Beverskog, M. Bojinov, P. Kinnunen, T. Laitinen, K. Mäkelä, T. Saario, Corros. Sci., 44 (2002) 1923. 3. M. Bojinov, P. Kinnunen, T. Laitinen, K. Mäkelä, T. Saario, P. Sirkiä, Electrochem. Commun., 4 (2002) 222. 4. Y. Takeda, M. Bojinov, H. Hänninen, P. Kinnunen, T. Laitinen, K. Mäkelä, T. Saario, T.Sakaguchi, T. Shoji, P. Sirkiä, A. Toivonen, Key Eng. Mater., 261-263 (2004) 925. 5. D.H. Lister, G. Venkateswaran, Nucl. Technol., 125 (1999) 316. 6. G. Repphun, A. Hiltpold, I. Mailand, B. Stellwag, Power Plant Chem., 1 (1999) 18. 7. H. Inagaki, A. Nishikawa, Y. Sugita, T. Tsuji, J. Nuclear Sci. Technol. 40 (2003) 143. 8. B. Beverskog, in Proc. Int. Conf. "Water Chemistry of Nuclear Reactor Systems", San Francisco, CA, 2004, paper 2.7. 9. B. Beverskog and I. Puigdomenech, Corrosion, 55 (1999) 1077. 10. Y.-J. Kim, Corrosion, 55 (1999) 81. 11. S.E. Ziemniak, M. Hanson, Corros. Sci., 44 (2002) 2209.

Section F Mechanical Properties of Passive Films

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Differences in Mechanical Properties of the Passive Metal Surfaces Obtained in Solution and Air Masahiro Seo*, Daisuke Kawamata, Makoto Chiba Graduate School of Engineering, Hokkaido University Kita-13 Jo, Nishi-8 Chome, Kita-ku, Sapporo 060-8628, Japan E-mail : [email protected]

Abstract - Nano-indentation or nano-scratching in solution and in air were performed for the single-crystal iron (100) and single-crystal tantalum (100) surfaces passivated at a constant potential in pH 8.4 borate solution to investigate the differences in mechanical properties of the passive metal surfaces obtained in solution and air. The dichromate treatment increased the hardness of the passive iron surface kept at 0.25 V (SHE) for 1 h. The effect of dichromate treatment on hardness was more significant in solution than in air. The friction coefficients obtained with nano-scratching in solution for the passive iron surface kept at various potentials for 1 h were always larger than those obtained with nano-scratching in air after passivation and increased with increasing formation potential of passive film. In contrast, the friction coefficients obtained with nano-scratching in air were not sensitive to the formation potential. The stick-slip like phenomenon was observed during nano-scratching in solution for the tantalum surface anodically oxidized at 5.0 V (SHE) for 1 h. The average friction coefficients obtained with nano-scratching in solution were significantly larger than those obtained with nano-scratching in air after anodic oxidation and increased with increasing logarithm of time, log t, required for scratching. On the other hand, the friction coefficients obtained with nano-scratching in air were almost constant independent of log t. The differences in hardness or friction coefficient of the passive metal surfaces obtained with nano-indentation or nano-scratching in solution and air were discussed from the difference in repassivation rate at ruptured sites of passive film between solution and air. Keywords: Nano-indentation; Nano-scratching; Passive metal surface; Hardness; Friction coefficient

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1. Introduction Nano-indentation and nano-scratching techniques have been applied to passive metal surfaces to evaluate their mechanical properties such as hardness, elastic modulus and friction coefficient. Most of the studies, however, have been made in air after passive film formation. In the case where the indenter tip is penetrated into the specimen substrate deeper than the thickness of passive film, the mechanical breakdown and repair of passive film would take place during nano-indentation and nano-scratching in solution. The mechanical breakdown and repair of passive film may influence the mechanical properties of passive metal surface. We performed so far nano-indentation or nano-scratching in solution for the iron, titanium and tantalum surfaces kept at passive states to measure the hardness and friction coefficients of the passive metal surfaces and to compare those measured in air after passivation. In this paper, we reviewed our results [1-4] and discussed the differences in mechanical properties of passive metal surfaces obtained in solution and in air from the viewpoint of mechano-electrochemistry. 2. Experimental Single-crystal iron (100) and single-crystal tantalum (100) disc plates with a diameter of 10 mm were used in experiments. The specimens were mechanically polished with alumina abrasives or diamond paste and then finally electropolished to remove a worked layer. The electrolyte solution employed for nano-indentation and nano-scratching was pH 8.4 borate solution in which the specimen was kept at constant potential in the passive region. A small electrochemical cell made from Diflon was specially designed for nanoindentation and nano-scratching in solution. Platinum ring and wire were used as counter and reference electrodes, respectively. The potential of the platinum wire was always measured with an Ag/AgCl reference electrode for calibration before and after every experiment. The potential of the metal specimen was converted to the standard hydrogen electrode scale (SHE). The nano-indentation/ nano-scratching apparatus (Hysitron Co., Ltd., Triboscope) was combined with AFM (Digital Instruments, Nanoscope IIIa) in which the electrochemical cell was set. A pyramidal diamond tip (cube corner) with an included angle of 90o and a tip radius less than 0.1 µm was used for nano-indentation, while a conical diamond tip with an included angle of 90o and a tip radius less than 1 µm was used for nano-scratching. Both tips were attached to a tungsten rod for use in liquid.

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3. Results and discussion 3.1. Difference between hardness measured in solution and in air The nano-indentation in solution was performed under a maximum load of Lmax = 100 – 500 µN and for an indentation time of 10 s for the iron (100) surfaces kept at 0.25 V (SHE) for 1 h in pH 8.4 borate solution after and without dichromate treatment to measure the load-depth curves from which the hardness of the passive surface was evaluated. For dichromate treatment, the electropolished iron specimen was immersed in 5 x 10-2 M K2Cr2O7 solution for 24 h. It is seen from Fig.1 that the load-depth curves shift to the left side due to dichromate treatment [1]. The hardness, H, is defined by dividing maximum load, Lmax, with projected contact area, A, of the indenter at Lmax [5]. H = Lmax / A

(1)

The hardness, H, was determined from the measured load-depth curves in Fig.1 by using a fused quartz as a standard material and calibrating the projected contact area, A, as a function of contact depth, hc. 800 nano-indentation in solution

Load, L / µN

600

after dichromate treatment without dichromate treatment

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Fig. 1 Load-depth curves measured with nano-indentation in solution for the passive iron (100) surface kept at 0.25 V (SHE) for 1 h in pH 8.4 borate solution after and without dichromate treatment.

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Fig. 2 shows the hardness, H, as a function of contact depth, hc, for the passive iron (100) surfaces after and without dichromate treatment [2], indicating that the dichromate treatment increases the hardness of the passive iron (100) surface at hc > 30 nm in spite of large fluctuation in hardness. It should be noted that the contact depth, hc = 30 nm, is ten times as large as the thickness of passive film formed on iron [6]. The part of passive film in contact with indenter will be ruptured as the indentation depth increases during loading. The ruptured sites, however, will be repaired, depending on repassivation rate. The rupture and repair, therefore, will be repeated during loading. It appears that the film rupture promotes the indentation, while the film repair impedes the indentation. 5.0 nano-indentation in solution

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It is known that a chromium-enriched passive film is formed with chromate or dichromate treatment [7,8]. The high corrosion resistance of chromium-enriched film results not only from the high barrier property of film itself against ion transport but also the high repassivation rate at ruptured sites. The increase in hardness due to dichromate treatment may be brought by promotion of repassivation at the ruptured sites due to chromium enrichment in the passive film. It is expected that the load-depth curves obtained with indentation in air are different from those obtained with indentation in solution since the

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repassivation in air may be slow as compared to that in solution under electrochemically controlled condition. The nano-indentation in air was also performed on the iron surfaces passivated at 0.25 V (SHE) for 1 h in pH 8.4 borate solution after and without dichromate treatment. Fig. 3 shows the comparison between hardness obtained with nano-indentation in solution and in air for the passive iron surfaces after dichromate treatment [2]. The hardness obtained with nano-indentation in solution is always larger than that obtained with nano-indentation in air in the measured contact depth range, suggesting that the repassivation rate in solution is higher than that in air. In contrast, for the passive iron surfaces without dichromate treatment, the hardness obtained with nano-indentation in solution was almost the same as that obtained with nano-indentation in air [2]. These results suggest that the promotion of repassivation due to dichromate treatment is more significant in solution than in air. 5.0 nano-indentation in solution nano-indentation in air

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3.2. Difference between friction coefficients measured in solution and in air

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3.2.1. Friction coefficient of passive iron surface Nano-scratching in solution was performed under a constant normal force, FN = 100 – 1000 µN at a scratching rate of 0.2 µm between a distance of 2 µm for the iron surface kept at a constant potential in the passive region for 1 h in pH 8.4 borate solution to measure the friction coefficient. Fig. 4 shows a) normal displacement, DN, and b) lateral force, FL, as a function of lateral displacement, DL, obtained with nano-scratching in solution at a constant normal force, FN = 500 µN for the passive iron surface kept at 0.5 V (SHE) [3,4]. The direction of normal displacement, DN, in Fig. 4a) was conveniently taken positive from the surface level before scratching. The lateral force, FL, i.e., friction force increases rapidly at the initial stage (DL < 0.5 µm) of scratching and then becomes steady state at DL> 0.5 µm. -20 0

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The friction coefficient, µ’, is usually defined by dividing FL with FN.

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µ’ = FL / FN

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The steady state values of FL were employed for determination of µ’ of the passive iron surfaces. The nano-scratching in air after passivation was also performed to investigate the difference in friction coefficient of passive iron surface in solution and air. Fig. 5 shows the friction coefficients, µ’, obtained with nano-scratching in solution and in air at a constant normal force, FN = 100 µN as a function of formation potential of passive film for the passive iron surface. The friction coefficient, µ’, obtained with nano-scratching in solution increases with increasing formation potential, while that obtained with nanoscratching in air does not significantly depend on formation potential.

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Potential, E / V vs. SHE Fig. 5 Friction coefficients, µ’, obtained with nano-scratching in solution and in air at a constant normal force, FN = 100 µN as a function of formation potential of passive film for the passive iron surface.

The normal displacement, DN, during nano-scratching at a constant normal force of FN = 100 µN is about 30 nm, which is ten times as large as the thickness of passive film on iron. The passive film, therefore, is ruptured at the moving front of the indenter tip during nano-scratching. In the case of nanoscratching in air, the ruptured sites of the passive film would be repaired only by air-oxidation and the repassivation at the ruptured sites may be slow. In contrast, in the case of nano-scratching in solution, the repassivation may be promoted by the potential difference between the substrate iron at the ruptured

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L

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sites and solution, since the repassivation rate increases exponentially with increasing potential. The increase in repassivation rate at the ruptured sites would provide the high mechanical resistance against lateral movement of the indenter tip to contribute to the increase in friction coefficient. Consequently, the significant potential dependence of friction coefficient obtained with nanoscratching in solution may be explained in terms of the potential dependence of repassivation rate at the ruptured sites of the passive film.

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Fig. 6 Normal displacement, DN, and lateral force, FL, as a function of lateral displacement, DL, obtained with nano-scratching in solution at a scratching rate of 0.4 µm s-1 for the tantalum (100) surface anodically oxidized at 5.0 V (SHE) for 1 h in pH 8.4 borate solution.

3.2.2. Friction coefficient of tantalum surface anodically oxidized Nano-scratching in solution and in air was performed under a constant normal force, FN = 300 µN at different scratching rates between a distance of 2 µm for the tantalum (100) surface anodically oxidized at 5.0 V (SHE) for 1 h in pH 8.4 borate solution to investigate the difference in friction coefficient between solution and air. The thickness of anodic oxide film formed on polycrystalline tantalum at 5.0 V (SHE) was about 13 nm [9]. Fig. 6 shows normal

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displacement, DN, and lateral force, FL, as a function of lateral displacement, DL, obtained with nano-scratching in solution at a scratching rate of 0.4 µm s-1. In nano-scratching in solution, FL increases with increasing DL and does not reach a steady state. Particularly, significant oscillation of FL, i.e., stick-slip phenomenon, is observed at the lower scratching rate. DN has a tendency to increase with increasing FL. On the other hand, in nano-scratching in air, FL increased rapidly at the initial stage and then took a steady state. The fluctuations in FL and DN during nano-scratching in air were very small as compared to those during nano-scratching in solution. Moreover, the maximum value of FL during nano-scratching in solution was about 10 times as much as the steady state value of FL during nano-scratching in air in spite of the fact that the thickness of anodic oxide film on tantalum does not change significantly between solution and air. This means that the high mechanical resistance arises against the movement of indenter for the tantalum surface kept at 5.0 V (SHE) in solution. The AFM images of grooves produced by nano-scratching indicated that the protrusions are significantly accumulated at the moving front of indenter for nano-scratching in solution, which would provide the high mechanical resistance against the movement of indenter to increase FL. In contrast, no accumulation of protrusions was observed for nano-scratching in air. The anodic oxide film is ruptured during nano-scratching in solution or in air since the scratching depth (40 nm – 60 nm) is more than 3 times as large as the thickness (about 13 nm) of anodic oxide film. Rapid formation of anodic oxide at the ruptured sites would take place at the moving front of indenter during nano-scratching in solution. The anodic oxide may adhere to the moving front of indenter although its mechanism is unclear. As the indenter moves to the lateral direction, the anodic oxide would be accumulated at the moving front of indenter. On the other hand, for nano-scratching in air, oxide formation and growth may be slow since oxidation of ruptured sites is only brought by oxygen in air. Therefore, the amount of oxide adhered to the moving front of indenter would be negligibly small for nano-scratching in air. Fig. 7 shows the average friction coefficient, µ’ave, as a function of logarithm of time, log t, required for scratching. The average friction coefficient, µ’ave, was obtained by graphic integration of lateral force, FL, with lateral displacement, DL and then dividing by DL = 2 µm. The time, t, required for scratching is given by t = DL / v

(3)

where is the scratching rate. The average friction coefficient, µ’ave, increases linearly with increasing logarithm of time, log t, for nano-scratching in solution, while µ’ave is almost independent of log t for nano-scratching in air. The linear relation between µ’ave and log t may mean that the accumulation of anodic oxide at the moving front of indenter is given as a function of log t. The significantly large values of µ’ave obtained with nano-scratching in solution as compared to

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those obtained with nano-scratching in air would be ascribed to the accumulation of anodic oxide at the moving front of indenter.

Average friction coefficient, µ'ave

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Fig. 7 Average friction coefficient, µ’ave, as a function of logarithm of time, log t, required for scratching.

4. Conclusions The following conclusions were drawn from comparison between mechanical properties of passive metal surfaces obtained with nano-indentation or nanoscratching in pH 8.4 borate solution and in air after passivation. 1) The dichromate treatment increased the hardness of passive iron (100) surface kept at 0.25 V (SHE) for 1 h. The effect of dichromate treatment on hardness was more significant in solution than in air after passivation, which was ascribed to the promotion of repassivation at the ruptured sites of passive film due to enrichment of chromium in passive film. 2) The friction coefficients obtained with nano-scratching in solution for the passive iron (100) surfaces kept at various potentials for 1 h were always larger than those obtained with nano-scratching in air after passivation at various potentials. Moreover, the friction coefficient obtained with nano-scratching in solution increased with increasing formation potential of passive film, while the friction coefficient obtained with nano-scratching in air was not sensitive to the

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formation potential. It was inferred that the repassivation at the ruptured sites which takes place at the moving front of indenter gives the mechanical resistance against the lateral movement of indenter to contribute to the increase in friction coefficient. The large friction coefficient and its potential dependence obtained with nano-scratching in solution were ascribed to the increase in repassivation rate at the ruptured sites with increasing formation potential of passive film. 3) The stick-slip like phenomenon was observed during nano-scratching in solution for the tantalum (100) surface anodically oxidized at 5.0 V (SHE) for 1 h. The average friction coefficient obtained with nano-scratching in solution was significantly larger than that obtained with nano-scratching in air after anodic oxidation and increased with increasing logarithm of time required for scratching, which was associated with the accumulation of anodic oxide at the moving front of indenter. References 1. 2. 3. 4. 5. 6. 7. 8. 9.

M. Seo and M. Chiba, Electrochim. Acta 47 (2001) 319. M. Chiba and M. Seo, Corros. Sci. 44 (2002) 2379. M. Chiba and M. Seo, J. Electrochem. Soc. 150 (2003) B525. M. Chiba and M. Seo, Electrochim. Acta 50 (2004) 967. B. Bhushan (Ed.), Handbook of Micro / Nano Tribology, CRC Press. New York ,1995, p. 321. N. Sato, K. Kudo and T. Noda, Electrochim. Acta 16 (1971) 1909. J. B. Lumsden and Z. Szklarska-Smialowska, Corrosion 34 (1978) 169. M. Seo and N. Sato, Corrosion-89, New Orleans, (1989) NACE Paper No. 138. T. Sakon, Master Thesis, Graduate School of Eng., Hokkaido University (1982).

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Current Transients of Passive Iron during Microindentation in Solution Koji Fushimi, Ken-ichi Takase, and Masahiro Seo Graduate School of Engineering, Hokkaido University Kita-13 Jo, Nishi-8 Chome, Kita-ku, Sapporo 060-8628, Japan [email protected]

Abstract - Micro-indentation of passive iron surface in borate solution was carried out to investigate the rupture and repair of passive film. During downward and upward driving a micro-indenter, a couple of anodic current peaks were observed. Both current peaks were responsible for partial exposure of iron substrate to the solution due to rupture and repair of passive film. They were sensitively influenced by electrode potential or concentration of sulfate ions containing in solution. The model for a series of rupture and repair processes of the passicve film by micro-indentation was proposed. Keywords: Current transient; Passive film; Iron; Micro-indentation; Mechanoelectrochemistry

1. Introduction Our previous studies of nano-indentation and nano-scratching in solution to passive metal surfaces such as iron and titanium [1-3] have suggested that the hardness and friction coefficient are influenced by mechano-electrochemical reaction such as mechanical breakdown of passive films followed by repassivation at the breakdown sites. The measurement of current transients during indentation or scratching in solution is essentially necessary to investigate the mechano-electrochemical reaction. For the nano-indentation, however, the indented area is so small (10-14 – 10-12 m2) that it is difficult to measure the current transients in the order of 10-15 – 10-13 A. Here micro-

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indentation in solution was performed to a polycrystalline iron surface kept at passive state so as to measure the current transients in the order of 10-9 A. 2. Experimental The polycrystalline iron plate (12 x 12 x 2 mm3) with a purity of above 99.99% was used. The iron specimen was mechanically polished with a diamond paste (3 m) and ultrasonically cleaned in ethanol. The electrolytes were pH 8.4 borate buffer solutions without and with 5x10-3, 2x10-2, 5x10-2, or 0.1 mol dm-3 Na2 SO4, which were deaerated with purified Ar gas during experiments. A conical diamond indenter with an including angle of c.a. 110o was employed for micro-indentation. For the indentation in solution, the diamond indenter attached to a force sensor was normally moved downward and upward with a piezoelectric-driving system to the passive iron surface kept at constant potential. The indenter was moved at a rate of 10 mm s -1 with an intermission of 10 s, up to a maximum depth of 60 m. The maximum load corresponding to the maximum depth was less than 0.5 N. 3. Result and discussions

h / µm

b)

0 -20 -40 -60 0.3 0.2 0.1 0.0 200

L/N

L/N

a)

h / µm

Fig. 1 shows typical transients of load and current flowing through the passive iron surface polarized at 0.7 V vs. SHE in pH 8.4 borate buffer solution when the indenter drove downward and upward. The load jumps to 0.25 N by downward driving of the indenter and returns to zero, while the load decreases

150

I / nA

I / nA

150 100 50 0

0 -20 -40 -60 0.3 0.2 0.1 0.0 200

100 50

0 20 40 60 80 100

t / ms

9801000

0

0 20 40 60 80 100

9801000

t / ms

Fig. 1 Time-transients of (middle) load and (bottom) current flowing through iron electrode polarized at 0.7 V vs. SHE in deaerated pH 8.4 borate buffer solution for a) downward and b) upward driving of indenter as shown in (top) height changes. The intermission between downward and upward driving of indenter was 10 s.

by 0.05 N during the intermission as shown in load-depth like curve of Fig. 2a. This slight decrease in load may come from the stress relaxation of specimen

453

Current Transients of Passive Iron during Micro-indentation in Solution

surface due to movement of dislocations. The load becomes zero at a depth of ca. 10 m in the unloading curve, which corresponds to plastic deformation. Two current peaks appear during downward and upward driving. The first peak appears at the initial stage of downward driving, i.e., at the moment when the indenter penetrates into the passive iron surface. The second peak emerges at the final stage of upward driving, i.e. at the moment when the indenter is removed from the surface. Each current peak decayed within 1 s, which was sufficiently short as compared to the intermission. In Fig. 1, the current peak height for upward driving is 3 times as much as that for downward driving. Moreover, the electric charge corresponding to the current peak areas for for down ward and upward driving were 2.3 and 6.3 nC, respectively. b)

c) 200

200

0.2

150

150

0.1

I / nA

0.3

I / nA

L/N

a)

100

50

100

50

0.0 0

-20

h / µm

-40

-60

0

0

-20

h / µm

-40

-60

0

0.0

0.1

0.2

0.3

L/N

Fig. 2 Relation between load and depth during indentation (Fig. 1). Current as a function of b) indentation depth or c) load.

Figs. 2b and 2c are the currents plotted versus indenter-displacement in depth and change in load, respectively. The current during downward driving increases with increase in penetration depth or load after the indenter contacts with the passive surface until depth or load reaches a maximum, and then decreases independent of depth or load. However the current during upward driving starts to increase from the intermediate point, dramatically increases with decrease in depth or in load until the plastic deformation depth at which the indenter tip releases from the surface, and then decays to the passivitymaintaining current. It is clear that the current flowing from the specimen surface during the indentation is strongly influenced by the relative position of the indenter tip and deformed surface. The appearance of a couple of current peaks during downward and upward driving throws light on a series of processes during the indentation including contact of indenter with passive surface, film and substrate deformations, film rupture and crack formation, exposure of the substrate to solution, and film repair. Each process is not only mechanical, chemical or electrochemical phenomenon but also synergistically related each other. The calculated ratio of contact area with indenter to its projected area is larger than 1.22. The indenter is penetrated by ca. 60 m into the iron substrate. The ellipsometrical thickness [4,5] of passive film

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formed on iron at 0.7 V vs. SHE is c.a. 4 nm. The penetration depth of the indenter, therefore, is larger by 4 decade than the film thickness, suggesting that significant area of passive film is ruptured. As the indenter drives downward, the indenter tip contacts with the surface and the load to the surface increases progressively. At the very initial stage of loading, the film as well as the substrate deforms elastically. The load, however, exceeds the yielding stress of the passive film to deform plastically and eventually the film is ruptured or cracked, since the passive film formed on iron is harder than the substrate [1,3]. The indenter penetrates more deeply, the total ruptured area becomes larger. On the other hand, upward driving gradually releases the load, a part of which is rivaled against elastic deformation of the substrate. As far as the load is rivaled against elastic deformation, the indenter tip keeps contact with the specimen surface, which means that most of the film area ruptured during downward driving is not yet exposed to solution. The exposure of ruptured area to solution is allowed at the later stage of upward driving, where the load as well as the elastic deformation decreases and the contact area is so sufficiently reduced that the ruptured area is exposed to solution and then current flows to repair the ruptured area. At the intermission stage, the indenter tip covers the ruptured area and no current flows. The electric charge for current peak, e.g., 6.3 nC in Fig. 1b, is too small to repair the ruptured area if the ruptured area is bare and free from oxide film since, assuming Fe2 O3 film, 580 nC is required for formation of oxide film with a thickness of 4 nm on the contact area corresponding to the penetration depth, 60 m, of indenter. It means that the area of passive film contacted with indenter during indentation is not completely ruptured to create a small number of cracks (only 1 % of contact area). 300 250

upward

Ip / nA

200 150 100 downward 50 0 0.0

0.2

0.4

0.6

0.8

1.0

E / V vs. SHE

Fig. 3 Relation between peak current, Ip, and electrode potential, E, during downward and upward driving of the indenter in case where the indentation was made for the passive iron surfaces polarized at various potentials in deaerated pH 8.4 borate buffer solution.

When the indentation was performed for the passive surfaces polarized at various potentials in pH 8.4 borate buffer solution, a maximum load of Lmax =

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455

0.25 ± 0.03 N was applied to the surfaces and the similar current transients for downward and upward driving were obtained as shown in Fig. 1. Fig. 3 shows the relation between peak current, Ip, which was defined as a maximum of the current subtracted by passivity-maintaining current, and electrode potential, E, during downward and upward driving. Both Ip for downward and upward driving increase linearly with increase in E, indicating that the film repair is accelerated by a potential drop at the interface between bare iron substrate and solution which is almost proportional to the increase in potential. Peak current was mainly due to charging of the double layer, in which current is proportional to potential difference applied. Although the hardness of passive iron surface determined by nano-indentation [1] increased slightly with increase in film formation potential, the effect of potential to the hardness might be negligible because the indent-depth was significantly deeper than the passive film thickness. -5.8 unloading

log(Ip' / A N-1)

-6.0 -6.2 -6.4 loading -6.6 -6.8

-inf -3

-2

-1 -3

log (csulfate / mol dm )

Fig. 4 Relation between logarithm of peak current per maximum load, log (Ip’), and logarithm of sulfate concentration, log (csulfate) where the indentation was made for the passive iron surfaces polarized at 0.7 V vs. SHE in deaerated pH 8.4 borate buffer solution containing Na2SO4.

After the passivation of iron electrode in pH 8.4 borate buffer solution, solution was exchanged to the solution containing Na2SO4 up to 0.1 mol dm-3 and then the indentation was made under polarization at the same potential employed for passivation. The loads applied to the surfaces were a little scattered and maximum load Lmax = 0.26 ± 0.05 N was obtained. Two anodic current peaks were observed by the indentation similar to Fig. 1. Compared to pure borate buffer solution, the value of Lmax is scattered and Ip has a poor reproducibility in sulfate containing solution, especially at higher sulfate concentration. Moreover, Ip had almost linear relationship with Lmax in each solution. Thus, the peak current per maximum load, Ip’, is defined by dividing Ip with Lmax. Fig. 4 is the relation between logarithm of peak current per maximum load, log(Ip’), and logarithm of sulfate concentration, log(csulfate), when the electrode was polarized at 0.7 V vs. SHE. Both Ip’ during downward and upward driving increase linearly with increase in log(csulfate). It is suggested that sulfate ion in solution

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accelerates some processes during indentation including rupture of passive film, dissolution of substrate, and repair of the ruptured sites. Assuming that Ip’ is composed of two terms; one is Ip,0’ which is consumed during the indentation, independent of sulfate concentration and other is that directly related to sulfate concentration, csulfate; Ip’ = Ip,0’ csulfate,, (1) where n is a constant corresponding to the slope in Fig. 4. In extremely high concentration, the linear relationships in Fig. 4 would intersect because n for downward driving is larger than that for upward driving. This suggests that the film repair or repassivation during penetration of indenter is difficult in solution of high sulfate concentration. Furthermore, the larger value of n for downward than upward means that some process in downward driving is more sensitively influenced by sulfate ion. It seems that the increase in Ip’ due to sulfate ion includes both mechanical and electrochemical, i.e., mechano-electrochemical effect. The passivity-maintaining current in sulfate-containing solution was slightly larger than that in the solution without sulfate. This comes from an electrochemical effect that the film in sulfate-containing borate solution is less stable. However sulfate ion may also exert an influence on mechanical effect to the passive surface since the reduction in load during intermission decreased with increase in sulfate concentration, e.g., 0.03 N in 0.1 mol dm-3 sulfate containing solution, indicating that stress relaxation of the specimen was prevented by the existence of sulfate ion although the role of sulfate ion was not clear. The area of ruptured film, i.e., the crack density may increase as stress relaxation of film becomes difficult. 4. Conclusion Micro-indentaion was applied to the passive iron surface in deaerated pH 8.4 borate solution. A couple of current peaks were observed during downward and upward driving of the diamond indenter. The increase in potential accelerates the peak currents during the micro-indentation, due to electrochemical effect. The influences of sulfate ions in solution on the current peaks were also investigated and it was stressed that the mechanical and electrochemical interactions, i.e., mechano-electrochemical effects worked for repassivation at the ruptured sites of passive film during indentation. References [1] [2] [3] [4] [5]

M. Seo and M. Chiba, Electrochim. Acta 47 (2001) 319. M. Seo and Y. Kurata, Electrochim. Acta 48 (2002) 3221. M. Chiba and M. Seo, Electrochim. Acta 50 (2004) 967. N. Sato, K. Kudo and T. Noda, Electrochim. Acta 16 (1971) 1909. K. Azumi, T. Ohtsuka and N. Sato, DENKI KAGAKU 53 (1985) 306.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Published by Elsevier B.V.

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Mechanical Properties of Single-Crystal Tantalum (100) Surface Covered with Anodic Oxide Film Daisuke Kawamata, Masahiro Seo* Graduate School of Engineering, Hokkaido University Kita-13 Jo, Nishi-8 Chome, Kita-ku, Sapporo 060-8628, Japan E-mail : [email protected]

Abstract - Nano-indentation and nano-scratching were performed to the single-crystal tantalum (100) surface covered with a thick anodic oxide film to compare the mechanical properties of anodic oxide film and tantalum substrate. The hardness of the tantalum surface covered with anodic oxide film was about three times as much as that of the electropolished tantalum surface. The friction coefficient of the tantalum surface covered with anodic oxide film as well as that of electropolished tantalum surface increased with increasing normal displacement due to the tip geometry. The normalized friction coefficient was defined by taking the tip geometry into account and it was independent of normal displacement. The normalized friction coefficient of the tantalum surface covered with anodic oxide film was less by about 20 % than that of the electropolished tantalum surface. Keywords: Nano-indentation; Nano-scratching; Hardness; Friction coefficient; Tip geometry

1. Introduction It is very important to investigate the relationship between mechanical and chemical properties of passive metal surface for better understanding of local breakdown of passive film and initial stage of pitting corrosion. There have been many studies on chemical properties of passive film such as composition, structure and film thickness. However, there have been few studies on mechanical properties of passive film such as hardness and friction coefficient.

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Nano-indentation and nano-scratching techniques have been recently developed to evaluate the mechanical properties of local sites of materials surfaces such as hardness and friction coefficient. In our previous study [1 - 3], the mechanical properties of passive iron and titanium surfaces were measured. However, the thickness of passive films on iron and titanium are so thin (2 - 20 nm) that the measured mechanical properties are the composite of passive film and metal substrate. In order to measure the mechanical properties of oxide film itself, the oxide film much thicker than the indentation depth, has to be formed on metal. In this study, nano-indentation and nano-scratching were performed to the single-crystal tantalum (100) surface on which thick oxide film was anodically formed. 2. Experimental A single-crystal tantalum (100) disk plate with a diameter of 10 mm and a thickness of 1 mm was employed in this experiment. The tantalum surface was mechanically polished with -alumina abrasives (0.3 µm), and then ultrasonically washed with acetone. Finally, to remove worked layer, i.e., Beilbly layer, the tantalum surface was electropolished at a current density of 78.4 mA cm-2 in a mixed solution of 47 % HF + 96 % H2SO4 (1 : 10). The tantalum surface was anodically oxidized at a constant current density of 150 µA cm-2 in pH 8.4 borate buffer solution up to a cell voltage of 200 V to form anodic oxide film with a thickness of about 310 nm [4]. The nano-indentation / nano-scratching apparatus (Triboscope, Hysitron Inc.) was combined with AFM (Nanoscope IIIa, Digital Instruments). A conical diamond indenter with a tip radius less than 1 µm and an included angle of 90o was used for nanoindentation and nano-scratching. For nano-indentation, load-depth curves were measured at a maximum load of Lmax = 100 to 1800 µN. Hardness, H, is defined by dividing maximum load, Lmax, with projected contact area, A. H = Lmax / A

(1)

For nano-scratching, lateral force, FL, and normal displacement, DN, were measured under a constant normal force of FN = 100 to 1800 µN at a scratching speed of v = 0.2 µm s-1 as a function of lateral displacement, DL, between lateral distance of 2 µm. Friction coefficient, µ’, is defined by dividing FL with FN.

µ’ = FL / FN

(2)

3. Results and discussion Fig.1 shows the hardness, H, as function of contact depth, hc for the tantalum surfaces electropolished and anodically oxidized up to 200 V. The hardness (8.6 GPa) of the tantalum surface covered with anodic oxide film was three times as much as that (2.8 GPa) of the electropolished tantalum surface. The hardness of

Mechanical Properties of Single-Crystal Tantalum (100) Surface…

459

the tantalum surface covered with anodic oxide film increased with increasing contact depth, hc. On the other hand, the hardness of the electropolished tantalum surface was independent of the contact depth, hc. The thickness of airformed film on electropolished tantalum surface is about 4 nm [4], which is less than 10 % of contact depth. Therefore, the hardness of electropolished surface in the contact depth range larger than 100 nm would be regarded as that of the substrate. On the other hand, the thickness of anodic oxide film is more than

three times as much as the contact depth. The hardness of the tantalum surface covered with the anodic oxide film in the contact depth range less than 30 nm would correspond to that of the anodic oxide film itself. The increase in hardness with contact depth for the tantalum surface covered with anodic oxide film may be associated with underestimation of contact depth due to pile-up [5]. Fig.1 Hardness, H, as a function of contact depth, hc, for the tantalum surfaces electropolished and anodically oxidized up to 200 V.

Fig.2 shows the friction coefficient, µ’, as a function of normal displacement, DN, for the tantalum surfaces electropolished and anodically oxidized up to 200 V. These friction coefficients are averaged values measured 20 times at each normal force. Both friction coefficients of the tantalum surface electropolished and anodically oxidized increase with increasing normal displacement. According to Amontons and Coulomb’s law, lateral force is proportional to normal force, i.e., friction coefficient should be independent of normal force. Since a normal displacement depends on a normal force, the results of Fig.2 mean that the friction coefficient does not obey the Amontons and Coulomb’s law.

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Fig.2 Friction coefficient, µ’, as a function of normal displacement, DN, for the tantalum surfaces elecropolished and anodically oxidized up to 200 V.

In general, a friction coefficient, µ’, consists of an adhesion term, µ’a, and a ploughing term, µ’p, [6].

µ ’ = µ ’a + µ ’p

(3)

As previously reported by Chiba and Seo [3], the ploughing term has the major contribution to the friction coefficient obtained with nano-scratching. The ploughing term of friction coefficient, µ’p, is given by Eq.(4)

µ ’p = k

f

/H

(4)

where k is the parameter associated with tip geometry, f, the plastic flow pressure of materials against lateral movement of the diamond tip during nanoscratching and H is the hardness. Furthermore, k is given by Eq.(5) k = Av / Ah

(5)

where Av and Ah are the vertical and horizontal cross sectional areas of the conical diamond tip in contact with the specimen surface, respectively. If the conical diamond tip with an included angle of 90o has an ideal geometry, k takes a constant value of 0.64. The real diamond tip, however, has not ideal but round shape. In the case of non-ideal shape, k is less than 0.64 and approaches 0.64 as the normal displacement increases. The value of k can be estimated from the AFM observation of the indent obtained by nano-indentation of fused quartz as a standard specimen [3]. If the value of k is estimated as a function of normal

Mechanical Properties of Single-Crystal Tantalum (100) Surface…

461

displacement, the normalized friction coefficient, µ’’, may be newly defined by dividing µ’ with k to eliminate the dependence of k on normal displacement.

µ’’ = µ’p / k =

f

/H

(6)

In Eq. (6), it is assumed that the adhesion term is negligible. In this study, the values of k were estimated as a function of normal displacement from the AFM observation of indents on fused quartz. Fig.3 shows the normalized friction coefficients, µ’’, as a function of normal displacement, DN for the tantalum surfaces electropolished and anodically oxidized. As seen from Fig.3, the values of µ ’’ are independent of

Fig.3 Friction coefficient, µ’’, as a function of normal displacement, DN, for the tantalum surfaces elecropolished and anodically oxidized up to 200 V.

normal displacement within some scatters. This means that the dependence of friction coefficient, µ’, on normal displacement in Fig.2 is attributed to the tip geometry. The normalized friction coefficient of the tantalum surface covered with anodic oxide film is less by 20 % than that of the electropolished tantalum surface. On the other hand, the hardness of tantalum surface covered with anodic oxide film is about three times as much as that of the electropolished surface. The plastic flow pressure, f, can be estimated from Eq.(6). The plastic flow pressure (8.2 GPa) of the tantalum surface covered with anodic oxide film was 2.2 times as much as that (3.6 GPa) of the electropolished surface. This difference in plastic flow pressure may be mainly caused by the difference in resistance against the lateral movement of indenter tip between anodic oxide film and tantalum substrate.

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4. Conclusions Nano-indentation and nano-scratching were performed to the tantalum (100) surfaces electropolished and covered with anodic oxide film. The following conclusions were drawn. 1) The hardness (8.6 GPa) of the tantalum surface covered with anodic oxide film was about three times as much as that (2.8 GPa) of the electropolished surface. 2) The friction coefficient of the tantalum surface electropolished and covered with anodic oxide film increased with increasing normal displacement due to the tip geometry. 3) The normalized friction coefficient was newly defined by dividing friction coefficient with geometrical factor to eliminate the effect of the tip geometry on friction coefficient. As a result, the normalized friction coefficients were independent of normal displacement. 4) It was suggested from the difference in normalized friction coefficient that the plastic flow pressure of anodic oxide film was larger than that of the tantalum substrate. References 1. 2. 3. 4. 5. 6.

M. Seo, M. Chiba, Electrochim. Acta, 47 (2001) 319. M.Seo, Y.Kurata, Electrochim. Acta, 48 (2003) 3221. M. Chiba, M. Seo, Electrochim. Acta, 50 (2004) 967. T. Sakon, Master Thesis, Graduate School of Eng., Hokkaido University (1982). A.Bolshakov, G.M.Pharr, J.Mater.Res, 13 (1998) 1049. D.F. Moore, “Principles and Applications of Tribology”, Pergamon Press, Oxford (1975) p.44.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Study of the mechanical effects on passivity breakdown by local probe techniques V. Vignal, N. Mary and R. Oltra LRRS, UMR CNRS 5613, Université de Bourgogne, BP 47870, 21078 Dijon France Email : [email protected] Abstract - This paper aims at demonstrating that a relationship exists between surface stress and pitting corrosion. The surface stress field generated by polishing was first calculated using a thermo-mechanical model and a finite element code. Pitting corrosion tests performed at the microscale along the austenite/ferrite interface using the electrochemical microcell technique were then analyzed considering the microstructure and the residual surface stress field. Mechanochemical criteria are proposed leading to an enhancement of pitting corrosion of duplex steels. Keywords : Duplex steels, local probes, pitting, residual stress

1. Introduction Duplex stainless steels are being increasingly used as structural materials in chemical and petrochemical applications, in power plants, etc. This is mainly due to their high resistance to localised corrosion as well as to their high strength and toughness. The two phases of duplex stainless steels have different mechanical properties and coefficients of thermal expansion and high stress gradients can originate from these differences under straining conditions or after various surface treatments [1-2]. The presence of high stress gradients may also affect the corrosion resistance of each phase and consequently of the entire metallic alloy. To our knowledge, no works have been devoted to study at the microscale the influence of residual surface stresses on the electrochemical behaviour and the pitting susceptibility of duplex stainless steels. However, the evolution of the time to pit was analysed at the macroscale on the same duplex steel under straining conditions in 0.5M NaCl, pH6.5 at 50°C for various potentials applied in the passive range [3]. The results obtained for applied stresses below the apparent yield strength of the duplex stainless steel indicated

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that its electrochemical behaviour can be correlated to the surface stress state measured by XRD techniques within the austenite. For an applied stress below 34% of the apparent yield strength (Rp0.2), compressive surface stresses were measured within this phase and the corrosion tests show mainly no pitting. Between 34 and 54% of Rp0.2, tensile surface stresses were determined and the time to pit was mainly located between 0.08 and 53 hours. Above 54% of Rp0.2, microplasticity occurred and the time to pit was found to drop (between 0.02 and 0.32 hours). 2. Experimental 2.1. Material and surface preparation Experiments were performed on a duplex stainless steel UNS S31803. The chemical composition of this material is (wt.%): Cr: 22.616%; Ni: 5.45%; Mo: 2.886%; Mn: 1.653%; Si: 0.471%; N: 0.1725%; Cu: 0.133%; P: 0.021%; C: 0.02% and S: 0.0009%. After hot rolling, the plates (thickness of about 20 mm) were solution annealed at 1050°C for 15 minutes and quenched in water. The grain structure was heavily banded and the grain size in the short transverse direction was about 10 µm, as shown in Fig. 1(a). This grain size was not suitable for performing local electrochemical measurements within the austenite and the ferrite and a second heat treatment was carried out. It consists of a homogeneisation treatment at 1300°C for 1 hour, followed by a slow cooling down to 1080°C (formation of the austenite) and a water quenching. The material showed a globular microstructure with a grain size of about 75 µm, as shown in Fig. 1(b).

Fig. 1 : Optical micrographs of the (a) lamellar microstructure and (b) globular microstructure of the duplex stainless steel.

Specimens were mechanically ground by using SiC papers (down to 4000 grade) and polished with several grades of diamond pastes (from 6 µm down to

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1 µm). Between each polishing, they were cleaned in ethanol under ultrasonics for 5 minutes. 2.2. Local electrochemical measurements Local measurements were carried out in 15M LiCl using the electrochemical microcell technique [4]. This technique consists of a glass microcapillary which is filled with electrolyte. The tip of the microcapillary (12 µm in diameter) is sealed to the specimen surface with a layer of silicon rubber. The microcell is mounted on a microscope for precise positioning of the microcapillary on the surface and the entire setup is placed in a Faraday cage. A modified high resolution potentiostat is used to have a current detection limit of 20 fA (Jaissle 1002T-NC-3). A cathodic prepolarisation of –300 mV/SCE was applied for 2 min (in order to reach a steady-state in current and to reduce the passive film) after which the current was plotted versus time for 15 minutes at +150 mV/SCE with a measurement frequency of 18.2 Hz. All potentials were measured versus a saturated calomel electrode (SCE) and the counter electrode was a platinum wire. 3. Results and discussion 3.1. Surface stress field induced by polishing The numerical calculation of the surface stress field involves the use of a new procedure in the finite-element code for meshing the real microstructure, as shown in Figs. 2(a) and (b), and a material model (including the real texture and the real mechanical characteristics of each phase) capable of describing the anisotropic behaviour of these materials. Mechanical polishing using 1 µm diamond paste induces a small and homogeneous variation of the lattice plane spacing with respect to the stress-free lattice plane spacing such that the elasticity theory can be used. The generalized Hooke’s law was then considered and the terms of the stiffness tensor expressed in the crystallophysic coordinates were found in Refs [5-6]. The coefficients of thermal expansion between 25°C and 600°C were found out in Refs [7-8]. The surface stress field resulting from mechanical polishing was then calculated using a 2D thermomechanical model [9], as shown in Fig. 2(c). In order to verify the validity of the numerical method, the values of the average total stress within both phases (values determined at the macroscale) were compared to the experimental values derived from standard XRD measurements [9]. The numerical method described above was applied to the globular microstructure of the duplex steel UNS S31803. The austenite was found to be more in tension than the ferrite after mechanical polishing, leading to a large stress step across the austenite-ferrite interface, as shown in Fig. 2(c). Some

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large stress gradients which extend over an average distance of about 15 µm were also detected on both sides of the austenite-ferrite interface. This result permits the determination of a characteristics length of about 15 µm which dictates the scale of analysis in mechano-electrochemistry.

Fig. 2 : (a) Optical micrograph of austenite grains surrounded by the ferrite, (b) meshing of this microstructure and (c) surface stress field along the X1-axis (rolling direction), s11.

3.2. Local mechano-electrochemical analysis The mechanical parameter in equation 1, which takes into account the influence of the local average stress <s>loc and the local stress amplitude (smin-smax)loc, was calculated within the sites where microelectrochemical measurements were performed. The diameter of the sites selected was choosen equal to the characteristics length defined previsouly (15 µm in diameter). w = 2.<s>loc.

(smin-smax)loc A

(1)

where smin and smax is the lowest stress and the highest stress within the selected site of area A, respectively. Regarding the duplex steel UNS S31803, the values of w was found to range from 5 to 50 MPa/µm along the austenite-ferrite interface and the sites defined by w > 40 MPa/µm constituted a minority group.

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467

A tentative criterion was used to classify these sites : low stress gradients were defined by w £ 15 MPa/µm whereas medium and high stress gradients were defined by 20 £ w £ 32 MPa/µm and w ³ 32 MPa/µm, respectively. It has been demonstrated that w and this criterion are relevant parameters for quantifying the role of stress gradients on pit initiation at the microscale [10]. 28 corrosion tests were carried out for low and medium stress gradients (w £ 15 MPa/µm and 20 £ w £ 32 MPa/µm, respectively) and 15 corrosion tests were performed for high stress gradients (it was difficult to locate sites defined by w ³ 32 MPa/µm), as shown in Fig. 3. A tentative criterion was proposed to define the time to pit as the time of appearance of the first current transient characterized by an electrical charge of 4 pC. The time to pit was then plotted vs. the parameter w, as shown in Fig. 4.

Fig. 3 : Current vs. time evolution during corrosion tests at +150 mV/SCE in the case of (a) no pitting, (b) stable pitting and (c) metastable pitting.

Fig. 4 : Evolution of the time to pit at +150 mV/SCE vs. the parameter w. The selected sites contain an austenite/ferrite interface (surface ratio : 50/50).

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When w £ 15 MPa/µm, no pitting was detected within 15 minutes, as shown in Figs. 3(a) and 4. By contrast, stable pitting (Fig. 3(b)) or metastable pitting (associated with current transients as those observed in Fig. 3(c)) were systematically observed for w ³ 20 MPa/µm (Fig. 4). 4. Conclusion The influence of microstress gradients generated by polishing on the pitting susceptibility of a duplex steel was investigated in 15M LiCl at 25°C under potentiostatic control using the electrochemical microcell technique. The following conclusions can be drawn : (i) below a stress gradient of about 15 MPa/µm, no pits initiated for 15 minutes immersion whereas above 20 MPa/µm, metastable or stable pits were systematically observed. (ii) stable pits were mainly encountered for stress gradients between 20 and 32 MPa/µm 5. Acknowledgements The authors are grateful to CEA-DRN/DMT/SEMT (Saclay, France) which has designed and developed the FE code Cast3M and to J. Peultier and C. Lojewski (Industeel, Arcelor group) for providing materials and for stimulating discussions. One of the author (N.M.) would like to thank the Conseil Régional de Bourgogne (France) for the financial support. 6. References 1. V. Vignal, J. Favergeon and R. Oltra, Phil. Mag. Lett., 82, 503 (2002). 2. J. Johansson, M. Oden and X.H. Zeng, Acta Mater., 47, 2669 (1999). 3. V. Vignal, N. Mary, C. Valot, R. Oltra and L. Coudreuse, Electrochem. Solid-State Lett., 7, C39 (2004). 4. H. Bohni, T. Suter and F. Assi, Surf. and Coat. Techn., 130, 80 (2000). 5. Y.D. Wang, R.L. Peng and R. Mcgreevy, Scripta Mater., 41, 995 (1999). 6. E.A. Brandes, in Smithless metals reference book, 6th Edition, Edited by E.A. Brandes and the Butterworth & Co publisher, p. 15-5 (1983). 7. T. Siegmund, E. Werner and F.D. Fischer, J. Mech. Phys. Solids, 43, 495 (1995). 8. F.D. Fischer, F.G. Rammerstorfer and F.J. Bauer, Metall. Trans., 21A, 935 (1990). 9. N. Mary, V. Vignal, R. Oltra and L. Coudreuse, Phil. Mag., 85, 1227 (2005). 10. N. Mary, V. Vignal, R. Oltra and L. Coudreuse, J. Mat. Res., 19, 3688 (2004).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Mechanical properties of metastable r.f magnetron sputter-deposited Al1-xCux thin films Mohamed Draissia , Mohamed Y. Debili LM2S, Physics department, Badji-Mokhtar University, BP12 Annaba, 23200 Algeria [email protected]

Abstract- The mechanical properties of a family of homogeneous Al1-xCux (0 < x < 0.22) thin films made by radiofrequency (13.56 MHz) cathodic magnetron sputtering from composite Al-Cu targets have been investigated. The as-deposited microstructures for all film compositions consisted of a mixture of the two expected face-centered-cubic (fcc) Al solid solution and tetragonal (Al2Cu) phases. The microhardness regularly increases and the grain size decreases both with copper concentration. This phenomenon of significant mechanical strengthening of aluminium by means of copper is essentially due to a combination between solid solution effects and grain size refinement. Keywords : Aluminium alloys, Intermetallic compounds, Sputtering, Microstructure, Microhardness.

1. Introduction The addition of an alloying element to a metal reinforces the mechanical hardening of this one. The hardening observed in an alloy depends primarily on the nature of the base metal, the additional element characteristics and the elaboration method and parameters. The mechanical hardness of a metallic material increases typically with decreasing the grain size [1,2]. Microcrystalline and nanocrystalline materials can be currently produced by several elaboration methods and the resulting metal has a polycrystalline structure without any preferential crystallographic grain orientation. Gas-phase deposition processes using physical or chemical methods (PVD or CVD) are currently used to produce thin films coatings for mechanical engineering industries. Aluminium and its alloys with their low density and easy working occupy a significant place in the car industry, aeronautics and food

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conditioning. The on-glass slides sputter-deposited aluminium-based alloys thin films, such Al-Mg [3], Al-Ti [4], Al-Cr [5] and Al-Fe [6], exhibit a notable solid solution of aluminium in the films and microhardness higher than that of corresponding traditional alloys. This work describes firstly the preparation of the Al-Cu targets and the elaboration of the films and secondly the structure and microhardness characterizations. Analogies with other aluminium-based thin films systems based on aluminium are displayed. 2. Elaboration procedure The targets used in the elaboration of the aluminium-copper thin films are made from bulk aluminium crown of 70 mm of diameter in which is inserted a bulk copper disc. Using bulk material minimizes oxygen in the films. This target shape enables an easy control of additional element composition in the films, this one is parabolic with insert diameter and then it is linear with surface fraction (Figure 1). The equation which governs this variation was determined from geometrical target form and aluminium and copper characteristics, so it is [7] : at.%Cu » 100 ´ F/(0.71+0.28 ´ F)

(1)

where F is the insert surface fraction defined as the ratio between the insert copper surface and all the target surface (70 mm of diameter).

Fig.1 Aluminium-copper target shape and evolution of the atomic copper composition with surface fraction in sputtered films.

The films were on glass 75mm´25mm´1mm slides radiofrequency (13.56 MHz) sputter-deposited under low pressure of 0.7 Pa and a substrate temperature which doe’s not exceed 400 °K (130 °C). The substrate-target distance is 80mm. The sputtering is carried out with a constant power of 200 W, an autopolarization voltage of -400 V, that acquired by the plasma is –30 V, a regulation intensity of 0,5 A and a argon flow of 30 sccm. After a 1 hour and 30

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Mechanical properties of metastable sputter-deposited Al1-x Cux thin films

minutes, the deposition velocity is 2.5 thickness were then obtained.

m/h and films of about 3 to 4

m

3. Chemical and Microstructural characterizations The chemical analysis of atomic Cu composition in the Al-Cu films was made by X-ray dispersion spectroscopy. The microstructure of the Al-Cu films was studied by X-ray diffraction (XRD) (Figure 2) and transmission electron microscopy (TEM) (Figure 3).

Fig.2 X-ray diffraction diagrams of the cosputtered Al-Cu films : (a) Al0.982Cu0.018, (b) Al0.9279Cu0.0721 and (c) Al0.7805Cu0.2195.

Fig.3 Bright field TEM micrograph (a) and its associated electron diffraction pattern(b) of the Al0.982Cu0.018 films.

The figures 2 and 3 show that the Al-0 to 21.95 at.%Cu films microstructures consisted of a mixture of a gAl solid solution of aluminium and the expected tetragonal Al2Cu phase (Table 1). The solubility of copper in aluminium is lower than 1.8 at.%Cu in agreement with the Al-Cu equilibrium diagram, this is due to the difference between valences as expected by Hume-Rothery laws and

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the strong tendency to intermetallic compounds formation [7]. However, the sputtered Al-based films [3-6] proceeded under similar conditions exhibit a notable gAl solid solution of substitution of copper in aluminium as shown in table 2, their authors attribute this to difference in size between solvent and solute atoms (Table 2). Composition Insert (mm) Structure

Al 0 gAl

Al0.982Cu0.018 10 gAl + Al2Cu

Al0.9279Cu0.0721 15 gAl + Al2Cu

Al0.7805Cu0.2195 20 gAl + Al2Cu

Tab.1 Relation between the composition and the structure of the co- sputtered Al-Cu thin films.

Characteristics rX (nm) rAl - rX (nm) Structure Valence gAl (at.%) References

Base Al 0.143 0 fcc 3

Alloying element Mg Ti 0.160 0.147 -0.017 -0.006 hcp hcp 2 3 20 27 [3] [4]

Cr 0.127 +0.016 bcc 2 5 [5]

Fe 0.126 +00.17 bcc 2 5.5 [6]

Cu 0.128 +00.15 fcc 1 <1.8 [8]

Tab.2 Solubility extends and characteristics of aluminium and its alloying elements.

4. Microhardness investigations

Fig.4 Microhardness evolution with atomic alloying element in the aluminium-based thin films.

Fig.5 Evolution of the microhardness of the cosputtered Al-Cu thin films with grain size.

The Vickers microindentation was used to determine the intrinsic mechanical hardening of the Al-Cu thin films. The applied load was 1 g corresponding to

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473

films thickness lower than 10 m. The microhardness measurements were converted to that of 10 g applied load and were compared to other Al-based films (Figure 4). This one increases regularly with copper enrichment. We note that chromium and iron increase the hardness of aluminium much more than copper or titanium. This is due to difference in size between the solvent (aluminium) and solute (alloying elements) atoms (Table 2). The mean grain size does not exceed 120 m in the Al0.982Cu0.018 films. The extrapolated Debye-Scherrer [9] grain size measurements from the XRD diagrams are respectively 10 nm, 15 nm and 30 nm for the three studied samples. The figure 5 shows that there is grain size refinement by adding copper with lower lattice volume of 0.0088 nm3 in aluminium with greater lattice of 0.0122 nm3. The straight evolution as expected by Hall-Petch rule [1,2] shows that the films act as solid solution where the copper is incorporated substitutionaly in aluminium matrix and this makes possible to elevate the copper solubility in aluminium over the equilibrium Al(rich)-Cu diagram [8]. 5. Conclusion A family of radiofrequency sputtered aluminium-copper thin films have been elaborated from composite targets where the composition is easily controlled. Their structure consisted of a polycrystalline fcc solid solution aluminium matrix and the expected tetragonal phase. The microhardnes increased and the main grain size decreased regularly with copper contents in the films. This is due to a combination effects of the strong tendency to intermetallic compounds formation and the difference in size between the solvent and solute atoms. Aknowledgments - The elaboration and measurements were made at LSGS de Nancy, the authors are grateful to Dr. Tran Huu loi for technical assistance and Pr. C. Frantz for useful discussions. This work is supported by Research project : D2301/07 /2005

References [1] E. O Hall, Proc. Phys. Soc., London, Vol. 643 (1951), p. 747. [2] N. J. Petch, J. Iron Steel Inst., London, Vol. 173 (1953), p. 25. [3] R.D. Arnell, R.I. Bates, Vacuum, 43 (1992), p. 105. [4] F.Sanchette, Tran Huu Loï, C. Frantz, Surf. Coat. Technol., 74-75 (1995), p. 903. [5] F.Sanchette, Tran Huu Loï, A. Billard, C. Frantz, Surf. Coat. Technol., 57 (1993), p. 179. [6] M. Y. Debili, Tran Huu Loï, C. Frantz, La revue de Métallurgie-CIT/Science et Génie des Matériaux, Décembre (1998), p. 1501. [7] M. Draissia, M.Y. Debili, Accepted publication in issue3 volume 3 (2005) of Central European Journal of Physics. [8] M. Draissia, M.Y. Debili, Crystal Growth, Vol. 270 (2004), p. 348. [9] A. Guinier, Théorie et Technique de la radiocristallographie, 3ème éd. Dunod Paris (1964) p. 461.

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Section G Passivity Issues in Stress Corrosion Cracking and Tribocorrosion

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

477

Passivity Issues in Tribocorrosion D. Landolt Institut des matériaux,Ecole polytechnique fédérale de Lausanne, CH-1015 Lausanne, Switzerland, email: [email protected]

Abstract - Passivity issues play a crucial role for the tribocorrosion behavior of metals and alloys. Rubbing between solids or impingement of particles can damage the protective passive film and thus accelerate the rate of material loss. In this overview the role of repassivation and of depassivation rate is discussed in relation to mechanical film damage mechanisms, materials properties and mechanical variables. Galvanic interactions between the wear scar and the surrounding passive film and the effect of potential distribution are also considered. Finally, areas for future research are mentioned. Keywords: Tribocorrosion, wear, repassivation, depassivation

1. Introduction Most metals and alloys used in engineering are protected against corrosion by passive oxide films. Rubbing between contacting surfaces or impingement of solids or fluids can damage the protecting passive film and thus accelerate corrosion. The field of tribocorrosion deals with the study of the effects of simultaneous corrosion and wear on the rate of degradation of materials and on the performance of tribological systems. In addition to damaging the passive film, rubbing or particle impingement also lead to changes in surface topography and to structural changes of the metal just underneath the surface. Furthermore, the formation of solid corrosion and wear products alters the mechanical conditions in the contact and therefore the mechanical wear rate. All these phenomena can have a positive or negative effect on the rate of mechanical wear and of film damage. In other words, the rates of corrosion and

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wear are strongly interdependent, often in complicated ways. The following discussion shall mostly focus on sliding contacts, but parallels with particle impingement under erosion-corrosion conditions will also be pointed out. When a passive metal rubs against a metallic or non metallic material the passive film at the contacting surfaces is continuously damaged and will spontaneously repair itself by oxidation of metal. The rate of film damage is essentially governed by the mechanical parameters of the contact, the materials properties and the surface topography of the contacting bodies. The rate of film repair, on the other hand, depends critically on the electrochemical properties of the metal and the prevailing potential, which usually is imposed by the nature of the corrosive environment. For modeling the tribocorrosion behavior of passive metals two limiting situations are of particular interest: (a) Transient conditions, where a single depassivation event is followed by film repair and repassivation. Single scratch experiments, nano indentation and single particle impingement studies fall into this category. (b) Steady state conditions where the rate of film damage is equal to the rate of film repair resulting in a constant rate of material loss. Pin-on-disk experiments are a typical example. During rubbing of rough surfaces as well as in erosioncorrosion, microscopic depassivation and repassivation events take place randomly on a surface. The sum of these events then gives the average rate of material loss. In a similar way, in an alternating motion tribometer with a dead time between strokes the average rate of material loss is the sum of the losses during each stroke. 2. Repassivation Rate Let us assume that a hard pin slides over a passive metal surface damaging the passive film. As the pin has passed repassivation occurs. According to Faradays law the rate of film growth dL/dt (cm/s) is proportional to the anodic growth current density, igr (A/cm2): dL/dt = kF igr . Here kF = M/rnF, where M and r are, respectively, the atomic mass and the density of the metal and n is the charge number of the anodic reaction. The functional dependence of dL/dt, and hence of igr, on potential and on film thickness depends on the film growth mechanism. Film growth models have been studied by many authors and a recent discussion can be found in reference [1]. Two limiting cases can be identified. High field conduction models (HFM) assume that the transport of ionic species in the oxide limits the growth rate. At constant potential the electric field diminishes with increasing oxide thickness and so does the growth rate. Interface models (IFM) assume that the electrochemical reactions at either the metal-oxide or the oxide-solution interface limit the growth rate. The electric field in the oxide is constant and therefore the potential difference at the interface, and hence the growth rate, diminishes with increasing film thickness.

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Both models predict a similar decrease of growth current with anodizing time and with increasing film thickness and for this reason they can not readily be distinguished from current transient measurements [2]. In terms of predicting repassivation rates under tribocorrosion conditions both models therefore yield comparable results. Jemmely et al [3] simulated repassivation transients in tribocorrosion experiments where an aluminum oxide pin rubbed on a stainless steel surface in an alternating motion tribometer with a dead time between strokes. The metal was kept at constant potential in the passive potential region. The authors assumed high field film growth and they took into account the ohmic resistance between reference electrode and wear scar. All current was assumed to serve for film growth neglecting dissolution, which would reduce the film growth rate by about 20-30%. A more important uncertainty resulted from the unknown real contact area, which differs from the geometrical area. More recently, Wu and Celis applied a high field conduction model to the study of current transients observed under open circuit conditions where the electrode potential varies during an experiment [4]. Olsson and Stemp [5] simulated tribocorrosion current transients for Fe-17Cr alloy in 0.5M H2SO4 assuming interface controlled growth (IFM) and taking into account film dissolution. For anodic film growth on stainless steel in sulfuric acid the measured anodic current, Ia, includes the film growth current and a dissolution current that results from selective iron dissolution during film growth [6]. The growth current can be expressed as Igr = Rg Ia where Rg is the growth fraction, a quantity that can be experimentally determined by means of an EQCM or by solution analysis [7,8]. The described repassivation models were developed for pin-on-plate experiments assuming that the surface depassivated by the sliding pin repassivates once the pin has passed. Similar models have previously been proposed for single scratch experiments by a number of authors [9-11]. Recent experiments by Stemp et al [12] suggest that for typical tribocorrosion conditions in sliding contacts depassivation rather than repassivation is the limiting step. These authors studied the effect of the contact arrangement on tribocorrosion transients at constant potential for Fe-17Cr in 0.5M H2SO4 in an alternating motion tribometer with a dead time between strokes. They observed very similar current transients and metal loss rates (taking into account slight differences in wear scar area) whether a metal pin was sliding on a ceramic disk or whether a ceramic pin was sliding on a metal plate (Fig. 1). This result is surprising because the physical situation is quite different: when a ceramic pin slides on a metal plate the metal surface has time to repassivate when the pin moves on, but when a metal pin slides on a ceramic plate the rubbing metal surface is continuously in contact with the counter body. The fact that on the

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metal pin/ceramic plate contact repassivation is observed during the dead time suggests that surface roughness or slight misalignments permit electrolyte access the metal surface in the contact permitting thus passive film formation. The results cast some doubts on tribocorrosion models that consider repassivation of the surface to occur after the pin has passed. They suggest that the anodic current results from electrochemical reactions that take place in the contact zone itself. One possibility is that equal rates of mechanical film thinning and anodic film growth determine the measured current, but it is more likely that during rubbing microscopic asperity contacts depassivate and repassivate in a statistical way. The rate of depassivation therefore plays a crucial role in tribocorrosion.

Fig. 1 Three contact arrangements and corresponding current transients measured on anodically polarized Fe-17Cr stainless steel in alternating motion tribometer with a dead time between strokes (after ref. [12])

3. Film Damage Mechanisms Four distinct wear mechanisms may lead to film damage in sliding contacts: adhesion, abrasion, ploughing and fatigue. Which mechanism predominates depends on the properties of the contacting materials, their surface topography and on the operational parameters of the contact. Adhesion is of particular importance in metal-metal contacts. Extensive plastic deformation of the contacting metal surfaces cause passive film damage and

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Passivity Issues in Tribocorrosion

permits local metal-to-metal contacts to be formed, that can lead to seizure. The worn surfaces therefore exhibit the typical aspect of local seizures. Abrasion occurs when a hard body slides over a softer metal. The hard body may be a ceramic pin or sphere often used in tribocorrosion experiments or it may be an abrasive powder. Worn surfaces usually exhibit scratches in sliding direction [12]. Mild abrasion can in principle lead to film thinning, but because in most cases the roughness of the contacting surfaces by far exceeds the passive film thickness, a more likely scenario is that the passive oxide film is destroyed at the contacting asperities and then reforms. Ploughing means strong plastic deformation and displacement of a ductile material as a result of scratching with a hard body. The scratch test is a classical example. The plastic deformation of the substrate can lead to passive film damage by film thinning and film rupture [13]. Fatigue wear is mostly observed with brittle materials and results in formation of cracks that start from the surface or form underneath the surface in parallel direction [14,15]. Crack propagation eventually leads to the detachment of a metal particle that subsequently must be eliminated from the contact. Release of a particle from a passive metal locally exposes the naked metal surface to the corrosive environment. Rapid repassivation and/or active dissolution will occur on both the damaged metal surface and the released particle [16].

90o impact 50o impact

Current from the wear trace / µA

5 4 3 2 1 0

0

1

2

3

4

5

6

7

8

Energy / J

Fig. 2 Threshold Energy for Passive Film Rupture by Particle Impingement measured with silica sand on 304 stainless steel in 0.6M NaCl by Sasaki and Burstein [20]

Fig. 3 Effect of dissipated mechanical energy on anodic current during tribocorrosion of anodically polarized (0.5V vs Ag/AgCl) Ti6Al4V in 0.9% NaCl [22]

Similar damage mechanisms are to be expected in erosion corrosion due to particle impingement. Sharp particles impinging at oblique angle may lead to abrasion [17]. Impingement of particles without sharp edges may lead to plastic deformation and material displacement similar as in ploughing, especially at relatively high impact angles with respect to the surface [18]. Finally, repetitive impacts can lead to cold hardening and fatigue damage. According to Tan et al [19] cutting is another mechanism. To produce surface damage, the kinetic energy of impinging particles must exceed a certain critical energy as is

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illustrated by the data of Fig. 2 from Sasaki and Burstein [20]. Recently it has been proposed to use indentation measurements to determine the critical energy [19]. Barril et al [21,22] recently studied the effect of dissipated energy on anodic currents measured on Ti6Al4V in 9% NaCl under fretting corrosion conditions, varying the amplitude and the applied normal force (Fig. 3). The wear volume measured under anodic polarization were found to increase steadily as a function of dissipated energy, while under cathodic polarization the wear volume did not change [23]. In Fig. 3 the anodic current resulting from passive film damage increases in a similar way with dissipated energy. At amplitudes below about 10 micrometer, only elastic deformation took place that did not lead to a measurable anodic current. 4. Depassivation Rate For rubbing surfaces mechanical wear theory predicts that in the absence of chemical effects the material loss is given by Vmech = kWj vs FN/H, where kWj is a wear coefficient the value of which depends on the wear mechanism, vs is the sliding speed, FN is the applied normal force and H is the hardness of the material. One may expect that the rate of passive film damage follows a similar law. Landolt et al [24] studied the effect of normal force for rubbing contacts between a hard and a ductile material of different relative surface roughness. When a smooth hard pin slides over the rough surface of a ductile metal the contact between the two bodies takes place at asperities and the real contact area varies with the applied normal force and the hardness of the metal: Ar = S Aj = nj Aj,av = FN / H Here Aj is the area of an asperity contact, nj is the number of asperity contacts and Aj,av is the average area of asperity contacts. The surface of the depassivated area is given by Rdep = k' Rn Aj,av , where Rn is the rate of renewal of asperity contacts during sliding which is proportional to the sliding speed vs. The proportionality factor k' represents the probability that a contact actually leads to depassivation. The value of Rn depends on the surface topography of the two contacting bodies as has been discussed elsewhere in detail [24]. One finds thus that the rate of generation of depassivated area, Rdep, is proportional to sliding speed and to (FN/H)b where 0.5 < b < 1 depending on the surface topography of the contacting bodies. At steady state the average rate of metal oxidation in the contact is equal to the average rate of generation of depassivated surface multiplied by the charge needed for repassivation: Ia = Rdep Qp The measured anodic current therefore should be proportional to sliding speed and vary with applied normal force and hardness as predicted by the above equations. Experimental results obtained with stainless steel in sulfuric acid agreed well with this prediction [25].

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483

In principle the repassivation charge for a given metal-environment combination can be measured in an independent experiment performed with a depassivated metal surface. Such surfaces have been prepared by breaking [26], diamond scratching [11], guillotining [27] or cathodic reduction [28] among other. Such measurements are delicate, but in tribocorrosion there is a further complication. Indeed, complete repassivation of a depassivated microscopic surface element requires that the time between two subsequent depassivation events is larger than the time to reform the passive film. This may not always be the case. For example, if a pin moves at a sliding speed of 5 mm/s and the distance between asperity contacts is 50 micrometer the time between depassivation events would be on the order of ten millisecond. Typical time constants observed for the current decay between strokes in Fig. 1 are on the order of one hundred milliseconds but passive film growth will continue for even longer times. During rubbing the microscopic depassivated surface areas therefore will not necessarily repassivate completely before being depassivated again. In spite of this uncertainty, repassivation charge may be a useful criterion for comparing the tribocorrosion rate of different metals and alloys. Mischler et al. found a good correlation between the repassivation charge determined electrochemically and the tribocorrosion rate for a variety of metals [28]. 5. Third Body Effects Most physical modeling in the literature concerns two body contacts. It is well known, however, that mechanical wear produces solid wear products that remain in the contact for a certain period before being ejected or dissolved. In a tribological contact these particles form a third body [29,30]. It has been found for steel that the presence or not of a passive film had a strong effect on third body formation and mechanical wear [31]. The presence of a third body changes the prevailing mechanical conditions and hence the wear rate: hard particles tend to be abrasive, increasing wear, while soft particles tend to act more like a lubricant, reducing the wear rate. Since under tribocorrosion conditions third body particles are most often oxidized, the physical properties of the relevant metal oxides are expected to play an critical role. For the understanding of tribocorrosion mechanisms it is important to know whether wear particles are released from the rubbing metal surface in oxidized or metallic form. The anodic current measured in a tribocorrosion experiment includes the oxidation of the first body as well as of wear particles. Current measurements alone therefore do not permit to identify the nature of the wear particles released from the first body. A detailed analysis of these effects and of the different material fluxes in tribocorrosion contacts involving a third body has recently been presented elsewhere [16].

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6. Galvanic Interactions and Potential Gradients Two types of tribocorrosion experiments are typically found in the literature: those performed at constant applied potential and those performed at open circuit. For given conditions the tribocorrosion rate was found to be the same whether the potential was imposed electrochemically with a potentiostat or by the corrosive environment [16, 32]. In the first case the electrons liberated by the anodic reaction in the wear scar flow to a counter electrode, in the second case they react with an oxidizing agent present. The latter reaction could take place at the active surface in the wear scar or, more likely, at the surrounding passive surface leading to formation of a corrosion cell. Besides electrolyte conductivity the electronic properties of the passive film and the geometrical arrangement of the contact should play a crucial role, but no systematic studies are available at present.

Fig. 4 Galvanic effects due to tribocorrosion: rubbing of Ti6Al4V and nitrided Ti6Al4V (N1,N2) against alumina ball in 0.9M NaCl. (a): Potential shift, (b): Galvanic current between rubbed surface and identical non rubbed surface. After reference [34].

At open circuit, rubbing of a passive metal leads to a potential shift in the negative direction, due to the damaging of the passive film [33-35]. As an example, Fig. 4 shows the potential change induced by rubbing on Ti6Al4V alloys with different surface treatment in aerated 0.9M NaCl [34]. Also shown is the galvanic current between the rubbed surface and a passive surface of the same material connected through a zero resistance ampere meter. The current flow results from a galvanic cell established between the activated wear scar and the surrounding passive surface. The indicated current values are qualitative only, because part of the cathodic reaction may have taken place on the passive surface of the rubbed sample adjacent to the wear scar rather than on the counter electrode. The actual current in the wear scar therefore is expected to be larger than the measured current. Recently, Wu and Celis reported similar current transients during rubbing of stainless steel under open circuit conditions [4]. To measure the galvanic currents a platinum electrode was connected to the working electrode through a zero resistance ampere meter forming a galvanic

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Passivity Issues in Tribocorrosion

cell. Not surprisingly, the measured current was found to depend on the size of the platinum electrode. In principle, galvanic effects could also play a role during potentiostatic tribocorrosion experiments, if part of the current is drawn to the surrounding passive film rather than to the counter electrode. In such a case the measured current would represent only a part of the total anodic current in the wear scar. However, up to this date no evidence for such a mechanism has been found. Galvanic and capacitive effects in repassivation experiments with bare metal surfaces have been discussed by Bastek et al [11] who found them to be unimportant in well conducting electrolytes. In poorly conducting media a broadening of current transients and a slight diminishing of their height was observed. Current, mA

0.14 0.12

0.05 M H2SO4

0.10 0.08 0.06 0.04 0.02 0.00 0.0

0.2

0.4

0.6

0.8

1.0

0.6

0.8

1.0

50

Current, mA

40

0.005 M H2SO4

30 20 10 0 0.0

0.2

0.4

Time, s

Fig. 5 Current transients measured on Fe-17Cr pin sliding on alumina plate in alternating motion with a dead time between strokes. Applied potential 0.25V MSE. After reference [36]

Several authors showed that repassivation transients in the absence of rubbing are governed to a large extent by ohmic effects [3,9,11]. Therefore, as the current decreases during film growth the potential at the surface changes even if the applied potential is kept constant. In tribocorrosion experiments ohmic effects are even more important because of the unfavorable geometry of the wear zone which resembles that of a crevice. A potential drop is expected to exist from the outer boundary to the interior. Under some conditions the potential in the interior of the wear zone might even be lower than the passivation potential. In such a case, instead of repassivation, active dissolution dominates. Stemp et al [25] studied current transients for a stainless steel pin sliding in alternate motion on an alumina plate in sulfuric acid of different concentration (Fig.5). The transients in 0.005M H2SO4 suggest that the sample

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did not repassivate during the off-time. Photographs of surfaces after experiments suggested crevice corrosion like phenomena resulting from a potential gradient in the contact zone [37]. At this time no quantitative data on the potential distribution in the contact zone and its role for the tribocorrosion rate of passive metals are available. 7. Concluding Remarks Tribocorrosion presents many challenging passivity issues. For several systems physical models have been proposed and experimentally tested and the mechanical and electrochemical factors that determine passive films damage and repair have been identified. Many aspects require further study. They include a better understanding of the electrochemical conditions at the rubbing surfaces in relation to the prevailing potential distribution, more quantitative information on how passivation charge and critical current density affect the tribocorrosion rate, and experimentally confirmed models on how the mechanical and electrical properties of oxide films influence galvanic cell formation during rubbing. The role of passive oxide films in the formation of third body particles and the transformation and dynamic behavior of these are other topics for future research. While passivity issues are the key for understanding tribocorrosion behavior of metals, they should always be considered in context of the tribocorrosion system as a whole, including contact design, mechanical stability, operating conditions, materials properties of work piece and antagonist and the corrosivity of the environment. Acknowledgements The author wishes to acknowledge the major contributions to the field of tribocorrosion by S. Mischler , P. Jemmely, M. Stemp and S. Barril. References 1. C.-A.O. Olsson, D. Hamm, D. Landolt, J. Electrochem. Soc. 147 (2000) 4093 2. C.-A.O. Olsson, M.-G. Vergé, D. Landolt, J. Electrochem. Soc. 151 (2004) B652 3. P. Jemmely, S. Mischler, D. Landolt, Wear 237 (2000) 63 4. P.Q. Wu, J.-P. Celis, J. Electrochem. Soc. 151 (2004) B551 5. C.-A.O. Olsson, M. Stemp, Electrochim. Acta 49 (2004) 2145 6. P. Schmutz, D. Landolt, Corr. Sci. 41 (1999) 2143 7. D. Hamm, K. Ogle, S. Weber, C-O.A. Olsson, D. Landolt, Corr. Sci. 44 (2002) 1443 8. D. Hamm, C-O.A. Olsson, D. Landolt, Corr. Sci. 44/5 (2002) 1009 9. G.T. Burstein, A.J. Davenport, J. Electrochem. Soc.136 (1989) 936

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10. F.J. Graham, H.C. Brookes, J.W. Bayles, J.Appl.Electrochem. 20 (1990) 45 11. P. D. Bastek, R. C. Newmann, R. G. Kelly, J. Electrochem. Soc. 140 (1993) 1884 12. M. Stemp, S. Mischler, D. Landolt, Corr. Sci. 45 (2003) 625 13. M. Chiba, M. Seo, J. Electrochem. Soc. 150 (2003) B525 14. S.C. Lim, M.F. Ashby, 35 (1987) 1 15. J. Jiang, M.M. Stack, A. Neville, Tribology Intl. 35 (2002) 669 16. D. Landolt, S. Mischler, M. Stemp, S. Barril, Wear 256 (2004) 517 17. R.E. Winter, I.M. Hutchings, Wear 29 (1974) 181 18. I.M. Hutchings, R.E. Winter, Wear 27 (1974) 121 19. K.S. Tan, A.W. Hassel, M. Stratmann, Mat.-wiss.u.Werkstofftech. 36 (2005) 13 20. K. Sasaki, G.T. Burstein, Phil.Mag.Lett. 80 (2000) 489 21. S. Barril, S. Mischler, D. Landolt, Wear 256 (2004) 963 22. S. Barril, Ph.D. thesis No 2751, EPFL Lausanne 2003 23. S. Barril, S. Mischler, D. Landolt, Wear 259 (2005) 282 24. D. Landolt, M. Stemp, S. Mischler, Electrochim. Acta 46 (2001) 3913 25. M. Stemp, S. Mischler, D. Landolt, Wear 255 (2003) 466 26. G.S. Frankel, C.V. Jahnes, V. Brusic, A.J. Davenport, J. Electrochem. Soc. 142 (1995) 2290 27. G.T. Burstein, K. Sasaki, J. Electrochem. Soc. 148 (2001) B282 28. S. Mischler, S. Debaud, D. Landolt, J. Electrochem. Soc. 145 (1998) 750 29. M. Godet, Wear 100 (1984) 437 30. M. Godet, Wear 136 (1990) 29 31. S. Mischler, A. Spiegel, M. Stemp, D. Landolt, Wear 251(2001) 1295 32. S. Barril, S. Mischler, D. Landolt, Tribol. Intl. 34 (2001) 599 33. N. Latona, P. Fetherston, A. Chen, K. Sridharan, R.A. Dodd, Corrosion 57 (2001) 884 34. F. Galliano, E. Galvanetto, S. Mischler, D. Landolt, Surf. Coat.Technol. 145 (2001) 121 35. P.Q. Wu, J.-P. Celis, Wear 256 (2004) 480 36. M. Stemp, Ph.D. thesis No 2292, EPFL Lausanne 2001 37. M. Stemp, S. Mischler, D. Landolt, Tribology Research: From Model Experiment to Industrial Problem, G. Dalmaz et al. Editors, Elsevier Science B.V. (2001 ) p. 539

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Wear in nuclear power plants, a tribocorrosion approach A. Beaudouin* , P. Combrade* , D. Kaczorowski* , J-P. Vernot* * FRAMATOME-ANP Technical Center, Porte Magenta 71205 Le Creusot, France Corresponding author : [email protected]

Abstract - Wear of sliding 304L stainless steel couples in PWR primary water is studied to understand the environment effect in the degradation process and quantify the corrosion wear process with mechanical and electrochemical measurements. Keywords : Tribocorrosion, Wear, PWR environment, Stainless steel

1. Introduction In pressurized water reactors (PWR), some components are submitted to relative motions due to necessary operational processes or to unwanted effects such as flow induced vibration. Complex combinations of wear can occur [1], [2], [3] from a large range of contact kinematics (sliding, fretting, impact [4]) and from PWR environment (temperature (290-325°C), pressure (154 bars), water chemistry (deaerated, low conductivity water)) resulting in specific corrosion phenomenon [5], [6]. The combined effect of mechanical removal from the contact and wear accelerated corrosion on a wear process, is called tribocorrosion [7]. Wear accelerated corrosion arises from the fact that an asperity rubbing on a metal surface produces a bare surface track which is usually more sensitive to corrosion [8]. It is usually assumed that on passive material the anodic current variation due to the oxidation of mechanically damaged surfaces occurs during a limited time period required for the surface film to recover its protectiveness. Accordingly, the so called "film repair transient" is believed to be a controlling process of tribocorrosion. Indeed experimental results [9] tend to show that, in PWR environment, the wear rate of Stellite material sliding on stainless steel is likely to be controlled by such a film repair process. The influence of environment in 304L stainless steel wear is examined and electrochemical measurements were performed during wear tests

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to quantify the effect of oxidation processes. Film repair kinetics were measured by tensile tests under controlled potential [17]. 2. Experimental 2.1. Tribometer - AURORE bench Understanding of wear mechanisms in nuclear power plant requires the development of specific "parametric tribometers" [10], reproducing the contact geometry, the high temperature and pressure water environment and the displacements. The FRAMATOME-ANP “AURORE” test bench, was originally designed to realize tests in primary PWR coolant conditions [11] (control rod against support plate, fuel rod against fuel assembly grids…). The test bench is housed in an autoclave working at temperatures up to 320°C +/5°C and pressure about 154 bars. Fig.1 presents a sketch of the tribometer. The contact partners consist of two coaxial cylinders. The stationary ring (height 3 mm, internal diameter 10.7 mm), is mounted on a flexible support of rigidity K. The moving tube (length 5 mm, diameter 8.7 mm and thickness 0.95 mm) is placed around a mast, clamped at its ends. Four electro-magnets, create a controlled motion of the tube. The deflection of the ring support, measured with load sensors, gives the normal and tangential forces. Eddy current sensors are used to measure the differential displacement of the tube. The entire system is computer controlled. Contact kinematics is an important Fig.1: Schematic cross section of the parameter in wear phenomenon [12], thus tribometer (upper view) the AURORE apparatus allows the use of four main kinematics: pure sliding, fretting, impact and impact plus sliding. During sliding, the continuous contact between tube and ring results from revolutions described by the tube’s axis. Sliding plus impacts is generated by adding a higher frequency disturbing the orbital sliding. It leads to a succession of oblique impacts with sliding. 2.2. Electrochemical measurements during wear The tribometer is equipped to perform electrochemical measurements (Fig. 1b). The rubbing electrodes are insulated but when both samples are metallic, they are electrically connected to

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each other. An Ag/AgCl electrode is used as reference. According to MacDonald [13], its potential (at 300°C) is 20mV lower than Standard Hydrogen Electrode. Platinum disk is used as an auxiliary reference electrode in the tests involving potential control. The counter electrode also is a platinum disk. Electrical leads are insulated with PTFE, allowing a maximum test temperature about 300° C. The tests under controlled potential presented in this paper were performed at a potential slightly above the corrosion potential (Ecorr+30 mV) in order to minimize the cathodic reduction of H+ contribution to the measured current. 2.3. Experiments Water chemistry

Mass variation

(mg) Experiments were performed on a tube (ppm) pH at 300°C tube ring versus ring sliding couple made of 304L. [B] : 1200, The test environment is a fully deaerated [Li] : 2 (A) -1.7 -0.8 primary water solution containing boric 6.9 [B] : 1200, acid and lithium hydroxide. Boric acid (B) (B) [Li] : 0 -13.3 -4.2 takes part in the control of the nuclear 5.3 reaction and lithium oxide (Li) regulates Table 1- Conditions and results of the tests the pH. Table 1 indicates the contents of performed with two pHs solutions. both species for the tests environments (dissolved O2 < 5 µg/kg). described in this paper. The applied kinematics are reactors representative movements known from scale 1 simulators and components observations.

3. Results 3.1. Surface oxides in HT water In Pressurized High Temperature Water (PHTW), the oxide layers formed on 304L are much thicker than passive films grown at ambient temperatures [5], [6]. After 1400 hours in primary water at 300°C the tube surface is oxidized and a cross section of the duplex layer structure obtained is shown in Fig. 3. It appears that the oxide grains, which apparent hardness is between 10 and 30 Gpa, may cleave during the indentation loading [15]. External layer

Figure 3 - Cross section of the oxidized tube in PHTW. There are two oxide layers. The external layer is made of magnetite Fe3O4 crystals, (100 to 1500 nm thick). The internal layer is a rich chromium spinel mainly formed by solid state growth process (50 to 300 nm thick).

External layer Internal layer Internal layer

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3.2. Wear results It has already been proved that the wear damage morphology observed on control rods in service can be reproduced in laboratory tests by a combination of sliding and impacts [10], in water at 300°C but not at 80°C [15]. In PWR (at 300°C), the pH environment effect on the wear rate was investigated in the sliding plus impact configuration. The reference test “A” was performed in the nominal primary environment (pH 6.9 at 300° C). The wear scars observed on the tube are made of elongated indents, with a typical length of 100 µm and a width about 20µm (Fig. 4 and 5). Lowering the pH to 5.3 causes the wear rate increase (5 to 7 times the test “A” wear rate). However, the scars’ morphology was very similar to the test “A” one. Fig. 4 - Wear scars on the tube surface after test. Fig. 5 – Section of one scar.

100µm

20µm

20µm

3.3. Electrochemical measurement during wear test The pH effect on the electrochemical behavior of rubbing 304L was investigated by using wear tests under controlled potential. The tests scheme was either 600s continuous friction or 60s friction sequences taken in turns with 180 or 600s waiting sequences where the two rubbing surfaces were separated. Variations in anodic current during tests are given by Fig.6 and Fig.7.

50

pH 7

60

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pH 5 Current (µA)

Current variation (µA)

pH 7 40

30

20

40

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0 0

1000

2000

3000

4000

5000

6000

7000

Time (s)

0 0

500

1000

1500

2000

2500

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Time (s)

Fig. 6 - Current variations for 2 pH solutions during continuous friction. (pHs at 300°C)

Fig. 7 - Current variations during 5 sequences of 60s of friction + 600s of unloading

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As soon as the friction starts, there is a quick current increase which does not aim to stabilize within the 60s of friction. As soon as separating the samples, the current starts to decrease. At the lowest pH, the increase of anodic current is lower than at pH 6.9 but the decay towards the initial value is slower. This behavior is typical of the protective film degradation and repair: during friction the partial film repair is responsible for the progressive stabilization of the anodic current and when there is no more friction it is responsible for the current decay. Fig. 6 and 7 suggest that the corrosion rate is lower at pH 5.3 than at pH 6.9. The metal volume dissolved by corrosion during friction is calculated by applying Faraday’s law to the increase of anodic current, as shown in table 2. Dh (nm/cm²) pH6.9 60+600 60+180 29 16

Dh (nm/cm²) pH5.3 60+600 60+180 24 13

Table 2 – Calculation of the metal loss due to corrosion during the five sequences of friction and unloading shown in Fig. 7.

It shows that the pH influence on the wear rate is not the result of corrosion phenomena, and that wear is mainly due to the particular properties of oxide in PHTW (the hardness of friction films). 4. Conclusion. The effect of environment on wear rate was clearly demonstrated in the PWR primary environment: ̶ Temperature plays a major role: representative wear is observed only at high temperature. This can be related to the oxide layers nature: in high temperature water, the oxide layers are much thicker and made of very hard magnetite crystals of micrometer size. Those could make up very abrasive friction films when separated from the stainless steel surface. ̶ The solution pH appears to have a large effect on the wear rate but this effect is not correlated with a dissolution (or oxidation) process. The pH has opposite effect on wear rate and oxidation processes. Thus the weight loss during wear of 304L Stainless Steel in primary PWR water is not dominated by an oxidation process at least at the lower pHs. ̶ The pH effect on the oxidation rate can be mainly considered as an effect of the solution conductivity, which is always low, but much lower at pH 5.3 that at pH 6.9. In the pH 6.9 solution, bearing about 2 ppm of Li ions, the upper conductivity range allows higher electrochemical reaction kinetics. ̶ If not directly due to film repair transients, the effect of pH on the wear rate of 304L SS in PWR primary water must be due to a mechanical parameter. A comparison of the pH effect on the anodic current during tensile and wear tests show that the local solicitation inside the contact is less severe at pH 6.9 than at

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pH 5.3 [17]. According to the tensile tests, if the surface damages created by rubbing were identical at both pHs, the increase of anodic currents due to rubbing should be more different when decreasing the pH from 6.9 to 5.3. This means that the same external solicitation causes a lower surface damage at pH 6.9 than at pH 5.3. Two causes can be invoked to explain this difference: - a difference in the number of abrasive particles produced by the surface oxidation, but this is unlikely due to the lower oxidation current at pH 5.3. - a difference in the electrostatic forces due to the electrical duplex layers at the oxide/solution interface that modify the local pressure in the contact area. ̶ Of course more work is required to ratify the above analysis which currently relies on a very limited set of results. A work including a parametric study of the effect of environment on the wear damages is presently in progress. It aims at detailed analysis of the impact energy, characterization of the damaged surfaces, and if possible, measurement surface charges on 304 L in PWR primary water. ̶ A study to develop modeling of the electrochemistry cell and then of depassivation and repassivation currents is also going on. 5. References [1] R. Cauvin, Proceedings Fontevraud 4, (1998), pp 1289-1299 [2] P.L Ko, M.C Taponat, M. Zbinden, ASME publ PVP, vol 328, Flow induced vibration, (1996), pp 211-218 [3] S. Hogmark, A. Öberg, B. Stridh, Proceedings of the J.S.L.E International Tribology Conference in Tokyo, (1985), pp 723-728 [4] P.J Hoffman, T. Schettler, D.A Steininger, ASME, Pressure Vessel and Piping Conference in Chicago, PVP-2, (1986) [5] B. Stellwag, Corrosion Science, vol 40, (2-3), (1998), 337-3710 [6] DH. Lister, RD. Davidson, E.Mc Alpine, Corrosion science, vol 27, n°2, (1987), pp113140 [7] Mischler et al, J. Electrochem. Soc.,Wear, Vol 145, n°3, (1998), pp750-758 [8] S. Mischler, Wear, vol 251, (2001), pp1295-1307 [9] M. LeCalvar, E. Lemaire, Wear, 8656, (2000), pp 1-7 [10] Kaczorowski et al., Proceedings of the International Symposium on Fretting Fatigue 4, Ecole Centrale de Lyon, (2004), to be published. [11] J.Ph Vernot, D. Kaczorowski, C. Crenn, Proceedings of Eurocorr, Nice, France, (2004). [12] P.L Ko, Wear, vol 106, (1985), pp 211-281 [13] Mac Donald, Journal of Electrochemical Society, vol 6, (1979), p126 [14] R.W. Bosch, B. Schepers, M. Vankeerberghen, Electrochemical Methods in Corrosion Research, Nieuwpoort, Belgium, Electrochimica acta, vol 49, (2004) pp 3029-3038 [15] D. Kaczorowski, Wear of a stainless steel in high temperature and high pressure water, Ecole Centrale de Lyon, PhD Thesis, (2002). [16] P.L Ko, M.C Taponat, M. Zbinden, ASME, PVP-vol 298, Flow induced vibration, (1995) pp 99-110. [17] P. Combrade, P. Scott, Corrosion Deformations Interactions, The institute of Materials, London, (1996) pp 394.

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Comparison between tribocorrosion mechanisms of Stellite 6 and Zircaloy 4 in LiOH-H3BO3 solutions Viorel-Eugen Iordache1, François Wenger1, Pierre Ponthiaux1, Antoine Ambard2, Jean Peybernès3, Joëlle Vallory3. 1

LGPM, Ecole Centrale Paris, 92295 Châtenay-Malabry Cedex, France Corresponding author : F. Wenger ; [email protected] 2 EDF - R&D, Les Renardières. 77818 Moret sur Loing, France 3 DEN/CAD-DTN, CEA Cadarache 13108 Saint-Paul-lez-Durance cedex, France

Abstract - Tribocorrosion mechanisms of Stellite 6 and Zircaloy 4 alloys were studied by means of tribocorrosion tests carried out in conditions of unidirectional sliding friction in a solution of boric acid and lithium hydroxide, with same chemical composition as water in nuclear Pressurized Water Reactors. With Stellite 6, a "soft wear" process is obtained, controlled by mechanical removal and subsequent restoration of a very thin passive film. With Zircaloy 4, in the same tribological conditions, a major influence of abrasion by zirconia particles, pulled out from the zirconia surface layer, was observed. These major differences in the tribocorrosion mechanisms explain the differences found in wear laws. Keywords : tribocorrosion, wear, PWR, Stellite 6, Zircaloy 4

1. Introduction Many component parts of nuclear pressurized water reactors (PWR) are submitted to the combined actions of friction and corrosion (tribocorrosion). Passivating alloys with high mechanical properties were used for the manufacture of such parts. The gripper latch arms (GLAs) of the control rod drive mechanisms, made of stainless steel, are protected from wear and corrosion by a thick layer of Stellite 6, a hardfacing and passivating cobalt-base alloy, with 28 wt. % chromium as main alloying element. Based on field experiences, Lemaire and Le Calvar [1] have found that the wear loss W

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increases with the number N of contacts (impact + friction) between GLA and control rod, and also, for a given value of N, with the average latency time t between two successive steps (or contacts), following the empirical law : W = Nw0 (t / t0)1-n

(1)

where w0 and t0 are constants and where n was found to be around 0.65. This law was interpreted by the occurrence of a tribocorrosion mechanism consisting in a periodical destruction by friction of the passivating Cr2O3 film grown during t. The fuel rods of the fuel assemblies are made of Zircaloy 4 (zirconium-base alloy), a hard and passivating material. In few cases, friction of the rod against the spacer grids, can lead to perforations of the cladding, difficult to predict on the basis of simple wear laws. The aim of this study is therefore to analyze, in laboratory conditions, the tribocorrosion mechanisms of Stellite 6 and Zircaloy 4. First, wear-latency time dependence was examined and then, the relative contributions of the mechanical and electrochemical components of wear in the overall tribocorrosion process were investigated. 2. Experimental conditions Tribocorrosion experiments were performed by using a pin-on-disc tribometer fully described elsewhere [2, 3], and permitting to carry out tests in corrosive aqueous media under controlled electrochemical conditions. The working surface was the disc-shaped upper cross section of a cylindrical specimen (25 mm in diameter and 25 mm high), installed in an electrolytic cell mounted on the tribometer. The counterbody was an inert cylinder-shaped corundum pin, drawing a circular wear track of 16 mm diameter on the disc. The pin had a spherical end of 100 mm radius (sphere-on-plane contact). The normal forces used in the tests were chosen in order to induce average static contact pressures of 120 Mpa (actual pressure between the GLAs and the command rods), 60 MPa, or 30 MPa. 30 Mpa is close to the order of magnitude of the contact pressure between fuel rods and spacer grids. For both alloys, these values are smaller than the yield strength. All tests were performed at 120 rotations per minute (rpm) and stopped after 10000 laps. Tests under intermittent friction conditions were carried out by applying 2500 steps of friction at 120 rpm during 4 laps (duration = 2 s), two successive steps being separated by motionless steps (latency time t = 20 or 200 s). The Stellite 6 specimens were prepared by plasma spraying deposition of a 1 mm thick layer on the cross section of a cylinder of stainless steel. The Zircaloy 4 specimens were machined from a rod, 26 mm in diameter. Before the tests, the working surfaces were polished with abrasive paper up to grade 1200, and then rinsed with water and alcohol. At the end of the tests, the worn volumes were estimated by two independent methods: weight loss measurements (0.2 mg accuracy), and 3-D

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microtopographic surveys of the wear track with a STIL® High-Resolution Optical Measurements Station (1 µm lateral and 30 nm vertical resolutions). For Stellite 6, the tests were conducted at 85°C in a pH 7.7 solution of H3BO3 (1000 ppm of B) and LiOH (12 ppm of Li). For Zircaloy 4, the tests were conducted at room temperature in a similar solution (1000 ppm of B and 130 ppm of Li, pH = 8.7). The concentrations of B and Li were chosen to obtain the same pH / pKe ratio in the solutions as with the water of PWR, ensuring therefore similar chemical properties. A three-electrode arrangement was used for the electrochemical measurements. The disc-shaped end of the specimens acted as working electrode and a circular platinized titanium grid was used as counter-electrode. An Ag-AgCl reference electrode (Ag/AgCl/saturated KCl solution, +360mV/NHE at 85°C) was chosen for the measurements at 85°C on Stellite 6, and a SSE (Hg/Hg2SO4/saturated K2SO4 solution, +670mV/NHE) was chosen for the room temperature measurements on Zircaloy 4. Electrochemical measurements were carried out with a PAR 273A potentiostat controlled by a PC through the CorrWare® software. 3. Electrochemical behavior of the alloys With no applied friction, the potentiodynamic polarization curves of the alloys given in Fig 1a show that both alloys are passivating materials in the aqueous medium used in this study, as in PWR environment.

Figure 1. Potentiodynamic polarization curves at 100mV/min scan rate, scanning from cathodic to anodic range, for Stellite 6 and Zircaloy 4, under no friction conditions (a) and under continuous friction at 120 MPa initial average pressure and 120 rpm rotation speed (b).

The Stellite 6 alloy shows a wide passivation plateau (0.8 V width; passivation current in the order of 2 µA/cm-2), starting immediately after the zero-current potential, limited at the cathodic end by hydrogen evolution and by transpassive dissolution at the anodic end. Potentiodynamic curve of Zircaloy 4 are similar to those of Stellite 6: a wide ( > 1 V) passivation plateau is observed, and no active dissolution peak is observed around the zero current potential (Fig. 1a). On the contrary, when friction is applied, a continuous increase of the anodic current

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with the potential is observed on both curves (Figs. 1b); this indicates that fast oxidation of the bare metal occurs in the wear track, on the areas where the passive film was removed by friction. To note also that in the potential domain where passivation occurs, most of the anodic current measured during friction is flowing from the depassivated areas of the wear track (out of the wear track, the alloy remains in a passive state). 4. Wear laws – intermittent friction tests The relationship linking up the total wear Wt to the latency time t was determined by using intermittent friction tests in open-circuit potential (OCP) conditions. Continuous friction tests at 120 rpm were considered as intermittent friction tests of 0.5 s latency period (periodicity of the contact with the pin for every place of the wear track). Each test was repeated at least three times and the average values of Wt versus latency time t are presented in Fig. 2. It is worth noticing that, for both materials, the difference between the wear values estimated by weighing and by topographic surveys was less than 5 %. It means that corrosion did not develop out of the wear track, and that the behaviour of the specimens is only determined by the tribocorrosion process occurring in the wear track.

Figure 2. Total wear Wt versus latency time t (intermittent friction tests). The continuous lines (for Stellite 6) and the dashed lines (for Zircaloy 4) represent linear regressions. Average error values of 10% for Stellite 6 and 60 % for Zircaloy 4 were found.

For Stellite 6, a well-defined linear relationship was found in log-log scale between the total wear Wt and t. The parameter n (given by the slope of the line) has the same value of (0.70 ± 0.05) for both values of the contact pressure. This indicates that the kinetic law of the oxidation reaction involved in the repassivation of the surface after friction is the same for both contact pressures. Moreover, to note that the value n is close to the one of 0.65 declared for the wear of the GLAs in PWR environment. For Zircaloy 4, Wt is much higher than for Stellite 6, and almost independent of t, indicating that the tribocorrosion

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mechanisms of the alloys are significantly different. Complementary investigations were carried out to explain these differences. 5. Complementary investigations Micrographs in Fig. 3 show the typical aspects of the wear track after a tribocorrosion test in the same conditions, with Stellite 6 and Zircaloy 4. For the latter, the influence of abrasive wear in the wear process seems much stronger. This interpretation is confirmed by the observation of the solution after the tests: after a test with Stellite 6, the solution was limpid and no wear debris were found in the cell, whereas the solution was cloudy after the tests with Zircaloy 4, and many particles were found on the bottom of the cell. 200 µm

50 µm

a)

b)

Figure 3. Aspect of the wear track after tribocorrosion tests (intermittent friction with t = 20 s; 10000 laps; 120 rpm; 120 MPa) in OCP conditions, for Stellite 6 (a) and Zircaloy 4 (b).

In tribocorrosion tests, the total wear Wt can be expressed as the sum of We (electrochemical component of wear) and Wm (mechanical component) : Wt = We + Wm whereWe corresponds the amount of oxide formed during the latency times and removed by mechanical action, whereas Wm corresponds to the amount of bare metal removed by friction. To evaluate the order of magnitude of We for Stellite 6, tests in OCP conditions and continuous or intermittent friction in distilled water (taken for its weak oxidation properties) allowed us to show that Wm had always a minor contribution in Wt: Wm < Wt / 5, and that Wm/Wt decreased as t increased. With Zircaloy 4, the contribution of Wm in Wt seems of major importance as shown by Fig. 4: continuous friction tests were performed under various values of the applied potential. The anodic current was recorded and integrated on the duration of the test. By applying the Faraday's law, the mass of oxidized metal WF was calculated, giving an assessment of We. Contrary to the case of Stellite 6, WF << Wt at all applied potentials. In addition, an increase of Wt with the potential is found which can not be explained by the increase of WF (or We): that means that Wm increases with the potential.

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V.-E. Iordache et al.

Figure 4. Zircaloy 4: total wear Wt and electrochemical wear WF (or We) versus potential E after friction tests under potentiostatic conditions. Continuous lines are only for guidance.

To explain the differences between the mechanisms of the alloys, it seems necessary to consider first the differences between structures and mechanical properties of the passive films of Cr2O3 and ZrO2, the latter being much thicker hard and brittle than the chromium oxide one. Obviously, the mechanical properties of the alloys must also be considered. Surprisingly, in the case of Zircaloy 4, the mechanical component of wear Wm depends on the potential. The electrochemical conditions could perhaps influence the production, the structure and the properties of the abrasive particles. These complex aspects of the tribocorrosion mechanism of Zircaloy 4 will be the subject of our next research work. 6. Conclusions The results of this study revealed major differences between the tribocorrosion mechanisms of Stellite 6 and Zircaloy 4 in the H3BO3-LiOH solution. In the first case, the wear process could be explained mainly by the removal of the thin passive film of Cr2O3 on the contact areas and following restoration of this film through electrochemical oxidation of the bare metal areas, with a minor contribution of the pure mechanical component of wear affecting the bare metal. On the contrary, in the case of Zircaloy 4, the oxidation process leading to the restoration of the damaged surface film of ZrO2 seemed to have a negligible weight in the total wear. An intense abrasive process was observed probably due to the production of numerous hard zirconia particles (wear debris) in the wear track, some of them being trapped in the contact. References [1] E. Lemaire, M. Le Calvar, Wear, 249 (2001) 338. [2] L. Benea, P. Ponthiaux, F. Wenger, J. Galland, D. Hertz, J.Y. Malo, Wear, 256 (2004) 948. [3] F. Wenger, P. Ponthiaux, L. Benea, J. Peybernès, Proceedings of EUROCORR 2004, France.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

501

Detrimental effect of lead on the passivity of UNS N06690 alloy B.T. Lua, J.L. Luo a,* and Y.C. Lub a

Department of Chemical and Materials Engineering, University of Alberta, Edmonton, Alberta, Canada T6G 2G6 [email protected] b Atomic Energy of Canada Ltd., Chalk River Laboratories, Chalk River, Ontario, Canada K0J 1J0

Abstract - The pitting induction times under the potentiostatic conditions indicate that the detrimental impact of lead on the passivity degradation is markedly enhanced by calcium ions present in alkaline crevice solutions. The mechanism is discussed based on the results of X-ray photoelectron spectroscopy (XPS) analyses, photoelectrochemical and Mott-Schottky measurements. Keywords: lead-induced corrosion, steam generator tube alloy, passive film

1. Introduction The Pb-induced stress corrosion cracking (SCC) of steam generator (SG) tubes is considered to be related to the degradation of the passive films. 1 It was reported that the presence of lead could incorporate into the passive films2 and reduce the passive film thickness,3 resulting in a significant deleterious impact in the passive region. 4 The evidence has indicated that the other corrosive species coexisting may affect the Pb-induced corrosion1 but the interaction mechanisms are still unknown. In this study, an attempt is made to gain some insight into the detrimental effects of lead contamination on the passivity of the SG tubing alloy in simulated SG crevice chemistries. A special attention is paid to the role of interaction between lead contamination and other species.

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2. Experimental Procedures The nominal composition of UNS N06690 alloy in weight percent is 0.02C, 30Cr, 10Fe, 0.5Mn, 0.5Si and bal. Ni. The ion concentrations in test solutions are listed in Table 1, where No. 1 and No. 4 were derived from the simulated CANDUÒ SG crevice chemistries. CANDU SG water chemistry management maintains the SG bulk water at pH 9.5. The impurities in SG water are concentrated in the SG crevices because of boiling, resulting in alkaline condition.5 No. 2 and No. 3 were the modified versions of No. 4 to investigate the roles of Ca2+ and SO42- in the passivity breakdown. The pitting induction times were measured at 40°C under a constant potential of 0.45 VSCE (vs. the Saturated Calomel Electrode) in order to initiate pits in reasonable test durations. The Mott-Schottky measurements were performed using the potential swept from the open circuit potential in the positive direction with a scanning rate of 0.01 V/step. In accordance with Ferrieria et al.,6 an ac signal of 3,160 Hz with the peak-to-peak magnitude of 10 mV was selected superimposed on the scanning potential. In the photoelectrochemical experiments, the light with a 150 W Xenon arc lamp with the wavelength range of 200-900 nm was provided by a scanning digital monochromator controlled at a scan rate of 0.8 nm/s and a lock-in amplifier was used to increase signal-to-noise ratio. The photocurrent was defined as the difference in the currents under the illuminated and dark conditions. The changes in compositions of the passive films prepassivated in the test solutions at 40°C under a constant potential 0.45VSCE for 30 min were determined by XPS analyses. Table 1. Concentrations (mM) of various species in the test solutions and pH* ID No.

PbO

Na+

K+

Cl-

SO42-

Ca2-

pH27oC

pH40oC

1 2 3

0 2.2 2.2

812 718 699

50.7 50.0 50.0

659 650 650

50.9 50.0 0

1.83 0 0.5

12.7 12.7 12.7

12.2 12.3 12.2

4

2.2

812

50.7

659

50.9

1.83

12.7

12.2

*Calculated with ChemSolv

3. Results and Discussion 3.1. Pitting Induction Time Pits could initiate in less than 20 min when specimens were immersed in the alkaline crevice solution containing PbO (Solution No. 4) while no local attack Ò

CANadian Deuterium Uranium is a registered trademark of Atomic Energy of Canada Limited.

Detrimental effect of lead on passivity of UNS N06690 alloy

503

was observed in the Pb-free solution (Solution No. 1) after immersed in the test solution for more than 12 h, indicating a detrimental effect of lead contamination on the pitting resistance. Further experiments showed that, when the specimens were polarized at the same potential and temperature in the leadcontaminated solutions without calcium ions (Solution No. 2), the minimum incubation time for pit initiation was longer than 12 h, implying that the impact of SO42- on the passive film breakdown is insignificant. By reducing the concentration of the calcium ions from 1.83 mM (Solution No. 4) to 0.5 mM (Solution No. 3), the minimum pitting incubation time increased from 15 min to about 120 min, suggesting that an increase in the calcium ion concentration enhanced the detrimental effect of lead on the passivity of the alloy. Since calcium ions may concentrate in the SG-crevices, the synergism of the lead contamination and calcium ions may play a role in the breakdown of the passive films in the lead-induced SCC processes. The main difference between the compositions of Solutions No. 2 and No. 3 are that the former contains SO42but is Ca2+-free, and the latter is free of SO42- but contains Ca2+. The results of the pitting induction times in these two solutions suggest that, under the present alkaline chemistry conditions, the presence of calcium ions is necessary for Pbinduced passive film breakdown. 3.2. Compositions of Passive Films

The composition-depth profiles for passive films were analyzed by XPS and the results for the specimens prepassivated in Solution No. 1 and No. 4 are shown in Figure 1. It is clearly that the passive film formed in a lead-contaminated alkaline solution containing calcium is more complicated than that formed in lead-free solution. The outer layer with a thickness of several tens micron comprises mainly PbO. It should be noted that a detectable amount of lead, in agreement of the previous reports,2,3 and calcium as well, can be found in the inner layer of the film (Figure 1b), whereas, without lead contamination, calcium ions could hardly enter in the passive film (Figure 1a). The lead contamination also reduces the chromium concentration in the passive film. These results imply that the detrimental effect of lead contamination on the stability of the passive films may be exacerbated by certain synergisms between Pb and Ca ions. When lead contamination is absent, calcium ions in the solution do not incorporate into the passive film and will not affect the passivity of the alloy. The synergism of Pb and Ca ions in the passivity degradation has been well demonstrated by the pitting induction times.

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50 Ni2p

Fe2p

Cr2p

O1s

Ca2p

Pb4f

60 Ni2p

O1s

Cr2p

Pb4f

40

Atomic percentage

Atomic percentage

Fe2p Ca2p

50

40

30

20

30

20

10

10 0

0

20

40

60

80

100

Etch depth from surface (nm)

(a) Prepassivated in Solution No.1

120

0 0

20

40

60

80 100 120 140 160 180 200 220 240 260 280 300

Etch depth from surface (nm)

(b) Prepassivated in Solution No.4

Figure 1. Composition-depth profiles of prepassivated UNS N06690 alloy 3.3. Capacitance and Photoelectrochemical Response of Passive Films Figure 2 shows the Mott-Schottky plots for passive films formed in solution No 1 and No. 4. The negative slopes of the plots indicate that the passive films are p-type semiconductors. The slope of linear section of Mott-Schottky plot is given by -2/(ii0qNA),8 where e is the dielectric constant of the passive film, e 0 the vacuum permittivity, q the electron charge and NA the acceptor density. Figure 2 shows that the absolute value of the slope measured for the passive film formed in the lead-contaminated solution is about 25 times as large as that obtained for the passive film formed in the lead-free solution. This must result from the changes in the acceptor density and the dielectric constant of the passive films. Because the dielectric constant of PbO (~26) is about twice as large as those of other oxides (NiO: ~11, Cr2O3: ~12, FeO: ~14, CaO: ~11),7 it is expected that the presence of lead oxide in the passive film will not result in a decrease in the dielectric constant. As a result, the passive film formed in the lead-contaminated solution probably has a lower acceptor density than that formed in the lead-free solution. The shape of the photocurrent spectra obtained from the passive films formed in the lead-free and lead-containing solutions are similar. A detectable photocurrent first appears at incident photo energy of approximately 1.8 eV, subsequently rises with decreasing light wavelength and peaks at around 2.87 eV. The peak photocurrent density of the passive film formed in lead-contaminated solution is about 20 times greater than that formed in the lead-free solution. Photocurrents can be detected in the incident light energy range h <Eg, which is known as the “Urbach tail”. Eg is the band-gap energy between the conductive band (CB) and valence band (VB). The appearance of the Urbach tail indicates that both passive films are highly defective, or nonstoichiometry.8 For a heavily doped semiconductor, local electronic states

Detrimental effect of lead on passivity of UNS N06690 alloy

505

are introduced between the CB and VB, which makes the passive films more defective and reduces the photon energy needed to excite the charge carriers in the semiconductor. The values of Eg can be determined from ( hv )0.5 vs. h plots in Figure 3 using the Gartner-Butler equation. 8 As indicated by Figure 3, the band-gap energy of the passive film formed in the lead-contaminated solution (No.4) is slightly lower than that of the film formed in lead-free solution (No.1). However, the change in the band-gap energy is not significant.

-2

0.5

(h.hv)

4 -1

(F m V )

4 -2

2

2/(ee 0qNA )

-2

-0.5

500

80

Prepassivated in Solution W ith Pb Pb-free 2.35

0.025

3

1000

C (F m )

0.030

Prepassivated in solution W ith Pb Pb-free

1500

1

0.020 0.015

Eg (eV)

2000

0.010

60

2.25

40

0.005

0 No.1 No.4 Solution

20 0 -0.4

2.3

-0.2

0.0

0.2

0.4

0.6

0.8

E (V SCE)

Figure 2. Mott-Schottky Plots

0.000 1.0

No.1 No.4 Solution

1.5

2.0

2.5

hv (eV)

Figure 3. Plots of (hhv)0.5 vs. hv

The XPS analysis has indicated that lead ions in the solution may enter the passive film and form lead oxide. The photoelectrochemical tests have indicated that the PbO is a p-type semiconductor with relative higher photoactivity. 9 Therefore, the deposit of PbO in the outer layer may contribute the stronger photocurrent. In line with the generalized Gartner-Butler equation,8 the photocurrent density at a given wavelength is proportional to [ii0/(qNA)]0.5[St/(St+Sr)], where St is the minority carrier’s transfer reaction rate and Sr the total recombination rate. The Mott-Schottky measurements indicated that the prepassivation in the solution with lead-contamination would result in a larger dielectric constant and a lower acceptor density. This change will enhance the photocurrent because it increases [ii0/(qNA)]0.5. In a p-type semiconductor, holes (h) are the major charge carriers and electrons (e) are minor ones. It is reasonable to assume that concentrations of holes, [h], and electrons, [e], are proportional to the densities of acceptor, NA and donor, ND, respectively. Obviously, St = kt’[e] = ktND, where kt and k t' are the experimental constants. Note that the recombination reaction results in the annihilation of a pair of hole and electron. Therefore, the total recombination rate Sr = kr’[h][e] = krNAND, where kr and kr’ stand for the reaction constants for the recombination. Hence, we have St/(St+Sr) = kt/(kt+krNA). It suggests that a decrease in the acceptor density will result in an increase in the photocurrent. Pb2+ is likely to replace the Cr3+ in the oxide films by occupying the vacancy of Cr3+ through the 2+ defect reaction: Pb 2 + + VCr''' 3+ + h + ⇒ PbCr 3+ , which will consume hole and

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reduce Cr-content in the passive films. This deduction agrees with the experimental and analytical results given above. It has been well documented that a decrease in the Cr-content reduces the stability of passive films. 4. Summary The coexistence of lead and calcium reduce the pitting induction time significantly, indicating that the detrimental effect of lead contamination on the passivity of UNS N06690 alloy is enhanced by the presence of calcium ions in solution. Calcium ions may enter the oxide layer with the aid of lead contamination during the corrosion process, resulting in certain changes in both composition and the structure of the passive films. Lead-contamination also causes certain changes in the electronic properties of the passive films, which may contribute to their low stability. Acknowledgements This work was supported by Atomic Energy of Canada Ltd. Life Management Technology program Manager, R. L. Tapping is acknowledged for supporting this work. The authors would like to thank C. W. Turner and B. W. Hocking for reviewing this paper and provided many helpful suggestions. G. Strati’s assistance on calculating the test solutions using ChemSolv code is appreciated. References 1 R. W. Staehle and J. A. Gorman, Corrosion, 2004, 60, 115-180. 2 J. B. Lumsden and P. J. Stocker, in Proc. 4th Intern. Symp. Environmental Degradation of Materials in Nuclear Power Systems, - Water Reactors, Jeklly Island, GA, 1989, pp.6-38~6-51. 3 M. Garcia-Mazario, A. M. Hernandez and C. Maffiotte, Nuclear Engineering and Design, 1996, 167:155-167. 4 H. Radhakrishnan, R. C. Newman and A. Carcea, Presentation in Corrosion 2005, Houston, TX, USA, April 2-7, 2005. 5 M. D. Wright, G. Goscynski and F. Peca, 7th Intern. Symp. Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Vol. 1, 1998, La Grange Park, Illinois, 11. 6 M.F. Montemor, M.G.S. Ferreira, N.E. Hakiki, M. D. C. Belo, Corrosion Sci., 2000, 42:16351650. 7 A. B. Kunz, J. Phys. 1981, C14: L455. 8 F. Di Quarto and M. Santamaria, Corr. Eng. Sci. Tech., 2004, 39,71-81. 9 P. Veluchamy and H. Minoura, J. Electroanal. Chem., 1997, 440, 41-49.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

507

The role of Borides in the Passivity and SCC of Alloy 600 Yongsun Yi, Hongpyo Kim, Joungsoo Kim Korea Atomic Energy Research Institute, 150 Deogjin-dong, Yuseoung-gu, Daejeon, 305-353, Korea, [email protected]

Abstract - The effects of an addition of borides, NiB and CeB6, to a NaOH solution as the reference solution on the passivity and stress corrosion cracking (SCC) of Alloy 600 was investigated. Polarization curves of the alloy were measured in the solutions at 315 °C. The addition of the borides to the reference solution resulted in a decrease in the peak current and the active potential range, implying that passive films could be formed more easily in caustic solutions containing the borides. Samples of the alloy were exposed to the solutions for 120 hours and the surfaces of the samples were examined using Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS). The results showed that much thinner passive films were formed on the samples exposed to the solutions containing the borides than that in the reference solution. Using slow strain rate tests (SSRT) the SCC susceptibility of the alloy was investigated in the same solutions. The SSRT test results showed that the borides suppressed the SCC of the alloy in the caustic solution. It was concluded that the borides enhanced the passivation of the alloy in the solution, resulting in a decreased SCC susceptibility. Keywords: Alloy 600, passivity, SCC, boride, caustic solution

1. Introduction Alloy 600 had been used as steam generator tubing materials in many nuclear power plants constructed in the late 1960s and early 1970s. However, experience with Alloy 600 in the early pressurized water reactors (PWRs) revealed numerous corrosion problems [1]. The causes of a steam generator tube plugging have evolved with time and in recent years, outer-diameter stress

508

Y. Yi et al.

corrosion cracking (ODSCC) and intergranular attack (IGA) have become the predominant identifiable causes that lead to a plugging of the steam generator tubes [2]. The ODSCC occurring in the secondary side of steam generators is known to be explainable by the metallic dissolution/oxidation [3] which is directly related to the oxide properties. In this paper the effects of borides as candidate inhibitors on the passivity and SCC behavior of Alloy 600 are discussed. 2. Experimental 2.1. Material and test solutions Archived YGN 3&4 steam generator tubing, Alloy 600, with an nominal outer diameter of 19.05 mm and a thickness of 1.07 mm was used. Chemical compositions of the tubing used are given in Table 1. The reference solution was a dearated 40 % NaOH aqueous solution. To investigate the effects of borides on the electrochemical and SCC behaviours of Alloy 600, two kinds of boride, NiB and CeB6 (2g/l), respectively, were added to the reference solution. All the tests were performed at 315°C. 2.2. Electrochamical and SCC tests 2.2.1. Potentiodynamic polarization For the polarization curve measurements, coupons of 5 mm x 10 mm x 1 mm were machined from the above S/G tubes. The coupon surface was polished up to 0.1 m using alumina paste and it was then spot-welded to an Alloy 600 wire. The polarization curves were measured in a 1 liter volume static autoclave made of Ni using the Model 273 Potentiostat / Galvanostat (EG&G PAR). An external Ag/AgCl and the autoclave itself were used as the reference and counter electrodes, respectively. First, at -1.5 V vs. the open circuit potential (OCP), the samples were cathodically treated for 30 min to remove the oxides on the surface and then polarized from -1.5 V to 1.5 V vs. the OCP at a scan rate of 1 mV/sec. In the SSRT tests and the polarization measurements, all the solutions were deaerated by a bubbling with pure nitrogen gas for 24 hours. Table 1 Chemical compositions of the tubing used in this study C

Ni

0.026 72.4

Cr

Fe

16.81 9.01

Ti 0.36

Al

Mn

0.16 0.83

Si

Co

Cu

N

B

P

S

0.33 0.01

0.01

0.018

0.001

0.007

0.001

The role of Borides in the Passivity and SCC of Alloy 600

509

2.2.2. Surface oxide examination Samples were exposed to the solutions for 5 days and the surface oxides formed on the samples were examined using AES and XPS. AES depth profiles were measured with a PHI 680 Auger nano-probe operated at a primary beam energy of 5kV and an electron current of 15 ~ 20 nA. Depth profiling of the surface oxides was performed using 1-4 keV argon ions. Surface oxides were also analyzed by an XPS using a PHI 5800 ESCA System. The system background pressure was approximately 3 x 10-10 mbar. An AlKg (1486.6 eV) radiation generated by using an anode source operating at 15 kV and an emission current source of 24 mA was used in all the measurements. The signal from the adventitious carbon (284.6 eV) was used for calibrating the XPS data. The analyzed spot size was 400 m x 400 m. 2.2.3. SCC tests To investigate the effects of the borides for the reference solution on the SCC behavior of Alloy 600, SSRT tests with tensile samples were performed. Tensile samples with a gage length of 25.4 mm, a gage width of 4 mm, and a thickness of 1.07 mm were used. SSRT tests were conducted in a static autoclave made of Alloy 625. Before straining the samples, the solution was continuously purged with pure nitrogen gas for 24 hours. The tests in the aqueous solutions were performed at a strain rate of 1 x 10-6 s-1. Surface appearance in the gage section and that of the fracture surfaces after the SSRT tests were examined with SEM. 3. Results and Discussion 3.1. Polarization curves Fig. 1 shows the polarization curves of the Alloy 600 obtained at 315°C in 40% NaOH solutions with and without borides. By the addition of the borides, the active peak current density decreased and the active potential range became narrower, implying that a passive film could form more easily in the boride containing solutions. 3.2. Surface oxide analysis Fig. 2 shows the AES depth profiles of the samples exposed to three kinds of solutions for 120 h. It is known that Cr is depleted at the outer layer of oxide of Alloy 600 formed in caustic solutions [4,5] In this study, Cr was not depleted in the reference solution [See Fig. 2(a)]. McIntyre et al. [6] examined the effects of dissolved oxygen (DO) on the surface compositions of Alloy 600 in caustic solutions containing LiOH. According to their results, the surface oxide is

510

Y. Yi et al. 1.0 Potential (V vs. OCP)

40% NaOH + CeB6

0.5

40% NaOH + CeB6 40% NaOH + NiB

0

-0.5 10-5

10-4

10-3

10-1

10-2

Logarithmic current density

1

(A/cm2)

Fig. 1 Polarization curves of Alloy 600 obtained in 315°C 40% NaOH solutions with and without borides.

abundant in Cr at a DO of 5 ppb while Cr is depleted when the DO is higher than 150 ppb. In the boride-containing solutions, the concentration of the Cr decreases at the most outer surface of the oxides. The most distinct difference in the oxides between the reference and boride containing solutions appeared in the thickness of the surface oxides. The thickness of the surface oxide formed in the reference solution was about 1 m, but that of the oxides in the boridecontaining solutions was extremely thin, only a few nm.

O

Atomic conc. (%)

190 Å/min (SiO2)

Cr

Fe

35 Å/min (SiO2)

Ni

100

100

80

80

80

60

60

60

40

40

40

20

20

20

0

0 0

20

40

60

Sputtering time (min) a) Deaerated 40% NaOH at 315 C

43 Å/min (SiO2)

100

0 0

2

4

6

8

10

Sputtering time (min) b) Deaerated 40% NaOH + 2g/l CeB6 at 315 C

0

2

4

6

8

10

Sputtering time (min) c) Deaerated 40% NaOH + 2g/l NiB at 315 C

Fig. 2. AES depth profiles of samples exposed for 5days to (a) 40% NaOH solution (reference solution), (b) CeB6, and (c) NiB containing solutions at 315°C.

511

XPS intensity (Arbitrary unit)

The role of Borides in the Passivity and SCC of Alloy 600 O 1s

Cr2O3

Cr 2p3/2

Cr2O3

Ni 2p3/2

Ni(OH)2

Ni(OH)2

Ref. Ni

CeB6 NiB 537

533

529

590

582

865

574

857

849

XPS intensity (Arbitrary unit)

Binding energy (eV) (a) O 1s

Cr 2p3/2

Ni 2p3/2 Ni

537

533

529

590

582

574

865

857

849

Binding energy (eV) (b) Fig. 3. Narrow scan XPS depth spectra of samples exposed for 5days to 40% NaOH solution (reference solution), CeB6, and NiB containing solutions at 315°C (a) before cleaning and (b) after cleaning.

(a)

(b)

(c)

Tensile direction

The XPS spectra of O 1s, Cr 2p, and Ni 2p for the samples exposed to the caustic solutions for 5 days are shown in Fig. 3. The spectra obtained for the as-received surfaces are shown in Fig. 3(a). Fig. 3(b) shows the spectra after a cleaning which was done by an ion sputtering by about 2 nm. The peak of Cr 2p2/3 around 576.8 eV corresponds to the chromium oxide, Cr2O3. In the spectra of Ni 2p2/3 two peaks before a cleaning are observed which correspond to the nickel hydroxide, Ni(OH)2, and metal nickel, but after a cleaning the peak of the metal nickel becomes dominant in the three conditions (Fig. 3). Although a big difference in the oxide thickness was observed, no discernable differences were found in the XPS results.

Fig. 4. SEM micrographs of the gage lengths of the samples tested in (a) 40% NaOH solution (reference solution), (b) CeB6, and (c) NiB containing solutions at 315°C.

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3.3. SCC tests Fig. 4 shows the SEM micrographs of the samples strained in the reference and boride containing solutions. No cracks were observed for the samples tested in the boride containing solutions while a severe IGSCC occurred in the sample tested in the reference solution. The decrease in the oxide thickness in Fig. 2 and the inhibition of the SCC of the same alloy were also observed for Alloy 600 in 315°C ammonia solutions (pH = ~9.6) by the addition of the borides [7]. This suggests that the borides could be applicable to solutions with different pH values. Concerning the inhibition mechanism, the cracking of Alloy 600 occurring in the secondary side is known to be associated with metallic dissolution/oxidation [8], which is explained by a slip dissolution model. This could draw a speculation that no IGSCC occurred when a sample was tested in the solutions containing borides since the corrosion rate was greatly suppressed by the additives. 4. Summary The polarization measurements and surface oxide analysis showed that the borides enhanced the passivation of Alloy 600 in the caustic solution, which resulted in an inhibition of the IGSCC of the alloy References [1] R.W. Staehle and J.A. Gorman, in Proceedings of the 10th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, (2001). [2] D.R. Diercks, W.J. Shack, and J. Muscara, Nuclear Engineering and Design, 194 (1999), pp. 19-30. [3] P.M. Scott, In proc. of the 9th Int. Symp. on Environmental Degradation of Materials in Nuclear Power systems – Water Reactors, The Metallurgical Society, Warrendale, PA (1997), pp. 3-13. [4] M.F. Montemor, M.G.S. Ferreira, M. Walls, B. Rondot, and M. Cunha Belo, Corrosion, Vol. 59, No. 1 (2003), pp.11-21. [5] U.C. Kim, K.M. Kim, J.S. Kang, E.H. Lee, and H.P. Kim, Journal of Nuclear Materials 302 (2002) 104-108. [6] N.S. McIntyre, D.G. Zetaruk, and D. Owen, Journal of Electrochemical Society, Vol. 126, No. 5, (1979), pp. 750-760. [7] J. Kim, H. Kim, Y. Yi, SMiRT 18 (Beijing, China, August 7-12, 2005), to be presented. [8] P.M. Scott, P. Combrade, Proc. 8th Int. Symp. on Env. Deg. of Mat’ls in Nuclear Power Systems-Water Reactors (La Grange, IL: American Nuclear Society [ANS], 1997), p. 65.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

513

Effect of Nickel and Tungsten on Repassivation Rate of Stainless Steels in Chloride Solution by Potential Step Chronoamperometry Sung-Yu Kima*, Heesan Kimb, Hyuk-Sang Kwonc a,b

School of Materials Sci. & Eng. Hongik University 300 Sinan-Ri, Chochiwon-Up, Youngi-Gun, Chungnam, 339-701, KOREA c Dept. of Materials Sci. & Eng. KAIST 373-1 Guseong-Dong, Yuseong-Gu, Daejeon, 305-701, KOREA

Abstract - The recent development of the method for analyzing repassivation rate by a rapid scratch test and high field conduction model makes it possible to predict SCC susceptibility more quantitatively but it impossible to distinguish the repassivation rates with a little difference due to the inaccurate measurement of scratched surface area. From the scratch test and the potential step chrono- amperometry (PSC) test, the function of nickel and tungsten affecting repassivation rate is explained as blocking the passivation by metal not participating in the formation of passive film and stabilizing the film by molybdenate (MoO42-) with accelerating the dissolution rate of iron by tungsten, respectively. In addition, a PSC test can be approved by a technique for simulating repassivation rate by the comparison of the results from two tests. The details in test conditions are mentioned in the paper. Keyword : repassivation rate, high field conduction model, tungsten, nickel, ferritic stainless steel

1. Introduction SCC phenomena of materials having passive film such as stainless steels or nickel alloys have been explained well by slip dissolution model where SCC susceptibility depends on both repassivation rate and dissolution rate of materials at sites where passive film is broken [1]. Hence, the SCC susceptibility has been evaluated by comparisons of repassivation rates using

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S. Y. Kim et al.

various methods [2-4] under an environment meaning that the dissolution rate of bare materials is constant [5]. The repassivation kinetics of stainless steels conforms to the place exchange model [6] at the initial stage of repassivation ( less than 10 milliseconds after the film rupture), but, at the next stage ( less than 100 milliseconds after the film breaks down), the high field conduction model [7] which is known to affect SCC susceptibility [8,9]. According to the high field conduction model, current density (i(t)) across a metal/film interface is inversely proportional to charge density (q(t)). log i (t ) = log A +

zF 1 1 BV = log A + cBV 2.3Mf q(t ) q(t )

(1)

,where A and B are parameters related to energy barrier, z is the charge number of mobile ion, F is Faraday’s constant, is the density of a passive film, M is a molecular weight, f is a ratio of charge density due to forming the passive film to total charge density and c is a constant(=zF /2.3M f). However, the rapid scratch test still has difficulties in evaluating cBV with respect to reproducing a scratched area sensitively affecting cBV. In this study, passive film is removed by electrochemical method [10] instead of mechanical method and the repassivation rate was measured from the log i(t) - 1/q(t) plot. Hence, the goal of this study is to get a better understanding of the effects of the alloying elements (W, Ni) on the repassivation rates in both rapid scratch test and the electrochemical method and to prove the validity of the electrochemical method.

In order to examine the effect of alloying elements (W, Ni) on the repassivation rate by a scratch test and a potential step chronoamperometry(PSC) test, three experimental alloys with the basic composition of 30%Cr and 4%Mo (Table 1) and nickel (purity 99.9%, NILACO corp.) were used. The specimens with cross section of 2 mm wide and 8 mm were soldered copper-lead wire, were cold-mounted and were polished with diamond paste to 3 µm finish. The deaerated chloride solutions (0.8M HCl+ 3.2M LiCl (pH 0), 4M LiCl (pH 6.8)) were used as test environments.

0.8M HCl + 3.2M LiCl at 25oC 30Cr-4Mo 30Cr-3MO-2W 0.2 30Cr-3Mo-2W02Ni Ni(pH=0.8) 0.0 Ni(pH=6.8) 0.4

SCE ) Potential (V

2. Experimental

-0.2 -0.4 10-7

10-5

10-3

101 2

Current density (A/cm )

Figure 1 Effects of W and Ni on the anodic polarization behaviors of stainless steels in deaerated 25Ł 4M-chloride solution

515

Effect of Nickel and Tungsten on Repassivation Rate of Stainless Steels Table 1 Chemical composition of experimental alloys No 1 2 3

Cr 30.1 29.9 29.7

Mo 4.0 3.2 3.2

C .008 .011 .015

N .01 .01 .01

Ti .05 .06 .06

Nb .26 .26 .26

Ni NA NA 2.15

W NA 2.15 2.18

P 0.03 0.03 0.03

Si 0.33 0.3 0.29

Mn 0.3 0.3 0.3

Al 0.05 S 0.002 O 0.006

For the polarization test, the specimen was potentiodynamically polarized from the potential more active to the open circuit potential by 50mV at the scan rate of 20mV/min after the corrosion potentials had been stabilized. For the rapid scratch test, the repassivation currents were measured every one millisecond during and after the passive film was removed by a rapid scratch method [8,9], to calculate cBV value from the log i(t) - 1/q(t) plot. Prior to scratching, the potential of the specimen was held 200mVSCE for 500 seconds to to form a stable passive film. A PSC test was performed with the cell and the specimens used in the polarization test. The potential of the specimen was held 200mVSCE for 500 seconds after the corrosion potentials had been stabilized (first stage) and stepped to the active potentials (-287mVSCE, -400mVSCE, in Fig. 1) for three different time (second stage); one second after potential transition, eight seconds after observation of oxidation current, two hundred seconds which is a minimum time required for a stable oxidation current after the dissolution of passive film. After the potential was stepped back to 200mVSCE (third stage), current flow and repassivation rate were measured like the scratch test. XPS analysis was performed using ESCA LAB 2000VG Scientific Co. equipped with a mono-chromate Al anode (Al-K : 1486.6eV) in order to examine compositions on the surfaces exposed to the acidic chloride solution for 500 seconds at various potentials (-287mVSCE, -400mVSCE). 3. Results and Discussion Fig. 2 and Fig. 3 show that cBV was decreased by tungsten but increased by nickel. However, the sensitivity of alloy elements depended on pH as listed in Table 2. The reasons are as follows: The function of tungsten promoting repassivation rate by accelerating the dissolution rate of iron [11,12] does not work effectively in the acidic solution due to high dissolution rate of iron in the acidic solution. In the meanwhile, nickel, not participating in the formation of passive film, puts blocks on the formation of passive film to decrease repassivation rate [13]. Table 2 shows that the sensitivity of dependency of the alloying elements to cBV on after potential transition from an active potential to a passive potential in the PSC test. The cBV’s measured after transition from -287mVSCE provide that the repassivation rate is almost independent of the alloying elements. In the mean while, the cBV’s measured after transition from -400 mVSCE, show that the repassivation rate from the PSC test is almost independent of tungsten but is

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S. Y. Kim et al.

more sensitive to nickel than that from scratch test. In the PSC test, the inconsistent sensitivity of the alloying elements to cBV is due to the difference of metallic compositions/components of the surface placed at an active potential prior to transition to a passive potential as shown in XPS analysis (Fig. 4 and Fig. 5). Like to the scratch test, the insensitivity of repassivation to tungsten in two PSC tests is not due to effectively playing a role of promoting the dissolution rate of iron as mentioned in previously. The inconsistent dependency of nickel on the active potential is due to whether the enrichment of nickel on the surface occurred or not. In other words, nickel is enriched on the surface after exposure at –400mVSCE (Fig. 5), lower than the corrosion potential of nickel but is not after exposure at –287mVSCE (Fig. 4). 3.21x10-4 1.67x10-4

10-1

2

100

101

1.00x10-4

10-2

4M LiCl at 25oC 30Cr-4Mo 30Cr-3Mo-2W 30Cr-3Mo-2W-2Ni

10-3 10-4

5.0x102

1.0x103

1.5x103

2.0x103

Current density (A/cm )

2

Current density (A/cm)

101

5.40x10-3

100

5.12x10-3 4.61x10-3

10-1

0.8M HCl + 3.2M LiCl at 25oC 30Cr-4Mo 30Cr-3Mo-2W 30Cr-3Mo-2W-2Ni

10-2

10-3

0

100

Charge density-1 (cm2/c)

200

300

Charge density-1 (cm2/c)

Figure 2 Log i(t)-vs.-1/q(t) plots of ferritic stainless steels after scratch test at 200 mVSCE in deaerated 25Ł 4M LiCl solution

Figure 3 Log i(t)-vs.-1/q(t) plots of ferritic stainless steels after scratch test at 200 mVSCE in deaerated 25Ł 0.8M HCl + 3.2M LiCl solution

Table 2 cBV values of ferritic stainless steels measured by scratch test in the deaerated 25Ł 4M chloride solutions and PSC test after the potential was held to the active potential for 200 seconds in 25Ł 0.8M HCl + 3.2M LiCl solution Solution containing 4M chloride Relative cBV

Active Potential (mVSCE) : PSC test

4M LiCl

0.8M HCl + 3.2M LiCl

-287

-400

cBV of 30Cr – 3Mo – 2W cBV of 30Cr – 4Mo

0.59

0.90

0.95

0.89

cBV of 30Cr – 3Mo – 2W – 2 Ni cBV of 30Cr – 4Mo

1.91

1.05

0.89

1.94

Hence, the pre–exposure at –400mVSCE for 200 seconds provides test condition of the PSC test to simulate the repassivation rate in the acidic solution depending on alloying elements (W, Ni). However, with this condition, the repassivation rate of stainless steels in neutral pH is not simulated by PSC test. It may due to the difference of repassivation rate between acidic pH and neutral pH reaction though they are not listed in Table 2. The other PSC test was

517

Effect of Nickel and Tungsten on Repassivation Rate of Stainless Steels

Atomic Concentration (%)

Atomic Concentration (%)

conducted for the specimen with the passive film, whose thickness was controlled by the time spent at active potential (-287mVSCE) after passivation. The PSC test after 1-second exposure at –287mVSCE, provides that log i(t) 1/q(t) plot did not depend on alloying elements (Fig. 6). In the mean while, the PSC test after eight-second exposure (Fig. 7) gives not only that log i(t) - t plots is affected by the alloying elements but also that the dependency of cBV on the alloying elements is similar to the dependency of cBV from the scratch test in the neutral pH solution rather than in the acidic solution(Table 2). Unlike onesecond exposure, the similarity in the dependency on alloying element between two tests is thought to be explained as follows: The eight-second exposure is not enough to remove the passive film, whose presence is indicated by the current 60 Deaerated 0.8HCl + 3.2M LiCl solu. at Room temp. 50

At -287 mVSCE for 500 sec.

CrMe FeMe WMe MoMe NiMe

40 30 20

CrOx FeOx WOx MoOx NiOx

10 0 0

10

20

30

40

60 Deaerated 0.8HCl + 3.2M LiCl solu. at Room temp. 50

At -400 mVSCE for 500 sec.

CrMe FeMe WMe MoMe NiMe

40 30 20 10 0 0

50

2

10-3

30Cr-4Mo 30Cr-3Mo-2W 30Cr-3Mo-2W-2Ni

-5

10

10-6

o

0.8M HCl + 3.2M LiCl at 25 C stripping cond.: -287mVSCE for 1 s 4.0x104

8.0x104

1.2x105 -1

20

30

40

50

Figure 5 XPS analysis of ferritic stainless steels (No. 3) polarized to – 400mVSCE for 500 seconds in deaerated 25Ł 0.8M HCl + 3.2M LiCl

Current density(A/cm )

2

Current density(A/cm )

Figure 4 XPS analysis of ferritic stainless steels (No. 3) polarized to 287mVSCE for 500 seconds in deaerated 25Ł 0.8M HCl + 3.2M LiCl

10

10

Depth from surface (nm)

Depth from surface (nm)

-4

CrOx FeOx WOx MoOx NiOx

1.6x105

2

10-2

5.21x10-5

3.08x10-5

2.29x10-5

10-3

10-4

0.8M HCl + 3.2M LiCl at 25oC stripping cond.: -287mVSCE for 8 s 30Cr-4Mo 30Cr-3Mo-2W 30Cr-3Mo-2W-2Ni

-5

10

0.0

4.0x104

8.0x104 -1

1.2x105

1.6x105

2

Charge density (cm /C)

Charge density (cm /C)

Figure 6 Log i(t)-vs.-1/q plots of ferritic stainless steels at 200 mVSCE which the potential is stepped to after the potential was held at -287mVSCE for one second in deaerated 25Ł 0.8M HCl + 3.2M LiCl solution

Figure 7 Log i(t)-vs.-1/q plots of ferritic stainless steels at 200 mV which the potential is stepped to after the potential was held at -287mVSCE for 8 seconds in deaerated 25Ł 0.8M HCl + 3.2M LiCl solution

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density (10 µA/cm2), much less than primary passive current density (10 mA/cm2, in Fig. 1). Hence, residual passive film blocks the dissolution of iron and makes tungsten effectively enhance repassivation rate by the stabilization of molybdenate. In addition, the presence of nickel at the interface of metal in contact with passive film [13] is thought to cause it to difficult to transport across passive film and retard repassivation rate. 4. Conclusions 1. By the scratch test and PSC test, how nickel and tungsten affect repassivation rate is explained as follows: - Blocking the passivation by metal not participating in the formation of passive film such as nickel - Stabilizing the film by molybdenate (MoO42-) with accelerating the dissolution rate of iron by tungsten. 2. By the comparison of cBV’s in between the scratch test and the PSC test, a PSC method turns to be valid as a technique to measure repassivation rate and provides the test conditions as follows: - In neural pH solution: PSC test in passive zone after eight-second exposure at –287mVSCE - In acidic pH solution: PSC test in passive zone after 200-second exposure at –400mVSCE. Acknowledgements The work was supported by Korea Science and Engineering Foundation (KOSEF) and Ministry and Science and Technology (MOST) (Grant No. R01-2003-000-10836-0).

References 1. R. N. Parkins: Corrosion (Ed. L. L. Shreir), 2nd ed., Newnes-Butterworths, London 1976 p.83 2. X. G. Zhang and J. Verrecken: Corrosion, 45 (1989) 57. 3. S. I. Pyun and M. H. Hong: Electrochim. Acta., 37 (1992) 2437. 4. T. R. Beck: Electrochim. Acta., 18 (1973) 815. 5. H. J. Engell: The Theory of Stress Corrosion Cracking in Alloys (Ed. J. C. Scully), N.A.T.O. Scientific Affairs Division, Brussels 1971 p.86 6. N. Sato and M. Cohen: J. Electrochem. Soc., 111 (1964) 512. 7. N. Cabrera and N. F. Mott: Rep. Prog. Phys., 12 (1948) 163. 8. E. A. Cho: M. S. Thesis, KAIST 1998 9. H.S. Kwon, E. A. Cho and K.A. Yeom: Corrosion, 56(1) (2002) 32. 10. I. Garcia, D.Drees, and J.P. Celis, Wear, 249 (2002) 452. 11. P. Jixiang, J. S. Kim and K. Y. Kim: to be published on Corrosion, 61 (2005). 12. Y. H. Lee: Ph. D. Thesis, Pohang University of Sci. & Tech. 1998 13. J. E. Castle and J. H. Qiu: J. Electrochem. Soc., 137 (1990) 2031.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

519

Dislocations effect on kinetic of passivation of polycrystalline nickel in H2SO4 medium Mohamed Sahal, Catherine Savall, Juan Creus, R. Sabot, Xavier Feaugas LEMMA, EA3167, Université de La Rochelle, Pôle Sciences et Technologie, Bâtiment Marie Curie, 25, rue Enrico Fermi, 17000, La Rochelle, France

Abstract : Metallurgical state is an important parameter which affects the passivation process. In this study, the first experimental results on the effect of dislocation density and their distribution on the surface on the process of oxide growth are reported for polycrystalline nickel in H2SO4 media at 291 K. It is demonstrated that this effect can not be investigated without tacking into account the grain size and composition of medium. Dislocation density promotes the adsorption of hydroxide and modifies the properties of the oxide layer. Keywords : nickel, passivation kinetic, dislocation density, dislocation patterns, grain size.

1. Introduction According to previous works [1-3], structure of the oxide NiO depends on the crystallographic orientation and on the composition of the solution. It has been clearly demonstrated that when passivation takes place the magnitude of the residual current (jp) depends on the orientation of single crystals, in direct relationship with epitaxy of the oxide formed [1]. In polycrystalline nickel the process of oxides growth and their distribution on metallic surfaces is less documented. In present work, the electrochemical technique is used to evaluate the effect of plastic strain on passivation process on polycrystalline nickel. A particular attention has been provided on the influence of grain size and composition of medium on the interaction between passivation mechanism and dislocation density. 2. Experimental setup Experiments were carried out on polycrystalline high purity (99.99%) nickel with a relatively weak texture and a dislocation density below 1010 m-2 [4]. All the tensile tests were performed at room temperature (300K) with a constant

520

M. Sahal et al.

strain rate of 5.10-4 s-1. For each test, different features of the dislocation density and distribution were studied over 50 grains using a Jeol JEM2011 transmission electron microscope operating at 200 kV [4,5]. Each sample used for the electrochemical analysis was first mechanically then electrochemically polished in a solution of 50 ml of H2SO4 and 350 ml of CH3OH at a current density of 260 mA.cm-2 and temperature 292 K for 10 min. It was then rinsed in ultra pure (UP) water and dried with filtered compressed air. Finally, it was transferred to the electrochemical cell [5] and immediately, a polarisation curve was performed at 0.5 mV/s in the anodic domain. The electrolyte solutions were xM H2SO4 ( x Î [0.25 - 2] mol.L-1) prepared with high-purity H2SO4 and UP water (resistivity of 18.2 W.cm-1) which were adjusted to pH 0.06-4 and deaerated with pure nitrogen for at least 1h 10 min at a pressure of 1.4 bar. Conventional three electrodes potentiostatic method was used with a sulphate saturated electrode (reference electrode), and a platinum grid as counter electrode. 3. Experimental results and discussion Quasi-steady state polarisation curves of the nickel (figure 1a) are comparable to those obtained on monocrystalline [2, 6] and polycrystalline [7,8] samples. The evolution of the logarithm of current density versus potential is a good diagnostic of the mechanisms which occur (figure 1b). In region where the current decreases with potential in relation with the formation of the passive layer, these curves are used to identify potentials range where the logarithm current density is a decreasing linear function of voltage. This specific behaviour can be well described by the phenomenological law: j = j 0- Ad exp[- b Ad E ] where j 0- Ad and b Ad are two parameters which characterise the kinetics of oxide formation [9]. The region located beyond a critical potential exhibits a passive behaviour where the current (jp) is low due to the formation of a protective film composed by hydroxide and oxide species. 1.2E-01

1.0E+00 2

-1

j (A/cm )

dln(j)/dE (mV )

0%

1.0E-01

0% Ad

1.0E-01

34%

34%

D-II 8.0E-02

1.0E-02

(a)

1.0E-03

6.0E-02

D-I

D-II

1.0E-04

4.0E-02

P

(b)

P

Ad

1.0E-05

D-I

2.0E-02

P

E (mV /ESH) 1.0E-06 -250

0

250

500

750

E (mV /ESH)

1000

1.0E-06 -250

0

250

500

750

1000

Figure 1: Quasi- potentiostatic curves (a) and logarithmic slope |dLn(j)/dE| vs E (b) of a strained (ep=0.34) and unstrained nickel in 1M H2SO4 medium .

521

Dislocations effect on kinetic of passivation of polycrystalline nickel

Grain size effect. Pre-passive region is affected by grain size (figure 2). b Ad stays quite constant (0.1 mV-1) but the current density j0- Ad increases as function of grain size. In other term, the formation of hydroxide in relation with adsorption of OH - is promoted by the density of grain boundaries rf following a simple relation: j0 - Ad = Krf -2.6 with rf = 4 ( 3f ) and K equal 1030 and 1028 for an unstrained and a strained (ep=34%) nickel respectively. The passive region is observed for lower potentials when grain size decreases (480 mV/SHE for 18 µm and 510 mV/SHE for 168 µm) but the current density jp is not affected by grain size (figure 2b). Consequently, grain boundary seems do not modify the physical characteristics of oxide layer. However, irrespectively grain size considered, the plastic strain clearly favors the formation of oxyde layer (K[34%] < K[0%]) and affects the stability of this one (jp[34%]>jp[0%]). The effect of the anion concentration (M). The pre-passive stage is clearly affected by an increase of the concentration of ion HSO4- (figure 3a). j0- Ad increases as a function of M for unstrained material. The adsorption of HSO4probably reduces the number of adsorption sites of OH - and consequently delays the formation of hydroxide. This effect is less important for strained metal (figure 3a). Assuming that dislocation is a preferential site for OH adsorption, this fact is well explained in term of dislocation density which promotes the passive process. jp is not affected by the anion concentration (M). 1.E-03

1.E+20

2

2

A/cm

A/cm

1.E+19

j0-Ad j0-Ad-34%

1.E+18

jp jp-34%

1.E-04

1.E+17 1.E-05

1.E+16 1.E+15

(a) 1.E-06

1.E+14 1.E+13

(b)

f(µm)

f(µm)

1.E-07

1.E+12 0

50

100

150

200

0

50

100

150

200

Figure 2: The effect of grain size f on current density j 0- Ad (a) and on residual current density jp (b) of an unstrained and a strained (ep=34%) nickel in 1M H2SO4 medium.

522

M. Sahal et al.

In other term, the properties of oxide layer do not depend on the HSO4concentration which suggests a desorption of this specie before the formation of a well established oxide layer in accordance with previous work [9]. 1.E-03

1.E+23

2

A/cm

2

1.E+22

A/cm

jp

j0-Ad

1.E+21

jp-14%

j0-Ad-14% 1.E+20

1.E-04

1.E+19 1.E+18

(a) 1.E-05

1.E+17 1.E+16 1.E+15

(b)

M

1.E+14 0.00

0.50

1.00

1.50

2.00

2.50

1.E-06 0.00

M 0.50

1.00

1.50

2.00

2.50

Figure 3: The effect of anion concentration (M) on current density j 0- Ad (a) and on residual current density jp (b) of an unstrained and strained (ep=14%) nickel.

pH effect. Independently to plastic strain, b Ad stays quite constant (0.07 ± 0.01 mV-1 for ep=0% and 0.06 ± 0.005 mV-1 for ep=14%) in pH range studied. j0- Ad decreases as a function of pH (figure 4a). This result, in accordance with previous works [7,9], shows that OH - promotes the formation of hydroxide on polycrystalline nickel. Irrespective of pH, plastic strain decreases j0- Ad . 1.E-03

2

1.E+18

2

A/cm

A/cm

j0-Ad

1.E+16

jp

j0-Ad-14%

jp-14%

1.E+14

1.E-04

1.E+12 1.E+10 1.E+08 1.E-05

1.E+06

(a)

1.E+04

(b)

1.E+02

PH

PH

1.E-06

1.E+00 0

1

2

3

4

5

0

1

2

3

4

5

Figure 4: The effect of pH on current density j 0- Ad (a) and on residual current density in passive region jp (b) of an unstrained and strained (ep=14%) polycrystalline nickel.

Dislocations effect on kinetic of passivation of polycrystalline nickel

523

The slope of j0- Ad vs pH is lower for strained metal than for unstrained metal, consequently the difference between strained and unstrained nickel decreases as a function of the pH increase. The passive region is observed for lower potentials when pH increases (450 mV/SHE for pH=4 and 550 mV/SHE for pH=0.6) and the current density jp decreases as a function of pH (figure 4b). Plastic strain increases residual current density jp for lower pH than 2 which confirms the fact that dislocations modify physical properties of oxide film at this level of pH. This last effect seems to be less clear for higher pH than 2. The effect of plastic strain. The effect of plastic strain level on the formation of passive layer has been studied in electrolyte solution 1M H 2 SO4 (figure 5).

b Ad is constant in the strain range studied (0.08 ± 0.01 mV-1) and j0- Ad is lower in strained material than in unstrained one (figure 5a). Consequently, dislocations promote the formation of oxide film. A rapid decreases of the current density j0- Ad is observed at early stage of plastic strain (ep<0.01), followed by an increase at higher strain. This evolution do not correspond to the one observed for total dislocations density r (figure 5b) as function of plastic strain. However, it seems that a correlation exists between the fraction of walls (fW) and the current density j0- Ad (figure 5a). In a first stage of plastic strain, the distribution of dislocations is homogeneous and the j0- Ad decrease can be related to an increase of dislocations density, leading to the increase of the number of adsorption sites, and then to the increase of the hydroxide adsorption sites. From approximately ep=0.024, a heterogeneous distribution of dislocations appears in the nickel (fw increases). Dislocations re-arrange and form progressively a cell structure [5]. This new distribution leads to the concentration of dislocations, and then of adsorption sites in the walls of the cells. It seems that over ep=0.024, the additional sites, created by increasing the strain, develop steric repulsions in the dislocation walls that lead to the increase of j0- Ad . This behaviour is quite similar to the one related to hydrogen adsorption on strained nickel [11]. When passivation takes place, residual current (jp) is an increasing function of plastic strain before ep=0.024, and a decreasing function over this level of plastic strain. Consequently, jp evolution can be correlated to one of dislocations density as a function of plastic strain (figure 5b). The effect of plastic strain on residual current can be interpreted as dislocations inhibiting or reduction of the degree of crystallisation and the epitaxy between NiO and Ni. It is commonly accepted that the residual current in passive region is controlled by field-assisted diffusion of the cations throughout the oxide. Consequently, dislocation density and distributions

524

M. Sahal et al.

promote the diffusion of the cations probably in relation with an increase of vacancies concentration [10]. This feature necessitates further works. 1.E+21 1.E+20

50% 2

A/cm

j0-Ad

1.E+19 1.E+18 1.E+17

40% fw

35%

1.E+16 1.E+15 1.E+14 1.E+13 1.E+12

45%

1.E-04

ep

1.E-03 1.E-02

7.E-05

-2

m

1.0E+15

25%

4.E-05

jp

20%

3.E-05

r

5%

0% 1.E-01 1.E+00

1.2E+15

6.E-05 5.E-05

10%

(a)

1.4E+15 2

A/cm

30%

15%

1.E+11 1.E+10 1.E-05

8.E-05

8.0E+14 6.0E+14 4.0E+14

2.E-05 1.E-05 0.E+00 0.E+00

(b) 1.E-01

ep 2.E-01

3.E-01

2.0E+14

0.0E+00 4.E-01

Figure 5: The effect of plastic strain ep on current density j 0- Ad and on residual current density in passive region jp of polycrystalline nickel in 1M H2SO4 medium Also are shown the fraction of walls (fw) and the density of dislocation (r).

4. Conclusion The first investigations of the effects of dislocations density and distribution on formation and properties of passive film seem to demonstrate that plastic strain promotes the adsorption of OH - and modifies the properties of the oxide layer. Both aspects depend on medium composition but do not depend on grain size. References [1] J. Oudar and P. Marcus, Application of Surface Science, 1979, 3, 48. [2] D. Zuili, V. Maurice, P. Marcus, J. of the Electrochem. Society, 2000, 14 (4), 1393. [3] J. Scherer, B.M. Ocko and O.M. Magnussen, Electrochimica Acta, 2003, 48, 1169. [4] X. Feaugas and H. Haddou, Metalalurgical. Transaction. A, 2003, 34A, 2329. [5] M. Sahal, J. Creus, R. Sabot and X. Feaugas, Scripta Materialia., 2004, 51, 869. [6] V. Alonzo, F. Berthier and L. Priester, Mat. Sci. and Eng. A, 1998, A249, 158. [7] M. Itagaki, H. Nakazawa, K. Watanabe, K. Noda, Corr. Science., 1997, 39 (5), 901. [8] N. Sato, G. Okamoto, Journal of the Electrochemical Society, 1963, 110 (6), 605. [9] A. Jouanneau and M.C. Petit, Journal de Chimie Physique, 1976, 73 , 878. [10] B. Pieraggi, B. MacDougall, R. A. Rapp, Corrosion Science, 2005, 47, 247. [11] H. El Alami, J. Creus, and X. Feaugas, Electrochimica Acta, submitted, 2005.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

525

Electrochemical potential oscillations during galvanostatic passivation of copper in NaNO2 solution and their role in TGSCC mechanism Y. Yagodzinskyya, P. Aaltonenb and H. Hänninena a

Laboratory of Engineering Materials, Helsinki University of Technology P.O. Box 4200, FIN-02015 HUT, Finland b Technical Research Centre of Finland (VTT), Manufacturing Technology P.O. Box 1704, FIN-02044 VTT, Finland; e-mail:[email protected]

Abstract - Auto-oscillations of the electrochemical potential around its average value close to +100 mVSCE under galvanostatic polarization of pure copper in 0.3M NaNO2 solution were studied in annealed state of copper and during slow strain rate tensile tests of cold-deformed copper. The oscillations occur in the potential range where the passivation current increases dramaticaly due to Cu++ cation generation in the semicoherent oxide film on the copper substrate. Parameters of the oscillations were measured depending on the polarization conditions. It was supposed that the observed oscillations represent a nonlinear relaxation process in the passive oxide film interacting with the metal substrate at the semi-coherent oxide/metal interface. Based on the idea that cation vacancy balance at the Cu2O/Cu interface depends on the stresses induced in the metal substrate by the interface conditions, a model of the observed nonlinear phenomena was formulated. Keywords: Copper, Cuprous oxide, Cation vacancy, Galvanostatic polarisation, Electrochemical potential oscillation

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1. Introduction Auto-oscillations of the electrochemical potential under galvanostatic anodic polarization conditions have been observed and well documented for several metal-electrolyte systems [1-3]. This phenomenon was studied in detail for iron in sulfuric acid [1] and was attributed to fluctuation of hydrogen ion concentration in solution at the iron/solution interface in dilute acid or to mass transfer of ferrous ions from corroding steel surface into solution for concentrated sulfuric acid [3]. Auto-oscillations of the potential under galvanostatic conditions were observed in polarisation of Ag in NaOH solution just approaching the potential corresponding to equilibrium between Ag2O and AgO oxides [2]. The observed auto-oscillations of potential seem to be a generic process which can provide useful information on oxidation/dissolution mechanisms. As an auto-catalytic process, the potential oscillation needs some nonlinear feedback for its kinetics. Studies of the origin of feedback are based on the analysis of the oscillation parameters, which can result also in more detailed information on oxide film properties and interactions between the oxide film and the metal substrate. Aim of the present study was to examine the potential oscillations of pure copper galvanostatically polarised in NaNO2 solution under slow strain rate tensile test (SSRT) conditions. Main attention was paid to the origin of the oscillations and their possible role in the oxidation/dissolution mechanisms of pure copper. 2. Experimental Pure copper (99.998 wt.%) containing less than 75 ppm oxygen was used in the study. Billets of copper produced by Outokumpu Poricopper Oy (Finland) were cold-rolled to 1 mm thick strips with intermediate heat treatments. Flat tensile test specimens with 0.5 x 5 mm cross-section and gauge length of 15 mm were prepared from the copper strip. Main part of the specimens were studied in the cold-deformed state. Some of the strips were annealed in vacuum at 800 oC for 20 min and used as a reference state in the anodic oxidation/dissolution studies. The galvanostatic polarisation measurements were carried out for the annealed copper and during SSRT of the cold-deformed specimens with strain rates of 5 x 10-8 and 5 x 10-10 s-1. SSRT tests were carried out by a pneumatic loading device equipped with an electrochemical cell. Slow straining of 0.2 % preloaded specimens was initiated with the potential scan from -700 mVSCE with sweep rate of 1 mV/min. In the range of potentials where the passivation current starts to increase due to Cu++ generation (between +50 and +100 mVSCE), the potential scan was switched to galvanostatic polarisation with the certain value of the

527

Electrochemical potential oscillations during galvanostatic passivation of copper

passivation current. Similar polarisation procedures were used for the annealed copper in a separate electrochemical cell. Polarisation measurements were performed at room temperature in 0.3M NaNO2 solution (pH 8.3) deaerated by nitrogen gas purging. Saturated calomel type electrode and platinum plate were used as reference and counter electrodes in both electrochemical cells. 3. Results Typical potential sweep curves for copper in annealed state and during SSRT of cold-deformed copper with strain rate of 5 x 10-8 s-1 are shown in Fig. 1 a. The second polarisation curve was completed at the tensile strain less than 1 %. It can be seen that the passivation current exhibits two peaks at about +50 and +150 mVSCE separated by a passive range at about +100 mVSCE. It was also found that this pasive range can be observed on the polarisation curve at slow sweep rates only and the passive current density increases with strain rate and the passive range vanishes at about 10-6 s-1 [4]. The sharp increase of the current density at the potential of about +40 mVSCE is caused by the dissolution of the passive Cu2O film due to massive Cu++ cation generation and their transfer from the cuprous oxide to the nitrite solution [4]. Auto-oscillations of the electrochemical potential during galvanostatic polarisation of copper appear at least in the range of negative differential electrical conductivity, namely, between +50 and +100 mVSCE. Initial stages of the electrochemical potential variations at galvanostatic polarisation with current density of 0.5 mA/cm2 for annealed and cold-deformed copper during slow deformation with strain rate of 5 x 10-10 s-1 are shown in Fig. 1 b. 10 2

1

0,14

POTENTIAL, VSCE

CURRENT DENSITY, mA/cm

Cold-deformed copper, strain rate 5 x 10- 10 s- 1

0,16

Annealed copper -8 -1 Cold-deformed copper, strain rate 5 x 10 s

0,1

0,12 Annealed copper

0,10

0,01 0,08

b

a 1E-3 -0,20 -0,15 -0,10 -0,05

0,00

0,05

POTENTIAL, V SCE

0,10

0,15

0,20

0

1000

2000

3000

4000

TIME, s

Figure 1. a) Slow scan rate (1 mV/min) polarisation curves for copper in 0.3M NaNO2 solution at room temperature in annealed state and during SSRT of cold-deformed copper with strain rate of 5 x 10-8 s-1. b) Initial stages of the potential oscillations during galvanostatic polarisation of annealed and cold-deformed copper at current density of 0.5 mA/cm2.

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It can be seen that annealed and cold-deformed copper show rather different potential behavior after application of the constant current: the potential increases for cold-deformed copper, while it decreases in the case of the annealed specimen. Cold-deformed copper specimen needs about 50 mV higher potential in comparison to that for annealed copper to keep the same current density of 0.5 mA/cm2, and regular oscillations of the potential around its average value start to occur in the deformed specimen earlier than in the annealed state. Only a few potential pulses with fast decay were observed in the cold-deformed copper, while the oscillation in the annealed specimen lasts more than 2 h. As it can be seen in Fig. 2 a, oscillations of the annealed specimen finish with an abrupt increase of the potential, up to about +550 mVSCE. The similar behavior of the potential takes place in the cold-deformed copper, but the potential jump begins earlier than in the annealed specimen. Observed auto-oscillations of the electrochemical potential manifest at least quite regular periodicity, while their amplitude varies irregularly in a wide range, as shown in Fig. 2 b. 0,12

0,6 0,11

POTENTIAL, V SCE

POTENTIAL, VSCE

0,5

0,4

0,3

0,2

0,10

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0,0

b 0,06

0

4000

8000

12000

TIME, s

16000

20000

0

3000

6000

9000

12000

15000

TIME, s

Figure 2. a) Time dependence of the potential under galvanostatic polarisation (current density 0.5 mA/cm2) of annealed copper in 0.3M NaNO2 solution at room temperature. b) Enlargement of a).

Tensile deformation and applied current density have a remarkable effect on the observed auto-oscillations of the potential. As can be seen in Fig. 3 a, increase of the strain rate from 5 x 10-10 to 5 x 10-8 s-1 results in an increase of the average period of oscillations from 166.5 to 130.5 s (208 s for annealed state) and a slight increase of their average amplitude from about 22 to about 30 mVSCE (19 mVSCE for annealed copper) at polarisation current density of 0.5 mA/cm2 for cold-deformed copper. Reduction of the current density leads to remarkable changes of the average parameters of oscillations (see Fig. 3 b). When the polarisation current density of cold-deformed copper strained with the strain rate of 5 x 10-8 s-1 was changed from 0.5 to 0.05 mA/cm2 during the tensile test, the average amplitude of oscillations decreased from about 30 to 10 mVSCE, but their period increased from 130.5 to 171.5 s, respectively.

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Electrochemical potential oscillations during galvanostatic passivation of copper

0,17

a 0,15

0,15

POTENTIAL, VSCE

POTENTIAL, VSCE

b

0,16

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0,14 0,13 0,12

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Strain rate 5 x 10- 1 0 s- 1

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Current density 0.5 mA/cm 2

0,10 0

1000

2000

3000

4000

2000

TIME, s

2200

2400

2600

2800

3000

TIME, s

Figure 3. a) Effect of SSRT strain rate on the electrochemical potential oscillations under galvanostatic polarisation with current density of 0.5 mA/cm2. b) Effect of the current density on the potential oscillation during SSRT with strain rate of 5 x 10-8 s-1.

4. Discussion Observed electrochemical potential oscillation is a nonlinear relaxation process which is governed by feedback due to the rate-determining process of the galvanostatic polarisation. Cuprous oxide, Cu2O, is a metal-deficient p-type semi-conductor with electrical conductivity resulting from migration of negative cation vacancies and electron holes [5]. It has been argued [6] that the rateconrolling process of the galvanostatic transient attributes to the metal-oxide interface, where surface charge jn is collected due to the flow of current: j = j0 exp(B(E+jn /i0ir )),

(1)

where B is a constant, E is applied electrical field strength and i0ir is permittivity of the oxide. In the studied case of pure copper the surface charge jn is induced by cation vacancies accumulated at Cu2O/Cu interface. The kinetics of the surface cation vacancy concentration nV´ can be described by the balance equation: dnV´ /dt = j – WnV´jM( ,nV´),

(2)

where W is the capture cross-section of copper ions captured from the metal substrate to the cation vacancy at the oxide side of the Cu2O/Cu interface and jM( ,nV´) is the copper diffusion flux to the interface, which can depend on the strain rate of the copper substrate and on the cation vacancy concentration in the oxide at the interface. The last dependency originates from the crystallographic semi-coherent coupling of the oxide and metal at the interface and any change in nV´, which results in elastic deformation in the oxide film, results also in elastic deformation of the substrate metal due to the conditions of

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elastic-plastic compatibility [7]. Resulting elastic stresses induced in the metal substrate at the Cu2O/Cu interface can markedly affect the enthalpy of the metal vacancy nucleation, which is the key parameter of jM( ,nV´). Misfit and misorientation dislocations, and disconnections [7] are interfacial defects which play a role in the elastic-plastic compatibility at the Cu2O/Cu interface. All these defects can effectively consume or inject vacancies and interact with the lattice dislocations generated during plastic deformation of the copper substrate resulting in the jM dependence on the strain rate . The second term in the nonlinear Eq. (2) provides the feedback for the cation vacancy kinetics at the semi-coherent Cu2O/Cu interface by adjusting the elastic-plastic compatibility conditions and the feedback can result at certain parameter values in oscillating time-dependency of the cation concentration nV´. At galvanostatic polarisation it accommodates according to Eq. (1) with the applied potential oscillations. Detailed analysis of the cation vacancy kinetics at the semi-coherent Cu2O/Cu interface needs, however, further efforts. 5. Conclusion Electrochemical potential oscillations arise in pure cooper during its galvanostatic anodic polarisation in 0.3M NaNO2 solusion in the range of stadystate potentials above +50 mVSCE, where an intensive disollution of the cuprous oxide starts to proceed by Cu++ cations at the oxide/solution interface. Potential oscillations are observed in annealed as well as in cold-deformed copper under SSRT. The average period of oscillations depends on the strain rate of copper substrate. Observed oscillations originate from the nonlinear relaxation oscillations of the cation vacancy V´Cu concentration at the semi-coherent Cu2O/Cu interface. Elastic-plastic compatibility conditions at the interface and vacancy generation and consumption play the key role in the feedback of the nonlinear relaxation mechanism. References 1. J. Wojtowitz, Modern Aspects of Electrochemistry. Eds. Bockris J. O`M. and Conway B., Vol. 8, Chapter 2. Plenum Press, New York, 1972. 2. M. J. Dignam, H. M. Barret, and G. D. Nagy, Canadian Journal of Chemistry, 47 (1969), 4253. 3. M. Tingjiang, X. Naixin, Corrosion Science, 34 (1993), 915. 4. P. Aaltonen et al., this issue. 5. P. Kofstad, Non-Stochiometry, Diffusion, and Electrical Conductivity in Binary Metal Oxides. Wiley Inter-Science, New York, 1972. 6. H. J. De Wit, C. Wijenberg, and C. Crevecoeur, J. Electrochem. Soc., 126 (1979), 779. 7. J. P. Hirth, B. Pieraggi, and R. Rapp, Acta Metall. Mater., 43 (1995), 1065.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

531

Effects of straining on oxide films and passivity of copper in nitrite solution at ambient temperature P. Aaltonena, Y.Yagodzinskyyb, H. Hänninenb a

VTT Technical Research Centre of Finland, Industrial Systems P.O. Box 1704, FIN-02044 VTT, Finland, [email protected] b Department of Mechanical Engineering, Helsinki University of Technology, P.O. Box 4200, FIN-02015 HUT, Finland

Abstract - The effects of strain rate on interactions between copper and its oxide films have been studied. The electrochemical oxidation process of copper is accompanied by the generation of vacancies in the copper substrate. Diffusion of vacancies from the oxide/metal interface or annihilation of vacancies by dislocation reactions is essential for oxidation to continue. Sufficiently slow straining without breaking the passive film on copper leads to a re-arrangement of the dislocation sub-structure at the interface, which helps to consume the oxidation-generated vacancies. The balance of slow straining and oxidation/dissolution produces environmentally-enhanced plasticity in the copper substrate, and simultaneously, accelerated corrosion. These conditions occur as long as the strain rate is low, i.e., on order of less than 10-8 s-1. The effects of slow straining on oxidation/dissolution of high purity copper were studied at various strain rates, even as low as 10-10 s-1. The electrochemical oxidation behaviour of colddeformed copper was recorded by dynamic polarization scans. The synergistic effects of oxidation/dissolution process and straining on the mechanical properties of metal, i.e., dislocation sub-structure of copper were studied by TEM. The results are discussed in conjunction with a TGSCC model (SDVC, Selective Dissolution – Vacancy Creep). Central to the TGSCC model is the concept of vacancies generated in the process of oxidation/dissolution taking part in substrate recovery and playing a key role in the stress corrosion crack growth behaviour. Keywords: Copper, TGSCC, oxidation/dissolution, strain

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1. Introduction One option for the disposal of spent nuclear fuel, is packing it in cast iron canisters having a corrosion shield made of copper, and subsequently embedding the canisters in a deep bedrock repository. Adequate evaluation of the expected 100,000 years lifetime of a 50 mm thick corrosion shield requires the ability to measure environmentally-assisted cracking (EAC) growth rates of -11 less than 10 mms-1. An alternative approach is to better understand the mechanisms of EAC and to improve EAC models and therefore increase the reliability of canister lifetime predictions. There is evidence confirming that excessive amounts of vacancies are generated in copper during oxidation/dissolution reactions, as compared to the amount of vacancies at thermal equilibrium. Continuous generation of vacancies causes changes in the dislocation behaviour at the oxide/substrate interface which further affects the deformation behaviour of copper. Selective anodic dissolution of zinc in brass exposed to tap water has also been shown to provide excessive amounts of vacancies in the substrate [1].The reactions involved in the oxide film and at the interfaces during oxidation/dissolution reactions have been schematically described by a point defect model [2]. The present SCC models utilize electrochemical data gathered from so called “fast-straining” electrode measurements, which measure the repassivation kinetics of films ruptured in specific environments upon rapid, dynamic straining [3]. Peak current densities of copper and brass measured in this manner are orders of magnitude higher than those recorded for unstrained surfaces [3, 4]. Crack growth rates calculated using such “fast straining” current density data in a slip-step dissolution model correlate well with the crack growth rates measured in slow strain rate (SSRT) experiments at relatively fast -6 strain rate (>10 s-1) [4, 5]. However, the crack growth rate data obtained with very slow SSRT of copper demonstrate that EAC crack growth rate does not approach zero, even if the number of film rupture events becomes very small. Thus, assuming that dissolution is an essential part of EAC, there must be additional sources of dissolution current than slip-step dissolution in order to account for the crack growth rates measured in very slow strain rate or static conditions. Crack tips are surrounded by plastic zones having a high dislocation density. The wake surfaces of EAC cracks in copper are always covered with dark cupric oxide all the way up to the tip. For geometrical reasons, the strain rate in the crack wakes is reduced as compared to the nominal strain rate in a SSRT specimen. There is little information available regarding the electrochemical properties of the oxide films formed on the surfaces in the crack wake. The aim of this study is to measure the corrosion-deformation interactions by straining copper and its oxide films without rupture of the films, simulating the exposure of oxidized crack wakes.

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533

2. Experimental In order to first quantify the expected distribution of equivalent von Mieses strains in the wake behind the crack tip a 200 µm deep side crack in SSRT specimen was modelled using finite element analysis (FEA). The crack tip was described by a small notch of a finite radius. The analysis was performed using a quarter of the full geometry by applying symmetrical boundary constraints. The material was modelled as elastic-plastic with isotropic hardening, deformations following an updated Lagrangian scheme, and finite strain implementation. Material for electrochemical tests was pure copper (99.998 wt. %). In order to simulate the advancing stress corrosion crack tip, the material was cold deformed by rolling with intermediate heat treatments to produce a 1.00 mm thick strip. The room temperature yield strength of the resulting cold-deformed copper was 160 N/mm2. Small flat tensile test specimens having a 5 mm wide and 15 mm long gauge region were fabricated by EDM. Slow strain rates were produced by a pneumatic loading device capable of generating low linear displacement rates. Thermal relaxation and dynamic recovery during slow straining of the prestrained copper were also recorded. Dynamic polarization curves using 1 mV/min scan rate were generated while straining the bars at strain rates of 10-6 to 10-10 s-1 in room temperature 0.3M NaNO2 solution of pH 8.3, deoxidized by nitrogen gas purging. A platinum plate was used as a counter electrode, while the reference electrode was a saturated calomel electrode (SCE). Slow straining of 0.2 % pre-strained specimens was initiated simultaneously with potential scan from -700 mVSCE. Transmission electron microscopy (TEM) was used to study the dislocation sub-structure of the copper in the as received, cold-deformed condition, and following thermal relaxation and dynamic recovery in air at room temperature and in conjunction with electrochemical polarization. 3. Results The FEA results show that crack tip plastic zone extends to the crack wake but at distances greater than 15 µm behind the crack tip, only residual plasticity is present and the oxide film experiences slow deflection. In general, the strains, and correspondingly the strain rates experienced by the crack wake are about two orders of magnitude smaller than the nominal value applied in the SSRT. In electrochemical polarization curves of copper in 0.3M NaNO2 solution, Fig 1a, the first peak (at about -160 mVSCE) is related to the thermodynamic cuprous oxide film formation potential. The other (at about +160 mVSCE) is related to cupric oxide film completion, which passivates the specimen surface at the same potential at which the TGSCC in NaNO2 solutions ceases for pure copper [3]. However, reduced scan rate polarization reveals a third peak, at about

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+40…60 mVSCE, which is related to the emergence of Cu++ ions in the oxide film. The potential range between these two last peaks is within the potential range where TGSCC has been reported for pure copper exposed to NaNO2 solutions.

a)

b)

Figure 1. (a) Fast (60 mV/min) and slow (1 mV/min) scan rate polarization curves for unstrained copper at room temperature in 0.3M NaNO2 solution. Also shown are the thermodynamic initiation potentials for oxide formation reactions as well as the experimental potential range for TGSCC in copper in NaNO2 solutions. (b) Slow scan rate polarization curves for copper strained at room temperature in 0.3M NaNO2 solution. Strain rates were 10-6, 10-7 and 10-8 s-1.

Strain rates up to 10-7 s-1 produced continuous film rupture on the specimen surface as indicated by the absence of passivation after the cuprous film formation, and a steadily increasing current density, Fig. 1b. The increase in current density continued in these specimens until the completion of cupric oxide film passivated the specimen surface. Strain rate 10-8 s-1 does not rupture the passive film of copper, and provides a passive potential range following the cuprous oxide film formation. With this slow strain rate the peak at about +40 mVSCE, for initiation of cupric oxide formation, moves closer to the thermodynamic equilibrium of Cu++ emergence potential than was measured with a slow scan rate under unstrained conditions, Fig. 1a. During the scan from +40 to +160 mVSCE, the current density gradually increases and is about two orders of magnitude higher than the current density measured for an unstrained specimen, Fig. 2a. The total strain of a specimen during the potential scan through the TGSCC potential range was much less than 0.1 % at strain rate of 10-8 s-1, which was not enough to initiate visible cracks in the specimens. Increasing scan rate resulted in overshooting of the corrosion potential measured for copper to some extent as compared to the thermodynamically calculated Cu2O formation potential. In the cathodic part of the potential scan reduction of the atmospheric oxide film on the specimen occurred. Mechanical straining consistently produced an elevated corrosion potential value, Fig. 1b. However, reduction of the strain rate resulted in the corrosion potential approaching the equilibrium value.

Effects of straining on oxide films and passivity of copper

a)

535

b)

Figure 2. (a) Comparison of polarization curves for unstrained and strained (10-8 s-1) copper specimens obtained at room temperature in 0.3M NaNO2 solution after normalizing corrosion potentials. A strain-induced difference is visible in the Cu++ emergence potential and in the current density after that for cupric oxide formation. (b) Experimental observations of reduction of area for copper and brass in SSRT conducted in 1M NaNO2 (pH~9) solution at room temperature. Decrease in reduction of area is due to the TGSCC [3].

Thermal relaxation and dynamic recovery take place in cold-worked copper when slowly straining (10-8s-1) at room temperature. Figure 3a shows the reduction in load observed due to thermal relaxation and dynamic recovery in the constant strain rate test in room temperature air alone and for the test conducted with potentiodynamic polarization in the 0.3M NaNO2 solution. TEM examination of the cold-worked copper, slowly strained and polarized in 0.3M NaNO2 solution, revealed areas of dislocation cell formation typical for recovery, Fig. 3b.

a)

b)

Figure 3. (a) Load reduction in constant strain rate tests due to thermal relaxation and dynamic recovery in cold-worked copper at room temperature. The dynamic recovery is slightly accelerated due to oxidation/dissolution reactions, when the potential is high enough, i.e., above +40 mVSCE, for emergence of Cu++ ions in the oxide film. (b) TEM micrograph of cold-worked copper, after recording dynamic polarization curve in 0.3M NaNO2 solution at room temperature while straining the tensile specimen slowly (b).

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4. Discussion Strain rates near the SCC crack tip are low along the crack wake. Slow scan rate dynamic polarization reveals formation and structural transformations of the oxides. Slow straining in the potential range of oxide transformation increases the dissolution/oxidation current density as compared to the current density of the unstrained specimens. The observed increased dissolution/oxidation of copper under slow straining results from enhanced vacancy injection into the substrate enabled by the accelerated consumption of the vacancies by vacancydislocation interactions. Thus, oxidation/dissolution related vacancy generation promotes strain, which in turn makes continuation of oxidation/dissolution possible. Annihilation of vacancies in the substrate is the rate controlling step for oxidation/dissolution reactions at the cupric oxide formation potential range. Thermal relaxation and dynamic recovery at room temperature in slowly strained cold-worked copper reduce the loads to about 20% of the original yield strength. These effects are visible by TEM even in the bulk metal. 5. Conclusions

· Oxidation/dissolution reactions inject vacancies in the base metal. · Annihilation of vacancies controls the reaction rate in cupric oxide formation. · The vacancy surplus is annihilated by slow deformation making continuation of oxidation/dissolution possible and enhancing simultaneously recovery of the deformed material simulating the TGSCC crack tip plasticed zone. · Thermal relaxation, dynamic recovery and corrosion enhanced recovery localise plastic deformation at the crack tip. Acknowledgements This work was performed under the framework of European Fusion Technology Programme by association Euratom/Tekes. The contributions of P. Moilanen, W. Karlsen, E. Hosio, M. Bojinov and E. Arilahti are gratefully acknowledged. References 1. P. Aaltonen, Y. Jagodzinski, A. Tarasenko, S. Smouk and H. Hänninen, Acta Materialia, 46 (1998), 2039. 2. D. D. McDonald, Journal of the Electrochemical Society, 139 (1992), 3434. 3. J. Yu and R. N. Parkins, Corrosion Science, 27 (1987), 159. 4. M. Alvarez, C. Manfredi, M. Giordano and J. R. Galvele. Corrosion Science, 24 (1984), 769. 5. P. Ford, Environment-induced Cracking of Metals, eds. R.P. Gangloff and M.B. Ives (Houston, TX: NACE, 1990), 139.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

537

Repassivation Kinetics of Al-Alloys for Aircraft Structures Joachim Wlokaa, Theo Hackb, Sannakaisa Virtanena a

Department of Materials Science, LKO, University of Erlangen-Nuremberg, Martensstr. 7, Erlangen D-91058, Germany, e-mail: [email protected] b LG-SC, EADS Corporate Research Center Germany GmbH, Munich D-81663, Germany

Abstract - Two different approaches to measure repassivation kinetic parameters of Alalloys at open circuit potential and under polarization were studied. A definition of the term “repassivation time” is given and differences in repassivation behavior of various alloys are discussed. It was found that repassivation is much faster at the lower pH end of the passivity region than at the upper end but the protection of passive layers formed at pH 9 solution was found to be better. Keywords: corrosion fatigue, repassivation kinetics, aluminium alloys, scratch technique, microstructure

1. Introduction Stress Corrosion Cracking (SCC) and Corrosion Fatigue (CF) are serious problems to cope with in aircraft bearing structures. Due to the take-off and landing cycles as well as some resonant oscillations of e.g. the wings, CF is inevitable because an electrolyte film originating from the fallout of the atmosphere (e.g. rainfall, dew, atmospheric moisture) and chloride ions contained in industrial waste fumes or ocean spray is almost always present. [1] The severity of CF is such that it consists of mechanical fatigue with an influence of corrosion and supplementary detrimental synergistic effects. With each cycle a growing fatigue crack is opened and therefore bare metal surface is exposed to electrolyte which is pumped inside by vacuum forces. To minimize

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the influence of corrosion on CF and therefore the detrimental synergistic effects of both, it might be very useful to know if and how fast different materials can repassivate after exposing them to the corroding environment. In order to determine if a thermodynamically stable system suffers from considerable corrosion until the crack tip is fully repassivated during a crack opening and closing cycle, kinetics of repassivation have to be considered. In this work we compare different methods of determination of repassivation parameters. Results of measurements carried out under crack-like conditions on various aluminum aircraft alloys are presented and discussed. 2. Experimental Methods Prior to each experiment, the samples consisting of different aluminum alloys used for aircrafts (pure Al (99,5 %), AA 2024-T351, AA 6013-T4, AlMgSc (with Mg2Al3 intermetallics, Al3Sc dispersoids), AlMgLi 1424 (with Mg2Al3 and LiMgAl2 intermetallics)) were ground on abrasive cloth to remove surface strain and then polished with diamond polishing slurry with grain sizes of 6, 3 and 1 µm, respectively. The polished sample was mounted in an O-ring cell with an opening of 0.78 cm2 in a conventional three electrode assembly with the sample as working electrode, a platinum gauze as counter electrode and an Ag|AgCl reference electrode. The electrolytes were prepared with pure water and reagent grade chemicals and consisted of 0.1M and 0.5M NaCl solution; the pH was adjusted to 3, 4, 5, 6, 9 by adding diluted HCl or NaOH, respectively. The electrodes were connected to an Autolab PGSTAT30 potentiostat and controlled by GPES 4.8 software. Out of the various possibilities to evaluate repassivation properties of metals (enumerated in [2]), we chose a rapid scratching technique. While applying a constant normal force, a diamond tip was rapidly pulled across the sample’s surface by means of a current pulse through a solenoid coil to locally activate the sample. To evaluate the influence of the test method on repassivation kinetic parameters, transients under potentiostatic and at open circuit potential (i = 0) conditions have been measured, Fig. 1.

i

i q(t)

i=g+ ·logt t0

trepass

t

t

Fig. 1 Potential transient and current transient under potentiostatic conditions after scratching.

Due to repassivation, the current density or potential transient recovers to the value which prevailed previously to the activation event [3-9]. As the scratching

Repassivation Kinetics of Al-Alloys for Aircraft Structures

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time is fast compared to the recovering time, a mathematical deconvolution of the single activation and repassivation events could be disregarded [3]. The repassivation behavior under open circuit conditions (= recovery of OCP) can be described by the following empirical equation (1):

e (t ) = a + b × log t

(1)

where i(t) represents the open circuit potential (OCP), g and are constants which depend on the system. Unfortunately g does not represent a minimum value as one could expect from the equation. Stoudt et al. [5] defined as a repassivation rate (as dV/dt) at the maximum of the potential transient. It turned out that this method is inappropriate for the present case because in early stages of the repassivation starting effects (convolution of single activation and repassivation events) show great influence on the shape of the curve. For long scratches the OCP remained on a minimum value for a longer period. Therefore this equation is only applicable for intermediate repassivation stages. Current density transients obey the following empirical rule (2): i (t ) = imax × t -a

(2)

where i(t) is the measured current density, imax is the maximum current density of the transient, t is the time and g is a decay constant. Kwon et al. [9] found for SS 304 that this method is inappropriate to compare the repassivation rate as there are almost no changes for different applied potentials. For aluminum we were able to confirm this conclusion. For varying chloride concentrations, pHvalues or potentials we found no significant differences. To find appropriate terms to characterize repassivation kinetics, we defined a repassivation time t10 which means that the transient declines to 10 % of its peak value. This method is applicable for both the open circuit (OC) and the potentiostatic measurement technique. Although this repassivation time is a relative measure it turned out to describe the repassivation kinetics reliably. As the rise and fall times of the transients are long compared to the sampling intervals, an influence of the sampling rate on the repassivation time t10 could be neglected. 3. Results and Discussion

As is shown schematically in Fig. 1, the repassivation curve under OC conditions follows for the present experiments equation (1). To verify that behavior, an OCP vs. logt plot was drawn for five scratches on AlMgLi, Fig. 2. The OCP decrease portion of the curve was discarded and – with respect to the logarithmic time scale – the minimum potential was shifted to 0.01 s which is the incremental measuring time. Clearly three different regions can be distinguished: (I) early stage (< 0.1 s) where convolution phenomena of single activation and passivation events are present, (II) intermediate stage (0.1-1 s) where equation (1) is valid and (III) final stage (>1 s) where secondary

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OCP in V vs. Ag|AgCl

processes become visible. The total scratching time is about 30 ms. With increasing number of scratches, the repassivation retards. This might be due to the weak passive layer of previously repassivated scratches. With a new potential drop during activation, a freshly repassivated area suffers potential changes which might re-activate them partially. -0,8

III

-0,9 -1,0

II Scratch 1 Scratch 2 Scratch 3 Scratch 4 Scratch 5

-1,1 -1,2

I -1,3 0,01

0,1

1

10

100

time in s Fig. 2 Repassivation curves of AlMgLi alloy under open circuit conditions in 0.5M NaCl at pH 6.

0,24

30

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25

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0,18 15 0,16 10

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repassivation rate in V/s

The repassivation time t10 as it was defined for comparing both methods is in agreement with previously utilized [4, 5] repassivation rate (= from eq. (1)), Fig. 3.

pH of electrolyte

Fig. 3 Repassivation rates in comparison to repassivation times of AA 2024-T351 in 0.1M NaCl solution of different pH.

As it can be seen by comparing both curves in Fig. 3 the repassivation rates are high for low pH-values whereas the repassivation times are low. This shows that repassivation rates and repassivation times are reciprocal values. Another interesting feature can be observed in Fig. 3: Repassivation is retarded for high pH-values (here pH 9) whereas according to the Pourbaix-diagram of pure Al, Al shows stable passivity. For low pH values repassivation is very fast whereas pH 3 designates the passive-active transition of Al. This is a general behavior

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Repassivation Kinetics of Al-Alloys for Aircraft Structures

for all tested aluminum alloys, Fig. 4. As the measured conductivity of all solutions show no significant differences because of the high chloride concentration, an ohmic effect could be ruled out. It is more likely that the prevailing cathodic counter reaction dominates the repassivation kinetics. [10] This finding is of great importance for the SCC and CF behavior of aircraft structures as aluminum repassivates faster in acidified gaps and crevices and cracks than it would do on open non-acidified surface. From Fig. 4 it is obvious that for alloys which contain Mg as a major alloying element, repassivation is strongly retarded for intermediate pH-values (here pH 6). For alloys which contain no Mg the repassivation time increases almost linearly with the pH-value. The lower the concentration of the alloying elements, the better the linear dependence of the repassivation time from the pH-value. repassivation time t10 in s

140

Al AA2024 AlMgLi AlMgSc AA6013

120 100 80 60 40 20 0 -20

3+

+

2 Al + 3 H2O = Al2O3 + 6 H (pH 3.9) 3

4

5

6

7

8

9

pH of electrolyte (0.1M NaCl) Fig. 4 Repassivation times for all tested alloys depending on the pH of the electrolyte.

Under potentiostatic polarizations (up to 100 mV cathodic and anodic to OCP) the measured repassivation times are for e.g. AlMgSc in 0.1M NaCl at pH 9 only 42 ms ± 2 ms which is about 3 oders of magnitude lower than those measured under OCP conditions. Generally, the repassivation times under potentiostatic conditions are in all cases much lower than repassivation times under OCP conditions. Under potentiostatic conditions the activated system is always polarized anodically (compared with OCP curve) and therefore a high anodic current can be measured. This may help to provide the necessary Al3+ ions for forming alumina on the surface. Moreover, under OCP conditions the recharging of the passive film capacity (which is rapidly discharged during scratching) is overlapping the repassivation process and thus only the recharging behaviour is measured [11, 12]. Assessing the two different repassivation times has to take into account what is considered as “repassivation”: The re-establishment of the status antea (OCP) or the re-formation of a protective oxide layer (potentiostatic).

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Evaluations of the passive current density under potentiostatic conditions revealed that the passive current is significantly higher for pH 3 and 4 than for pH 6 or 9. This result shows that the protection by the oxide layer in acidic environments is weaker than in neutral or slightly alkaline environments. This finding is in good agreement with general considerations of the stability of Al passivity. 4. Conclusions

1. Repassivation under potentiostatic conditions appears to be much faster than under open circuit conditions because different physico-chemical effects are measured. Generally open circuit conditions are closer to real systems than polarization of the sample. 2. Repassivation times strongly depend on pH-value; fast repassivation was found in acidic environments, slower repassivation in neutral or slightly alkaline environments due to the prevailing cathodic counter reaction. 3. The protection effect of the oxide layer is larger in higher pH environments than in acidic environments. Therefore we conclude that in low pH environments a weakly protecting oxide layer is rapidly formed, whereas in neutral and slightly alkaline environments the oxide layer forms slower but exhibits more protection. 5. Acknowledgement

This work was supported by Dr. Tempus, Airbus Bremen and EADS CRC, Munich. F. Keller and B. Rackl are gratefully acknowledged for performing experiments and interesting discussions. References 1. M. Tullmin and P. R. Roberge in: Uhlig’s Corrosion Handbook, Chapter 18, Ed. R. W. Revie, John Wiley & Sons, Inc., New York, 2000, 2nd Edition. 2. G.T. Burstein, KSCS 2004, Mechanisms of Corrosion & Corrosion Protection Proceedings, Espoo, Finland 2004. 3. T.A. Adler, R.P. Walters, Corrosion 49 (1993) 399-408. 4. E. Cavalcanti et al., Materials Science Forum 44 & 45 (1989) 235-246. 5. Stoudt, M. R., A. K. Vasudevan, et al., International Symposium of Corrosion Testing of Alumium Alloys, San Francisco, California, ASTM special technical publications, 1990. 6. J.-D. Kim and S.-I. Pyun, Electrochimica Acta 40 (1995) 1863-1869. 7. S.-I. Pyun and E.-J. Lee, Electrochimica Acta 40 (1995) 1963-1970. 8. D. Chidambaram et al., Surface and Interface Analysis 35 (2003) 226-230. 9. H.S. Kwon, E.A. Cho, and K.A. Yeom, Corrosion 56 (2000) 32-40. 10. G. T. Burstein and C. Liu, Corrosion Science 37 (1995) 1151-1162. 11. H. S. Isaacs, Y. Ishikawa, J. Electrochem. Soc. 132 (1985) 1288-1293. 12. M. Hashimoto, S. Miyajima and T. Murata, Corrosion Science 33 (1992) 885-904.

Section H Passivity Breakdown and Localized Corrosion

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Localized Corrosion Growth Kinetics in AA7178 G.S. Frankel, Tsai-Shang Huang, and Xinyan Zhao Fontana Corrosion Center, The Ohio State University, 477 Watts Hall, 2041 College Rd., Columbus, OH 43210, USA ([email protected])

Abstract - The growth kinetics of exfoliation corrosion (EFC) and sharp intergranular corrosion (IGC) fissures in high strength AA7178 were determined and the effects of several factors on the growth kinetics have been studied: alloy, temper, relative humidity (RH), orientation, and stress. Samples were cut in specific orientations relative to the plate rolling direction and were pretreated electrochemically to initiate the attack and generate localized corrosion environments. They were then placed in constant humidity chambers with controlled humidity in the range of 50-96%. The attack proceeded either as IGC or EFC depending on the sample configuration. The growth of EFC depended strongly on RH, and EFC was not found at RH values below about 60%. On the other hand, the growth rate of sharp IGC fissures was independent of RH for test times of up to 1000 h. The difference is likely that sharp IGC fissure growth occurs inside the material with much less connection to the environment than EFC, which pries open the microstructure. Interestingly, the average sharp fissure growth rate was constant with time over this period. Some observations indicate that the sharp fissures propagate as a result of hydrogen embrittlement, and then the crack is filled with electrolyte from the prior IGC region, allowing subsequent corrosion of the crack flanks. Keywords: Al alloy, intergranular corrosion, exfoliation, hydrogen embrittlement

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1. Introduction Wingskins on airplanes are commonly fabricated from AA7xxx alloys such as AA7178. These high strength Al alloys are susceptible to various forms of localized corrosion, including pitting, exfoliation corrosion (EFC), and intergranular corrosion (IGC) in the form of extremely sharp fissures [1]. The structural integrity of aircraft typically is ensured using programs based solely on fracture mechanics. Critical flaw sizes for unstable cracking and rates of fatigue crack growth are known. Corrosion is not part of the calculation, and any corrosion that is found during inspection must be fixed or replaced. This approach has been very successful at maintaining integrity, but it is very expensive. A structural integrity program that includes corrosion would be able to predict the growth of any corrosion and its effect on structural integrity so that a decision might be made to leave the corrosion until the next maintenance cycle and therefore save on maintenance costs. Standardized procedures exist for assessing alloy susceptibility to EFC and IGC. For example, tests for intergranular and exfoliation corrosion in Al alloys are described in ASTM standards G34, G66, G110, and G112 [2-5]. However, these are all simple immersion tests, with visual observation of the results, and they do not provide a measure of growth kinetics. The focus of this study was to develop techniques that can provide information on the growth kinetics of EFC and IGC, and to determine the effects of various important parameters. The foil penetration technique is a simple and direct non-electrochemical method that provides localized corrosion growth kinetics [6,7]. The penetration time for the fastest-growing localized corrosion site is determined for foil samples of varying thickness, which generates a relationship that can be inverted to provide an expression for the depth of the fastest growing site as a function of time. There is no need to make assumptions regarding the shape of the corrosion feature or the extent of hydrogen evolution current, as is required for most electrochemical methods. Zhang and Frankel used this approach to show that the IGC of wrought AA2024-T3 was extremely anisotropic owing to the elongation of the microstructure along the working direction [8,9]. In this study, a modification of the foil penetration method was used to study the growth of sharp IGC fissures in AA7178. EFC was studied using a newly developed method. 2. Experimental Samples were taken from an AA7178 plate (nominal composition Al-6.8Zn2.8Mg-2.0Cu) that was removed from the wingskin of a retired airplane. The samples were cut in specific orientations relative to the plate rolling direction

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and were pretreated electrochemically to initiate the attack and generate localized corrosion environments. The plate microstructure was severely elongated and unrecrystallized. The electrochemical pretreatment was performed in 1 M NaCl. The IGC samples were potentiostatically polarized for 2 h at -725 mV SCE and the EFC samples were galvanically polarized for 7 h at 5.5 mA/cm2, which results in a similar potential range above the second breakdown potential. These pretreatments create localized corrosion of various forms, but primarily selective grain attack where certain grains are selectively attacked and dissolve from the sample, leaving a void. IGC can also be seen on the sides of the cavities [10]. However, this corrosion morphology is not the same as what is observed in real structures exposed to service conditions. In order to obtain realistic forms of attack, the pretreated samples were placed in constant relative humidity (RH) chambers with controlled relative humidity in the range of 50-96%. The samples were simply rinsed with water and dried prior to insertion in the humidity chambers, so that the corrosive solutions within the localized corrosion sites remained. The attack proceeded either as IGC or EFC depending on the sample configuration and area exposed during pretreatment. EFC was tested using through-plate slices that were completely exposed (unmasked, even on the edges) during the pretreatment, and the attack proceeded inward from the edges of the samples, corresponding to the plate surfaces. The progress of EFC was documented by digital photography through the transparent walls of the humidity chamber, and was reported as the depth of attack or thickness of the unattacked region in the center of the sample. IGC in the form of sharp fissures was generated by exposing only a part of the face of a slice during pre-treatment; the edges were not exposed. The fissures propagated from the cavities generated by the pretreatment. The time for IGC to penetrate the slice was determined by visually monitoring the backside of the samples, and the exact length of the sharp IGC fissure at the penetrated site was assessed by serial sectioning. 3. Results and Discussion 3.1. Sharp IGC Fissures As mentioned above, the very sharp fissures observed in real wingskin samples exposed to service environments could not be reproduced by electrochemical treatments, which is perhaps not surprising since real structures are not submersed in liquid solutions. Figure 1a is a cross section of AA7178-T6 sample after the potentiostatic pretreatment of polarization in 1 M NaCl. The dark regions are cavities representing grains that have been selectively dissolved. It is not clear why this form of localized attack in AA7xxx stops at grain boundaries rather than developing into a hemispherical pit. It is possible that solute depleted zones surrounding the grain boundaries are less corrodible

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owing to a depletion of Zn [10,11]. Figure 1b shows a similar sample that was pretreated and then exposed to 96% RH. A sharp fissure can be seen to extend from the bottom of the grain that was dissolved during the pretreatment. This type of sharp IGC fissure is very similar to what is observed on real structures.

(a)

(b)

Figure 1. Cross sections of AA7178-T6 samples. For both samples, the top surface was polarized and the longitudinal or rolling direction is vertical. (a) after polarization pretreatment showing selective grain attack (b) after pretreatment and then exposure to 96% RH showing selective grain attack and sharp IGC fissures.

The kinetics of sharp fissure growth were difficult to assess. The amount of solution in the sharp fissures was insufficient to trigger the sensing mechanism used for immersed foil penetration experiments [8,9]. Therefore, it was necessary to visually monitor the back side of the samples to determine the time of the first fissure penetration. In order to determine the fissure growth kinetics, it was also necessary to know the exact length of the IGC fissure, which is roughly equal to the total sample thickness minus the depth of the pretreatment selective grain attack, which varied from sample to sample and from site to site in a single sample. Serial polishing was used to determine the length of the penetrated fissure. Figure 2 shows a few images from a series of cross sections of a sample that had penetrated from the exposed surface on top to the back surface on the bottom Reconstruction of such sections indicates that the sharp fissures grow intergranularly in three dimensions along grain boundaries and around grains. The origin of the sharp IGC fissure was taken as the point furthest away from the pretreated surface where it was connected to the selective grain attack. Figure 2c shows the deepest selective grain attack associated with the sharp IGC fissure. Therefore, the depth of the sharp IGC fissure was determined from this image, i.e. from the endpoint of the selective grain attack to the bottom of the sample.

Localized Corrosion Growth Kinetics in AA7xxx Alloys

(a)

(b)

(c)

(d)

(e)

549

(f)

(g)

Figure 2. Serial sections of an L sample. A sharp IGC fissure grew from the end of a selective grain attack. This sharp IGC fissure was determined to be 267 mm in length

Measurements such as these were made on many samples. Figure 3 shows the results for fissure depths in AA7178 wingskin material during exposure to relative humidities from 58-96%. It is clear that there is no measurable dependence of sharp fissure growth rate on RH. It seems that the environment built into the corrosion sites by the pretreatment is isolated from the bulk environment. This is very different than the EFC case described below. Furthermore, the rate of growth is practically linear with time over a period of 1000 h; a fitted line gives a slope of 0.33 mm/h. Linear growth kinetics for localized corrosion is also unusual, as localized corrosion usually slows down with time owing to increases in ohmic and diffusional resistance. However, this technique measures a range of fissure growth rates, not only the fastest growing ones as does the foil penetration method. For example, examination of samples with multiple penetration sites has indicated that, in some cases, the site that penetrated first to the backside just initiated closer to the backside, and actually grew slower than other sites that grew longer distances and penetrated later. The upper envelope of the data in Figure 3 represents the fastest growing fissures, and a line fitted to these data points results in an equation for the fastest fissures, d(mm) = 45t(h)1/3. However, applicability at longer times of this relationship or the linear growth law is unclear at this time, so extrapolation should be done with care. Penetrated samples were pulled apart in tension and the matching sharp fissure surfaces were studied by SEM. Some of the sharp fissure areas were found to be covered by a layer of corrosion product on only one side, and this layer exhibited oriented “mud crack” lines that were roughly parallel and regularly spaced at about 1.5-4 mm. Other areas had matching sides both with oriented mud crack lines, and some had matching sides both without mud crack lines. The areas without mud crack lines had matching morphology on both sides.

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One explanation for these observations is that the IG cracks propagate discontinuously. They might propagate initially by a mechanism other than corrosion (possibly hydrogen embrittlement), and then the freshly-opened cracks are filled with electrolyte from the prior IGC region and corrosion of the crack flanks proceeds, which generates more hydrogen. The stress to cause the cracking is either residual or a small tensile force exerted by the corrosion product. However, so far there is no evidence providing clear proof of this interpretation.

Fissure depth (mm)

600 96% RH 83% RH 64% RH 58% RH

500 400 300 200 100 0

0

200 400 600 800 1000 1200 Time t (hr)

Figure 3. Sharp fissure depths of AA7178 wingksin in the longitudinal direction in different humidities.

3.2. EFC Growth Figure 4 shows the pictures of through-plate slices exposed to constant humidity conditions for 11 days following a galvanostatic electrochemical pretreatment. Unlike the sharp fissure growth, there was a strong dependence of exfoliation extent on humidity and a critical RH between 56-66% below which exfoliation seemed to cease. Above that critical RH, the extent of EFC on the edges of the slices increased with increasing humidity. Also with higher humidity, more corrosion product oozed out of the face of the sample. The width of the central unexfoliated zone was measured as a function of time and is given in Figure 5. The rate of exfoliation (equal to ½ the slope of these curves since EFC occurs on both sides) increased with RH and slowed with time. Note that no EFC was found at 56%, so there is no line on this plot at that RH.

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Figure 6 shows the effects of humidity cycling between 56 and 96% RH. There seems to be a rather fast response of EFC to the external environment. It is possible to fit the data in figure 5 to an equation describing the combined effects of RH (fractional) and time t (days) on exfoliation depth d (mm):

d = 1.64E-5 · t · exp(6.35 · RH) Because of the rapid response of the EFC to RH, this equation can be used to predict the exfoliation extent for a sample exposed to varying RH conditions. However, the responsiveness of this system to changes in the environment might result from the specific sample geometry and configuration used in the experiment. The full section of the slices was exposed to the environment, whereas a real structure is typically only exposed on the top surface. Nonetheless, the prying open of the microstructure by the EFC results in greater communication between the bulk environment and the local environment than in the sharp fissure case.

a.

b.

c.

d.

Figure 4. AA7178 wingskin slices after pretreatment and then 11 days in constant humidity. The rolling direction is vertical. a) 56% RH b) 66% RH c) 77% RH d) 96% RH.

Given that the environment is uniformly exposed to the plate section, it is interesting that the rate of EFC slows with time. It is possible that this slowing is caused by the increasing restraint imposed by the previously exfoliated material on the attack at the interface of the exfoliated edge and unexfoliated center of the sample.

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Width Change (mm)

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standard error in whole population

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Time (days) Figure 5. Width change of central unexfoliated region as a function of time for AA7178 wingskin in various constant RH environments.

Cylic Test (96%-56%)

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Figure 6. Results of cyclic exfoliation test for AA7178 wingskin samples.

AA7178 wingskin samples were re-heat-treated to the T6 and T7 tempers and tested for exfoliation. As has been reported previously from EXCO tests [12], the T6 temper was found to be much more susceptible to EFC. For example, the rate of exfoliation for the T7 temper in 96% RH was similar to that of the T6 temper in 76% RH, Figure 7.

Localized Corrosion Growth Kinetics in AA7xxx Alloys

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T6, 96% RH

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Time (days) Figure 7. Width change of central unexfoliated region as a function of time for AA7178 wingskin re-heat-treated to T6 and T7 temper in various constant RH environments.

4. Conclusions The growth rates of sharp intergranular fissures and exfoliation corrosion in AA7xxx alloys was studied, and the following was found: 1. Samples that were electrochemically pretreated to initiate corrosion and then exposed to constant humidity conditions generated primarily either sharp intergranular fissures or exfoliation corrosion depending on the orientation of the sample and the pretreatment exposure configuration. 2. Sharp intergranular fissures grew with a rate that was constant with time over 1000 h, about 0/33 mm/h, independent of the RH level. 3. Parallel mud-crack lines on the sharp fissure surfaces suggest that the sharp IGC fissures grew discontinuously, perhaps by a hydrogen embrittlement mechanism. 4. The rate of exfoliation was found to depend strongly on humidity, and to cease below a critical relative humidity. 5. The exfoliation growth rate reacted quickly to changes in humidity, as there seemed to be good communication between the local and bulk environments. 6. The rate of exfoliation decreased with time, perhaps as a result of an increase in constraint by previously exfoliated material 7. The exfoliation rate for AA7178-T7 was lower than that of AA7178-T6.

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5. Acknowledgements The authors acknowledge funding from the Aging Aircraft Division of ASC in support of the Aeronautical Enterprise Structures Strategy with a contract through S&K Technologies. References 1. R. Kinzie, "Corrosion Suppression: Managing Internal & External Aging Aircraft Exposures," Proceedings of 6th FAA/DoD/NASA Aging Aircraft Conference, San Francisco, 2002. 2. "G112-92, Standard Guide for Conducting Exfoliation Corrosion Tests in Aluminum Alloys," in Annual Book of ASTM Standards Vol. 3.02, ASTM, Philadelphia, PA, 2000, p. 489. 3. "G34-99, Standard Test Method for Exfoliation Corrosion Susceptibility in 2XXX and 7XXX Series Aluminum Alloys (EXCO Test)," in Annual Book of ASTM Standards Vol. 3.02, ASTM, Philadelphia, PA, 2000, p. 124. 4. "G66-99, Standard Test Method for Visual Assessment of Exfoliation Corrosion Susceptibility fo 5XXX Series Aluminum Alloys (ASSET Test)," in Annual Book of ASTM Standards Vol. 3.02, ASTM, Philadelphia, PA, 2000, p. 261. 5. "G110-92, Standard Practice for Evaluating Intergranular Corrosion Resistance of Heat Treatable Aluminum Alloys by Immersion in Sodium Chloride + Hydrogen Peroxide Solution," in Annual Book of ASTM Standards Vol. 3.02, ASTM, Philadelphia, PA, 2000, p. 481. 6. F. Hunkeler and H. Böhni, Corrosion, 37, (1981) 645-650. 7. A. Sehgal, G.S. Frankel, B. Zoofan, and S.I. Rokhin, J. Electrochem Soc., 47, (2000) 140148. 8. Weilong Zhang and G.S. Frankel, Electrochem. Solid-State Lett., 3, (2000) 268-270. 9. Weilong Zhang and G.S. Frankel, J. Electrochem Soc., 149, (2002) B510-519. 10. Qingjiang Meng and G. S. Frankel, J. Electrochem. Soc., 151, (2004) B271-283. 11. T. Ramgopal, P. Schmutz, and G.S. Frankel, J. Electrochem Soc., 148, (2001) B348-356. 12. S.Lee and B.W.Lifka, “New methods for corrosion testing of aluminum alloys”, STP 1134. (ed. V.S. Agarwala and G.M.Ugiansky), ASTM, Philadelphia, PA, 1992.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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An in situ AFM study of the first steps of localised corrosion on a stressed 304L stainless steel in chloride media Frantz Martin a, Sophie Fréchard a, Christian Bataillon b and Jacques Cousty a a b

CEA de Saclay, DSM/SPCSI, 91191 Gif sur Yvette cedex, France; [email protected] CEA de Saclay, DEN/SCCME/LECA, 91191 Gif sur Yvette cedex, France

Abstract – We use a set-up combining an AFM and an electrochemical cell to study in situ the local corrosion of a 304L stainless steel in an aqueous chloride-containing solution. We focus on the sites where pits initiate under controlled potential. We show the influence of the mechanical history of the material on the location of the first pits: it appears that 70% of the pits initiate at strain hardened areas resulting from mechanical polishing. Keywords – in situ AFM, chloride solution, strain hardening, pitting corrosion

1. Introduction It is well known that pitting corrosion of stainless steels is promoted by the addition of chloride ions to aqueous solutions. This type of corrosion has long been investigated for the last decades, and recently the propagation steps of localised corrosion have been well defined and modelled [1]. The uprising of the local probe microscopy in the 90’s, and especially the atomic force microscope (AFM) and its derivatives (SKPFM, EFM…), has brought a new insight on pit initiation: in situ observations with very high resolution compared to the only previously existing in-situ technique, which was optical microscopy, permit at least to link sites of localised corrosion with surface defects. The case of materials containing inclusions has recently been studied: stainless steels with MnS inclusions (ex-situ and in-situ studies) [2,3,4,5], or chromium carbide precipitates [6,7]. In all cases, pits initiate in the neighbourhood of the inclusions or precipitates or at the matrix/impurity interface. However, to our

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knowledge, no study has been carried out on the pit initiation when the stainless steel does not present these chemical defects. The motivation of our study is to understand what generates pitting in chloride solution in absence of chemical defect like inclusions or precipitates. The pitting corrosion of an austenitic 304L stainless steel in chloride borate buffer solution was monitored using an in situ electrochemical AFM. The surface was directly observed before and after an anodic potential scan, and at pitting potential with time. This work is part of a study of the very first stages of stressed induced corrosion of stainless steels in WPR nuclear plants environment. 2. Experimental The set-up used was an AFM Picoscan II from Molecular Imaging, mounted in a “stand-alone” configuration. Underneath, a Teflon electrochemical cell has been installed on a fixed support. All the experiments have been carried out in AFM contact mode with Si3N4 pyramidal tips. The 4 cm3 electrochemical cell had merely little room for a reference electrode, so the measurements were carried out with a 2 electrodes cell: the working electrode was the sample and the counter electrode a silver wire Ag/AgCl, which also played the role of pseudo-reference electrode (+120 mV versus NHE). All potentials given thereafter are referenced to the Ag/AgCl electrode. The exposed sample area in contact with the solution was approximately 2 mm², the rest of the sample being varnished. The potential was controlled using a GAMRY. The solution used was a borate buffer (pH=7.56, [borax]= 0.1 M, [boric acid]= 0.002 M) with diverse chloride concentrations, provided by NaCl salt. All the experiments were carried out at room temperature and atmospheric pressure. The composition of the polycrystalline 304L stainless steel studied was (w%): Cr: 18.68, Ni: 10.14, Mn: 1.72, Mo: 0.35, Cu: 0.15, N: 0.072, C: 0.018 and Fe balanced. The samples were prepared by mechanical polishing with diamond paste (til 1 µm), then submitted to at least 20 hours of chemical-mechanical polishing with colloidal silica, and finally rinsed and sonicated in distilled water and ethanol. The average corrugation (RMS) of the surface as polished was of few angstroms for areas of 1-10 µm². This finishing is supposed to flatten very softly the surface and to suppress even the tiniest scratch coming from the former mechanical polishing. The grain structure was clearly visible because of preferential dissolution processes, depending on grain orientation, occurring during the chemical-mechanical polishing [8]. The average grain size was 10 µm. Electropolishing was performed on some samples in a traditional mixture of ethylene glycol butyl ether (760 mL) and perchloric acid (40 mL) for 10 minutes at 30V, 1A.

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3. Results and Discussion The in situ observations were performed either at free corrosion potential (open circuit) or at fixed controlled potential. In the second case, in order to get a good reproducibility of the surface state of the specimens, the passive films were formed by immersing the samples in solution and polarising to -400 mV until activation. Then the potential was raised from –400 mV until pitting occurred. The 304L stainless steel in these operating conditions exhibits a free corrosion potential (vs. Ag/AgCl ref. electrode) at –230 mV and a passive range of about 510 mV before pitting at +280 ±2 mV (see Fig.1). Because of the sensitivity of the AFM set-up to corrosive solutions, the duration of the experiments has been limited to 48 hours. This short time range restricts the study to the very first steps of localised corrosion, which are incubation and nucleation of pits. At free corrosion potential (-230mV to -50mV) and for various chloride concentrations, from 9g.L-1 to 70 g.L-1 respectively, the evolution of the surface was observed continuously in situ while AFM scanning the surface in contact mode. After 48 hours, no surface modification was noticed, despite the good vertical resolution of the set-up which allows the detection of relief variations of

Figure 1: Polarisation curve of a 304L stainless steel in borate buffer solution, 9 g.L-1 NaCl. a) corrosion potential is –230 mV, b) in the passive range, anodic peaks corresponding to metastable pitting are observed before stable pitting at +280 mV for this sample c). d) shows the reverse part of the scan (rate: 1 mV/s).

the order of few angstroms. The reason is that no detectable localised corrosion events occured, at least in the area observed, which is at maximum 80x80µm². Statistically, according to Shibata et al. [9], incubation time for pit initiation at

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open circuit potential is long and thus the pitting probability quite low. This result indicates that pitting must be initiated via the control of the potential in order to be observed with in situ AFM. Figure 2.a. shows the surface of the sample just after chemical-mechanical polishing. The aspect of the surface is general to the entire sample. One can notice that the surface is very flat (note the z vertical range of only 10 nm) with no ray nor scratch. One notices the presence of slip lines due to plastic deformation, which occured while removing the sample from the polishing plate. The effect of such gliding lines will not be discussed here. A wavelet substructure can be observed as well (encircled): this substructure is due to hydrodynamical processes occuring during the polishing with colloidal silica (periodicity and amplitude depend on the pressure applied by the polishing disc). Figure 2.b. shows the surface after the complete potential scan to +280 mV (different area, same sample). The image shows evident signs of pitting. The pits are very small: 200 nm diameter at the opening for a typical depth of 10-15 nm. This indicates that one observes the first stages of nucleation and growth of the pits. One also notices that the pits are not disposed randomly. A large number of pits is indeed aligned along preferential directions, which can be linked to the apparent scratches on the surface. A statistical study made on ten different areas reveals that at least 70% of the pits are aligned along these scratches, which origin remains to be determined. a

10.00 nm

0.00 nm

b

30.55 nm

0.00 nm

Figure 2: AFM images of a 304L SS sample first mechanically polished and then chemicallymechanically polished. a) Surface as prepared. A grain boundary is underlined by the black dashed line, and wavelets are encircled; b) after polarisation scan to +280 mV in borate buffer solution with 9 g.L-1 NaCl. Scratches have been revealed and pitting occurs along theses rays.

The scratches or lines observed on the surface after the polarisation scan cross different grains without changing direction, which means that their origin is not crystalline. Furthermore, their depth is approximately 10 nm, which is larger than the total vertical range of the images presented in figure 2.a. A possible

In situ AFM study of the first steps of localised corrosion on a stressed 304L SS

559

explanation for the presence of these rays is that they result from the mechanical polishing treatment. In order to support this interpretation, one may eliminate the outermost layer of the material, in which the eventual rays or residual strains due to mechanical stresses might be concentrated. Electropolishing was therefore performed on the samples after mechanical polishing. After that, the samples were submitted to the soft chemical-mechanical polishing treatment like before. An image of the surface obtained at this stage is presented figure 3.a. One can notice the typical wavelet substructure and the absence of ray or scratch. The surface obtained after the polarisation scan to pitting potential is shown in figure 3.b. Pits are observed but there is no evident alignment of the pits. No rays are revealed. The absence of rays and the non-alignment of pits confirms that: 1) the rays observed on the samples studied earlier were only due to the mechanical polishing and 2) pitting occurs preferentially on these rays, since they are aligned along them. In fact, the mechanical polishing induces strain hardening in depths of the order of 5-20 microns. The reliefs made by the mechanical polishing are thereafter suppressed by the chemical-mechanical polishing in surface but the strain hardened zones remain present in the bulk close to the surface. We observe that a cathodic scan followed by an anodic scan can reveal the presence of strain hardened zones, forming rays at the surface. a

38.72 nm

b

0.00 nm

Figure 3: AFM images of a 304L SS sample mechanically, electrochemically and chemicallymechanically polished taken before (a) and after pitting (b). At OCP (a) only wavelets are visible. At pitting potential (b), pits (in black) are visible and randomly distributed on the surface.

The rays revealed in fig.2a correspond thus to areas with remnant strains in the material due to former mechanical polishing. The strains on the surface are thus not uniform. Li et al. [10,11] have shown on brass that the work function, measured using macroscopic Kelvin probing techniques, differs when the material is under compressive, relaxed or tensile deformation conditions. It implies that, at least in the case of brass, the surface potential is not uniform if

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the sample presents compressive or tensile stressed areas. It is assumed that this argument can be applied to our case. This difference in surface potentials could explain both the formation of the scratches during the potential scan and the preferential initiation of pitting inside the scratches. In potentiostatic conditions, this would result in an inhomogeneous spatial distribution of anodic current density. Higher anodic current density would thus be expected at strain hardened areas, at which localised corrosion could be preferentially initiated. 4. Conclusions It has been shown that in-situ AFM imaging of a 304L surface under corrosive conditions can reveal the initiation of localised corrosion at surface heterogeneities other than inclusions. The results emphasise the effect of sample preparation, particularly mechanical polishing. Mechanical polishing induces local strain-hardening of the material, which is proposed to modify substantially the local electrochemical properties on the surface. It is shown that if the sample contains superficial strain-hardened zones (induced by a mechanical polishing treatment), pits initiate preferentially in these areas. References 1. H-H. Strehblow, Mechanisms of Pitting Corrosion in Corrosion Mechanisms in Theory and Practice, Second edition, P. Marcus and J. Oudar, Editors, Marcel Decker, Inc., New York (2002) 2. B. Vuillemin, X. Philippe, R. Oltra, V. Vignal, L. Coudreuse, L.C. Dufour, E. Finot, Corros. Sci. 45 (2003) 1143-1159 3. T. Suter, E.G. Webb, H. Böhni, R.C. Alkire, J. Electrochem. Soc. 148, 5 (2001) B174B185 4. E.G. Webb, T. Suter, R.C. Alkire, J. Electrochem. Soc. 148, 5 (2001) B186-B195 5. R. Ke, R. Alkire, J. Electrochem. Soc. 139, 6 (1992) 1573- 1580 6. R.E. Williford, C.F. Windish Jr, R.H. Jones, Mater. Sci. Engineer. A 288 (2000) 54-60 7. I. Reynaud-Laporte, M. Vayer, J.P. Kauffmann, R. Erre, Microsc. Microanal. Microstruct. 8 (1997) 175-185 8. J.T. Dickinson, R.F. Hariadi and S.C. Langford, Finishing of Advanced Ceramics and glasses, (Ceramics Transactions, Volume 102), American Ceramics Society, Westerville, OH, 213-232 9. T. Shibata and T. Takeyama, Corrosion 33, 7 (1977) 243-251 10. D.Y. Li, L. Wang, W. Li, Mater. Sci. Engineer. A 384 (2004) 355-360 11. W. Li and D.Y. Li, Appl. Surf. Sci. 240 (2005) 388-395

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Initial stage of localized corrosion in artificial pit formed on zinc coated steels by photon rupture SAKAIRI Masatoshi, UCHIDA Yoshiyuki and TAKAHASHI Hideaki Graduate School of Engineering, Hokkaido University, Kita-13, Nishi-8, Kita-ku, Sapporo, 060-8628, Japan. e-mail; [email protected]

Abstract - Photon rupture method, oxide film and metal removal by focused pulsed Nd - YAG laser beam irradiation, has been applied to form an artificial micro pit in zinc coated steels. It takes about 2 s to remove the zinc coated layer with a thickness of about 20 µm by continuance laser irradiation in this experiments. The rest potential transient was measured during laser irradiation. During zinc covered on the steel substrate, the rest potential was changed to noble direction just after laser was irradiated and then returned to previous value. However, after the steel substrate was exposed to the solution, the rest potential move to positive direction just after laser was irradiated and then returned to previous value. The amplitude and duration of the potential change after laser irradiation increased with increasing irradiation period which is related to pit depth and exposed area ratio of zinc / steel substrate. These rest potential fluctuation difference can be explained by the galvanic reaction change in the artificial pit formed on zinc coated steels during irradiation. Key words: zinc coated steel, localized corrosion, artificial pit, photon rupture, chloride ion

1. Introduction Because of their excellent corrosion protection characteristics in atmospheric environments, zinc and its alloy coated steels are widely used. The corrosion protection of the coated layers are ascribed to cathodic protection by a galvanic reaction between the coated layer and substrate [1-3], and the formation of stable and compact corrosion products which have high corrosion resistance. The composition of corrosion products formed on the steels caused by

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atmospheric corrosion has been investigated in a range of exposure conditions [4-8]. Analysis of abruptly destroyed passive oxide films and their repair are important to understand the localized corrosion of metals. Analysis of this behavior has been carried out by monitoring potential- and current-transients after mechanical stripping of the oxide films [9 - 13]. The mechanical film stripping poses problems in the film stripping rate, contamination from stripping tools, and stress and strain on the substrate. Recently, there are reports of film stripping by a photon rupture method (focused pulses of pulsed YAG-laser irradiation), which resolves many of the problems caused by mechanical film stripping. The irradiation of a pulsed laser beam is able to strip oxide film at extremely high rates without contamination from the film removing tools. This technique has been applied to iron by Oltra et al. [14] and by Itagaki et al. [15], to aluminum electrodes [16] and to zinc and its aluminum alloy coated steel by Sakairi et al. [17 - 22]. And this technique is applied to form micro pattern on metals [23, 24]. The purpose of this study is to investigate the effect of area ratio of zinc and steel of artificial pit, which was formed in zinc coated steels by photon rupture, on initial stage of localized dissolution in 0.5 kmol m-3 H3BO3 - 0.05 kmol m-3 Na2B4O7 ( pH = 7.4 ) with 0.01 kmol m-3 NaCl solutions. 2. Experimental The zinc coated steel sheets ( coated layer thickness of about 20 µm, Nippon Steel Co.) were cut into 20 x 20 mm coupon. After cleaning ultra sonically, samples were dipped in nitrocellulose / ethyl acetate solution two times to form an approximately 30 µm thick protective nitrocellulose film on the samples. After formation of the nitrocellulose film, the specimens were immersed in 0.5 kmol m-3 H3BO3 - 0.05 kmol m-3 Na2B4O7 ( pH = 7.4 ) with 0.01 kmol m-3 NaCl solutions, and irradiated by a pulsed Nd - YAG laser ( Sepctra Physics GCR 130 ) through a lens and quartz window at open circuit condition. The laser beam was the second harmonic wave, wave length 532 nm, wave duration 8 ns, and frequency 10 s-1, the laser power was adjusted to 30 mW before the lens. The irradiation interval was 0.1 s. The rest potential transients of the specimens after the laser irradiation were measured by a computer through an A/D converter. The laser irradiation time was also detected by a photo-detector to investigate the irradiation time. A saturated Ag/AgCl electrode was used as a reference electrode to measure the electrochemical data. After the tests, specimen surfaces were examined by a con-focal scanning laser microscope (CSLM). The pit depth was measured by depth analysis function of CSLM.

563

Initial stage of localized corrosion in artificial pit

3. Results and Discussion 3.1. Formation of artificial pit The CSLM contrast images after different periods, ti = 0.1 and 10 s, laser irradiation are shown in Fig. 1. Not only the zinc coated layer, center black area, but also the nitrocellulose film are removed by one pulse of laser irradiation. nitrocellulose film removed area

100 µm

100 µm ti = 10 s ti = 0.1 s Fig. 1 CSLM contrast images after continuous laser irradiation.

However, the shape and the size of the removed area at ti = 0.1 s are almost same as these of the removed area at ti = 10 s. This result means, that we can make artificial pit on coated steel and also can change area ratio between zinc and steel by this technique. 50

Pit depth, d/µm

40 30 20 10 0

0

Same depth as Zn coated layer 5

10

Irradiation duration, ta / s Fig. 2 Change in pit depth with irradiation

Figure 2 shows the change in the pit depth with irradiation duration in 0.5 kmol m-3 H3BO3 - 0.05 kmol m-3 Na2B4O7 solutions. The pit becomes deeper with

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irradiation time and the depth is almost the same size as the zinc coated layer thickness after about 2 s continuous irradiation. The slope also changes when the zinc layer was removed at the irradiated area because of a removal efficiency difference between zinc and steel, which depends on reflectivity and thermal property of the materials. From these results, about 2 s is enough to remove zinc coated layer. Hence, only zinc is exposed to the solution at stage I and both zinc and steel substrate are exposed to the solution at stage II.

Rest potential, E / V vs. Ag/AgCl

3.2. Rest potential change after laser irradiation -1.03 -1.04 Stage II -1.05 -1.06 -1.07 -1.08 -1.09

0.1 s

Stage I

-1.10

Fig. 3 Change in rest potential after laser irradiation with time at stage I and II

Figure 3 shows the change in the rest potential with time after laser irradiation at stage I and stage II. The rest potential suddenly decreases after laser irradiation and show peak and then return to previous value at stage I. However, the rest potential change direction at stage II is opposite at stage I. The amplitude and life time of each potential fluctuation becomes lager with irradiation duration which is related to pit depth and exposed area of metal. Schematic drawing of anodic and cathodic polarization curves, before (solid line) and after (dotted line) laser irradiation is shown in fig. 4, and cross section of artificial pit is shown in fig. 5. In this study, anodic reaction and cathodic reaction are as follows. M M2+ + 2eO2 + 2H2O + 4 e4 OHAt stage I, the anodic reaction site is the center of the irradiated area where coated layer removed and the cathodic reaction site is the outer of the irradiated area where coated layer does not removed by continuous laser irradiation. During laser irradiation, anodic reaction, metal dissolution, may be activated because of oxide film or corrosion product removed at the anodic reaction site.

565

Initial stage of localized corrosion in artificial pit Stage I: activated a nodic reaction

M

Mn+ + ne -

passive surfa ce

Stage I

Laser

nitroce llulose

film

Zn coated

layer

Activated anode site steel substrate

Potential Stage II

nitrocellulose

Zn coated

O2 +4H- +4e-

Laser film

layer

H2 O

Stage II: activated cathodic reaction

Fig. 4 Schematic drawing of anodic and cathodic polarization curves

Activated cathode site steel substrate

Fig. 5 Schematic drawing of artificial pit at each stage

This means the anodic polarization curve changes after laser irradiation. If cathodic reaction do not change at stage I, rest potential should change in the negative direction. After a certain period, the activated area may be covered by corrosion products or oxide film, this causes the rest potential return to previous value. The area of both reaction should be swap round at stage II. Because of this area change, the cathodic reaction is activated by the laser irradiation at stage II, and the rest potential changes in the positive direction after laser irradiation. After a certain period, the activated area may be covered by corrosion products or oxide film, this causes the rest potential return to previous value. 4. Conclusions The effect of area ratio of zinc and steel substrate in artificial pit, which was formed in zinc coated steels by photon rupture method, on initial stage of localized dissolution was examined, and the following conclusions may be drawn. 1) An artificial pit can be formed in a zinc coated layer by continuous focused pulsed YAG laser irradiation. The pit diameter remains constant and the depth increases with increasing time of irradiation.

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2) After laser irradiation, the rest potential changes to the negative direction during zinc coated layer exposed to the solution,. However, after steel substrate also exposed to the solution, the rest potential changes to the positive direction. These rest potential change direction difference can be explained by changing of activated reactions by laser irradiation. Acknowledgements The authors are indebted to The Iron and Steel Institute of Japan and wish to thank Nippon Steel Co. for providing the zinc coated steel sheets. References [1] G. X. Zhang, " Corrosion and Electrochemistry of Zinc ", Plenum Publication Co., New York (1996). [2] Y. Hisamatsu, Bull. Jpn. Instl. Mtals, 20 (1981) 3. [3] Y. Miyoshi, J. Oka and S. Maeda, Trans ISIJ, 23 (1983) 974. [4] T. E. Graedel, J. Electrochem. Soc., 136 (1989) 193C. [5] J. E. Svensson and L. G. Johansson, J. Electrochem. Soc., 143 (1996) 51. [6] J. E. Svensson and L. G. Johansson, Corrosion Sci., 34 (1993) 721. [7] J. J. Frei, Corrosion, 42 (1986) 422. [8] S. Oesch and M. Faller, Corros. Sci., 39 (1997) 1505. [9] F. P. Ford, G. T. Burstein, and T.P. Hoar, J. Electrochem. Soc.,127 (1980) 1325. [10] G. T. Burstein and P. I. Marshall, Corros. Sci., 23 (1983) 125. [11] G.T. Burstein and R. C. Newman, Corros. Sci., 21 (1981) 119. [12] G.T. Burstein and R. J. Cinderey, Corros. Sci., 32 (1991) 1195. [13] R. J. Cindery and G. T. Burnstein, Corros. Sc., 33 (1992) 493. [14] R. Oltra, G. M. Indrianjafy and R. Roberge, J. Electrochem. Soc.,140 (1993) 343. [15] M. Itagaki, R. Oltra, B. Vuillemin, M. Keddam, and H. Takenouti, J. Electrochem. Soc.,144 (1997) 64. [16] M. Sakairi, Y. Ohira and H. Takahashi, Electrochem. Soc. Proc., 97-26 (1997) 643. [17] M. Sakairi, K. Itabashi and H. Takahashi, Corros. Sci. and Tech., 31 (2002) 426. [180] M. Sakairi, K. Itabashi and H. Takahashi, Proc. of Japan-China Joint Seminar on Marin Corrosion (2002) 58. [19] M. Sakairi, K. Itabashi and H. Takahashi, Electrochem. Soc. Proc., 2002-24 (2002) 399. [22] M. Sakairi, K. Itabashi and H. Takahashi, Proc. of Int. Symposium Corrosion Science in the 21st Century (2003) C093. [21] M. Sakairi, K. Itabashi and H. Takahashi, Zaiyro-to-Kankyo, 52 (2003) 524. [22] M. Sakairi, K. Itabashi and H. Takahashi, Proc. of 13th APCCC (2003) 62. [23] T. Kikuchi, M. Sakairi and H. Takahashi, J. Electrochem. Soc. 150, (2003) C567. [24] T. Kikuchi, M. Sakairi, H. Takahashi, Y. Abe, and N. Katayama, Surface and coatings technology, 169 - 170 (2003) 199.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Influence of the chemical dissolution of MnS inclusions on the composition of passive films and the local electrochemical behaviour of stainless steels H. Krawieca,b, V. Vignala, R. Oltraa and O. Heintza a

LRRS, UMR CNRS 5613, Université de Bourgogne, BP 47870, 21078 Dijon, France AGH, University of Science and Technology, Department of Foundry, Reymonta 23 Street, 30-059 Cracow, Poland ; Email : [email protected] b

Abstract - Immersion of stainless steel containing MnS inclusions in aqueous electrolytes leads to the chemical dissolution of these heterogeneities. Chemical dissolution of MnS inclusions in 1M NaCl, pH=3 was studied using in-situ AFM and the dissolution rate of MnS was estimated between 0.04 and 0.19 mm3/min. The local electrochemical measurements reveal that the chemical dissolution of MnS inclusions promotes pitting corrosion. Similary, chemical dissolution of MnS inclusions in 1M NaClO4, pH=3 solution modified the surface close to the inclusions by the presence of FeSO4 in the passive film. Keywords: Pitting corrosion, in-situ AFM, microcapillary, XPS, SIMS

1. Introduction A survey of the literature indicates that numerous studies have been devoted to determine the nature of species released during MnS dissolution and to quantify some key-parameters such as the onset potential for electrochemical dissolution of the MnS inclusion, the amount of MnS dissolved and the pitting potential [13]. To our knowledge, no studies have been attempted to quantify the changes of the passive film composition and the local electrochemical behaviour of the metallic matrix induced by the presence of sulfur species released in the electrolyte during the chemical or electrochemical dissolution of MnS.

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H. Krawiec et al.

In the present paper, chemical dissolution of MnS inclusions was monitored in 1M NaCl, pH=3 at the OCP value using in-situ atomic force microscopy (AFM) and the dissolution rate was calculated from the volume of dissolved MnS. The effects of the chemical dissolution of inclusions on the passive film composition and the local electrochemical behaviour of the matrix were investigated both in 1M NaClO4, pH=3, and 1M NaCl, pH=3, using SIMS, XPS and the electrochemical microcell technique. 2. Experimental Measurements were performed on a 304L type resulfurized stainless steel (Ni: 8.75 wt.%, Cr: 18.3, Mn: 1.7, S: 0.17, Si: 0.5, P < 0.035, C: 0.05, Al: 0.02, B: 0.001 and N: 0.055). The specimens were machined from plates along the short transverse direction (grain size : 50 µm), mechanically ground with silicon carbide (SiC) emeri papers down to 4000 grit and polished with diamond pastes down to 1 mm. They were ultrasonically rinsed in ethanol between each step. The chemical composition of passive films was determined using X-Ray photoemission spectroscopy (XPS) with a SIA 100 Cameca-Riber apparatus and a nonmonochromated AlKa line (potential of 1486.6 eV and power of 300 W). A Mac 2 semi-imaging spectrometer was used with a resolution of 1.3 eV (width of Ag3d5/2 level). C1s peak from pollution (285 eV) was considered for the energy calibration. The surface distribution of chemical species around inclusions was mapped using secondary ion mass spectroscopy (SIMS) with a MIQ 256 Cameca Riber apparatus located in a ultra-high vacuum chamber (of about 10-7 Pa). Images (128´128 pixels) were acquired without preliminary sputtering and repartition of 55Mn+, 34S-, 53Cr+ and 68CrO+ ions was obtained using a Ar+ primary ion source (13 keV, 2 nA, incident angle of 45°). The chemical dissolution of MnS inclusions was investigated using a D3100 AFM Microscope (Veeco-Digital Instruments). The specimens were located in a glass cell containing about 50 cm3 of electrolyte. The local electrochemical behaviour of the matrix close to the inclusions was studied from potentiodynamic measurements carried out at 25°C using the electrochemical microcell technique (diameter of microcapillaries: 50 mm). All potentials were measured versus a saturated calomel electrode (SCE) and a platinum wire counter electrode. A cathodic potential of –500 mV vs. SCE was first applied for 3 minutes and the potentiodynamic polarization curves were determined at a scan rate of 1 mV s-1. 3. Results and Discussion MnS inclusions were found to be spheroidized and heterogeneously distributed on the surface (depth < 5 µm and size ranging between 5 and 40 µm). X-Ray microanalyses performed after polishing indicate that these inclusions are composed of sulfur (48.26±2at.%), manganese (43.4±1.5at.%) and a small

Influence of the chemical dissolution of MnS inclusions

569

amount of chromium (between 3 and 6.5at.%) and iron (between 0.5 and 6at.%). In addition, a depletion of oxygen was observed on the inclusions with respect to the matrix by SIMS, indicating that they were not passivated [4]. Regarding the matrix, the quantitative evaluation of XPS data yields 21.65±0.5% metallic iron, 19.2±0.5% Fe2O3, 48.7±0.5% FeO and 10.45±0.5% FeOOH in Fe2p band. The ratio of FeO to Fe2O3 was found to decrease down to 2.56, corresponding nearly to the inverse spinel structure of Fe3O4. No chromium hydroxide was detected in the Cr2p level and the quantitative evaluation yields 12.9% metallic chromium and 87.1±0.5% Cr2O3 [5]. 3.1. Chemical dissolution in 1M NaClO4, pH=3 for 25 minutes The chemical dissolution rate is low in 1M NaClO4, pH=3, and only a few inclusions were completely dissolved after 25 minutes of immersion. Chemical dissolution occured locally and no uniform dissolution along the interface was detected. This indicates that initiation processes were probably linked to the spatial distribution of defects (nanocracks, vacancies or dislocations) and/or residual stresses at the interface. Once a hole was initiated, dissolution occured continuously up to the absence of MnS. SIMS experiments showed that the surface of cavities left by the dissolved inclusions contains a small amount of sulfur and chromium (Figs. 1(a) and (b)). The peak related to the S2p level was identified in the XPS spectrum at 168.4 eV (Fig. 1(c)). According to the literature [5], this peak corresponds to ferrous sulphate FeSO4.

Fig. 1 : (a-b) Repartition of ionic species determined from SIMS experiments on sites containing some inclusions and (c) S2s and S2p levels on the XPS spectrum obtained after immersion of 304L-RES in 1M NaClO4, pH = 3, at the OCP value for 25 minutes.

However, the quantity of adsorbed sulfur is not homogeneous and grains located far from the cavities left by the dissolved inclusions exhibited a wide passive range, as shown in Fig. 2. On the other hand, the electrochemical behaviour of grains located close to the surface defects was strongly modified by the presence of FeSO4 in the passive film. Higher cathodic currents were measured and no passive range was observed in the anodic domain. Above 130 mV vs. SCE, the current increased sharply to reach a limiting current of about 2 mA/cm2.

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Fig. 2 : Local polarization curves determined in 1M NaClO4, pH=3 on sites containing grains located far from the surface defects (solid line) and close to the surface defects (dotted lines). The specimen was previously immersed in 1M NaClO4, pH=3, at the OCP value for 25 minutes.

3.2. Chemical dissolution in 1M NaCl, pH=3 for 15 minutes The chemical dissolution rate is significantly increased in chloride-containing media and values between 0.04 and 0.19 µm3/min were calculated in 1M NaCl, pH=3, by evaluating the volume of dissolved MnS on AFM images, Fig. 3, [5]. It can also be seen that dissolution does not propagate uniformly along the interface.

Fig. 3 : In-situ AFM images showing the different steps of chemical dissolution of an inclusion. The time indicated corresponds to the time elapsed from the immersion of the specimen.

Only a small quantity of FeSO4 was present in the passive film (small peak visible at 168.8 eV in the XPS spectrum shown in Fig. 4(a)). As a consequence, the beneficial effects of passivation on the electrochemical behaviour of grains located far from inclusions (lower currents) can be observed in Fig. 4(b). By contrast, cathodic reactions were significantly enhanced on grains surrounding inclusions where a high content of FeSO4 was formed.

Influence of the chemical dissolution of MnS inclusions

(a)

571

S2p

(b)

Fig. 4 : (a) S2p level obtained from the XPS spectrum and (b) local polarization curves determined on sites containing grains located far from the inclusions (solid black line) and close to the inclusions (dotted black line) after immersion in 1M NaCl, pH=3 at the OCP value for 15 min (the gray curve was obtained on the polished specimen).

When the ultrasonic cleaning process was performed using water instead of ethanol, chromium hydroxide was detected in the Cr2p level and the quantitative evaluation yields 67±0.5% Cr2O3 and 18.7±0.5% Cr(OH)3. After 15 minutes immersion in 1M NaCl, pH=3, two peaks related to the S2p level were identified in the XPS spectrum at 168.8 eV and 162.2 eV (related to chromium monosulfide), as shown in Fig.5. In this case, the regions surrounding inclusions were strongly affected, as shown in Fig. 6(a), and stable pitting was systematically observed in the polarisation curves at about 500 mV vs. SCE (Fig. 6(b)).

168.8eV

162.2eV

Fig. 5 : S2p level obtained from the XPS spectrum after 15 min immersion in 1M NaCl, pH=3.

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H. Krawiec et al. polished specimen

(b)

grains close to the inclusions

Fig. 6 : (a) Dark field image of a region containing inclusions after 15 min immersion in 1M NaCl, pH=3 and (b) local polarization curves in 1M NaCl, pH=3 on grains that appears bright.

4. Conclusion These results show that under certain conditions the electrochemical behaviour of stainless steels is strongly affected by the chemical dissolution of MnS inclusions. The composition of the passive film plays a significant role and it was found that the surface was less affected by the presence of sulfur-containing species when it was enriched in chromium oxide. Chemical dissolution starts at MnS/matrix interface leading to the formation of chromium and iron sulfide at the surface. In the case of solution containing chloride ions, the regions surrounding MnS inclusions were affected and stable pitting was observed in the polarization curves. After chemical dissolution of MnS inclusions in 1M NaClO4, pH=3, solution, no passive rage was observed on the grains located close to the inclusions. Acknowledgements One of the authors (H.K.) thanks Marie Curie fellowship (grant MEIF-CT2005-007363) for its financial support. References 1. 2. 3. 4.

T. Suter and H. Bohni, Electrochim. Acta, 42, 3275 (1997). E.G. Webb and R.C. Alkire, J. Electrochem. Soc., 149, B272 (2002). H. Krawiec, V. Vignal and R. Oltra, Electrochem. Comm., 6, 655 (2004). H. Krawiec, V. Vignal, E. Finot, O. Heintz, R. Oltra and J. M. Olive, Met. and Mat. Trans., 35A, 3515 (2004). 5. H. Krawiec, V. Vignal, O. Heintz, R. Oltra and J. M. Olive, J. Electrochem. Soc., 152, B213 (2005).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Electrochemical characterization of corrosion resistant alloys in chloride solutions Fabio Bolzonia, Patrizia Fassinab, Gabriele Fumagallia, Sara Goidanicha a

Dipartimento di Chimica, Materiali e Ingegneria Chimica "G.Natta" Politecnico di Milano, Via Mancinelli 7 - 20131 Milano, Italy. b ENI SpA E & P Division - Via Emilia 1 – 20097 S. Donato Milanese, Italy. [email protected]

Abstract - Seven corrosion resistant alloys CRA (six stainless steel and one nickel alloy) were tested in environments related to petroleum industry. Tests were performed in aqueous solution at different temperatures, pH 5.0 and 10, 35, 150 g L-1 NaCl content. The metallurgic, surface and electrochemical characterization was carried out by means of traditional chemical analysis, laser profilometry, SEM and cyclic potentiodynamic polarization. The cathodic process depolarises with increasing temperatures for all studied systems. Localised corrosion resistance is influenced mainly by temperature. A weaker effect is related to increasing chloride concentration. Materials which may undergo localised corrosion show an lectrochemical behaviour characterised by instability of the ip (metastability phenomena). Keywords: localized corrosion, breakdown potential, protection potential, stainless steel, petroleum industry

1. Introduction The paper presents the results obtained in the first phase of a research project dealing with probabilistic models for the localised corrosion prediction of metallic materials for the petroleum industry. The chemical and physical characterisation of the materials is an essential preliminary analysis for a correct approach to a statistical study of localised corrosion [1-3]. Films formed in environments, which may promote localised corrosion, could be unstable. Such film instability phenomena are relatively short-lived and take

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F. Bolzoni et al.

place by local depassivation, giving rise to what we could call a metastable state [4-5]. Considering the presence of metastabile phenomena, it is possible to formulate a three-state model of localised corrosion: (Fig. 1): 1) passivity, 2) metastability and 3) localized corrosion [6].

passivity

localized corrosion

metastability

Figure 1 – Three-state model for localised corrosion phenomena. 2. Metal finishing and metallurgical properties Seven CRA (25 ≤ PREN ≤ 47) were tested: six stainless steels and one nickel alloy. The chemical composition and the corresponding PREN values of the tested materials are reported in Tab. 1. The tensile strength Rm, the yield strength R0.2, the % elongation e, the Brinell hardness HB and the Grain size number G (ASTM E112-96,) are reported in Tab. 2. The materials indicated with [a] were 2 in (≈ 5 cm) diameter and 4 mm thickness pipes, those indicated with [b] were 6 mm thickness plates. The surface treatments performed by the producers were HNO3 and HF pickling and, only for [b], soft sand-blasting. In Tab. 2 the surface roughness revealed by laser profilometer is also reported as parameter Ra (DIN 4768).

390 1.66

29

33

Incoloy 825 *

N08825

6

320 0.8

14

2 23.7 39.3

34

AISI 904 L

N08904

8

390 1.62

23

38

duplex

S31803

23

480 0,78

39

superaustenitic

N08028

11

45

h Mo superaust. N08926

47

superduplex

S32750

0.9

17 11.2

2.1 75

- bal.

-

1.9 30.6

0,5 19.5 24.5

4.1 51

1.3 bal.

18

0.6 22.2

5.3

3.1 173

- bal.

480 1,63

17

0.5

27 31.1

3.4 58

0.9 bal.

9

290 0.88

13

1 20.9 24.7

6.4 210

1.1 bal.

14

280 0,48

18

3.9 288

0.1 bal.

0.7 25.4

6.5

3.2

%

% Mo

11

Fe

% Ni

S31603

%

% Cr

AISI 316 L

Cu

10-3 % S

25

N 10-3 %

10-3 % P

% Mn

Si 10-3 %

UNS number

10-3 %

alloy type

C

PREN **

Table 1 – Chemical composition for the metallic materials under investigation

** * other elements: Ti, Al e Co; austenitic stainlees steels PREN = x + 3.3 y + 16 z, duplex stainlees steels PREN = x + 3.3 y + 30 z (x = % Cr, y = % Mo and z = % N)

Electrochemical characterisation of corrosion resistant alloys in chloride solutions 575

The surface of the materials appeared homogeneous both by visual inspection and by stereomicroscope (50x). The microgeometry was also evaluated by SEM (Scanning Electron Microscope) observation: the pickling attack is more pronounced for the pipes than for the plates; on the latter there was no trace of the abrasive used for the sandblasting. Table 2 – Main mechanical and surface finishing characteristics. alloy type

UNS

Rm

R0.2

(MPa)

(MPa)

i伊 %

HB

G

Ra ( m)

AISI 316 L [a]

S31603

621

303

51

160

6.7

2.8

incoloy 825 [b]

N08825

602

265

45

210

7.9

2.8

AISI 904 L [a]

N08904

622

324

42

180

5.4

1.3

12.3

2.9

5.6

0.5

4.8

2.5

12.7

5.3

duplex [a]

S31803

761

601

29

260

superaustenitic [a]

N08028

659

366

40

150

h Mo superaust. [b]

N08926

713

337

56

200

superduplex [a]

S32750

875

662

31

280

*

*

[a] seamless pipe; [b] plate; * not normed

3. Electrochemical behaviour 3.1. Sample preparation and electrochemical methods Tests were performed at 20, 40 and 70 °C in aqueous solutions at pH 5.0 and with a 10, 35, 150 g L-1 NaCl content. All the materials were tested as received by the production. For the pipes, the internal surface was analysed. Samples preparation (Fig. 1) and electrochemical tests (Fig. 2) followed a procedure that did not damage in any way the surface. The specimens were carefully cleaned. The silicone sealant was acetic acid free and showed good performances under the tested conditions.

electric cable

specimen area = 2 cm2

silicone shield policarbonate tube Figure 1 – Specimen used for electrochemical tests

576

F. Bolzoni et al.

The reference electrode was a Ag/AgCl/KClsat.25°C (+198 mV vs SHE, temperature coefficient -0.6 mV K-1). Electrochemical tests were performed following the ASTM G 6-86 (approved again 2003): -700 mV start potential, 1 V h-1 scanning rate, 2 mA current of reverse potentials. 3.2. Potentiodynamic polarization results

a

irev

1,20

E vs Ag / AgCl / KClsat [V] sat sat [V]

E vs Ag / AgCl / KClsat [V] sat sat[V]

It is well known that potentiodynamic results depend on the scanning rate and pre-immersion time and are scattered on a quite wide range (ASTM G 6-86). Potentiodynamic tests give qualitative information on electrochemical and corrosion properties of the investigated systems. Potentiodynamic tests are of paramount importance in order to choose the correct conditions for the longterm potentiostatic tests which will be carried out in the future to characterise localised corrosion. In Fig. 2 the most significant results obtained from the “potential E-Log(current density i)” characteristics are underlined: · metastability phenomena, due to local breakdown and repassivation of the film (Fig. 2a); · breakdown potential Eb, corresponding to a sudden change of the anodic slope (increasing E) revealing further anodic processes which may be associated to localised corrosion (Fig. 2a) or to other oxidation processes (Fig. 2b), such as water or chloride oxidation; · protection potential Ep, is extrapolated in the return curve, corresponding to i < ip (passivity current density); after a large curve loop, electrode repassivation takes place (Fig. 2a) while, in the presence of a narrow loop (Fig. 2b), the competitive anodic processes disappear.

Eb

0,80

metastability 0,40

Ep

0,00 -0,40

cathodic

-0,80

-4

-2

0 -2

Log i [A m-2]

2

b b

0,80

irev

Eb

1,20

Ep EO2/OH-

0,40 0,00 -0,40

cathodic

-0,80

-4

-2

0

-2

2

Log ii [A [A m m-2]] Log

Figure 2 – E-Log (i) curve obtained in cyclic potentiodynamic polarisation tests in 3.5 g L-1, pH = 5, T = 40 °C: a) duplex SAF 2205, b) high-Mo superaustenitic stainless steel

Electrochemical characterisation of corrosion resistant alloys in chloride solutions 577 Table 3 – Breakdown potential Eb vs Ag/Ag Cl/KClsat (mV) environment pH=5, NaCl

T

10 g L-1 35 g L-1 150 g L 10 g L

-1

35 g L

-1

150 g L 10 g L

20°C

-1

40°C

-1

-1

150 g L

-1

70°C

AISI 316 L

incoloy 825

AISI 904 L

duplex

+ 380

>1V

>1V

-

-

-

-

+380

>1V

>1V

-

-

-

-

+290

>1V

>1V

>1V

>1V

>1V

>1V

-

+680

+1110

-

-

-

-

+210

+550

+540

>1V

+780

>1V

>1V

-

-

-

-

-

>1V

>1V

+120

+440

+480

-

-

+550

+990

0

+170

+140

+120

+260

+560

+480

superhigh-Mo austenic superaustenitic

superduplex

Table 4 – Protection potential Ep vs Ag/Ag Cl/KClsat (mV) environment pH=5, NaCl

T

10 g L-1 35 g L

-1

150 g L

20°C

-1

10 g L-1 35 g L

-1

150 g L 10 g L

40°C

-1

-1

150 g L

-1

70°C

AISI incoloy 316 L 825

AISI 904 L

duplex

superhigh-Mo austenic superaustenitic

superduplex

+50

≥ 600

≥ 600

-

-

-

-

+20

≥ 600

≥ 600

-

-

-

-

-40

≥ 600

≥ 600

≥ 600

≥ 600

≥ 600

≥ 600

-

+520

+250

-

-

-

-

+60

+160

+120

≥ 600

0

≥ 600

≥ 600

-

-

-

-

-

≥ 600

≥ 600

-200

+200

+160

-

-

+140

0

-200

-140

+100

-60

+80

-50

-100

In Tables 3 and 4 potentials Eb and Ep are reported; a conservative approach was choosen, so the less noble values are reported when more than one test was performed. In Table 5 corrosion morphologies detected by means of stereomicroscope (x 50) at the end of tests are reported. In some cases, when the test was repeated, different corrosion morphologies were observed (PCC, see notes in Table 5); in other cases, very weak attacks at the shield boundary of the specimen (CCI) were observed at high magnification by means of SEM. All the specimens subjected to localised corrosion showed presence of metastable phenomena; on the contrary, these phenomena were not observed on specimens that did not suffer localised corrosion, except incoloy 825 in 150 gL-1 NaCl at 20 °C. The trend of cathodic curves (increasing E) is analogous to the behaviour in Fig. 2 for all the materials and environments. The cathodic curves shift to higher

578

F. Bolzoni et al.

values of potential (-250 s +50 mV for i = 0) and to higher values of current density (0.1 s 1 A m-2 for E = - 600 mV) when temperature increase from 20 to 70 °C. This behaviour is confirmed also on platinum electrode. Table 5 – Corrosion morphology after cyclic potentiodinamic polarization tests environment pH=5, NaCl

T

10 g L-1 35 g L

-1

150 g L 10 g L

-1

35 g L

-1

150 g L

40°C

-1

10 g L-1 150 g L

20°C

-1

-1

70°C

superhigh-Mo austenic superaustenitic

superduplex

AISI 316 L

incoloy 825

AISI 904 L

duplex

PCC

NLC

NLC

-

-

-

-

PC

NLC

CCI

-

-

-

-

PC

NLC

NLC

NLC

NLC

CCI

CCI

-

PCC

CC

-

-

-

-

PC

PCC

PCC

NLC

CC

NLC

NLC

-

-

-

-

-

CCI

CCI

PC

PC

PC

-

-

CC

CC

PC

PCC

PC

PC

PC

CC

CC

- = not investigated CCI = Crevice Corrosion Initiation PC = Pitting Corrosion

NCL = No Localized Corrosion CC = Crevice Corrosion PCC = Pitting and Crevice Corrosion

4. Conclusions The cathodic process depolarises with increasing temperatures. Localised corrosion resistance is influenced mainly by temperature. A weaker effect is related to increasing chloride concentration (10 s150 g L-1). According to temperature and PREN, systems potentially susceptible to localised corrosion are: · at 70 °C – all tested materials (maximum PREN = 47); · at 40 °C – materials with PREN < 45; · at 20 °C – material with PREN < 33 (AISI 316 L). Materials which may undergo localised corrosion show an electrochemical behaviour characterised by instability of the ip (metastability phenomena). References [1] [2] [3] [4] [5] [6]

Z. Szlarska-Smialowska, Pitting corrosion of metals. Nace, Houston, Texas (1986), p.431. G. Salvago, G. Fumagalli, D. Sinigaglia, Corr. Sci., 23 (1983) 515. G. Salvago, G. Fumagalli, Corr. Sci., 33 (1992) 985. D.E. Williams, J. Stewart, P.H. Balkwill, Corr.Sci., 36 (1994) 1213. G.S.Frankel, J.Electrochem.Soc., 145 (1998) 2186. F. Bolzoni, P. Fassina, G. Fumagalli, E. Mazzola, proc. 30° convegno nazionale AIM, Milano, 2004, paper 79 (in italian).

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

579

Effect of HCl on Pickling of 304 Stainless Steel in Iron Chloride-Based Electrolytes Lian-Fu Li a,*, Mathieu Daerden b, Peter Caenen b, Jean-Pierre Celis a a

Department MTM, Katholieke Universiteit Leuven, Kasteelpark Arenberg 44, B-3001 Leuven, Belgium b UGINE-ALZ Belgium NV (Arcelor Group), Swinnenwijerweg 5, B-3600 Genk, Belgium * Corresponding author. E-mail address: [email protected] Abstract - Weight loss and corrosion potential measurements were performed to understand the role of HCl in the pickling of oxidized 304 stainless steel in HCl-FeCl3FeCl2 electrolytes. The surface finish was analyzed with SEM-EDX. The oxidized material is active on immersion, resulting in a low corrosion potential and a high weight loss. After certain duration the material either remains active or passive depending on the HCl content. At low contents, an ongoing active-to-passive transition results in localized corrosion at pits, grain boundaries and honeycombed recesses. The corrosion potential becomes high and the weight loss is suppressed. The weight loss decreases in initial stage and rises on extended pickling with adding HCl. Because of surface brightening, the material is always active at concentrated HCl. Keywords: 304 stainless steel, pickling, localised corrosion, uniform corrosion

1. Introduction Pickling is aimed at a total removal of the oxide scales and the Cr-depleted layer formed on stainless steels at elevated temperatures, and a least dissolution of the bulk metal [1,2]. An efficient picking can be achieved solely or successively with an anodic polarisation in a neutral electrolyte [3,4], and a treatment at the open circuit in an oxidizing acidic electrolyte [5-12]. The use of HNO3 as an oxidizing agent during open-circuit treatments is restricted for environmental concerns [10-12]. Electrolytes containing HF or HF-H2SO4 act effectively in the initial stage. However, an active-to-passive transition takes place after certain immersion time [8].

580

L.-F. Li et al.

The electrolyte containing HCl can efficiently remove the Cr-depleted layer, and result in a significantly smooth surface finish [7], attributing to the anodic brightening [13-15]. However, the HCl used in [5-7] is so concentrated that a practical application can be potentially problematic. Therefore, an electrolyte less concentrated with HCl is of more interest, and the study of the corrosion behaviour in such an electrolyte is needed. In this study, the corrosion potential, weight loss and surface finish of 304 stainless steel obtained in HCl electrolytes are investigated to unravel the role of HCl in a pickling. The ongoing localized corrosion and brightening are discussed. 2. Experimental Hot-rolled 304 stainless steel sheets were provided by UGINE & ALZ Belgium NV. The details of the material used were given in [5]. The surface finish consists of an outermost dense oxide layer that is partly broken, and the underlying porous oxide scales. That material is referred to as oxidized stainless steel. The sheets were cut into test samples of 25 mm ´ 25 mm. The samples were covered with an acid-resistant paint to leave only on one side an exposed area of 4 cm2 in the case of weight loss measurements, or 0.25 cm2 in the case of corrosion potential measurements. Analytical grade chemicals and deionizer water were used to prepare the electrolytes containing 0.54 M Fe3+, 0.27 M Fe2+ and up to 2.743 M HCl. The electrolytes were remained at 55 °C. Immersion of oxidized stainless steel was carried out for weight loss and surface finish investigations. After testing, the samples were slightly brushed to remove any reaction products. The surface finish was characterized by scanning electron microscopy (SEM Phillips XL30 FEG). Corrosion potential measurement was performed by using a potentiostat (Gamry CSM100). A three-electrode set-up was used with a test sample as working electrode, a platinum net as counter electrode, and an Ag/AgCl electrode as reference. 3. Results The corrosion potential of oxidized 304 stainless steel in the electrolytes containing 0.54 M Fe3+ and 0.27 M Fe2+ is dependent on HCl concentrations (Fig. 1). The evolution of corrosion potential can be classified into two groups, the first group at 0.137 M HCl or less, and the second group at 0.165 M HCl or more. In the first group, the corrosion potential stays at low levels for a certain time, and then abruptly increases up to a high value. This is due to an ongoing active-to-passive transition as noticed in HF or HF-H2SO4 electrolyte [5,8]. The material exposed to electrolyte is in active dissolution at the low corrosion potentials, thus, the weight loss is significant (Fig. 2). The weight loss stops

Effect of HCl on pickling of 304 stainless steel in iron chloride-based electrolytes

581

HCl content (M) 0 0.219

8 0.027 0.274

0.055 0.687

0.137 1.371

0.165 2.743 -2

0.05

Weight loss (mg cm )

Corrosion potential (V vs. Ag/AgCl)

increasing once the material is passivated. In the second group, the active-topassive transition is not noticed within the test duration (Fig. 1). The corrosion potential remains relatively constant once decreases to the low levels. That behaviour was ascribed to an anodic brightening as the material is immersed in a chloride-concentrated electrolyte [13,14]. As a result, the weight loss increases linearly related to immersion time (Fig. 2).

3+

2+

0.54 M Fe , 0.27 M Fe

0.00 -0.05 -0.10 -0.15

-0.25

4

2

0

-0.20

HCl (M) 0 0.137 0.165 0.274 2.743

6

0

100

200

300

400

Immersion time (s) 0

50

100

150

200

250

300

350

400

Immersion time (s)

Fig. 1 Corrosion potential of oxidized 304 stainless steel in the electrolytes containing 0.54 M Fe3+, 0.27 M Fe2+ and different HCl concentrations.

Fig. 2 Weight loss of oxidized 304 stainless steel in the electrolytes containing 0.54 M Fe3+, 0.27 M Fe2+ and different HCl concentrations.

With increasing HCl concentration, the weight loss becomes less significant in the initial stage, and more pronounced after extended pickling (Fig. 2). The total weight loss is mainly ascribed to the removal of the oxide scales undercut in the initial stage, and the dissolution of the Cr-depleted layer. Since the material remains active and any produced film is absent in the initial stage, the weight loss is simply related to the electrode potential, and the high corrosion potential obtained at low HCl contents results in high weight loss [5,8]. On extended pickling, the weight loss becomes dependent on the films formed. At low HCl contents, the material is subjected to active-to-passive transition. The resulting passive film suppresses the weight loss. On the other hand, anodic brightening takes place in chloride-concentrated electrolytes followed by a formation of a corrosion product film on the material. The high ionic conductivity of that film leads to a significantly high weight loss. SEM surface micrographs after pickling for 400 s depend greatly on HCl concentrations (Fig. 3). Narrow and deep intergranular corrosion takes place in the electrolyte free of HCl or containing 0.027 M HCl (Fig. 3a). The same surface finish was obtained in HF and HF-H2SO4 electrolytes and attributed to the active-to-passive transition [5,6,8]. In comparison to Fig. 3a, the attacks at grain boundaries and honeycombed recesses obtained at 0.137 M HCl are wider (Fig. 3b), indicating that a gradual transition takes place from localized

582

L.-F. Li et al.

corrosion to uniform corrosion. This consists with the extended duration of active state as HCl is added (Fig. 1). While the HCl content increases from 0.137 to 0.165 M, a significant transition in surface finish is achieved (Fig. 3c). The surface is significantly smoothened with limited intergranular corrosion and absence of honeycombed recesses. This tendency for uniform corrosion is confirmed by the active state of corrosion potential throughout immersion (Fig. 1). At 0.219 M HCl or higher, a brightened surface is induced (Fig. 3d), attributing to a compact solid film precipitated on the material [16]. It is suggested that the solid film is present at any HCl contents as the material is in active dissolution. The film formed at low HCl contents readily dissolves after the active-to-passive transition because the corrosion products are no longer replenished. However, at concentrated HCl the material does not undergo active-to-passive transition, the corrosion products are continuously precipitated so that the dissolution proceeds under mass-transport control.

(a)

(b)

(c)

(d)

Fig. 3 SEM surface micrographs of oxidized 304 stainless steel after pickling for 400 s in the electrolytes (a) free of HCl or containing 0.027 M HCl, (b) 0.137 M HCl, (c) 0.165 M HCl, and (d) 2.743 M HCl.

After the active-to-passive transition takes place, pitting corrosion is noticed in an electrolyte free of HCl or containing 0.027 M HCl in addition to the attacks at grain boundaries and honeycombed recesses (Fig. 4).

Effect of HCl on pickling of 304 stainless steel in iron chloride-based electrolytes

(a)

583

(b)

Fig. 4 SEM surface micrographs of pits induced after 400 s immersion in the electrolytes (a) free of HCl and (b) containing 0.027 M HCl.

4. Discussion At very low HCl contents, attacks at pits and grain boundaries take place simultaneously while the material on the other area is passivated after certain time. This is schematically represented by an Evan plot (Fig. 5). Curve 1 illustrates the potential-current relation at passive potentials, curve 2 the pitting corrosion, and curve 3 the anodic behaviour at grain boundaries. Because of the heterogeneity of composition, the material at grain boundaries breaks down at a lower potential than the remaining area. It was noticed in Figures 3a and b and in previous studies [5,8] that the local surface finish achieved at grain boundaries is much smoother than the other areas. Since the corrosion potential of the material is high after active-to-passive transition, the material at grain boundaries can be subjected to anodic brightening (curve 3), while the remaining areas are either passive (curve 1) or in pitting corrosion (curve 2). Because of the high corrosion rate at grain boundaries, the corrosion is controlled by the mass-transport in a corrosion product film precipitated, and the surface finish obtained at grain boundaries is therefore rather smooth. While HCl is increasingly concentrated, the localized corrosion is progressively replaced by a uniform corrosion. Pits are absent and only attacks at grain boundaries take place at the intermediate HCl contents. Moreover, pitting area becomes large and breakdown potential becomes low with increasing HCl content, thus curve 2 is finally merged into curve 3 (Fig. 5), and brightening takes place on the whole surface. 5. Conclusions Oxidized stainless steel is always active on immersion in the HCl-FeCl3FeCl2 electrolytes, the resulting corrosion potential is low and the weight loss is significant. After certain duration, the material is either remained active or

584

L.-F. Li et al.

passivated depending on HCl content. At low contents, an active-to-passive transition takes place, and localized corrosion at pits, grain boundaries and honeycombed recesses is induced. As a result, corrosion potential increases up to high levels, and subsequently the weight loss is suppressed. The corrosion potential of oxidized stainless steel is lowered with adding HCl, whereas the weight loss is decreased in initial stage and enhanced afterwards. When HCl is concentrated, the material always remains active and the weight loss increases linearly with time, because the surface tends to be brightening.

Fig. 5 Schematic Evan plot showing simultaneous pitting corrosion at pits, brightening at grain boundaries, and passivation on remaining area on oxidized 304 stainless steel in the electrolytes at low HCl contents.

References 1. D. Henriet, in Stainless Steels, P. Lacombe, B. Baroux, and G. Beranger, Editors, J.H. Davidson and J.B. Lindquist, English translator, p. 831, de Physique Les Ulis (1993). 2. L.-F. Li and J.-P. Celis, Can. Metall. Quart., 4 (2003) 365. 3. J. Hildén, J. Virtanen, O. Forsén, and J. Aromaa, Electrochim. Acta, 46 (2001) 3859. 4. E. Braun, Iron Steel Eng., 46 (1980) 79. 5. L.-F. Li, P. Caenen, M. Daerden, D. Vaes, G. Meers, C. Dhondt, and J.-P. Celis, Corros. Sci., 47 (2005) 1307. 6. L.-F. Li and J.-P. Celis, Scripta Mater., 51 (2004) 949. 7. L.-F. Li, P. Caenen, and J.-P. Celis, submitted to the J. Electrochem. Soc. 8. L.-F. Li, P. Caenen, and J.-P. Celis, J. Electrochem. Soc, in press. 9. B.S. Covino, J.V. Scalera, T.J. Driscol, and J.P. Carter, Metall. Trans. A, 17A (1986) 137. 10. S. Fortunati, E. Novaro, C. Pedrazzini, and A. Pollastrelli, in Innovation Stainless Steel, p. 2.119, Florence, Italy (1993). 11. N.J. Sanders, Anti-corros. Method. M., 44 (1997) 20. 12. J.P. Malingriaux and Ph. Morleghem, La Revue de Métallurgie, (2000) 1177. 13. T.P. Hoar, D.C. Mears, and G.P. Rothwell, Corros. Sci., 5 (1965) 279. 14. T.P. Hoar, Corros. Sci., 7 (1967) 341. 15. T.P. Hoar and J.A.S. Mowat, Nature, 165 (1950) 64. 16. M. Matlosz, Electrochim. Acta, 40 (1995) 393.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

585

Influence of thermal oxides on pitting corrosion of austenitic and duplex steels Vesna Alara, Vera Redeb, Ivan Juragab, B. Runjeb Faculty of Mechanical Engineering and Naval Architecture, University of Zagreb, Croatia, I.Lučića , 10000 Zagreb, [email protected] b Faculty of Mechanical Engineering and Naval Architecture, University of Zagreb, Croatia, I.Lučića, 10000 Zagreb

a

Abstract - The influence of thermal oxides on the pitting corrosion of austenitic and duplex (austenitic-ferrite) steels has been studied. Austenitic steel specimens AISI 316L were subjected to heating isothermally (200 to 1000 oC) in order to simulate conditions of the application in production. Duplex steel specimens AISI S31803 were heated isothermally at (1100 to 1300 oC) in order to simulate welding parameters. The characterization of thermal oxides formed by heating isothermally on austenitic steel was performed by AES, and XPS analyses. The microstructures and thickness of thermal oxides on austenitic specimens and microstructures on duplex specimens was determined. Pitting susceptibility of such specimens was tested by exposure to FeCl3 solution at 22 ±2 oC during 72 h (ASTM G 48-99a). Pitting criteria show that oxides developed at 600 oC cause the highest susceptibility to pitting corrosion for austenitic steel and 1300 oC for duplex steel. Keywords: thermal oxides, austenitic, duplex, pitting.

1. Introduction Thermal layers formed on austenitic and duplex steels by heating in air or in other gases containing oxygen, carbon dioxide, and/or water vapour affect considerably the susceptibility to pitting corrosion in chloride and in other locally depassivating solutions. That is technically very important since numerous manufacturing operations, like casting, rolling, drawing, forging, heat treatment, welding, and brazing of steels, are carried out at high temperature in contact with air or with other oxidising gases. As G. Bianchi [1] and collaborators have showed it already in the earliest publications on this topic thermal oxides may increase pits density. Later, many investigators studied pitting corrosion of austenitic and duplex steels on specimens with thermal oxides obtained by heating isothermally or in temperature gradient [2-8]. They used topographic, gravimetric, metallographic, electrochemical, and spectrometric techniques and concluded that the formation of thermal layers on steels causes often harmful

586

V. Alar et al.

changes of some important pitting susceptibility criteria while other criteria remain unchanged, change only slightly in favourable or in harmful direction, or even change significantly in favourable direction. These phenomena depend, of course, on austenitic and duplex steels type, on heating conditions, on circumstances causing pitting corrosion etc but the overall effect of thermal oxides is harmful. For that reason, several authors [5, 7, 8] recommend complete removal of thermal layers from steels surface by various mechanical, chemical, and/or electrolytic procedures with the purpose of restoring the original resistance of austenitic steel against pitting. As shown in several earlier papers, [9-11] such treatments may be useful but they have often opposite effects on certain topographic pitting observations so that their application has to be performed cautiously. In order to obtain deeper insight into the influence of thermal oxides on pitting corrosion of austenitic and duplex steels · specimens with thermal oxides were examined using AES, XPS and microscopic, techniques, · topographic criteria for pitting corrosion provoked in FeCl3-solution have been determined on specimens with thermal layers.

2. Experimental and results 2.1. Specimens preparation Rectangular specimens (300 mm x 30 mm) were prepared from 3 mm thick plate of AISI 316L austenitic steel by laser cutting and rectangular specimens (18 x18 x 10 mm) were prepared from AISI S31803 duplex steel. Before further treatment all specimens were degreased in ethanol and spontaneously dried. Experimental results relate only to one side of every specimen. Consequently, the observed surface area was equal to 90 cm2 for austenitic and 3.24 cm2 for duplex specimens. Following types of austenitic and duplex specimens were examined: · austenitic steels - type 1, isothermally heated in air for 20 min at 200, 400, 600, 800, and 1000 oC respectively (ie with thermal oxides) · duplex steels - type 2, isothermally heated in air for 30 min at 1100, 1200, 1250, and 1300 oC.

2.2. Examination of austenitic specimens Determination of phases on the surface of austenitic specimens type 1 has been by AES (Auger electron spectrometry) and X-ray photoelectron spectrometry (XPS). The penetration depth of both methods is about 5 nm so that the results,presented in Table 1. It is to point out that the comparison between bond energies for different peaks of Fe, Cr, and O in XPS spectrum proves the existence of trivalent state of Fe and Cr in their oxides. After heating at 800 and 1000 oC small quantities of Mo-oxides and traces of NiO have been detected. The heating at 200 oC causes no changes on austenitic steel surface due to the transparency of a very thin thermal oxide film. At higher temperatures the growth of semitransparent thicker oxide film brings about the

Influence of thermal oxides on pitting corrosion of austenitic and duplex steels

587

appearance of various interference colours on specimens surface (light yellow at 400 oC, brownish-violet at 600 oC and greenish-blue at 800 oC). The heating at 1000 oC produces thick thermal oxide film, which is opaque and dark gray. Table 1. Results of AES and XPS on the surface of specimens type 1

Attributed to

Types of specimens 1 Element Fe Cr Ni Mo O C

200 oC + + + +

400 oC + * + +

600 oC + ° + *

800 oC + ° ° * + *

1000 oC + * ° * + °

Fe2O3 Cr2O3 NiO Mo-oxides Metallic oxides Contaminations or carbides of Fe and Cr S ° ° * ° Sulphide inclusions (perhaps MnS) Note: - means no response, ° very weak response, * strong response, + very strong response

2.3. Examination of the microstructural of austenitic and duplex specimens Optical microscopy was also used for the examination of specimens type 1 and 2. Grain size of austenitic steel remains unchanged by heating at 200 and 400 oC but grows progressively by heating at 600, 800, and 1000 oC. The average grain sizes, determined by image analysis are approximately:16 µm for specimens type 200 oC, and 400 oC, 20 µm for specimens type 600 oC, 22 µm for specimens type 800 oC, 30 µm for specimens type 1000 oC. Optical micrographs of cross sections of specimens 600 oC, 800 oC, and 1000 oC reveal moreover the existence of zone with crushed grains close to specimens surface. On specimens type 600 oC and 1000 oC the thickness of this zone is about 7 µm and about 16 µm respectively. Ratio of ferrite/austenite in duplex steel is showed in Table 2. Increase of the temperature heating increase ratio of ferrite phase in duplex steel. Table 2. Ratio of ferrite/austenite in duplex steel

Types of specimens 2 F/A

1100 oC 61/39

1200 oC 78/22

1250 oC 86/14

1300 oC 91/9

2.4. Examination of the liability austenitic and duplex on pitting corrosion

F A The testing of the liability of austenitic steel and duplex specimens to pitting corrosion was performed according to ASTM G 48-92 in FeCl3 solution at 25 oC for 72 h. The specimens were immersed into the solution in horizontal position with the examined F

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surface turned upwards. Thereupon, the specimens were rinsed with water. At the same time loose particles of rust formed in the pits were removed by means of plastic brush. Finally, the specimens were rinsed in ethanol and dried in air. On the observed surface of each specimen pits number np has been determined and pits depths hp have been measured. Their orifice areas Sp have been calculated from measured dimensions. That was quite simple since the form of all orifices was almost ideally circular or elliptic. Pits depths measurements were performed using clock –like micrometer gauge with needle shaped feeler. Diameters of circular pits orifices as well as major and minor axes of elliptic orifices were determined microscopically. The changes on specimens type 1 and 2 after the exposure to FeCl3 solution are summarised in Table 3. Table 3.Phenomena observed on specimens type 1 and 2 after the exposure to FeCl3 solution

Specimens type

Phenomena

200 oC no visible changes 400 oC disappearance of yellow colour, many pits 1 600 oC considerable fading of colours, many pits 800 oC considerable fading of colours, many pits 1000 oC considerable fading of colours, many pits 1100 oC no pits 2 1200 oC no pits 1250 oC two pits 1300 oC four pits Colour changes are caused by total or partial dissolution of visible thermal oxides. As mentioned, loose rust particles appeared inside pits cavities. Pitting corrosion susceptibility of austenitic and duplex specimens has been evaluated on the basis of following criteria: pits density np/S (pits number over observed surface area), average pit depth p= ∑ (hp)/np ,average pit orifice area p= ∑ (Sp)/np The values of these pitting criteria for specimens type 1b-e and 2c,d are presented in Table 4. Micrographs with number pits on surface of austenitic and duplex specimens are shown in Figure 1.

a

b

Figure 1: Pits on surface after immersion in solution FeCl3 (a) austenitic specimen (b) duplex specimen

Influence of thermal oxides on pitting corrosion of austenitic and duplex steels

589

Table 4. Pitting observations for austenitic and duplex Specimen

pits density, cm-2 average pit depth, cm average pit orifice area, cm2

Type 1 600 oC 800 oC 1.333

0.889

1000 o C 0.278

0.096

0.099

0.093

0.143

1100 o C no pits -

0.021

0.027

0.015

0.018

-

200 o C no pits -

400 o C 0.555

-

Type 2 1250 1200 o o C C no 0.62 pits 0.11 -

1.64

1300 o C 1.23 0.065 0.135

3. Discussion Surface analysis by AES and XPS confirm the presence of trivalent iron and chromium oxides on specimens type 1 (Table 1) but it reveals also the existence of carbon, sulphur, nickel, and molybdenum compounds after heating of SS at certain temperatures. AES and XPS results refer only to the superficial layer about 5 nm thick on the top of much thicker thermal oxides films on austenitic steels. It is interesting that in this top layer iron oxides prevail considerably over chromium oxides after heating at 600 and 800 oC. Other authors [3,4] have established the same phenomenon but a satisfactory explanation is still missing. Carbon arises perhaps from contaminations or from carbide precipitates in metal. The presence of small quantities of sulphur is probably caused by sulphide inclusions (eg MnS) in the oxide film. After heating at 1000 oC no sulphur is detected possibly because of the transformation of sulphides to oxides by atmospheric oxygen. Optical micrography discloses normal grain growth inside austenitic and crushing of metal grains close to the surface after heating of specimens at 600 oC. The first phenomenon certainly contributes to stresses in specimens and the second one probably results from superficial tensile stress in metal owing to the growth of adhering oxide film with compressive stress. Both phenomena facilitate pits nucleation under certain circumstances. G.T. Burstein [11], local breakdown of passivity is caused by compressive stresses in the superficial film. In the case of thermal oxides such stresses arise primarily due to considerable volume increase going along with the transformation of metal phases (e. g. austenite) into new oxide. The influence of heating procedure on average pits depths and orifice areas, p and p, is ambiguous.

4. Conclusion Specimens of AISI 316L SS with thermal oxide films obtained by isothermal heating by were studied using AES and XPS and optical micrography, techniques. It has been established that the heating for 20 min in the range between 200 and 1000 oC produces visible oxide films adhering to the metal and containing iron and chromium oxides with unimportant quantities of NiO, Mo-oxides, carbon compounds, and sulphides. The

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oxidation of Cr is much faster than the oxidation of Ni and Mo that causes depletion of Cr and accumulation of Ni under the oxide film with accompanying appearance of crystal lattice defects. The presence of several phases in oxide film, metal grain growth at 600 oC and above, and partial transformation of austenite to ferrite about 1000 oC provoke also arising of new lattice defects. Besides, the conversion of metal into oxides increases the volume more than twice creating compressive stress in oxide films and corresponding tensile stress in the metal just underneath. Austenitic and duplex specimens with thermal oxides were tested on pitting susceptibility by immersion into FeCl3 solution at 25 oC for 3 days according to ASTM G 48-99a. · Thermal oxides on austenitic and duplex steels increase pits density, · Maximum increase of pit density is achieved after isothermal heating at 600 oC for austenitic specimens and 1300 oC for duplex specimens, · Thermal oxides affect ambiguously average pits depth and average pits orifice area, · Thermal oxide dissolution decline of oxide influence on pitting propagation.

References 1. 2. 3. 4. 5. 6.

G. Bianchi et al., Corros. Sci., 1970. 10, 19-23. G. Herbsleb et al., Werkst. Korros., 1989, 40, 651-660. S. Turner et al., Corrosion, 1989, 45, 710-716. P.K. Rastogi et al., Brit. Corros. J., 1994, 29, 78-80. J.R. Kearnset et al., Mater. Perform., 1994, 33, 57-61. E. Angelini et al., Proc. Eurocorr ´96, Nice, France, September 1996, EFC, Session VIII OR 13, 4 pp. 7. A.H. Tuthill et al., Mater. Perform., 1999, 38, 72-73. 8. C.P. Dillon, Mater. Perform., 1994, 33, 62-64. 9. I.Esih et al., Proc. Eurocorr ´99, Aachen, Germany, August/September 1999, EFC, Paper 7.10, 7 pp. 10. I. Esih et al., Proc. 15th International Corrosion Congress, Granada, Spain, September 2002, International Corrosion Council, paper 13, 6 pp. 11. 11.G.T. Burstein et al., Proc. 15th International Corrosion Congress, Granada, Spain, September 2002, International Corrosion Council, topic 9, keynote lecture, 14

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Modifications in the electrochemical behaviour of SAF 2205 in alkaline media induced by Cl ions C. M. Abreu, M. J. Cristóbal, R. Losada, X. R. Nóvoa, G. Pena*, M. C. Pérez E. T.S. E.I. e M., Campus Universitario, Universidade de Vigo, 36310 Vigo, (Spain) * Corresponding author,Tel.: +34986812229, e-mail:[email protected]

Abstract - The influence of chlorides on the films electrochemically formed on SAF 2205 in NaOH was studied as a function of the immersion time in NaOH 0.1M + NaCl 0.5M solution up to 1000 h immersion. No Cl incorporation into the passive film is observed. The OCP increases till about 500h immersion, as the outermost part of the passive layer gets oxidised. Keywords: Duplex steel, long-term test, Voltammetry, XPS, passive film ageing

1. Introduction One recognised cause of premature deterioration of concrete structures is the corrosion of the reinforcing steel result of the chloride-induced breakdown of the passive films formed in the alkaline medium of the concrete pores [1]. Therefore, an important number of investigations have been published related to the passive behaviour of mild steel in those media, trying to determine the chloride threshold and the initial stages of the corrosion process [2,3]. The duplex stainless steel (DSS) selected for this investigation has become a good alternative as reinforcement in new concrete structures designed for long service life, due to the similar proportions of austenite and ferrite in its microstructure that confers an excellent resistance to corrosion combined with good mechanic performance at relatively low cost. But this dual phase microstructure is also at the origin of the difficulties in the interpretation of the electrochemical and analytical measurements on these steels. For this reason more information is needed concerning the composition of the passive layers formed in strong alkaline solutions, as well as their stability and resistance to the chloride attack when this aggressive ion ingress in the concrete.

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In the present work the characterisation of the oxide film electrochemically developed on a SAF 2205 in alkaline media is achieved by cyclic voltammetry and electrochemical impedance spectroscopy, as well as XPS. Since the properties of the passive layer are established, the compositional changes induced on the film are analysed as a function of the immersion time into an alkaline chlorinated solution. The results are interpreted trying to enlighten the electrochemical behaviour of the material and the early stages of the film dissolution mechanism. 2. Experimental The chemical composition of the DSS used for this work is: 0.019% C, 22.39% Cr, 0.51% Si, 5.39% Ni, 1.68% Mn, 2.94% Mo, 0.024% P, 0.001% S, %Fe bal. The samples were polished with diamond paste to 6 µm finish and ultrasonically cleaned. The initial passive layer is electrochemically formed in NaOH 0.1M solution, on an exposed area limited to 0.28 cm2. The potential is scanned from hydrogen to oxygen evolution reactions at dE/dt=1mVs-1, for 11 cycles, at room temperature with continuous N2 bubbling, using a conventional three-electrode cell where the working electrode was the SAF 2205, and a Pt mesh was the large area counter electrode. Two reference electrodes were employed: a Hg/HgO 0.1M KOH, to assure the absence of chlorides during the film formation, and a saturated calomel electrode, SCE, when the specimens were submerged in chloride solutions. An AUTOLAB 30 potentiostat with FRA module (from EcoChemie, NL) was used for all electrochemical measurements. This solution was chosen because the redox processes in the passive layer are well differentiated and it is not far from the alkalinity found in cement pastes prepared from standard Portland cement. The low scan rate allows relaxation of the redox processes taking place in passive layers [4,5]. Using these experimental conditions an oxide layer thick enough to be analysed is formed. Once the passive film was generated, the specimens were removed from the NaOH solution and transferred into different recipients containing NaOH 0.1M + NaCl 0.5M solution [6], where the electrochemical impedance spectroscopy is measured at the corresponding open circuit potential (OCP) from 10 kHz down to 1 mHz for different immersion times (0, 100 h, 200h up to 1000h). After EIS one voltammetric curve is registered to follow the evolution of the films. In order to determine the chemical species in the oxide layers, and its distribution across them, the surface of the specimens is characterised by X-ray photoelectron spectroscopy (XPS) using a VG ESCALAB 250iXL spectrometer. The XPS data were collected using monochromatic Al K radiation at 1486.92 eV, and at constant analyser pass energy of 20 eV. Depth profiles analyses were performed using an EX05 Ar+ Ion Gun at 3 KeV. An initial survey spectrum is recorded and after, the high resolution regions of the detected elements (Fe 2p, Cr 2p, Ni 2p, Mo 3d, O 1s, C 1s, N 1s and Cl 2p) are analysed. The fitting of the spectra was accomplished using Shirley background

Modifications in the electrochemical behaviour of SAF 2205...

593

subtraction and Gaussian line shape with addition of an asymmetry factor for the metal peaks. Sensitivity factors for Fe, Cr, Ni, and Mo have been calculated using the bulk material as standard; the same values have been used for the ions of these elements, while for O, C, N and Cl the sensitivity factor was taken for the Wagner Library. Sputtering rate values were determined using perfilometry measurements. 3. Results and Discussion In Fig. 1 the last voltammetric cycle for SAF 2205 in NaOH 0.1M solution is compared to the curves recorded in NaOH 0.1M + NaCl 0.5M solution at four immersion times. Previous works by our group [7,8] explained the processes involved in each of the three potential zones considered. As it can be seen in the figure, as immersion time in alkaline chlorinated solution increases three main changes can be observed in the curves: 1) A decrease in the current density and a simultaneous shift of the peak corresponding to the magnetite formation (Fe2+/Fe3O4, region I) to more anodic potentials. 2) Decay in the Cr3+/Cr6+ current density. 3) Drop in the current density associated to Ni2+/Ni3+ (and consequently in the reduction peak) and O2 evolution.

-2

Current density µ /A cm

200 Region I

Region III

Region II

150 100

Fe2+/Fe3+

Cr3+/Cr6+

50 0

Ox

-50

-100 -150

2+

Red 3+

Ni /Ni t=0 t = 300 hrs t = 650 hrs

-200 -1.4 -1.2 -1.0 -0.8 -0.6 -0.4 -0.2 0.0 0.2 0.4 0.6

Potential / V vs SCE Figure 1: Cyclic voltammograms for SAF 2205 obtained at different immersion times in NaOH 0.1M + NaCl 0.5M. Potential range: –1.4 V to 0.5 V (vs. SCE) (dE/dt = 1 mVs-1 ).

The results of XPS analyses are summarised as follows: the high-resolution spectra of Cl 2p3/2 region (at 199 eV) recorded at different sputtering times show no evidence of chloride penetration into the film, even at the longest immersion time. This effect is in agreement with previously reported results [9].

C.M. Abreu et al.

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The comparison of the depth profiles from Fig. 2 denotes that there is only a small reduction in the film thickness of about 15% (from near 105 nm for t = 0 to 85 nm at t = 650 h), but an important dehydration experienced by the layer (as it can be seen in the OH- profile) as the exposition of chlorides increases. Another important fact derived from the compositional profiles is the enrichment in iron oxides of the external part of the film, and the slight rise in Cr oxides in the internal part. No remarkable differences are noted in the presence of Ni species, only detected as traces in outermost layers of the studied films. 80

80 OHO2Fe Oxide Cr Oxide Fe Metal Cr Metal Ni Metal

60

70

Atomic Conc (%)

Atomic Conc (%)

70

OHO2Fe Oxide Cr Oxide Fe Metal Cr Metal Ni Metal

50 40 30

60 50 40 30

20

20

10

10 0

0 0

20

40

60

Depth (nm)

80

100

0

20

40

60

80

100

Depth (nm)

Figure 2.- XPS depth profiles of a) SAF 2205 after cycling in NaOH 0.1M and b) after immersion in NaOH 0.1M + NaCl 0.5M solution for 650h

Voltammetric and XPS results seem to indicate that superficial adsorption of chlorides induces a modification in the film structure, from a Fe2+/Fe3+ magnetite type to a more oxidized structure. This transformation has been already reported to be enhanced on iron electrodes by the presence of chloride ions [10]. The decrease in the conductive character of the oxide derived from this change would explain the decay in the current density of magnetite formation, Cr oxidation and their shift to higher potentials. This effect in addition to the physical blockage of chloride ions on the surface can also explain the decay in the O2 evolution. In order to verify this hypothesis, the high resolution spectra of the Fe 2p doublet region (2p3/2 and 2p1/2 peaks ) recorded for several levels is displayed in Fig. 3.

Modifications in the electrochemical behaviour of SAF 2205...

595

The evolution in the intensity of the satellite Layer 10 2+ peaks at about 719 eV Fe (715 eV) (attributed to Fe3+) and 715 eV (assigned to Fe2+) [11] can be considered an Layer 1 evidence of a change in Fe2+/Fe3+ ratio, so in the magnetite type structure, to an increasing Fe3+ content towards the oxide-solution interface. The EIS spectra recorded in 3+ Fe (719eV) chlorinated solution at the OCP are in good agreement with the voltammetric results. In Fig. 4a, Nyquist diagrams at several immerFigure 3.- High resolution spectra of the Fe 2p region sion times are depicted. As recorded for several layers. The energies of the considered it can be seen on this figure, satellites is indicated. a continuous increasing in the low frequency limit of the impedance with immersion time is confirmed. 1.4 0.0

1.0 0.8 0.6 0.4 0.2

0 hrs 100 hrs 500 hrs

0.0 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4

OCP (VvsSCE)

Imaginary Part

1.2

-0.1 -0.2

OCP -0.3 -0.4

0

200 400 600 800 1000 1200

Time (h)

2

Real Part (MΩcm ) Figure 4.- a) Nyquist diagrams obtained at the OCP for different times in NaOH 0.1M + NaCl 0.5M, for SAF 2205 samples previously cycled 11 times in NaOH 0.1M solution. b) OCP evolution with immersion time in NaOH 0.1M + NaCl 0.5M

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The OCP values at which impedance measurements have been performed are represented in Fig. 4b. The progressive anodic shift of those potential also seem to confirm the formation of an external layer rich in Fe3+ acting as depolariser of the under layer corrosion reaction. 4. Conclusions The influence of chlorides on the films electrochemically formed on SAF 2205 in NaOH was studied as a function of the immersion time in NaOH 0.1M + NaCl 0.5M solution. Their behaviour was studied using cyclic voltammetry, EIS and XPS. From the discussed results, the following conclusions can be drawn: • Chlorides are adsorbed on the surface of the metal, no Cl incorporation to the passive film has been detected even at longer immersion times. • Ageing in chlorinated solution promotes a change in the film from a Fe2+/Fe3+ magnetite type to a more oxidized structure, enriched in Fe3+ species as immersion time is prolonged. This enrichment in Fe3+ ions in the external layer can explain the decrease in the current densities of the magnetite formation, and the Cr oxidation as well as the shifting the OCP towards more anodic values. • The generation of this less conductive external film produces an increase in the low frequency limit of the impedance spectra recorded at different immersion times. References 1. Corrosion of Reinforcement in concrete. Eds. C.L.Page, K.Wl Treadaway, P.B. Bamford, Elsevier Applied Science, London (1990). 2. M.Moreno, W. Morris, M. G. Álvarez, G. S. Duffó, Corrosion Sci, in press (2004). 3. S. Joiret, M. Keddam, X.R. Nóvoa, M. C. Pérez, C. Rangel, H. Takenouti, Cem. Conc. Comp., 24 (2002) 7. 4. C. Andrade, P. Merino, X. R. Nóvoa, M. C. Pérez, L. Soler, Mater. Sci. Forum, 192-194 (1995) 224. 5. S. Joiret, M. Keddam, X.R. Nóvoa, M.C. Pérez, H. Perrot, H. Takenouti,, Passivity of Metals and Semiconductors” , Ed. by M.B. Ives, J.L. Luo and J. Rodda, (2000) 785. 6. C. M. Abreu, M. J. Cristóbal, R. Losada, X. R. Nóvoa, G. Pena, M. C. Pérez, International Corrosion Congress. Granada, 2002. 7. C. M. Abreu, M.J. Cristóbal, X.R. Nóvoa, G. Pena, M.C. Pérez, V Congreso Nacional de Corrosión y Protección, Madrid, Spain, 2000. 8. C. M. Abreu, M.J. Cristóbal, X.R. Nóvoa, G. Pena, M.C. Pérez, Electrochim. Acta, 47 (2002) 2215. 9. S.V. Phadnis, A.K. Satpati, Muthe, K.P., J.C.Vyas, R.I. Sundaresan, Corros. Sci., 45 (2003) 2467. 10. C. Andrade, M. Keddam, X.R. Nóvoa, M.C. Pérez, C.M. Rangel, H. Takenouti, Electrochim. Acta, 46 (2001) 3905. 11. J.E.Castle, K.Ruoru, J.F. Watis, Corros. Sci., 30, 8-9 (1990) 771.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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In-situ Raman spectroscopy study of iron and carbon steel corrosion in mineral water Lise Lanardea,b, Suzanne Joireta, Xavier Campaignolleb, Michel Meyerb a

Laboratoire Interfaces et Systèmes Electrochimiques, UPR15 du CNRS, Université Pierre et Marie Curie, CP 133, 4 Place Jussieu, 75252 Paris Cedex 05, France. b Gaz de France, R&D Division, 361 av. du Prés. Wilson, BP33, 93211 Saint Denis La Plaine, France. [email protected]

Abstract - The corrosion behaviour of iron and carbon steel in mineral waters has been studied at open circuit potential or cathodic polarisation thanks to in-situ Raman spectroscopy. Corrosion process is not modified whatever the conditions (even with complete deaeration of the electrolyte). Pitting, associated with green rust formation, occurs in all cases at open circuit potential. Pitting is clearly associated with impurities in the metal and developed around sulfide (in carbon steel) or silicates (in iron). Evolution of green rust towards more stable iron compounds depends on the electrolyte conditions. Formation of an insulating barrier of a calcium (or magnesium) salt (silicate in acidic, carbonate in basic) has been analysed during cathodic polarisation. However dissolution of iron as Fe2+ ion is not totally stopped and siderite (iron II carbonate) or iron II sulfate has been founded underneath this layer. Keywords : Corrosion, cathodic polarization, steel, mineral waters

1. Introduction Onshore gas transmission lines are conjointly protected against external corrosion by cathodic protection (CP, from -775mV/SCE to -1125mV/SCE) and organic coatings. If both protection systems are simultaneously faulty, the pipeline steel may be subjected to a local lack of corrosion protection. Consequently, corrosions may develop due to the soil intrinsic corrosivity. This study is aimed to assess whether the integrity threat of a corrosion can be evaluated by the characterization and identification of its associated deposits. The corrosion behaviour of pure iron and X70 steel in either aerated siliceous water (Volvic®) or deaerated carbonated water (Evian®), have been studied at

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open circuit potential (OCP) and under CP (-925mV/SCE to -1125mV/SCE), at room temperature. 2. Experimental procedure 2.1. Characterization techniques The Raman spectrometer was a Jobin-Yvon Horida (LABRAM). The spectra were obtained with 632 nm radiation from an internal 10mW HeNe laser set at 0.1 or 1mW to avoid any thermal effect. Chemical/electrochemical synthetic green rust and in-situ analyses of corrosion products (formed on steel’s or iron’s surface), were analysed using, respectively, a ´100 objective allowing a 2µm2 spatial resolution, and a ´50 objective (5µm2). The confocal hole was kept at 200nm. X-ray EDS analysis (EDAX) was also performed. 2.2. References Raman spectra of green rusts GRs are hydroxy-compound based on a brucite structure, similar to layered double hydroxides (LDH). Metallic cations are located in M(OH)6 octaedra hcp stacked sheets. Some Fe(II) are oxidized to Fe(III). The electrical neutrality is insured by anions and water molecules trapped in the interlayer [1]. Preliminary experiments were run to characterize the Raman spectra of specific green rusts. GR(Cl), GR(CO3) and GR(SO4). GR(Cl) and GR(SO4) were precipitated by mixing acid solutions of 0.2M FeCl2 or 0.2M FeSO4×7H2O (pH 2) with 0.2M NaOH. Chemical GR(CO3) was synthesized by precipitation from a 0.1M Na2CO3 solution, in which 0.01M FeCl2 were dissolved. Electrochemical GR(CO3) was obtained from a 0.01M NaCl solution with 0.3M NaHCO3 at pH 8.4 introduced in an electrochemical cell containing a rod of carbon steel wrapped with filtering paper to ensure confined conditions. The system was maintained under Ar flow and was polarized at -770mV/SCE for 3 days. Then the sample was removed quickly dried and analysed by Raman spectroscopy (figure 4). Raman samples consisted of a drop of each precipitate placed between two glass windows to limit the GR oxidation by air. The Raman spectra of the considered GR are displayed in figures 1, 2, 3 and 4. The frequencies of the Raman bands (cm-1) observed from these four synthetics GRs are summarized in table 1.

In-situ Raman spectroscopy study of iron and carbon steel corrosion...

Relative Intensity

Relative Intensity

599

200

400

600

800

200

1000 1200

400

600

800

1000 1200 -1

-1

Raman Wavenumber (cm )

Raman Wavenumber (cm )

Figure 2: Raman spectrum of cleaned GR(CO3) precipitated at pH 11.

Relative Intensity

Relative Intensity

Figure 1: Raman spectrum of cleaned GR(Cl) precipitated at pH 7.

200

400

600

800

1000 1200

200

-1

400

600

800

1000 1200 -1

Raman Wavenumber (cm )

Raman Wavenumber (cm )

Figure 3: Raman spectrum of a cleaned GR(SO4) precipitated at pH 7.

Figure 4: Raman spectrum of electrochemical GR(CO3) formed on carbon steel.

Table 1: Raman bands (cm-1) observed from synthetics GR(Cl), GR(SO4) and GR(CO3).

Chemical GR(SO4) pH 7

Chemical GR(Cl) pH 7

Chemical GR(CO3) pH 11 1070

Electro-chemical GR(CO3) 1057

988 845, 920 506 433

845, 920, 990 497 429 366 324

220

237

845, 920, 1000 503 429

845, 935, 1020 510 432

260 156, 230

260 156, 220

Band Assignment n1 of CO32n1 of HCO3n1 of SO42Overtone or combination bands of Fe3+-OH/Fe2+-OH A1g Fe3+-OH [2,3] A1g Fe2+-OH Fe-Cl of [Fe(H2O)5Cl]2+ or [Fe(H2O)4Cl2]2+ ions [4] Eg Fe-OH

L. Lanarde et al.

600

2.3. Corrosion behaviour of carbon steel and iron in spring waters To simulate soil conditions, the atmosphere of aerated siliceous water and deaerated carbonated water were respectively controlled with 77%N2, 20%O2, 3%CO2 gas and 90%N2, 10%CO2 gas. A sealed electrochemical cell allowed insitu monitoring of the corrosion products with the Raman spectrometer. The specimens were cut either in X70 steel sheet of composition %wt Mn 1,45-1,66; C 0,08-0,17; V 0,08-0,10; Nb 0;05-0,08, or in a pure iron bar. The exposed surface was polished up to 2400 finish SiC emery paper and then rinsed with deionised water. At the end of experiment, the corrosion products were also characterized by SEM coupled with X-rays EDS analysis. 3. Results and discussion

Relative Intensity

As soon as the samples were inserted in the spectro-electrochemical cell, smallscattered dark spots appeared on the surface of the metal, some of them leading to pits. Those pits were associated with green rust formation, and occurred at both OCP or under CP. Under CP, the pits’ development was slower.

Table 2: X-rays EDS analysis recorded in a pit’s centre on X70 steel’s surface

200

400

600

800

1000 1200 -1

Raman Wavenumber (cm )

Element At% O 7,42

Fe 66,70 S 7,37

Ca 16,06 Mn 1,06

Figure 5: In situ Raman spectrum recorded in the pit’s centre of X70 steel after 1 hour of free corrosion (»-603mV/SCE) in deaerated carbonated water.

On the X70 steel’s surface at OCP/under CP, in both electrolytes, a similar insitu Raman spectrum (figure 5), not yet identified, was recorded in the pits’ centres. It displays a sharp peak at ≈205 cm-1 and a broad peak at ≈280 cm-1. According to X-rays EDS analysis (table 2), this compound could be CaxMnyS. On the iron’s surface, the in-situ Raman spectrum (figure 6), was recorded in pit’s centre at OCP/under CP, in both solutions. The two identified compounds were: fayalite Fe2SiO4 [5] with peaks at 157, 174, 187, 242, 284, 506, 817 coupled to 843 and 902cm-1, and magnetite Fe3O4 with a peak at 675cm-1. Their respective content varied from one pit to another.

601

Relative Intensity

Relative Intensity

In-situ Raman spectroscopy study of iron and carbon steel corrosion...

200

400

600

800

200

1000 -1

600

800

1000

Raman Wavenumber (cm )

Figure 6: In-situ Raman spectrum recorded in the pit’s center of iron after 4 hours of free corrosion (»-410mV/SCE) in deaerated carbonated water.

Figure 7: In-situ Raman spectrum characteristic of GR(CO3) recorded on iron’s surface after 1 day of free corrosion (»-673mV/SCE) in deaerated carbonated water.

first corrosion product formed. It occurred pits (faster in aerated siliceous water). With and thereafter developed on the whole metal 7) is similar to the synthetic electrochemical

Relative Intensity

Relative Intensity

In all conditions, GR was the preferentially in the vicinity of time, GR then covered the pits, surface. The GR formed at OCP (figure GR(CO3).

400

-1

Raman Wavenumber (cm )

200

400

600

800

1000 1200 -1

Raman Wavenumber (cm )

Figure 8: In situ Raman spectrum characteristic of GR(Cl) recorded on X70 steel’s surface after 1 day of cathodic polarization (-925mV/SCE) in aerated siliceous water.

200

400

600

800

1000 1200 -1

Raman Wavenumber (cm )

Figure 9: In situ Raman spectrum characteristic of GR(CO3) recorded on iron’s surface after 1 day of cathodic polarization (-1075mV/SCE) in deaerated carbonated water.

Under CP, pitting was also associated with GRs formation: GR(Cl) in aerated siliceous water and GR(CO3) in deaerated carbonated water. Indeed, in-situ Raman spectra given in figures 8 and 9 are similar to the reference Raman spectra displayed in figures 1 and 4. On the spectrum 8, the two peaks at 336 and 376cm-1 are shifted towards higher frequencies than the GR(Cl) reference Raman spectra. There is also a small peak at 1056cm-1 suggesting the presence of bicarbonate ions in the GR(Cl) structure. Insulating barriers of magnesium

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silicate salt in aerated siliceous water; and calcium carbonate salt (calcite or/and aragonite) in deaerated carbonated water, were identified. These salts covered the whole metal surface after a long-term exposure. Nevertheless, the dissolution of iron into Fe2+ ion was not fully stopped by these salts. Indeed, Siderite (iron II carbonate) or iron II sulfate were found underneath. Evolution of GR towards more stable iron compounds depended on the electrolyte. In aerated siliceous water, GR turned into ferrihydrite 5Fe2O3, 9H2O and in deaerated carbonated water, into magnetite Fe3O4. Both compounds lead to goethite a-FeOOH by slow oxidation.

4. Conclusions In our experiments, regardless to the potential, the solution type, and the metal (iron or steel), the observed metal oxidation was localized, associated with pit formation in the area of impurities. In the case of X70 steel, an unidentified, but reproducible Raman spectrum was recorded on top of the inclusions. EDAX analysis revealed the presence of sulfide with calcium and manganese. In the case of iron, the presence of silicates released by the polishing cannot be excluded. Fayalite formation in the pit centre has to be related to iron dissolution as Fe2+. In all experiments, Green Rusts were observed as transitory compounds. In our study, the main effect of the applied cathodic potential was to slow down the oxidation process. Nevertheless, Fe2+ compounds were still observed under the salt deposit that precipitated as the pH increased due to oxygen and/or water reduction. When cathodic polarization was applied, the deposits formation depleted locally the solution of carbonate, stabilizing the formation of the chloride GR. There is therefore a difference in the corrosion products composition depending on the application of cathodic polarization. References 1 R. Allman, Acta cryst. B24 (1968) p.972. 2 N. Boucherit, A. Hugot-Le Goff, Faraday Discuss., 94, (1992) p.134. 3 N. Boucherit, A. Hugot-Le Goff and S. Joiret, Corr. Sci., 32 (1991) p. 497. 4 P.M.L. Bonin, W. Jedral, M.S. Odziemkowski, R. W. Gillham, Corr. Sci., 42 (2000) p.1921. 5 A. Chopelas, American Mineralogist, 76 (1991) p1101.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Effect of Noble Element Alloying on Passivity and Passivity Breakdown of Ni Yeong Ho Kima*, G. S. Frankela, J. C. Lippoldb a

Fontana Corrosion Center, Dept. of Materials Science and Engineering Dept. Of Industrial, Welding and Systems Engineerig The Ohio State University Columbus, OH 43210, USA b

Abstract - Welding of stainless steels usually generates welding fumes containing carcinogenic hexavalent chromium (Cr+6). To mitigate this problem, a new Ni-Cu-Pd welding consumable alloy has been developed. The addition of small amount of Cu and Pd in Ni enhances the galvanic compatibility of Ni with stainless steels and improves its localized corrosion behavior. In this paper, the artificial pit electrode technique and xray photoelectron spectroscopy (XPS) were used to study the benefits of Cu and Pd alloying. Pd showed the catalyzed the reduction of Cu at the bottom of the artificial pit, which enhanced the cathodic reaction and thus ennobled the deactivation potential making stable pit growth more difficult. Keywords : Cu, Pd, Ni, Noble elements, Breakdown, Artificial pit

1. Introduction Stainless steels (SS) are widely used as construction materials because of their corrosion resistance. Components are often joined by arc welding when fabricated into complex structures. However, the evaporation and oxidation of Cr from the molten weld pools result in the generation of carcinogenic hexavalent chromium (Cr+6) in the welding fume1, 2. The US Environmental Protection Agency is considering a substantial reduction in the permissible exposure limit (PEL) for hexavalent Cr, which will make it difficult to weld SS in locations that are not extremely well-ventilated3. To reduce the emission of

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Cr+6 from weld fumes, we have developed a Cr-free consumable for welding austenitic stainless steel that provides mechanical properties and corrosion resistance comparable to the commonly-used Cr-bearing weld consumables4, 5. Ni-Cu alloys were selected initially based on their galvanic compatibility with types 304 and 316 SS in chloride environments4. It was found that 304L SS can be welded with Monel filler metal, which contains 28-34 wt% Cu, to create high quality welds with no cracks4 The welds pass bend tests and can survive long term exposure to mildly aggressive environments like 0.1M NaCl with no evidence of corrosion4. However, segregated regions rich in Cu were weak spots for corrosion susceptibility in the welds4. Subsequent work on alloys with lower Cu content led to the selection of Ni-10Cu-1Pd, which has improved corrosion properties as a Cr-free filler metal for welding of 304SS5. In this paper, the influence of Cu and Pd on the passivity and passivity breakdown of Ni is addressed. Passive film composition, as determined by XPS, and localized corrosion kinetics, as determined by artificial pit electrode experiments, are discussed to explain the effects of Cu and Pd. 2. Experimental Four Ni alloys were studied: pure Ni, Ni-10Cu, Ni-30Cu, and Ni-10Cu-1Pd (in wt.%). They were melted from pure materials and cast in the shape of small disk. By a series of processes, they were worked to wires with 0.17 or 0.34 mm diameter. The wires were degreased and annealed at 1100oC for 1 h in vacuumsealed quartz tubes. Finally, they were mounted in epoxy. The details of the artificial pit electrode are schematically shown in Figure 1. The samples were polished to #600 grit and then inserted facing upward into a cell that was filled with 0.1 M NaCl. Pt wire was used as a counter electrode and saturated calomel electrode (SCE) as a reference electrode. To evaluate the polarization behavior of artificial pits, three electrochemical steps were combined in series. First, to make an artificial pit of a specific depth, a constant potential of 0.8 VSCE for Ni-30Cu alloy and 0.5 VSCE for the rest of samples was applied. The time to generate a certain pit depth during the potentiostatic test was calculated using Faraday’s Law on preliminary experimental runs. Four different depths of pits on each sample were tested, where the ratio of the pit depth to the pit width was in the range of 0.5 to 4.0. The potential was then stepped down to -0.12 VSCE to dissolve the salt film in the pit6. Finally, potential was scanned from 0.3VSCE to -0.2VSCE at 10mV/min. The depths of the artificial pits were confirmed after the experiments by observation in an optical microscope and by integration of the curves. Samples of each alloy 10 x 10 mm in size were passivated at 0 VSCE in 0.1 M NaCl for 10 min and then analyzed using XPS at 13 kV with the a

Effect of Noble Element Alloying on Passivity and Passivity Breakdown of Ni

605

monochomated X-ray source from an Al target. The sample surface was intermittently sputter etched using Ar ions for 15 or 30 s. The sputtering rate was about 20 Å/min.

I

SCE Cell :Tygon Tube

I

III

0.5VSCE

V

0.3VSCE

CE : Pt

-0.12VSCE

Epoxy Mold

300~185Ksec 30sec

10mV/sec

-0.2VSCE

Time

WE : Wire(0.17mm dia.)

Figure 1. Schematics of the artificial pit electrode and the electrochemical test procedure.

3. Results and Discussion The O 1s spectra measured by XPS analysis indicated that the outermost surface of the passive film for all the tested Ni alloys consisted primarily of hydroxide, and that the inner part of the passive film was mixed oxide and hydroxide. The Pd 3d5/2 and the Cu 2p3/2 and Cu LMM Auger peaks were analyzed to assess the oxidation states of Pd and Cu, respectively. These elements mainly existed in the unoxidized metallic state in the outer layer, but small amounts of PdO and CuO were detected at the outermost layer, Figure 2. 5 10

7000

4

Sputtering Time (s) 4 10

3 10

4

Internsity (Counts)

Intensity (Counts)

180 4

150 120 90 75 60

2 10

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2p3/2

2p

6000

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CuO (Cu,Cu O) 0

920

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960

Binding Energy (eV)

970

Ni-10Cu-1Pd

Pd

Sputtering Time (s)

Ni-10Cu-1Pd

Cu

1000

330

3d5/2

Pd

335

3d3/2

PdO

PdO

2

340

345

350

Binding Energy (eV)

Figure 2. XPS results of Cu 2p and Pd 3d peaks of Ni-10Cu-1Pd alloy, showing some involvement of Cu and Pd in the outer layer

Y. H. Kim et al.

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3/2

OH

2000

500 1000

3/2

CuO

0

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100

(Counts)

Peak Intensity of Cu 2p

1000

0

150

and OH 2p

O

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5/2

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(Counts)

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2

and PdO 3d

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Peak Intensity of Pd 3d ,

2000

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PdO 3d

3/2

Cu

6000

3/2

and CuO 2p

2500

7000

Peak Intensity of O 2p

(Counts)

The heights of the O 1s, Cu 2p3/2, and Pd 3d5/2 peaks of Ni-10Cu-1Pd alloy were analyzed as a function of depth during sputter profiling, Figure 3. It is possible that some Cu +1 is included in the analysis for CuO, but the total amount of oxidized Cu was small. Figure 6 shows that both metallic Cu and Pd were present in the passive film, but in depleted concentrations relative to the content in the bulk alloy.

200 150

PdO

100 50 0

PdO2 0

50

100

150

200

Sputtering Time (sec)

Figure 3. XPS depth-profile results of Ni alloys after the passivation at 0VSCE for 10min.

The beneficial effects of noble Pd additions on the passivity of other Ni-alloys has been reported previously7, 8. However, the results of this study suggest that improvement in localized behavior obtained by the addition of Cu and Pd to Ni is more related to the effects on the anodic and cathodic kinetics in pits than to improvements in the protectiveness of the passive film. 0.4

0.2 0.1 0 -0.1 -0.2 -0.3

Aerated Deaerated

0.3

Potential (VSCE)

Potential (VSCE)

0.4

Pure Ni Ni-10Cu Ni-30Cu Ni-10Cu-1Pd

0.3

0.2 0.1 0 -0.1 -0.2

-8

10

-7

10

-6

10

-5

10

0.0001 0.001 0.01

0.1

1

2

Current Density (A/cm )

Figure 4. Polarization curves for artificial pit electrodes with 0.17 mm diameter and depths in the range of 0.3 to 0.4mm.

-0.3 -7 10

-6

10

-5

10

0.00010.001 0.01

0.1

1

Current Density (A/cm2)

Figure 5. Pit polarization curves for Ni-10Pd1Pd artificial pits of depth 0.34 mm in aerated and deaerated solutions

Figure 4 shows the polarization curves for artificial pit electrodes of 0.17 mm diameter and similar depths in the range of 0.3-0.4 mm, obtained by scanning downward from 0.3 VSCE. As the potential decreased, the pit current decreased

Effect of Noble Element Alloying on Passivity and Passivity Breakdown of Ni

607

until passing through a zero-current potential and then changing polarity. Rather than repassivation involving the formation of a protective passive film, the cessation of pit growth was a deactivation process. The addition of 10% Cu decreased the anodic current relative to pure Ni, resulting in an increase in the deactivation potential. The addition of Cu to 30% provided no further benefit. However, the addition of even 1 wt% Pd resulted in a large increase in deactivation potential as a result of a large increase in the cathodic reaction rate inside the pit. Noble metal alloys have been found previously to cause such increases in cathodic reaction rates9, 10.

Deactivation Potential (mV

SCE

)

The cathodic reaction in the artificial pit might come from reduction of dissolved oxygen molecules, protons, or metallic ions. Figure 5 shows pit polarization curves for Ni-10Cu-1Pd in deaerated and aerated solutions. The curves are identical, indicating that the cathodic reaction was not oxygen. Furthermore, the deactivation potential of Ni-10Cu-1Pd alloy was 0.15 VSCE, which is much higher than the reversible potential of H+/H2, but is very close to the standard potential of Cu2+/Cu, 0.337 VNHE. It is concluded that the primary cathodic reaction in the artificial pit bottom was reduction of Cu ions. The great enhancement in the rate of this reaction for the Ni-10Cu-1Pd alloy relative to the Ni-Cu binary alloys suggests that Pd catalyzed the Cu reduction, thereby ennobling the deactivation process. 300 Pure Ni Ni-10Cu

200

Ni-30Cu Ni-10Cu-1Pd

100 0 -100 -200

0

0.02

0.04

0.06

0.08

0.1

0.12

Attack Depth (cm)

Figure 6. Repassivation potentials of all alloys as a function of the artificial pit depth. Open symbols are 0.34 mm diameter and filled symbols are 0.17 mm diameter wires.

The deactivation potentials of all the samples are plotted in Figure 6 as a function of the artificial pit depth in the range of 0.1-1.1 mm. The deactivation potential tended to decrease as the pit depth increased. For pure Ni and Ni30Cu, the values were higher for the larger diameter wires. However, for Ni10Cu and Ni-10Cu-1Pd, the values were higher for the smaller diameter wires. The Ni-10Cu-1Pd alloy exhibited the highest deactivation potentials over the full range of pit depths. This corresponds to the improvement in breakdown and repassivation potentials measured on sheet samples in bulk solutions5.

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4. Conclusion The effects of Cu and Pd on the passivation and localized corrosion behavior of Ni alloys were investigated using XPS and the artificial pit electrode technique. The passive film on Ni-10Cu-1Pd alloy consisted of a mixture of oxide and hydroxide, with the outermost layer being mainly hydroxide. Most of the Cu and Pd in the passive layer existed in the metallic state. This small change in the composition of passive film cannot explain the improved resistance to the localized corrosion of Ni-Cu-Pd alloy. However, the addition of 1 wt.% Pd increased the deactivation potential considerably over the whole range of the pit depth studied. This effect was attributed primarily to enhanced reduction of copper ions in the presence of Pd. Acknowledgement Financial support of the Strategic Environmental Research and Development Program (SERDP) through project PP-1346 is greatly appreciated. References 1. P.-J. Cunat, Materiaux & Techniques 90, 1-2 (2002) p. 19. 2. S. B. Mortazavi, Welding in the World 39, 6 (1998) p. 297. 3. Welding : Fumes and Gases, Welding : Fumes and Gases, Nov., 1990, p. 1. 4. Y. H. Kim, G. S. Frankel, J. C. Lippold and G. Guaytima, accepted for publication to Corrosion. 5. Y. H. Kim, G. S. Frankel and J. C. Lippold, accepted for publication to Corrosion. 6. N. J. Laycock and R. C. Newman, Corrosion Science 39, 10-11 (1997) p. 1771. 7. P. L. Andresen, S. Hettiarachchi, Y. J. Kim and T. P. Diaz, Patent No. US5608766, (1997). 8. P. L. Andresen, S. Hettiarachchi, Y. J. Kim and T. P. Diaz, Patent No. US5768330, (1998). 9. J. Edwards, Coating and Surface Treatment Systems for Metals, A comprehensive guide to selection, ed. by translated by (Trowbridge, UK, Redwood Books Ltd., 1994) p. 102. 10. N. D. Tomashov and G. P. Chernova, Passivity and protection of Metals against Corrosion, Passivity and protection of Metals against Corrosion, ed. by translated by 1967) p. 82.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Analysis of Electrochemical Noise of Pure Aluminium in Sulphate and Molybdate IonContaining 0.1 M NaCl Solution Su-Il Pyun, Kyung-Hwan Na Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1 Guseong-dong, Yuseong-gu, Daejeon, 305-701, Republic of Korea, E-mail: [email protected] Abstract - In the present work, we investigated the effects of sulphate and molybdate ion addition on the pitting corrosion of pure aluminium as a function of anion concentration by using potentiodynamic polarisation experiment and electrochemical noise measurement. The experimentally obtained electrochemical noise was analysed based upon a stochastic theory. From the results, it was revealed that the localised corrosion process could be distinguished from the uniform corrosion process by the mean free time of events in the electrochemical noise. In addition, it was proposed that both sulphate and molybdate ions retards the metastable pitting or the pit initiation. Keywords: aluminium, sulphate, molybdate, electrochemical noise

1. Introduction Aluminium is susceptible to pitting corrosion in the presence of chloride ions. It has been generally known that pit initiation is regarded as a stochastic process at the native oxide film [1-3]. Since the fluctuations in potential or current associated with the pitting process can be easily observed in the experimental data, electrochemical noise measurement has been widely used for monitoring, as well as studying corrosion processes [4-8]. However, little attention has been paid on the analysis of electrochemical noise based upon a stochastic theory. Thus, the present work is aimed at investigating the effects of SO42- and MoO42ion additives on the pitting corrosion of pure Al as a function of anion concentration by using potentiodynamic polarisation experiment and electrochemical noise measurement. For this purpose, Weibull probability plot was constructed on the basis of a stochastic theory to resolve the noise signal.

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The change in linear slope in one Weibull probability plot with mean free time has been discussed in terms of the coupled processes of uniform and localised corrosion, which comprise the noise signal. 2. Experimental The test specimen used in this work was high purity Al rod of 6.35 mm diameter (99.999%, Aldrich Chemical Co., Inc.). The specimen was set in a block of polyethylene and the upper surface of the block was ground with silicon carbide papers to 1500 grit to expose the cross section of the electrode to the electrolytic solution. The electrolytes used in this study were 0.1 M NaCl solutions containing various Na2SO4 and Na2MoO4 concentrations of 0, 0.01 and 0.1 M at room temperature. The pitting potential (Epit) of Al in the absence of SO42- and MoO42- ions was compared with those in the presence by potentiodynamic polarisation measurement. The measurement was made at a scan rate of 0.5 mVs-1 in the applied potential range from -1.2 to +0.4 VSCE by using EG&G Model 263A Galvanostat/Potentiostat. A platinum gauze and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. Electrochemical noise measurement was carried out with a Zahner IM6e impedance measurement unit equipped with Zahner NProbe: Two identical Al specimens were galvanically coupled. After that, both the current between the two Al specimens and the potential of the Al specimens against a SCE were simultaneously recorded with time for 48h. The sampling interval used in this study was 1s. The cumulative probability of the frequency of events (fn) was calculated from the experimentally obtained noise for the stochastic analysis. 3. Results and Discussion Fig. 1 shows potentiodynamic polarisation curves for pure Al at a scan rate of 0.5 mVs-1 in aqueous 0.1 M NaCl solution containing various SO42- and MoO42ion concentrations of 0, 0.01 and 0.1 M. It is noted that Epit shifted to more positive values as SO42- and MoO42- ion concentrations increased. This indicates that pit initiation is suppressed by the addition of SO42- or MoO42- ions. If we assume that shot noise is produced in the present system [9, 10], the average corrosion current I corr is given by I corr = qf n (1) where q is the average charge in each event and f n the frequency of events. Since it is not possible to measure I corr , q and f n directly from the noise data, each parameter should be estimated from the following relations: I corr = B / Rn = Bs I / s E (2) q = s I s E / Bb (3) f n = I corr / q = B 2 b / s E2 (4)

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where B is the Stern-Geary coefficient, Rn the noise resistance, s I the standard deviation of current, s E the standard deviation of potential and b the bandwidth of measurement. 0.4

0.4

0.1 M NaCl 0.1 M NaCl + 0.01 M Na2SO4 0.1 M NaCl + 0.1 M Na2SO4

0.0

-0.2

Epit=-0.57 VSCE

-0.4

Epit=-0.64 VSCE -0.6

Epit=-0.70 VSCE

0.0

-0.2

Epit=-0.60 VSCE

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-1.2 -10 10

-1.2 -10 10

0

10

(b)

0.1 M NaCl 0.1 M NaCl + 0.01 M Na2MoO4 0.1 M NaCl + 0.1 M Na2MoO4

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Applied Potential / VSCE

0.2

A pplied Potential / VSCE

(a)

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-5

10

-4

10

-3

-2

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10

-1

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0

10

-2

Current Density / Acm

Current Density / Acm

Fig. 1. Potentiodynamic polarisation curves of pure Al with a scan rate of 0.5 -1 2mVs in aqueous 0.1 M NaCl solution containing various (a) SO4 and 2(b)MoO4 ion concentrations. 1.0

1.0

(b) Cumulative Probability of Events / -

Cumulative Probability of Events / -

(a) 0.8

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0.1M NaCl 0.1M NaCl + 0.01M Na2SO4 0.1M NaCl + 0.1M Na2SO4

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Frequency of Events / Hz

Fig. 2. Cumulative probability plots for the frequency of events, fn in aqueous 220.1 M NaCl solution containing various (a) SO4 and (b)MoO4 ion concentrations. Fig. 2 presents the cumulative probability plots for f n in aqueous 0.1 M NaCl solution containing various SO42- and MoO42- ion concentrations. It was found that the distribution of f n became wider with increasing SO42- ion concentration. On the other hand, the distribution of f n shifted to a lower

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frequency region by the addition of MoO42- ions. Considering that high frequency events will tend to occur all over the Al surface, the corrosion of Al will be more uniform as high frequency events become dominant. In contrast, the corrosion will be rather localised over the surface as low frequency events are dominant. Thus, f n is a useful parameter for the evaluation of the corrosion type. In this respect, it is conceivable that the corrosion of Al in the presence of MoO42- ions tends to be rather localised compared with that in the absence. However, in the case of SO42- ions, it was inferred that both uniform and localised corrosions occurred more actively than those in absence. 1.0

1.0

0.8

0.6

0.4

0.2

(b) Cumulative Probability of Events

Cumulative Probability of Events

(a)

0.1M NaCl 0.1M NaCl + 0.01M Na2SO4 0.1M NaCl + 0.1M Na2SO4

0

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Mean Free Time / ms

1.6

2.0

0

0.5

1.0

1.5

2.0

Mean Free Time / ms

Fig. 3. Cumulative probability plots for mean free time, t in aqueous 0.1 M 22NaCl solution containing various (a) SO4 and (b)MoO4 ion concentrations. In an attempt to investigate low frequency events associated with localised corrosion in more detail, the cumulative probability plots for mean free time t were reproduced from Fig. 2 and are given in Fig. 3. In the presence of SO42ions, it was found that the distribution of t became wider up to the SO42- ion concentration around 0.01 M and then narrower again with increasing SO42- ion concentration This means that localised corrosion is dominant at the low SO42ion concentrations around 0.01M. However, uniform corrosion becomes dominant with increasing SO42- ion concentration. In the presence of MoO42ions, it was observed that the distribution of t became wider with increasing MoO42- concentration. This indicates that the corrosion of Al is rather localised in the presence of MoO42- ions. In general, the cumulative probability F(t) of a failure system based upon a “weakest-link” model is described as a Weibull distribution function, which is expressed as [1-3] F (t ) = 1 - exp(-t m / n) (5)

Analysis of Electrochemical Noise of Pure Aluminium

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where m [-] and n [sm] are the shape and scale parameters, respectively. From rearrangement of Eq.(5) ln{ln[1 /(1 - F (t ))]} = m ln t - ln n (6) The two parameter m and n can be determined from the slope of the linear ln{ln[1 /(1 - F (t ))]} vs. ln t plot (Weibull probability plot) and from the intercept on the ln{ln[1 /(1 - F (t ))]} axis, respectively. 3

3

(b)

2

2

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1

0

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ln{ln[1/(1-F(t))]}

ln{ln[1/(1-F(t))]}

(a)

-1

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0.1M NaCl 0.1M NaCl + 0.01M Na2SO4 0.1M NaCl + 0.1M Na2SO4

-4

-5

-6 -25

-20

-15

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0

-1

-2

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0.1M NaCl 0.1M NaCl + 0.01M Na2MoO4 0.1M NaCl + 0.1M Na2MoO4

-5

-6 -25

5

ln (Mean Free Time t)

-20

-15

-10

-5

0

5

ln (Mean Free Time t)

Fig. 4. Weibull probability plots (plots of ln {ln [1/(1-F(t))]} vs. ln t) in aqueous 220.1 M NaCl solution containing various (a) SO4 and (b)MoO4 ion concentrations. Fig. 4 depicts the Weibull probability plots in aqueous 0.1M NaCl solution containing various SO42- and MoO42- ion concentrations. The plots showed satisfactorily good straight lines. The two slopes in one plot are likely to indicate that two failure modes exist, depending upon mean free time t. Considering that only uniform and localised corrosion will occur during the noise measurement, it is suggested that the slopes in the relatively shorter mean free time region are associated with uniform corrosion. On the other hand, the

100 (a)

100 (b)

100 (c)

Fig. 5. Optical micrographs of surface morphology of pure Al after the electrochemical noise measurement in: (a) 0.1 M NaCl; (b) 0.1 M NaCl + 0.01 M Na2SO4 and (c) 0.1 M NaCl + 0.1 M Na2SO4 solution.

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slopes in the longer mean free time region are responsible for localised corrosion such as metastable pitting or pit initiation.

100 (a)

100 (b))) )

100 (c)

Fig. 6. Optical micrographs of surface morphology of pure Al after the electrochemical noise measurement in: (a) 0.1 M NaCl; (b) 0.1 M NaCl + 0.01 M Na2MoO4 and (c) 0.1 M NaCl + 0.1 M Na2MoO4 solution.

Figs. 5 and 6 show the optical micrographs of surface morphology of pure Al after the electrochemical noise measurement in 0.1 M NaCl solution with various SO42- and MoO42- ion concentrations, respectively. It was observed that the number of pits was reduced with increasing SO42- or MoO42- ion concentrations. This means that both SO42- and MoO42- ions inhibit the metastable pitting or the pit initiation. 4. Conclusion From the analysis of electrochemical noise based upon a stochastic theory, it was found that localised corrosion process could be distinguished from uniform corrosion process in the electrochemical noise by the mean free time of events. It is suggested that a shorter mean free time of events may be associated with uniform corrosion. On the other hand, a longer mean free time of events is responsible for localised corrosion. In addition, it was revealed that both SO42and MoO42- ions retard the metastable pitting or the pit initiation of Al. References 1. S.-I. Pyun and E.-J. Lee, Surf. and Coat. Technol., 62 (1993) 480. 2. S.-I. Pyun, E.-J. Lee and G.-S. Han, Thin Solid Films, 239 (1994) 74. 3. J.-J. Park and S.-I. Pyun, Corros. Sci., 46 (2004) 285. 4. P.R. Roberge, S. Wang and R. Roberge, Corrosion, 52 (1996) 733. 5. J.W. Isaac and K.R. Hebert, J.Electrochem. Soc., 146 (1999) 502. 6. C.B. Breslin and A.L. Rudd, Corros. Sci., 42 (2000) 1023. 7. Y. Kobayashi, S. Virtanen and H. Böhni, J. Electrochem. Soc., 147 (2000) 155. 8. K. Sasaki, P.W. Levy and H.S. Isaacs, Electrochem. Solid-State Lett., 5 (2002) B25. 9. R.A. Cottis, Corrosion, 57 (2001) 265. 10. H.A.A. Al-Mazeedi and R.A. Cottis, Electrochim. Acta, 49 (2004) 2787.

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Localized corrosion of 2024 alloy: structure and composition of oxide films grown on model alloys representative of the different metallurgical phases Christine Blanc,a,* Georges Mankowski,a Corinne Dufaure,a Claude Mijoule,a and Yolande Kihn,b a

CIRIMAT – UMR CNRS 5085 – ENSIACET – 118 route de Narbonne 31077 Toulouse Cedex 04 - France b CEMES – UPR CNRS 8011 – BP 4347 – 29 rue Jeanne Marvig 31055 Toulouse Cedex 04 – France *corresponding author : [email protected]

Abstract - Structure and composition of passive films grown on model alloys representative of the metallurgical phases present in AA2024-T351 alloy were determined. It was found that copper enters the passive films formed on the different metallurgical phases (matrix or intermetallics). The other alloying elements (Fe and Mg) can also be introduced into the passive films which might explain the differences in the electrochemical behavior for all these phases. Keywords: localized corrosion, passive films, copper, aluminium, intermetallics

1. Introduction In Cu-containing Al alloys, Cu is present in solid solution in the matrix. Fine strengthening phases and coarse intermetallic particles are also enriched with Cu. In appropriate media, localized corrosion such as pitting can occur due to galvanic coupling between the matrix, the particles and the bordering region [12]. The nucleation step of pitting corrosion corresponds to passive film rupture and galvanic coupling are probably due to local differences between the passive film grown on the Al matrix and that on the intermetallics [3]. Thus, to understand pitting mechanisms, it is necessary to have a good knowledge of the corrosion properties of the passive films grown on the different metallurgical

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phases present in Al-Cu alloys. The electrochemical behavior of these oxide films is related to their structure and chemical composition. In order to understand further the electrochemical behavior of the different phases, examination of Al-Cu-X (with X = Mg or Fe) model alloys allows separation of the individual effects of the alloying elements [4-5]. Here, the electrochemical behavior of model alloys representative of the different metallurgical phases present in 2024-T351 alloy was studied. This last alloy is characterized by an Al matrix containing about 0.02wt%Cu and two types of coarse intermetallic particles i.e. Al-38wt%Cu-16wt%Mg and Al-27wt%Cu-7wt%Mn-11wt%Fe. Including the fine strengthening particles, the mean Cu content of the matrix is about 4wt% [6]. Thus, in this study, Al-4wt% Cu, Al-55wt%Cu-10wt%Mg and Al-35wt%Cu-10wt%Fe model alloys were deposited by magnetron sputtering. Polarization curves were plotted in a sulfate solution. The alloys were also polarized in sulfate solution and the oxide films grown on the surface were observed by Transmission Electron Microscopy (TEM) and analyzed by Electron Energy Loss Spectroscopy (EELS). In order to understand at the atomic level the influence of Cu incorporation in Al alloys on the oxide film growth, quantum calculations were also performed on the adsorption energy of atomic oxygen on pure Al clusters and bimetallic Cu/Al systems by using the density functional theory (DFT) approach. 2. Experimental Binary Al-4wt% Cu alloy, ternary Al-35wt%Cu-10wt%Fe and Al-55wt%Cu10wt%Mg alloys were prepared by magnetron sputtering with separate high purity Al (99.999%) and Cu (99.99%) targets. The substrate, onto which the alloys were deposited, consisted of mechanically polished (up to 1 µm diamond paste) 2017 Al alloy. The alloy layers were deposited at about 5 nm.min-1 to a final thickness in the range 150 – 250 nm. The chamber was first evacuated to 2.10-7 mbar, with sputtering then carried out at 5.10-3 mbar in 99.998% argon. The deposited alloys (without further polishing) were then potentiokinetically polarized at a scan rate of 15 mV.min-1 in a 0.1M sodium sulfate electrolyte at 293 K. All potentials quoted are relative to the saturated calomel electrode. Further experiments consisted in polarizing model alloy samples at 1000 mV/SCE for 1 hour in 0.1M Na2SO4 solution to make an oxide film grow. Suitable electron transparent samples of freshly deposited alloys and anodized alloys were examined by TEM using a Philips CM20T instrument with EDX and EELS facilities. The samples were first sectioned to a nominal thickness of about 30 nm by ultramicrotomy. Compositions of the alloys and anodic films were determined by using EELS. DFT computations were performed with the GAUSSIAN 98 package [7]. The metallic clusters used were three-layers systems containing 31 atoms, namely the Al31 and three different CuAl30 clusters (atomic copper concentration : 3.2%). They were built in order to simulate the (111) plane of the infinite

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surface. The O atom is adsorbed on the metallic surface in a threefold adsorption site. The O-surface distance was optimized perpendicularly to the surface and the O binding energy (BE) calculated as follows: BE = ( EM +EO ) -E M-O with M = Metallic cluster (Al31 or CuAl30) The influence of the spin multiplicity of the studied systems was taken into account. A Mulliken population analysis was performed in order to estimate the magnitude of the electronic transfers between the oxygen atom and the clusters. 3. Results and discussion TEM observations showed that model alloys representative of the different metallurgical phases present in 2024 alloy consisted of thin films homogeneous in thickness (150 nm thick for Al-Cu model alloy and about 250 nm thick for Al-Cu-Mg and Al-Cu-Fe model alloys). EDS analysis showed that, for the model alloy representative of the matrix, the chemical composition was Al4wt%Cu. For the model alloys representative of the intermetallics, the Cu content was too high in comparison with the Cu content in the particles; moreover, our experimental device for magnetron sputtering only allowed the preparation of ternary alloys. So, the intermetallic particles present in 2024 alloy were modelled by Al-55wt%Cu-10wt%Mg and Al-35%Cu-10%Fe model alloys. These alloy layers allowed to take into account the influence of the incorporation of Cu, Mg and Fe into the aluminum on the corrosion behavior of this material. Electronic diffraction experiments showed a nanocrystalline structure with a medium grain size less than 100 nm for Al-Cu alloy while an amorphous structure was obtained for Al-Cu-Mg and Al-Cu-Fe alloys probably due to the introduction of a third alloying element in significant proportions. Due to the influence of passive film properties (structure and chemical composition [3]) on the corrosion behavior of materials, passivity of model alloys was studied. Fig. 1 shows TEM observations of the passive films grown on model alloys polarized for one hour in 0.1M Na2SO4 at 1V/SCE. For Al-Cu alloy, the passive film observed was about 20 nm thick (Fig. 1(a)) and presented an amorphous structure as shown by the diffraction experiments. Both EDS analysis and further EELS analysis showed that this oxide film was predominantly composed with alumina slightly enriched with Cu with a Cu to Al ratio in the alumina film much lower than in the Al-Cu alloy. Recent studies showed that anodizing of Al-Cu alloys in ammonium pentaborate solutions proceeds in two stages [8]: in a first stage, formation of a Cu-free alumina film with Cu accumulating near the alloy/oxide interface and, in a second stage, when a sufficient concentration of Cu in the Cu-enriched interface has been achieved, introduction of Cu into the film with generation of oxygen gas near the alloy/film interface. This led to an increase of the resistance of the film to ion transport: thus, Cu incorporation in Al had a detrimental effect on the anodic film growth. Quantum calculations performed on the O adsorption on pure Al31 and bimetallic CuAl30 systems showed that the presence of the Cu atom affects

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the electronic transfer that occurs during the adsorption process (net atomic charge on O : -0.99 e and -0.82 e when O is adsorbed on Al31 and CuAl30 respectively), leading to a lower binding energy (Table I). Moreover, ELNES (Energy Loss Near Edge Structure - part of the EELS spectrum obtained at high resolution) profiles of the Cu L23-edge were plotted and compared with ELNES profiles obtained for reference samples of metallic Cu and Cu oxides. They showed that Cu was present as metallic Cu at the interface alloy/passive film and as oxidized Cu in the film. Metallic Cu present at the interface alloy layer/anodic film might correspond to the Cu enriched layer described previously but this Cu enriched layer was not observed in this work with TEM. ELNES analyses show that the Cu enriched layer might also exist in this case. (a)

Oxide film

200 nm (c)

50 nm (b)

Oxide film

Oxide film Alloy

Alloy 200 nm

200 nm

Fig. 1: TEM observations of the passive films formed by an hour's polarization at 1V/SCE in a 0.1M Na2SO4 solution of (a) Al-4wt%Cu (b) Al-55wt%Cu-10wt%Mg and (c) Al-35wt%Cu10wt%Fe model alloys. Table I : O binding energy (BE in eV) on Al31 and CuAl30 clusters (a : distance O-atom n°1, 2 or 3)

Clusters Al31 CuAl30 CuAl30 CuAl30

Cu position / 1 4 5

D O-atoma (Å) 1.855 1.894 1.874 1.856

BE (eV) 5.59 5.34 5.40 5.15

The passive films formed on Al-Cu-Mg and Al-Cu-Fe model alloys were much thicker than that formed on Al-Cu alloy with a thickness of about 150-200nm (Fig. 1 (b) and (c)). For both alloys, the passive film was found to contain a lower Cu content by comparison with the alloy layers. The passive film formed on Al-Cu-Fe model alloy was also enriched with Fe, with a higher Fe content (about 34wt%) than in the alloy (Fig. 2). Magnesium was also detected in the passive film grown on Al-Cu-Mg alloy with a lower content in the outer part of

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the film by comparison with the inner part of the film which was explained by the high diffusion rate of Mg in the alumina film and its dissolution. To sum up, whatever the model alloy, Cu oxidation was observed but the thicknesses of the anodic films were very different for Al-Cu alloy on one hand and for Al-Cu-Mg and Al-Cu-Fe alloys on the other hand. Maybe, the incorporation of another alloying elements, i.e. Mg and Fe, could explain the differences. One could also suggest that the crystallographic structure (nanocrystalline or amorphous) of the alloy layer could significantly influence the passive film growth. O-K Outer zone of the oxide film

Cu-L23

a.u.

Fe-L23

Al-K

substrat

400

600

800

1000

1200

1400

1600 eV

Fig. 2: EELS spectra for Al-35wt%Cu-10wt% Fe alloy polarized in 0.1M sulfate solution: oxygen K-edge, iron L23-edge, Cu L23-edge and Al-K edge

The differences observed in the chemical composition of the passive films formed on the three model alloys can explain the differences observed on the current-potential curves plotted in a 0.1M Na2SO4 solution. The currentpotential curve plotted for AA2024-T351 alloy was also reported for comparison. Al-Cu-Mg and Al-Cu-Fe model alloys presented similar electrochemical behavior both in the cathodic range and in the anodic range with cathodic current densities close to that measured for 2024 alloy (about 10-4 A.cm-2) and significantly higher than for Al-Cu alloy. Thus, the corrosion behavior of 2024 alloy was significantly influenced by the Al-Cu-Mg and AlCu-Fe intermetallic particles in the cathodic range and the cathodic current densities for the coarse intermetallics in 2024 alloy were high enough to determine the corrosion behavior of 2024 alloy. Thus, the corrosion potential of 2024 alloy was similar to those measured for Al-Cu-Mg and Al-Cu-Fe model alloys (Ecorr = -50 mV/SCE). By comparison with Ecorr measured for Al-Cu model alloy, the corrosion potential of Al-Cu-Mg and Al-Cu-Fe alloys was 150 mV shifted towards the more positive potentials. In the anodic range, the

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Current density (A/cm2)

electrochemical behavior of 2024 alloy was similar to that of its matrix since the anodic current densities measured for 2024 alloys and Al-Cu alloy were similar and significantly lower than those measured for Al-Cu-Mg and Al-Cu-Fe model alloys. This result suggested that the passive films formed on the intermetallics were less protective than that on the Al matrix, probably due to the chemical differences observed. However, the influence of these particles on the anodic behavior of 2024 alloy appeared much lower than in the cathodic range. 10-3

AlCuFe

10-4

AlCuMg Al 2024

10-5 AlCu

10-6 10-7 -800

-400 0 +400 Potential (mV/SCE)

+800

Fig. 3: Current density-potential curves for model alloys and 2024 alloy in 0.1M Na2SO4 solution

4. Conclusions Structure and composition of passive films grown on various metallurgical phases are significantly dependent on the chemical composition of the particles. For 2024 alloy, Cu enters the passive films formed on the matrix and on the intermetallics. The other alloying elements (Fe and Mg) can also be introduced into the passive films which might explain the differences observed in the electrochemical behavior for all these metallurgical phases. References 1. 2. 3. 4. 5. 6. 7. 8.

P. Schmutz and G.S. Frankel, J. Electrochem. Soc., 145 (1998) 2285. P. Schmutz and G.S. Frankel, J. Electrochem. Soc., 145 (1998) 2295. P. Schmutz, G.S. Frankel, J. Electrochem. Soc., 146 (1999) 4461. H. Habazaki, M.A. Paez, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros. Sci., 38 (1996) 1033. S. Garcia-Vergara, P. Skeldon, G.E. Thompson, P. Bailey, T.C.Q. Noakes, H. Habazaki, K. Shimizu, Appl. Surf. Sci., 205 (2003) 121. C. Blanc, B. Lavelle and G. Mankowski, Corros. Sci., 39 (1997) 495. M. J. Frisch et al., Gaussian 98, Revision A.3, Gaussian, Inc., Pittsburgh PA, 1998. S. Garcia-Vergara, F. Colin, P. Skeldon, G.E. Thompson, P. Bailey, T.C.Q. Noakes, H. Habazaki and K. Shimizu, J. Electrochem. Soc., 151 (2004) B16.

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Passivity Breakdown of Aluminum Alloys by Trace Element Lead Kemal Nişancıoğlu ([email protected]), Øystein Sævik and Yingda Yu Norwegian University of Science and Technology, Department of Materials Technology, N-7491 Trondheim, Norway

Abstract. Mechanisms of surface enrichment of trace element Pb as a result of heat treatment and ensuing passivity breakdown of the surface in chloride solution are investigated by use of electrochemical and surfaceanalytical characterization of model binary AlPb alloys containing 20 and 50 ppm Pb. Release of mobile Pb adatoms from a metallic sublayer, which is enriched in Pb as a result of heat treatment, appears to activate the surface anodically. The presence of chloride in the solution is necessary for activation, which occurs in the form of pitting potential depression relative to pure aluminum. Once segregated in the form of visible submicron and larger size particles, Pb no longer participates in the activation process. Keywords: Aluminum, lead, activation, dealloying, heat treatment

1. Introduction Thermomechanical processing has a significant role in the determination of the surface properties of aluminum alloys. Electrochemical activation of various alloys resulting from high temperature heat treatment has become a subject of attention recently1-3 because of its importance in galvanic and filiform corrosion. Anodic activation is characterized by a significant lowering of the pitting potential in chloride media in relation to the usual pitting or corrosion potential of aluminum of about - 0.75 VSCE. Surface repassivates by removal of the active layer, which is a fraction of a mm thick on rolled products, by mechanical or chemical means. The cause of activation was related to the enrichment of the material surface by lead up to a concentration of the order 1 wt%, while it is present in the bulk of the material as a trace element only at the ppm level. The present work was undertaken to obtain a more detailed chemical and microstructural information on model binary AlPb alloys for exploring the mechanism of activation.

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2. Experimental The binary AlPb alloys used in this work, containing 20 and 50 ppm Pb and designated as AlPb20 and AlPb50, respectively, were prepared from high-purity aluminum (99.98%) and lead (99.99%). The alloys were cast and subsequently scalped and cold rolled. Samples were metallographically polished through 1 µm diamond paste. This was followed by heat treatment at 600°C in air and subsequently quenching in distilled water at room temperature. Sample surfaces were characterized by glow discharge optical emission spectrometry (GD-OES), TEM analysis of cross-sectional foils and FEG-SEM. Electrochemical characterization was performed in 5% NaCl solution exposed to ambient air at 25 ±1°C and stirred by a mechanical stirrer at a fixed rate. Polarization measurements were performed at a potential sweep rate of 0.1 mV/s. Further details about the experimental apparatus and procedure for the electrochemical tests are given in reference 4. 3. Results Elemental depth profiles obtained by GD-OES for the annealed AlPb20 alloy verify earlier data2-4 insofar as significant Pb enrichment of the surface, including the oxide, oxide-metal interface and the metallic subsurface layer, as shown in figure 1. The data shown were obtained from two samples which were heat treated at the same temperature for different periods of time. Judging from the shape of Al and O profiles, the Pb enrichment was thought to be at a maximum at the metal-oxide interface, and the enriched subsurface region extended a certain distance into the metal. The only change detectable between 1 h and 4 h of annealing was the increase in the thickness of the Pb-enriched region in the metal phase with the time of annealing. The peak concentration, which was about 0.8 wt%, did not appear to increase significantly with annealing time.The wider Pb peak on the 4 h - annealed sample can be attributed to segregated Pbcontaining particles or nonuniform sputtering during depth profiling.

Figure 1. GD-OES elemental depth profiles for alloy AlPb20 heat treated in air at 600°C for different periods (1 h and 4 h). The Pb profile for the polished specimen is added for comparison.

Secondary electron images obtained by FEG-SEM at an accelerating voltage of 15 kV indicated a porous oxide structure on sample AlPb20 heat treated at 600ºC. The porous topogra-

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phy was typical for all samples heat treated at high temperatures including pure aluminum. The porous structure resulted from the gaps between the g-Al2O3 crystals, which contained amorphous aluminum oxide, as is well known from earlier studies of thermally-grown oxides on aluminum.5 Lead was not visible by backscattered electron FEG-SEM analysis of the 1 h - annealed specimen. On the 24 h - annealed specimen, however, the Pb particles filling the pores were readily visible. TEM nanoprobe EDS analysis (nominally 6 nm lateral resolution) of alloy AlPb50 indicated Pb enrichment close to the noise level of the instrument attributed to the presence of Pb in solid solution with the aluminum matrix in the subsurface region, since no segregated particles were detectable in most cases. Lead segregation in the form of 5-10 nm discrete particles was detectFigure 2. HRTEM images of the segregated Pb able only on a few samples as particles on AlPb50 specimen annealed for 1 h at shown in figure 2. The lattice 600°C. plane value was determined accurately as 0.286 nm by using the (111) lattice interplanar spacing (0.453 nm) of g-Al2O3 crystals in the same image as internal reference. This value agreed exactly with the (111) lattice interplanar spacing of metallic Pb. Inspection of Pb particles in several micrographs of type figure 2 revealed that a few were in contact with the aluminum metal, while many were detached from the metal surface with an oxide layer between the particle and the metal surface.

Figure 3. Polarization data for pure aluminum and alloy AlPb20 in 5 wt% NaCl solution and 5 wt% Na2SO4 solution.

Figure 3 shows typical anodic polarization curves for pure aluminum and alloy AlPb20 in 5% NaCl solution. The figure includes also polarization data for alloy AlPb20 in 5% Na2SO4 solution. The data for polished AlPb20 are without heat treatment, whereas all other curves represent heattreated samples. In chloride

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solution, pure Al and non-heat-treated AlPb20 exhibited similar polarization behavior insofar as anodic behavior above the pitting potential (> -0.75 VSCE) of pure aluminum is concerned. However, the AlPb alloy had a much higher apparent passive current than the pure Al in the range –0.9 to –0.7 VSCE. Observation of the surface under an optical microscope revealed slight pitting in this range, indicating that the surface was slightly activated anodically. Thus, the polished sample appeared to exhibit two pitting potentials, the lower one at about -0.9 VSCE corresponding to the active surface and the higher one at about 0.75 VSCE to the bulk. In the heat treated condition, the pitting potential, which was identical with the corrosion potential, was shifted in the negative direction relative to the pitting potentials of pure Al and non-heat-treated AlPb. The heat-treated AlPb sample exhibited a high oxidation current at the potential range where the other specimens were passive, indicating a significantly active surface relative to both the non-heat treated state and pure Al. Verifying an earlier more detailed study on the effect of chloride,1 activation due to Pb enrichment was not possible in chloride-free solutions of neutral pH, as demonstrated here by the sulfate solution. In this solution, no pitting was observed on the heat-treated AlPb20 alloy. Increasing heat treatment time above 1 h at 600ºC did not affect the polarization curve significantly. The results were not significantly different for Pb concentration in the range 5-50 ppm investigated in the binary model alloy system. More detailed studies of the electrochemical behavior of various alloys can be found in references 1 through 4. The present data verify the earlier results. The surface passivated as the active layer was etched away from the surface. The amount of charge passed to etch away the active layer by potentiostatic polarization, obtained by integrating the measured current-time transients, was commensurate with the thickness of the Pb enriched layer estimated by 10 mm GD-OES depth profile data (figure 1).

1 mm

Pb

1 mm

Figure 4. FEG-SEM images of alloy AlPb20 heat treated for 1 h at 600°C and water quenched after potentiodynamic polarization to -0.95 VSCE.

As revealed by controlled etching of the surface as shown in figure 4, corrosion was initiated at the grain boundaries (figure 4a) and in the form of small pits (figure 4b), distributed uniformly over the surface. Magnification of the corroded areas in figure 4a revealed trenched areas (figure 4c), corresponding to the grain boundaries, where corrosion of the active layer

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probably initiated. However, further propagation of corrosion appeared in the form of superficial, nearly two dimensional attack surrounding deeper trenched zones. Segregated Pb particles, which were not visible on the uncorroded surface, became visible at the outer boundaries of shallow corroded zones (figure 4c and d). Since no segregated Pb was observable on the heat treated and water quenched samples, the Pb observed at the corrosion front of the corroded samples was attributed to Pb segregated as a result of selective dissolution of aluminum during the polarization experiment. 4. Discussion The results demonstrated that metallic Pb can become enriched and segregated at the metal-oxide interface as a result of high temperature heat treatment in air. Independence of electrochemical activation on the bulk Pb concentration in the range 5 to 50 ppm for samples annealed at 600ºC indicated that the metal surface was saturated with Pb in solid solution with aluminum in a metallic subsurface layer of a fraction of a m thick on the average. On most 1 h - annealed (600ºC) samples, Pb enrichment could be verified by EDS analysis at the metaloxide interface and at the grain boundaries very near the metal-oxide interface, however, seldom observing any segregated Pb particles, except after prolonged heat treatment at 600ºC. We therefore believe that the Pb detected in the subsurface was in most cases in solid solution with aluminum. After prolonged heat treatment, i.e., 24 h, significant Pb segregation was observed. However, the electrochemical data did not indicate increased surface activation as a result of Pb segregation in the form of particles. This is additional evidence to the effect that Pb in solid solution with the Al metal in the metallic subsurface region, and not the segregated Pb particles, was responsible for activation. Annealing for 1 or 24 h at 600ºC did not affect the anodic polarization curves since the subsurface concentration of Pb, which controlled activation, was at the same saturation level in both cases. The Pb in solid solution was released again as mobile surface adatoms as a result of selective dissolution of aluminum by polarization in chloride media. Qian et al.6 claim that surface activator atoms increase the dissolution probability of adjacent aluminum atoms by breaking the Al-O-Al bonds. One activator atom can thus catalyze the dissolution of a large number of aluminum atoms as it moves around on the surface. This process must be enhanced by the presence of chloride ions in the solution. Increased enrichment of Pb by selective corrosion of Al presumably caused clustering and coarsening of surface Pb into particles. Since the surface Pb loses its mobility in this manner, the Pb particles could not be as effective in taking part in the activation process. Although corrosion of the active layer appeared as superficial etching, the mechanism was essentially pitting since corrosion of the active layer did not occur in the ab-

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sence of chloride. Anodic activation of the surface, therefore, occurred in the form of a reduction in the pitting potential. Pitting initiated at the grain boundaries, since the oxide was probably weakest and Pb was apparently most concentrated at the grain boundaries because these were the diffusion paths to the surface. Pitting, which was confined to the Pb-enriched sublayer, propagated faster in the lateral direction, giving an etching-type of morphology. Corrosion in the vertical direction was arrested by the Pb-lean bulk aluminum. The surface was passivated after the Pb-rich subsurface layer was corroded. A more localized, conventional type of pitting could be initiated again by only increasing the potential above the pitting potential of pure aluminum (-0.76 VSCE). 5. Conclusions Enrichment of the oxide-aluminum interface by lead as a result of high temperature heat treatment was demonstrated. Lead enrichment at the metal surface occurred both in solid solution with the aluminum matrix in a metallic subsurface layer and in the form of segregated metallic Pb nanoparticles. Pb in solid solution with aluminum in the metallic subsurface layer was responsible for surface activation and not the segregated Pb particles. Pb in solid solution was released as mobile Pb adatoms which reduced the passivity of the oxide in the presence of chloride in the solution and, thereby the surface was activated. Once segregated in the form of visible submicron size and larger particles, either by selective corrosion of aluminum or as a result of high temperature heat treatment, Pb no longer participated in the activation process. 6. Acknowledgments This work was part of a Norwegian national research program entitled "Light Metal Surface Science", supported by The Norwegian Research Council, Hydro Aluminium, Profillakkering AS, Norsk Industrilakkering AS, NORAL AS, Jotun Powder Coatings AS, Electro Vacuum AS, and DuPont Powder Coatings. References 1. Y. W. Keuong, J. H. Nordlien and K. Nisancioglu, J. Electrochem. Soc., 148, B497 (2001). 2. Y. W. Keuong, S. Ono, J. H. Nordlien and K. Nisancioglu, J. Electrochem. Soc., 150, B547 (2003). 3. J. T. B. Gundersen, A. Aytac, S. Ono, J. H. Nordlien and K. Nisancioglu, Corros. Sci., 46, 265 (2004). 4. J. T. B. Gundersen, A. Aytac, J. H. Nordlien and K. Nisancioglu, Corros. Sci., 46, 697 (2004). 5. K. Shimizu, R.C. Furneaux, G. E. Thompson, G. C. Wood, A. Gotoh and K. Kobayashi, Oxidation of Metals, 35, 427 (1991). 6. S. Qian, R.C. Newman, R. A. Cottis and K. Sieradzki, Corros. Sci, 31, 621 (1990).

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Passivity Breakdown of Aluminium Alloys by Surface Enrichment of Group IIIA - VA Trace Elements Brit Graver*, Øystein Sævik, Yingda Yu and Kemal Nisancioglu Norwegian University of Science and Technology, Department of Materials Technology, N-7491 Trondhem, Norway *[email protected]

Abstract - High temperature heat treatment of aluminium alloys gives rise to a curious anodic activation phenomenon, attributed to surface enrichment of the trace elements in Group IIIA - VA, specifically the low melting point elements Pb, Bi, In and Sn. The purpose of this paper is to investigate surface enrichment by heat treatment and anodic activation by electrochemical polarization, specifically of In in relation to Pb, by use of electrochemical, electron-optical and electron-spectroscopic techniques. Indium becomes easily segregated at the surface as a result of heat treatment, at much lower temperatures (300ºC) than that required for Pb. Passivity breakdown occurs by further surface enrichment of any of these elements by selective anodic dissolution of the more active aluminium component in chloride solution. Keynotes: Heat treatment, segregation, pitting, microstructure, depth profiling

1. Introduction Activation of aluminium alloys by the presence of trace elements, specifically in Group III-V, has been a recent subject of interest because of its significance in corrosion and surface engineering for corrosion protection [1]. Work in this laboratory showed that Pb in small concentration (ppm) levels in the bulk of the material can segregate or become enriched at the surface by heat treatment and activate the surface anodically in chloride environment [1,2]. Lead concentration in a submicron thick metallic sublayer near the surface could become as high as 1 wt%. The purpose of the present paper is to investigate whether the presence of small concentrations of In has a similar activating effect by enrichment at the aluminium surface and understand the mechanisms of enrichment and activation. Previous work has mainly focused on activation in the presence of high In concentrations in the metal, typically at about 0.1 wt%, or by adding

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In salts to the test solution [3,4]. The present interest is to investigate if lower concentrations of In, which may be present as a trace element in aluminium alloys, can become enriched and segregated at the aluminium surface as a result of heat treatment and thereby anodically activate the surface in a manner similar to Pb. Additionally, In was reported to activate aluminium by heat treatment at temperatures significantly lower (300ºC) than that required for Pb (600ºC) [5]. The cited temperatures are for 1 h heat treatment in air. 2. Experimental Binary AlIn specimens containing 20, 100, 500 and 1000 ppm indium were mixed from pure components and cast into 2.5 cm thick moulds. After cooling, the cast specimens were scalped and cold rolled down to 2.2 mm thickness. The materials were degreased in acetone, cut into 2x2 cm pieces and polished through 1 µm diamond paste. The specimens were thereafter heat treated for 1 h at various temperatures in an air circulation furnace and then quenched in distilled water. The specimens were dynamically and statically polarised in 5 wt% NaCl solution maintained at 25±1oC. The potentiodynamic sweep rate was 0.1 mV/s. The potentials are reported relative to the saturated calomel electrode. Surface morphology and microstructure were examined by a field emission gun scanning electron microscope (FEG-SEM), equipped with X-ray EDS capability, and a glow discharge optical electron spectroscope (GD-OES). 3. Results 3.1. Electrochemistry The polarisation curves for the as rolled conditions, shown in figure 1a, indicate that specimens containing 20 and 100 ppm were only slightly activated, as judged by the pitting potential depression relative to pure aluminium, reported as -0.76 V [6], while higher In containing specimens showed significant activation. Figure 1b shows that the samples heat treated at 600oC exhibited similar behaviour as the as-rolled variants, except for the 500 ppm In specimen, which exhibited reduced activation relative to the as-rolled condition. However, the specimens heat treated at 300oC exhibited clear activation for all In concentrations, with the possible exception of 1000 ppm, as shown in Figure 1c. Although not shown, the curves for 100, 500 and 1000 ppm specimens became similar to the curve for the as-rolled 100 ppm specimen (figure 1a), and the curve for 20 ppm sample was reduced to its corresponding as-rolled form, after polishing the 300ºC-annealed specimens. This indicated the presence of an active layer on the these specimens, which was at least partially removed by polishing.

Passivity Breakdown of Aluminium Alloys

629 -0.6

-0.6 a)

-0.7

20 ppm In

-0.8

Potential (VSCE)

Potential (V SCE )

-0.7

100

-0.9 -1

500 1000

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-1.3 1E-05 0.0001

b)

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0.1

1

10

100

2

Current density (mA/cm 2)

Current density (mA/cm )

-0.6 c)

Potential (V SCE )

-0.7

Figure 1. Potentiodynamic polarisation data for AlIn alloys a) in the as rolled condition, b) heat treated at 600oC and c) heat treated at 300oC

-0.8 -0.9

20 ppm In 1000

-1 500

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-1.1 -1.2 -1.3 1E-05 0.0001

o

300 C

0.001

0.01

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1

10

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Current density (mA/cm 2)

Figure 2 shows the effect of annealing temperature for 20 and 1000 ppm In specimens. Annealing temperature had a significant effect on the activation of 20 ppm In specimen, giving highest activation at the lower temperatures with maximum activation at 300ºC in agreement with [5] (figure 2a). Samples containing 100 and 500 ppm In exhibited a similar response to annealing temperature, although not shown. However, 1000 ppm samples did not show significant differences in anodic polarisation as a function of annealing temperature (figure 2b), although all variants except 1000 ppm In were significantly activated in relation to their corresponding as-rolled states (see figure 1a for comparison).

Figure 2. Effect of annealing temperature on potentiodynamic polarisation of AlIn samples containing a) 20 and b) 1000 ppm In.

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630

Figure 3 shows potentiostatic data obtained at -0.8 V for selected In concentrations and thermal conditions. The 20 ppm In was passive in the as-rolled condition and after heat treatment at 600 oC, and the potentiostatic current was therefore nearly zero (figure 3a). This result is as expected from the corresponding potentiodynamic polarisation curves in figure 2. The current density was however appreciable at the outset for the specimen annealed at 300ºC (figure 3a), indicating significant activation, also in accordance with the data shown in figure 2. However, the current density decayed with time indicating onset of passivation as an active surface layer was corroded away. In the case of the 1000 ppm In specimen, the current-time behaviour behaviour of the as-rolled and 600ºC heat-treated conditions were similar to one another (figure 3b). From an initial value of 40 mA/cm2 the current density increased to a nearly steady level of 50-60 mA/cm2. This behaviour indicated a uniformly active condition, the presence of a thick active layer for both thermal states or that enriched In cannot be etched away from the surface by anodic dissolution. The actual mechanism is not yet clarified. The 300ºC annealed condition exhibited a high current density of 80 mA/cm2 at the outset, indicating the existence of a highly activated surface sublayer. Current density decayed to a lower steady-state value of about 10 mA/cm2 with time as the thin active sublayer was corroded, and the steady-state current was controlled by a relatively less active bulk material condition. 90

20 ppm

a)

25 20 15 10

300 o C as rolled o 600 C

5

b)

80 Current density (mA/cm 2 )

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50 40 30 o

20

300 C

10

0

0

0

200

400 600 time (sec)

800

1000

0

2000

4000 6000 time (sec)

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Figure 3. Potentiostatic measurements for theAlIn specimens in as rolled condition and after heat treatment at 600 and 300oC for a) 20 ppm and b) 1000 ppm.

3.2. Surface analysis Based on the foregoing results, reporting of the surface analytical results are presently limited to 20 and 1000 ppm In specimens in as-rolled condition and after annealing at 300 and 600ºC. By FEG-SEM analysis, segregated In particles were observed only on the 1000 ppm specimen. In cases where indium was observed, it was in the form of particles located along the grain boundaries and inside the grains.

Passivity Breakdown of Aluminium Alloys

631

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200 a)

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0

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0.6

Figure 4. GDOES depth profiles for AlIn specimens heat treated at 300 oC for a) 20 ppm and b) 1000 ppm and heat treated at 600 oC for c) 20 ppm and d) 1000 ppm.

GD-OES measurements showed high concentration peaks of indium in an apparent surface sublayer of about 0.05-0.1 µm thick, as shown in figure 4. Since the oxide was much thinner, and by inspection of the O and Al profiles, In was deduced to exist in the oxide and at the oxide-metal interface. For samples heat treated at 600ºC, no appreciable enrichment of In was observed on the 20 ppm sample (figure 4c), while the 1000 ppm sample exhibited In peaks again at the oxide-metal interface region as shown in figure 4d. 4. Discussion and Conclusions Present work confirmed earlier studies of aluminium surface activation by In enrichment/segregation by heat treatment at temperatures 300ºC [5]. Most effective enrichment and activation occurred at 300ºC for all In concentrations investigated, as deduced by GD-OES data. With increasing annealing temperature, activation was reduced for the low In concentrations 20 - 500 ppm due to either homogenisation or evaporation of In. For 1000 ppm, heat treatment temperature had smaller significance. However, heat treatment at all temperatures appeared to cause enrichment of In on specimens with high In content, part of which was lost due to evaporation. This conclusion was deduced from the re-

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duction of activation of heat-treated 500 and 1000 ppm specimens as a result of polishing. Thus, the final state of surface enrichment/segregation resulted from the combined effect of In diffusion to the surface and In evaporation from the surface. It is not yet clear whether segregated In or In in solid solution with aluminium is responsible for activation. Two facts support the role of segregated In: a) Maximum activation occurs by heat treatment at 300ºC, at which the solid solubility of In in aluminium is limited [7], whereas liquid In droplets can easily segregate along the grain boundaries and diffuse to the surface, similar to the mechanism described for Pb segregation at higher temperatures [8]. b) Aluminium can easily be activated by the addition of In to the chloride solution [4]. The second mechanism is not applicable to activation by Pb, which was attributed to the presence of Pb in supersaturated solid solution with Al in a surface sublayer [2]. It is therefore suggested that the low melting point In forms a liquid alloy with aluminium, assisted by the heat generated by the oxidation process and melting point depression caused by the nanoscale size of the alloy film or particles. In combination with the presence of chloride in the solution, the passivity of the oxide is significantly reduced, culminating in an effective depression of the pitting potential and increased corrosion rates. 5. Acknowledgment Hydro Aluminium assisted with GD-OES measurements and preparation of model alloys. This work was supported by The Norwegian Research Council under contract no. 158585/431. References 1. Y. W. Keuong, S. Nordlien and K. Nisancioglu, J Electrochem. Soc. 150 (2003) B547. 2. J. T. B. Gundersen, A. Aytac, S. Ono, J. H. Nordlien and K. Nisancioglu, Corros. Sci. 46 (2004) 265. 3. W. M. Carroll and C. B. Breslin, Corros. Sci. 33(1992) 1161. 4. J. B. Bessone, D. O. Flamini and S. B. Saidman, Corros. Sci. 47 (2005) 95. 5. K. Fukuoka, Sumitomo Light Metal Technical Reports 42 (2001) 131. 6. K. Nisancioglu and H. Holtan Corros. Sci. 18 (1978) 835. 7. L. F. Mondolfo, Aluminium Alloys: Structure and Properties, p.304, Butterworths, London (1976). 8. H. Gabrisch, U. Dahmen and E. Johnson, Microsc. Microanal., 4 (1998) 286

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

633

Electrochemical behaviour of Al and some of its alloys in chloride solutions Sherif Zein El Abedin,a,b Frank Endresb a

Electrochemistry and Corrosion Laboratory, National Research Centre, Dokki, Cairo, Egypt.E-mail: [email protected] b Institute of Metallurgy, Robert-Koch-Str. 42, 38678 Clausthal-Zellerfeld, Germany.

Abstract - The present work is concerned with the activation of aluminium to be suitable for use as sacrificial anodes in the cathodic protection systems. The electrochemical behaviour of Al, AlIn and Al-Ga-In alloys in 0.6 M NaCl solutions was investigated. The study was performed by means of potentiodynamic polarization, potentiostatic current-time and electrochemical impedance spectroscopy measurements. It has been found that Al-In alloy exhibits the highest negative breakdown potential in 0.6 M NaCl and the corrosion resistance of the tested electrodes decreases in the following order: Al > Al-Ga-In > Al-In. The greater activity of Al-In alloy is interpreted on the basis of the autocatalytic attack by indium. The initial dissolution of Al-In alloy leads to increase the concentration of In3+ ions in the electrolyte, then the redeposition of In at active sites on the electrode surface occurs leading to the enhanced activity. Keywords: aluminium, aluminium alloys, activation, passivity breakdown, corrosion.

1. Introduction Aluminium and aluminium alloys are widely used in everyday practice, and have a remarkable economic importance in view of their characteristics of low cost, light weight and good corrosion resistance. Because of weight and cost advantages, aluminium is the most commonly used sacrificial material for cathodic protection of steel in seawater. The more important factor in the use of aluminium in aqueous electrolytes is that complex oxide films are formed irreversibly leading to the passivation of aluminium. The passivity of aluminium can be overcome by adding suitable alloying elements such as Sn [1-6], In [7-11], Zn [12,13] and Ga [14-16]. Activation of aluminium can be achieved also by adding small quantities of suitable metal cations, such as In3+, Ga3+, Hg2+, Sn4+ and Sn2+ [17-23].

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The objective of the present work is to study the electrochemical behaviour of Al, Al-0.77%In and Al-0.4%Ga-0.21%In alloys in 0.6 M NaCl solutions using potentiodynamic polarization, potentiostatic current-time and electrochemical impedance spectroscopy measurements.

2. Experimental Measurements were made on ultra pure Al (99.999%), Al-0.77%In and Al0.4%Ga-0.21%In alloys. Al-In and Al-Ga-In alloys were prepared from pure aluminium (99.99 %), ultra pure In (99.999 %) and Ga (99.999 %). The prepared alloys were used in the cast state. Prior to each experiment the working electrodes were successively abraded with metallographic emery paper of increasing fineness up to 800, then degreased with acetone and washed with running distilled water and then inserted immediately into the cell. The electrochemical cell was made of Pyrex glass fitted with a large area platinum auxiliary electrode and a saturated calomel reference electrode (SCE) connected through a salt bridge. All electrochemical measurements were performed using a Parstat 2263 Potentiostat/Galvanostat (Princeton Applied Research) controlled by a PowerCORR corrosion measurement software and a PowerSINE electrochemical impedance spectroscopy software. The impedance spectra were recorded under open circuit conditions in the frequency range between 100 kHz and 100 mHz with an amplitude of 10 mV. The tested electrodes were cathodically polarized at 2000 mV vs SCE for 3 min in the test electrolytes before potentiodynamic polarization and potentiostatic current-time measurements. In the potentiodynamic polarization measurements, a scan rate of 1 mV S-1 was employed and the potential was scanned from –2.0 V up to the breakdown potential.

3. Results and Discussions 3.1. Potentiodynamic polarization The potentiodynamic polarization curves of pure Al, Al-In and Al-Ga-In alloys in 0.6 M NaCl solution are shown in Fig. 1. The polarization curve of Al is characterized by a broad passive region extending for about 820 mV from –1580 to –760 mV (SCE) over which the passive current is about 6 µA cm-2. At the end of the passive region the current increases abruptly as a consequence of the pitting initiation process and hence the passivity breakdown. Al-In alloy exhibits different behaviour from that of pure aluminium, since the polarization curve shows an active behaviour without any passivity. The pitting potential takes a value close to the corrosion potential signifying the greater activity of the alloy compared to pure aluminium. This is attributed to the presence of In as an alloying element giving rise to increased Cl- adsorption at higher electronegative potential which, in turn, leads to the observed activity. The polarization curve of Al-Ga-In alloy is characterized by a major anodic peak followed by a gradual active to passive transition evidenced by more than one order of magnitude decrease in the current. This peak has been also observed by Tuck

Electrochemical behaviour of Al and some of its alloys...

635

et. al. [14] and represents the oxidation of Ga. Beyond the peak a small passive region of about 200 mV is observed followed by a rapid increase in the anodic current indicating the start of pitting attack. -1

-2

A l-G a -I n

Log I / A cm

-2

-3

-4

Al

A l-I n -5

-6

-7

- 2 ,0

- 1 ,8

- 1 ,6

- 1 ,4

- 1 ,2

- 1 ,0

- 0 ,8

- 0 ,6

E / V (S C E )

Fig. 1 Potentiodynamic polarization curves of Al, Al-In and Al-Ga-In electrodes in 0.6 M NaCl.

3.2. Current-time measurements The curves of Fig. 2a display the current-time behaviour of pure Al electrode polarized at different potential values, -700, -730, -740,-750, -800, -1000 and –1400 mV (SCE) in 0.6 M NaCl solutions. At E ≤ -750 mV, the curves exhibit a rapid decrease in the anodic current in the early stages and then becomes constant indicating the formation of a passive oxide film on the electrode surface. On the other hand, at E ≥ 740 mV the anodic current decreases swiftly within the first 20 seconds of polarization and then starts to fluctuate, with magnitude depending on the potential , revealing the initiation of the pitting attack. The initial and final current values depend on the potential, since they increase with increase in the applied potential. Taking into consideration that, the pitting potential of Al in 0.6 M NaCl is about –760 mV, obtained from polarization measurements, it can be concluded that the potentiostatic current-time results are in agreement with the potentiodynamic polarization ones. Figure 2b manifests a set of current-time curves for Al-In alloy in 0.6 M NaCl solutions, where individual specimens were potentiostatically polarized at –1050, -1100, -1150, -1250 and –1400 mV (SCE). It can be seen that, at the potentials of –1250 and – 1400 mV the current is first anodic and within the first 30 seconds becomes cathodic, then takes an approximately constant value. This indicates that Al-In alloy is not attacked at potentials more electronegative than –1150 mV. At E ≥ -1150 mV, the anodic current decreases at the early moments then increases exponentially reaching a maximum and initiates to be fluctuated. This can be ascribed to a combination of passive film formation, double layer charging and pit initiation. After reaching a maximum the anodic current decreases and the fluctuation becomes of higher magnitude revealing the increase in the extent of pitting attack. Here, the anodic current

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represents both the propagation of the first pits and the initiation of new pits and the oscillating nature of the current is due to the repeated birth and death of pits [24]. As shown in the curves, the initial, maximum and the final currents increase in proportion to the applied potential and the induction time prior to the fast increase in the anodic current decreases as the potential increases. The curves of Fig. 2c display the potentiostatic current-time profiles of Al-GaIn alloy in 0.6 M NaCl solutions. Here, it can be seen the active behaviour of the alloy at potentials more than –1100 mV, as revealed by the high anodic current , indicative of the severity of corrosion. At E ≤ -1200 mV the anodic current decreases at the early moments and attains approximately constant values after about 1 minute signifying the formation of a passive film. The initial and the final current values depend on the applied potential where, they increase by shifting the potential towards the anodic direction. 12 -700 mV

a) -2

8 -740 mV

6 -750 mV

4

b)

50

-730 mV

Current density / mA cm

Current density / mA cm

-2

10

-800 mV -1000 mV -1400 mV

2 0

- 1050 mV

40

30 - 1100 mV

20

10

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-1250 mV

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-1400 mV

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10

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-2

40

30 -1000 mV

20 -1100 mV

10

-1200 mV -1300 mV -1400 mV -1500 mV

0

-1600 mV

1

10

100

1000

Time / sec.

Fig. 2 Current-time curves for a) Al, b) Al-In, c) Al-Ga-In electrodes polarized at different potential values in 0.6 M NaCl.

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3.3. Impedance Spectroscopy In the present work, the impedance technique was used as a basis for a comparative study to show the corrosion behaviour of the tested electrodes in 0.6 M NaCl solutions. The electrochemical impedance spectroscopy of Al, Al-In and Al-Ga-In alloys was recorded under open circuit conditions in 0.6 M NaCl solutions. The measurements have been performed after immersion of the examined electrodes in the test electrolytes for 2 hours in order to get stationary conditions. After 2 hours the electrodes attain quasi steady state potentials of –810, -1110 and -1190 mV for Al, AlGa-In and Al-In alloys, respectively. The Nyquist plots of Fig 3 show an incomplete capacitive semicircle, for Al electrode, with a diameter of 8.5 k cm2 which represents the polarization resistance Rp. However, the capacitive semicircles obtained in the case of Al-In and Al-Ga-In alloys exhibit diameters of 1.35 and 2.85 k cm2, respectively. The capacitive semicircle for Al-In electrode is followed by a linear region at low frequencies which makes an angle of about 45o with the real axis. This signifies a Warburg type impedance corresponding to a mass transfer process involving ionic diffusion. The tested electrodes can be arranged according to the increase in the polarization resistance, as determined from the capacitive loop, in the order: Al > Al-Ga-In > Al-In This indicates that, Al-In alloy presents the most active behaviour compared with AlGa-In and pure Al electrodes which is in agreement with the aforementioned results. 4000

3000

Zim / W cm

2

100 mH

2000 1H

1000

1H 100 mH

0

Al Al-Ga-In Al-In

1H 100 mH

100 kH

0

1000

2000

3000

4000

5000

2

Zre / W cm

Fig. 3 Nyquist plots of Al, Al-In and Al-Ga-In electrodes in 0.6 M NaCl solutions.

4. Conclusions Among the tested electrodes, Al-In alloy presents the higher negative breakdown potential and lower polarization resistance in 0.6 M NaCl solutions. The corrosion resistance of the tested electrodes decreases in the order: Al > Al-Ga-In > AlIn. The greater activity of Al-In alloy is interpreted on the basis of the autocatalytic attack by indium. The initial dissolution of the Al-In electrode leads to increase the

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concentration of In3+ ions in the electrolyte, then the redeposition of In at active sites on the electrode surface occurs, leading to the enhanced activity.

References [1] D.S. Keir, M.J. Pryor and P.R. Sperry, J. Electrochem. Soc. 116 (1969) 319. [2] T. Valand and G. Nilsson, Corros. Sci. 17 (1977) 931. [3] D.R. Salinas and J.B. Bessone, Corrosion 47 (1991) 665. [4] M. Kliskić, J. Radosević and L.J. Aljinović, J. Appl. Electrochem. 24 (1994) 814. [5] S. Gudić, J. Radosević and M. Kliskić, J. Appl. Electrochem. 26 (1996) 1027.

[6] S. Gudić, J. Radosević and M. Kliskić, “Impedance transient study of barrier films on aluminium and Al-Sn alloys” Proceedings of the Symposium on Passivity and its Breakdown, Electrochemical Society, INC, New Jersey, 1997, p. 689. [7] S. Zein El Abedin and F. Endres, J. Appl. Electrochem. 34 (2004) 1071. [8] A. Venugopal and V.S. Raja, Corros. Sci. 39 (1997) 2053. [9] S.B. Saidman, S.G. Garcia and J.B. Bessone, J. Appl. Electrochem. 25 (1995) 252. [10] C.B. Breslin, L.P. Friery and W.M. Carroll, Corros. Sci. 36 (1994) 85. [11] C.B. Breslin and W.M. Carroll, Corros. Sci. 34 (1993) 1099. [12] M.C. Reboul, PH. Gimenez and J.J. Rameau, Corrosion 40 (1984) 366. [13] A. Tamada and Y. Tamura, Corros. Sci. 34 (1993) 261. [14] C.D.S. Tuck, J.A. Hunter and G.M. Scamans, J. Electrochem. Soc. 134 (1987) 2070.

[15] A.R. Despić, D.M. Drazić, M.M. Purenović and N. Ciković, J. Appl. Electrochem. 6 (1976) 527. [16] E. Aragon, L. Cazenave-Vergez, E. Lanza, A. Giround and A. Sebaoun, Br. Corros. J. 32 (1997) 121. [17] D. M. Drazić, S.K. Zecević and A.R. Despić, Electrochim. Acta 28 (1983) 751. [18] W.M. Carroll and C.B. Breslin, Br. Corros. J. 26 (1991) 255. [19] S.B. Saidman and J.B. Bessone, Electrochim. Acta 42 (1997) 413. [20] H.A. El Shayeb, F.M. Abd El Wahab and S. Zein El Abedin, Corros. Sci. 43 (2001) 655. [21] H.A. El Shayeb, F.M. Abd El Wahab and S. Zein El Abedin, Corros. Sci. 43 (2001) 643.

[22] H.A. El Shayeb, F.M. Abd El Wahab and S. Zein El Abedin, J. Appl. Electrochem. 29 (1999) 601. [23] H.A. El Shayeb, F.M. Abd El Wahab and S. Zein El Abedin, J. Appl. Electrochem. 29 (1999) 473. [24] R.G. Kelly, J.R. Scully, D.W. Shoesmith and R.G. Buchheit “Electrochemical Techniques in Corrosion Science and Engineering” (A. Philip and P.E. Schweitzer eds., Marcel Dekker, INC, New York (2003) p. 83.

Section I Modelling and Simulation

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

641

Cooperative Spreading of Pit Sites as a New Explanation for Critical Threshold Potentials J. R. Scullya, N. D. Budianskya, L. Organb, A.S. Mikhailovc, and J.L. Hudsonb a

Dept. of Materials Science and Engineering. bDept. of Chem. Engineering University of Virginia. Charlottesville, VA 22904 ; [email protected] c Abteilung Physikalische Chemie, Fritz-Haber-Institut der Max-Planck-Gesellschaft, Faradayweg 4-6, 14195 Berlin, Germany

Abstract - AISI 316 stainless steel rotating disk electrodes were pitted at applied potentials approaching pitting potentials in NaCl solutions. Point pattern analysis methods were utilized to elucidate whether pit sites occurred randomly or whether interactions were occurring and pit sites were clustered. Both the number of metastable pits increased explosively and their spatial proximity transitioned from random to highly clustered as the pitting potential associated with large increases in anodic current were approached. Similar behavior was seen by simulation using a pitting model that produced changes in metastable pitting rate based on environmental severity, potential, and surface oxide or inclusion damage. Strengthening of cooperative interactions triggered by surface damage promoted explosive growth in the number of pits and a transition from random to clustered pitting behavior with increasing potential. This transition potential could be shifted depending on rotation rate that controlled diffusional boundary layer thickness in both modeling and experiments. This diagnostic supports the notion that cooperative interactions between pit sites promote explosive growth in new pits across stainless steel surfaces as potential is raised. This phenomenon is proposed as a new explanation for critical threshold potentials in localized corrosion. Keywords : localized corrosion, pitting, cooperative interactions, critical thresholds

1. Introduction Threshold potentials, temperatures and corrodant concentrations are important factors in many, if not all localized corrosion phenomena.1-4 At these thresholds, significant increases in pit density are seen5, as well as growth of stable pits.1,6 Both can contribute to large increases in anodic current occurring abruptly at the threshold. Although corrosion engineers make significant use of

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such thresholds, their fundamental origins are not completely understood. Past studies have often focused on the stabilization of single pits and consider pit stabilization to be associated with a single weakest or most favorable site.7-9 Such frameworks generally do not consider explosive increases in the number of pit sites often observed at critical thresholds. In this study, we consider the role of cooperative interactions between such sites as a cause of explosive spreading of pit sites across metal surfaces at a critical threshold potential. Previous studies focusing on cooperative interactions have considered the processes that might promote or suppress cooperative growth of pit sites.10,11 These interactions could be explained within the context of the acid/halide pitting mechanism.8 Aggressive species accumulation adjacent to active pits causes clustering of future pits due to a temporary presence of a more aggressive local solution and subsequent surface damage on un-pitted surfaces. Such surface damage persist for long time periods even after initial pits had repassivated and associated locally aggressive chemical fields had dissipated.10 Surface damage in commercial 316 SS caused by partial dissolution of Mn(Cr,Fe)S inclusions that form occluded regions and promoting other phenomena such as sulfide ion release that enhance pitting susceptibility. In addition, local regions of suppressed pitting susceptibility, caused by ohmic potential shielding, were observed immediately adjacent to active pits.10,11 Point pattern analysis statistics (PPA) were used to investigate interactions on single electrodes.10 Cooperative interactions could be detected when PPA revealed clustered arrays of pits.12 In this work, we propose that such cooperative interactions cause explosive growth of the number of pit sites at a critical threshold potential associated with abrupt increases in global anodic currents. This theory is tested by investigating pitting over a range of applied potentials at various rotation rates. The intent is to suppress cooperative interactions by accelerating aggressive solution transport away from the electrode surface through control of the diffusion boundary layer thickness. The objective of this paper is to investigate whether the sudden onset of pitting corrosion as conventionally indicated by abrupt increases in anodic current at a given threshold potential coincides with an explosive increase in the number of pitting sites in conjunction with a transition from random to clustered spatial distributions of pit sites. 2. Methods Experimental studies were conducted using AISI 316 stainless steel (0.13 % C, 0.31% Co, 18.18% Cr, 0.38% Cu, 1.75% Mn, 2.40% Mo, 12.25% Ni, 0.016% S, 0.35% Si, balance Fe in wt.%) wire. The 250 mm diameter wire electrodes were flush mounted in epoxy resin and polished to 1200 grit

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polishing paper followed by a 1 mm and 0.3 mm alumina slurry polish. Experiments were conducted in 0.05 M NaCl (pH 6.8) held at 47oC. The rotation rate of the electrodes was controlled using a rotating disk electrode apparatus. The pitting potential was determined from a series E-Log i polarization curves at different rotation rates. Metastable pitting behaviour was investigated in a series of one hour potentiostatic holds conducted at various potentials under both stagnant conditions and at high rotation rates.12 After experiments, photomicrographs of electrode surface were utilized to determine the spatial proximity of pit sites. Ripley’s PPA methods were used to determine the spatial proximity of both inclusions and metastable pit sites before (inclusions) and after (inclusions and pits) potentiostatic testing.10,13,14 PPA describes the spatial proximity of specific observations (in this case metastable pit sites) enabling differentiation between aggregated (clustered), random, and periodic behavior. PPA results are shown as plots of analysis radius (i.e., distance from a selected pit) vs. L2 function.10,13,14 A region of complete spatial randomness is first calculated from computer generated random distributions of pit sites and shown in figures as the zone bounded by dashed lines delineating the region containing random patterns. Experimental data from patterns of pits that are located between the dashed lines can be described as spatially random. Sites are aggregated or clustered above the region of complete spatial randomness, while patterns below the random zone are regular or periodic. PPA results are also reported as normalized area beneath the L2 function vs. radius curves. In such normalized area curves, values at or below zero are random and values above zero are aggregated or clustered. A metastable pitting model, discussed previously, was used to explore the behavior of metastable pitting patterns associated with cooperative interactions.11,15 The core of this model is a memory function that determines the local pit generation rate depending on local environmental and electrochemical changes that can promote or suppress pitting interactions.10,11 In this study, the competing processes of ohmic potential shielding versus aggressive species accumulation and subsequent surface damage were contained in a two-dimensional model. The competition between OH- production at local cathodes and acidification by metal hydrolysis was not considered. Initial pitting sites were chosen at random but the subsequent local pit generation rates were decreased based on ohmic potential shielding or increased based on aggressive solution enhancement and subsequent surface damage. The conditions were varied between low, intermediate, and high initial pit generation rate simulating either high or low corrodant concentration or applied potential. Moreover, accumulation of aggressive species was mediated by changing the interfacial boundary layer thickness, simulating stirred conditions, which controlled transport of aggressive species into the bulk solution.15 The

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model results were expressed as x-y locations of metastable pit sites and memory function development. PPA was used to determine whether the spatial distributions of simulated pit sites were random or clustered.13 Increases in pit generation rate were brought about by increased local aggressiveness and surface damage. The increase in strength of interactions promoted an abrupt transition and explosive growth of pit sites at a critical condition. This condition can be altered in the model by changing the diffusional boundary layer thickness in direct agreement with experiments. 3. Results The median pitting potential of 316 SS in NaCl at 47oC increased with rotation rate from +0.488 to 0.571, and finally to 0.685 VSCE corresponding to no rotation, 1000 RPM, and 5000 RPM, respectively. An explosive growth in the number of metastable pit sites was seen as this conventional pitting potential is approached (Figures 1a-c). However, the potential associated with explosive growth in metastable pit sites was shifted to greater and greater potentials (Figure 1d) as the rotation rate is increased and the diffusional boundary layer thickness, d, decreased. At a selected potential near Epit, the number of metastable pit sites increases with decreasing rotation rate (Figure 1e-g). It has been shown that the number of metastable pits increases explosively with increased potential indicating a threshold potential for high activity.5,16 The critical potential associated with pit stabilization, Epit, has been previously shown to be dependent on the rotation rate of the electrodes.17,18 The dependence of rotation rate was believed to be caused by a reduction of the thickness of the diffusional boundary layer which increased the removal rate of corrosion intermediates and final products. This removal prevents accumulation of aggressive species at local sites. At potentials far below the critical potential (<0.3 VSCE) under both stagnant and rotating conditions, very few pit sites initiate and pit sites are randomly distributed and fall below zero on L2 integrated area curves. As discussed in previous publications12, PPA indicates that pit sites are randomly distributed when experimental curves fall between the bounds of complete spatial randomness. At slightly higher potentials (+0.35VSCE), electrodes in stagnant solutions exhibit weak interactions shown by PPA peaks at low analysis radii. Potentials (0.4VSCE) that begin to approach the critical potential produce bursts in current and strong interactions under stagnant conditions, as shown in Figures 2a and b. In contrast, under rotating conditions explosive bursts in current are not seen and weak interactions occur (Figures 2c-e) at the same applied potential. As the rotation rate is increased from no rotation to 5000 RPM to 10,000 RPM the mechanisms that control cooperative interactions are reduced. This is shown in PPA curves where the degree of clustering is reduced

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from highly clustered (Figure 2b) to clustering at small distances (Figure 2d) to nearly random behaviour (Figure 2f). Current-time series curves indicate that a rapid rise in the current is observed after a period of time and continuous metastable pit spikes are observed indicating strong cooperative interaction behavior under stagnant conditions (Figure 2a). The degree of cooperative interactions is reduced and the current-time series show fewer and fewer clustered metastable pit current spikes when solutions are stirred at 5000 RPM (Figure 2c) and 10,000 RPM (Figure 2e). Results from the spatial model indicate a strong dependence on Clconcentration, potential or other factor that increase pit generation rates. The model indicates that metastable pitting events are randomly distributed in the study area at low activity (low Cl- concentration or potential). As the generation rate is increased (representing increasing Cl- concentration or potential) aggregation of pit sites occurs at low analysis radii closely matching experimental results. At high generation rates, bursts in current and clustering occur over very large distances indicating that strong interactions are occurring across electrode surfaces, as shown in Figure 3a-c. In addition the model shows strong dependence on changes in the diffusion boundary layer thickness that simulates rotation rate. As theory predicts, decreasing the diffusion boundary layer thickness, simulating stirring, suppresses both significant bursts in current and strong interactions indicated by clustering which now only occurs at very small analysis radii (Figure 3d-f). Note that strong interactions occur under the same conditions when the boundary layer thickness is increased (Figures 3c and f). A summary of experimental results can be clearly seen in PPA integrated area curves vs. applied potential (Figure 4) where the data for stagnant conditions rise sharply at approximately 0.4VSCE while 5000 and 10,000 RPM data do not rise until higher potentials are applied. 4. Discussion The role of cooperative interactions in producing threshold potentials associated with large, rather abrupt, increases in pit sites was investigated by looking at the spatial distribution of pitting sites in the vicinity of critical potentials associated with global increases in pit current. Explosive growth in the number of metastable pits occurred proximate to these potentials (Figures 1a-c). At or just prior to critical thresholds, cooperative interactions occur with greater intensity than at more benign conditions causing explosive growth in the number of metastable pits. It is believed that a finite amount of surface damage is required to sustain such autocatalytic corrosion. This occurs at a critical threshold potential. However, as shown in Figure 1, this can be affected by stirring solutions which alters the extent of cooperative interactions at each

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potential by minimizing surface damage. Stirring is seen to shift the onset of explosive increases in the number of pit sites to higher potentials and the transition from randomly distributed to highly clustered spatial distributions of pits (Figures 1 and 4). At low potentials (+0.3 VSCE) where only a few sites initiate a metastable pit, such a threshold is not reached. Consequently, cooperative interactions are too weak to cause subsequent explosive lateral growth of pits under stagnant and stirred conditions. The behavior begins to change when the potential is raised closer to the threshold potential. At +0.35 VSCE the number of metastable pits begins to increase (Figure 1) and clustering of sites occurs (Figure 4). Interactions are occurring at small distances and cooperative spreading of pit sites may be occurring. The higher driving force for metastable pitting at this higher potential combined with more frequent aggressive local solution formation causes cooperative interactions to occur with greater intensity. At potentials (+0.4 VSCE) that are near or exceed the critical potential, the number of metastable pit sites increases explosively under stagnant conditions (Figure 1). If cooperative interactions occur at great enough intensity to spread pit sites across surfaces at some critical threshold, then the spatial proximity of metastable pit sites should be highly clustered, as seen in Figure 2b. However, under stirred conditions the mechanisms that control cooperative interactions are suppressed.15 Stirring causes a positive shift in the critical threshold potential by reducing the surface damage responsible for cooperative interactions.15 Both the applied potential where the rapid increase in pit sites occurs (Figure 1) and the potential associated with the transition from random to clustered behaviour are both observed to shift due to stirring (Figure 4). The increase in rotation rate decreases the diffusional boundary thickness causing an increased rate of intermediate and final corrosion product removal preventing local aggressive solution accumulation. PPA results indicate that cooperative interactions at 5000 RPM (Figure 2d) have been reduced as shown by the reduction in the analysis radius peak and the intensity of the peak. As the rotation rate increases, cooperative interactions are reduced significantly such that pit sites are nearly random, as shown in Figure 2f. As the rotation rate is increased the transition from random to clustered behaviour is clearly shifted to higher potentials, as shown in Figure 4. These experimental results strongly suggest that an intense increase in cooperative spreading of metastable pits occurs to form new pits across the entire surface at or near the critical threshold potential and a transition from random to clustered behavior occurs. The metastable pitting model that incorporates three sources of interactions shows similar results to experiments.11,15 The spatial proximity of metastable pits is random at low pit generation rate representing either low chloride concentrations or low applied potentials because surface damage is less

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extensive and may heal before the next event occurs.15 When the pit generation rate is increased, interactions between metastable pits begins to occur as surface damage accumulates causing small interaction patterns that show a peak in PPA results. Metastable pits first initiate at a few random sites but begin to cluster after a short time. The clusters promote local regions of surface damage that grow rapidly outward across the electrode area as time progresses. Model results for a high pit generation rate show that metastable pits spread across the electrode surface over time by a mechanism of cooperative positive interactions triggered by a build up of surface damage (Figure 3b). PPA results indicate that metastable pit sites are quickly clustered at distances that approach the size of the study area (Figure 3c). Model results predict that reducing the diffusional boundary thickness, or simulated stirring, will increase diffusion of aggressive species away from electrode surfaces preventing local aggressive solution accumulation and subsequent surface damage. At the same pit activity where metastable pit sites grew across an electrode surface under stagnant conditions (Figure 3b) the strength of interactions were greatly suppressed when highly stirred as shown by a drastic reduction in the peak radius on L2 statistics curves (Figure 3f). Given the notion that cooperative spreading is caused by local aggressive solution accumulation and surface damage, stirring should suppress clustering. Not only do these results suggest that a critical transition exists due to an explosive cooperative spread of local corrosion sites across metallic surfaces, but they also suggest a second explanation for abrupt thresholds in the pitting processes. Threshold potentials and Cl- ion concentrations may also be associated with the conditions that enable explosive lateral growth of pits sites across metal surfaces. If the factors that control cooperative spreading can be better understood then they can be used to design materials that are more intrinsically resistant to the explosive growth of pits sites. 5. Summary 1. Threshold potentials are not only characterized by an explosive increase in the density of metastable pits but also by a transition from random to clustered spatial patterns of pits. The clustered behavior is caused by cooperative interactions between metastable pit sites driven by local chemistry changes and subsequent surface damage near initial pit sites. 2. Cooperative interactions can be suppressed in stirred solutions that decrease the local accumulation of aggressive species. Such stirring produces positive shifts in critical threshold potentials associated with the explosive growth of pits.

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3. Cooperative interactions and subsequent explosive growth across electrode surfaces under certain conditions may provide a new explanation for critical threshold potentials in pitting corrosion. Acknowledgements The United States Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering supported this project under contract DEFG02-00ER45825 with Dr. Harriot Kung as contact monitor. References [1]. Z. Szklarska-Smilalowska, Pitting Corrosion of Metals, NACE, 1986. [2]. A. Broli, and H. Holtan, Corros. Sci. 17 (1977) 59-69. [3]. R.J. Brigham, and E.W. Tozer, Corrosion 29 (1973) 33-36. [4]. G.K. Glass, and N.R. Buenfeld, Corros. Sci. 39 (1997) 1001-13. [5]. G. Herbsleb, and W. Schewenk, STAHL EISEN 87 (1967) 709-13. [6]. N.J. Laycock, and R.C. Newman, in: G.S. Frankel, and J.R. Scully, Eds., Proc. of Corrosion/2001 Research Topical Sym., NACE, 2001, p. 165-90. [7]. H. Boehni, and F. Hunkeler, in: H.S. Isaacs, U. Bertocci, J. Kruger, and S. Smialowska, Eds., Proc of Adv in Localized Corrosion, NACE, 1990, p. 69-76. [8]. J.R. Galvele, J. Electrochem. Soc. 123 (1976) 464-74. [9]. G.T. Gaudet, W.T. Mo, T.A. Hatton, J.W. Tester, J. Tilly, H.S. Isaacs, and R.C. Newman, Amer Inst Chem Engineers Journal 32 (1986) 949-58. [10]. N.D. Budiansky, J.L. Hudson, and J.R. Scully, JECS. 151 (2004) B233-B43. [11]. T.T. Lunt, J.R. Scully, V. Brusamarello, A.S. Mikhailov, and J.L. Hudson, JECS 149 (2002) B163-B73. [12]. N.D. Budiansky, L. Organ, A.S. Mikhailov, J.L. Hudson, and J.R. Scully, in, Proceedings of 3rd Int. Sym. on Pits and Pores. Honolulu, HI, ECS, 2004, [13]. B.D. Ripley, Spatial Statistics, John Wiley & Sons, New York, 1981. [14]. N.D. Budiansky, L. Organ, J.L. Hudson, and J.R. Scully, JECS, 152 (2005) B152-B60. [15]. L. Organ, J.R. Scully, A.S. Mikhailov, and J.L. Hudson, Electrochim. Acta Accepted for Publication (2005) [16]. C. Punckt, M. Bolsher, H.H. Rotermund, A.S. Mikhailov, L. Organ, N.D. Budiansky, J.R. Scully, and J.L. Hudson, Science 305 (2004) 1133-36. [17]. F. Mansfield, and J.V. Kenkel, Corrosion 35 (1979) 43-44. [18]. F. Franz, and P. Novak, in: R.W. Staehle, B.F. Brown, J. Kruger, and A.K. Agrawal, Eds., Localized Corrosion 3, NACE, 1971, p. 576-79.

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Atomistic simulation of the passivation of ironchromium alloys using calculated local diffusion activation barriers Boubakar Diawara, Yves-Alain Beh, Philippe Marcus Laboratoire de Physico-Chimie des Surfaces, CNRS/ENSCP (UMR 7045), Ecole Nationale Supérieure de Chimie de Paris 11 rue Pierre et Marie Curie, 75005 Paris, France

Abstract - A 3D model has been developed to model the selective dissolution and passivation of alloys, and apply to simulate the passivation of iron-chromium alloys. The real structure of the alloy is taken into account. The passivation is modelled by considering the formation of "oxide" nuclei, resulting from the presence on the surface of local chromium-rich clusters. The dynamic evolution is based on the Kinetic Monte Carlo (KMC) method allowing us to take into account realistic kinetic evolution of the system, with simulation time related to the real time. Using the Modified Embedded Atom Method (MEAM), we have calculated the activation energies for the dissolution and the surface diffusion steps. The calculated surface diffusion probabilities are found to be in agreement with the empirical values used previously. In particular they confirm that chromium preferentially diffuses toward the chromium clusters on the surface while iron atoms show no preferential diffusion. The simulation leads to kinetics of passivation with a KMC time of the order of a few seconds, corresponding to the initial stage of passivation. The simulation exhibits the qualitative evolution of the corroded/passivated surface (Cr content, surface roughness). For a series of Fe-xCr alloys (x in the range 15-22%), the results of the simulation confirm our previous finding that the transition from incomplete or no passivation to complete passivation is continuous. Keywords : atomistic modelling, passivation, FeCr alloys, KMC, MEAM

1. Introduction While extensive experimental work has been published on the passivation of alloys, including Fe-Cr alloys, there is a very limited number of papers dealing with modeling and/or computer simulation of alloy dissolution and passivation. The reported simulations used macroscopic models [1-4, 9] or atomistic models [5-8, 10-16]. For atomistic modeling, 2D models [5-8, 11-13] and 3D models [10, 14-16] have been used.

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2. Our 3D model For the atomistic modeling of the initial stage of passivation of FeCr alloys we need structural models for the alloy and the passive film, and data on the probabilities of dissolution and surface diffusion of Cr and Fe atoms. With these data, an appropriate algorithm can be used for the simulation of the dynamic evolution of the system. 2.1 – Structural model for the alloy

Fig. 1: A model of FeCr alloy oriented along the [120] direction

We have used a 3D model taking into account the real lattice of alloys (body centered cubic for the Fe-Cr alloys investigated here). The alloy model is occupancy of the alloy elements (Fe-xCr with 15<x<22). The generated model (Fig. 1) is oriented along the [hkl] direction ([110] in this work). Depending on the size of the model the number of atoms varies from 1764 to 33600.

2.2 – Structural model for the passive layer It was assumed that the lattice of the "passive layer" remains the same as the lattice of the substrate. In the structure of a-Cr2O3, a chromium cation has 3 neighbors at a distance of 2.88 Å and 3 neighbors at 7.2Å. The mean distance is Fe Cr « oxidized » Fe « oxidized » Cr 4.2Å, which corresponds to the distance Fig. 2 The model of passivation nucleus of 3rd nearest neighbors in the lattice of the iron-chromium alloy. This critical size was evaluated from the experimental data on the structure of the passive films, measured by STM on Fe-22Cr(110) [17]. Therefore it was considered in the model that a passive layer nucleus is formed when a critical number of Cr atoms (denoted NCr(nuclei)) are present on the surface in an area around a Cr atom defined by an hemisphere of radius 4.2 Å (Fig. 2). O 4.2 A

2.3 – Evaluation of diffusion and dissolution This evaluation requires the calculation of the cohesive energy of clusters representing local areas of the alloys surface. FeCr alloys being chemically disordered it is necessary to consider a large number of configurations in order to cover the diversity of structural and chemical environments. In a previous approach we have used quantum chemistry calculations for the evaluation of the

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extraction energy of atoms. The cost of these calculations limits the size of the configuration space to be explored. An alternative is the use the Modified Embedded Atom Method (MEAM) potential [18]. The MEAM considers that each atom is embedded in the electronic density ri of the other atoms. The energy of the system is the sum of the energy of each atom Ei, which is a sum of a functional of the electronic density ri and a sum of all the pair potentials involving atom i :

1 ⎤ ⎡ Etot = ∑ ⎢Fi ( r i ) + ∑ fi j ( Rij ) ⎥ 2 i i ⎣ ⎦

were

Fi ( r i ) is the

embedded energy of atom i and fi j the pair potential between atom i and j.

The parameters used in the calculation of Fi ( r i ) are fitted using experimental values as the cohesive energy, the bulk modulus, the vacancy formation energy. The MEAM has been extended to alloys and parameters for FeCr alloys have been evaluated by Lee et al. [19]. 2.3.1 – Evaluation of the extraction energy of atoms (dissolution) The dissolution probability is related to the energy necessary to extract an atom (Fe or Cr) from a cluster of atoms with a given topology and chemical composition. The energy of extraction has been calculated as the difference between the sum of the energy of the cluster after the extraction of an atom and of the energy of the free atom, and the energy of the full cluster prior to the atom extraction: Eextrac = (E cluster n-1 atoms + E 1 atom) ) - E cluster n atoms (with E 1 atom= 0 in the MEAM method ) 7 Extraction energy (eV) 6 Fe Cr

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For the extraction of Cr from pure Cr clusters or Fe atom from pure Fe clusters, the calculated energies are in the range 0.5 eV - 6.0 eV and do not vary beyond the 6th neighbor (Fig. 4). Thus the size of the cluster was limited to the 6th neighbors. For the alloy, the topology and the chemical environment of the extracted atoms are taken into consideration (Fig. 5).

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654 Extracted atom (Fe)

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2.3.2 – Evaluation of the surface diffusion barrier The surface diffusion barrier was evaluated by considering a cluster containing up to 6th neighbors for the initial position and the final position of the diffusing atom (Fig. 6). The energy of the initial configuration (E1) and the final configuration (E2) are then calculated using the MEAM potential. The transition state is calculated by a conjugated gradient algorithm (Fig. 7). As in the Environment of the initial position

Environment of the final position

Fig. 6: Definition of the local environment of the initial and final positions

Fig. 7: Evaluation of the diffusion barrier

preceding part on extraction, the diffusion barrier was calculated for various topologies and chemical environments. The calculated values are in the range 1.4 eV - 2.7 eV for the diffusion Cr Fe barrier of iron atoms and in the range 1.84 eV 2.88 eV for the Iron diffusion with increasing Cr content diffusion of first nearest Barrier (eV) second nearest third nearest chromium atoms. neighbors neighbors neighbors An examination of Fe Fe Fe 2.08 the results shows Cr Fe Fe 2.10 that for iron atoms, Cr Cr Fe 2.40 the diffusion Cr Cr Cr 2.70 barrier height Fig. 8: Evolution of iron’s diffusion barrier with the local chromium increases with content

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increasing Cr content (Fig. 8), causing blocking of iron in Cr-rich areas. The same blocking behavior is observed for the diffusion of Cr in Cr-rich areas. A preferential diffusion of Cr atoms toward Cr-rich areas is observed (lowering the energy of the system). In contrast it is found that the diffusion of Fe atoms toward Cr-rich areas can lower or not the energy of the system, indicating that there is no preferential diffusion of Fe atoms toward Cr-rich areas. 2.4 – Kinetic Monte Carlo simulation of the dynamic evolution Our previous model used a classical Metropolis Monte Carlo algorithm [14,16]. The dynamic evolution is now based on the Kinetic Monte Carlo (KMC) method allowing us to simulate realistic kinetic evolution of the system, with simulated time related to the real time. In the KMC approach the dynamic evolution of the system is the result of n - Ei

processes, each process i being characterized by a rate Ri = v0 e kT probability pi =

Ri n

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∑Rj j =1

In the simulation of the passivation of FeCr alloys three processes are considered: surface diffusion, dissolution and blocking (passivation). The pre-exponential factor is set to the same value no=1013 for all the processes. In order to take into account the effect of the potential on the dissolution process, we applied a variable overpotential f which lowers the activation energy of the dissolution processes :

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3. Results of the KMC simulation The simulations have been performed with a model of 25 Å x 25 Å x 25 Å (11 atomic planes), wich contains 1764 atoms. The overpotential was fixed at 1.75 eV. The default value for NCr (number of Cr atoms forming a nucleus of passivating layer)) was 5. The results presented below are the mean values of five simulations for each set of parameters. 3.1 Simulation of passivation kinetics 300 Number of dissolved atoms

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Fig. 9: Amount of corroded material (number of atoms) vs. the KMC time (s)

The simulated kinetics are in qualitative agreement with experimental data. The KMC time is of the order of a few seconds, corresponding to the initial stage of passivation. The decreasing amount of corroded material prior to passivation with increasing Cr content in the alloy is clearly observed (Fig. 9). 3.2 Effect on passivation of the Cr content in the alloy

% Passivation

100

A continuous transition from non-passivated to passivated alloys is observed between 17 and 20% Cr (Fig. 10). This transition, which is not a percolation threshold, is consistent with the view that a nucleation and growth mechanism governs the initial stages of passivation (in agreement with experimental STM data).

80 60 40 20 %Cr 0 12 13 14 15 16 17 18 19 20 21 22 23

Fig. 10: Passivation probability vs. chromium content in the alloy

3.3 Simulated surface properties (chromium enrichment, surface roughness) 29 27 25

5

%Cr in the passive film

roughness

4

23 21

3

19 17

%Cr in the bulk

15 17

18

19

20

21

Fig. 11: Cr content in the passive layer vs. the Cr content in the alloy

22

%Cr

2 17

18

19

20

21

22

23

Fig. 12: Roughness of the passive layer vs. the Cr content in the alloy

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The passive film is enriched with chromium (Fig. 11) (in qualitative agreement with the experimental XPS data). For alloys with low Cr content, extensive iron dissolution is required to obtain passivation. This leads to increased surface roughness (Fig. 12). 4. Conclusions A KMC simulation has been performed with a 3D model of Fe-Cr alloy. Using a Modified Embedded Atom Method (MEAM): · dissolution rates have been calculated from extraction energy of atoms from clusters extended to the 6th neighbor, · diffusion rates have been evaluated from diffusion barriers corresponding to a transition state calculated by a conjugated gradient algorithm. An overpotential has been introduced for the dissolution process. “Real” kinetic data are derived from the results of simulations (equivalent to Q vs t). A continuous transition (not associated with a percolation threshold) is observed from non-passivated to passivated alloys for Cr contents between 15 and 20%, consistent with the view that a nucleation and growth mechanism governs the early stages of passivation. The simulation exhibits the qualitative evolution of the corroded/passivated surface (Cr content, surface roughness). References [1] H.W. Pickering and C. Wagner, J. Electrochem. Soc., 114, 698 (1967). [2] A.J. Forty, G. Rowlands, Phil. Mag. A, 43, 171 (1981). [3] M. Keddam, O.R. Mattos, and H. Takenouti, Electrochim. Acta, 31, 1147 (1986). [4] M. Keddam, O.R. Mattos, and H. Takenouti, Electrochim. Acta, 31, 1159 (1986). [5] R.C. Newman, Q. T. M. Foong, and K. Sieradzki, Corros. Sci., 28, 523 (1988). [6] K. Sieradzki, R. R. Corderman, K. Shukla, and R.C. Newman, Phil. Mag., 59, 4, 713 (1989). [7] S. Quian, R.C. Newman, R.A. Cottis, and K. Sieradzki, J. Electrochem. Soc., 137, 435 (1990). [8] Song Quian, R.C. Newman, R.A. Cottis, and K. Sieradzki, Corros. Sci., 31, 621 (1990). [9] J. Laurent, and D. Landolt, Electrochim. Acta, 36, 49 (1991). [10] J. Erlebacher, M. J. Aziz, A. Karma, N Dimitrov, and K Sieradski, Nature, 410, 450 (2001). [11] E. McCafferty, Electrochem. Solid State Lett. 3, 28 (2000) [12] E. McCafferty, Corros. Sci. 42, 1993 (2000) [13] E. McCafferty, Corros. Sci. 44, 1409 (2002) [14] M. Legrand, B. Diawara, J.-J. Legendre, and P. Marcus, Corros. Sci. 44, 773 (2002) [15] M. Legrand, B. Diawara, J.-J. Legendre, and P. Marcus, Chemometrics & Intelligent Laboratory Systems, vol 62/1, 1-15 (2002) [16] B. Diawara, M. Legrand, J.-J. Legendre and P. Marcus, J. Electrochem. Soc, 151 (3) B172 (2004) [17] V. Maurice, W. Yang, and P. Marcus, J. Electrochem. Soc. 143,1182 (1996) [18] B.-J. Lee and M. I. Baskes, Phys. Rev. B, vol 62, 8564 (2000) [19] B.-J. Lee, J.-H. Shim and H. M. Park. Calphad, Vol. 25, No. 4, pp. 527-534 (2001)

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DFT study of the interactions of Cl- with passivated Nickel surfaces: energetic and structural aspects B. Diawara a, , A. Bouzoubaaa, N. Pineaub, C. Minotb, V. Mauricea, P. Marcusa a

Laboratoire de Physico-Chimie des Surfaces, CNRS/ENSCP (UMR 7045), 11 rue Pierre et Marie Curie, 75231 Paris Cédex 05, France ; [email protected] b Laboratoire de Chimie Théorique, CNRS/UPMC (UMR 7616), boîte 137, 4 place Jussieu, 75252 Paris Cédex 05, France

Abstract - Density functional theory calculations of adsorption and penetration of chlorides on passivated nickel surfaces have been performed to investigate the mechanism of the breakdown of passivity and the localized attack of the substrate. The substitution of surface hydroxyl groups with chloride ions and the insertion of chloride in the first oxide layers were studied for structures with increasing Cl- surface coverage (25%, 50%, 75%, and 100%). In a first approximation, the relaxation of the four upper atomic layers was allowed only in the direction perpendicular to the surface. In order to evaluate the contribution of the in-plane relaxation of the atomic positions, full optimization of all the substituted and inserted structures have been investigated. While the full relaxation of the structures yielded no energy gain for the substitution step, a considerable gain was obtained for the optimized structure after insertion. An examination of the optimized structures shows a major reorganization in response to repulsive electrostatic interactions between adsorbed chloride ions at increasing chloride coverage. Depending on the coverage, the reorganized surface adopt local structure close to the structure of compounds as Ni(OH)Cl or NiCl2. Keywords: quantum chemistry, DFT, passivity breakdown, chloride, NiO

1. Introduction The interaction of Cl- with passivated metal surfaces can cause the breakdown of the passive film and the localized attack of the substrate, leading to localized corrosion by pitting. One of the mechanisms proposed to explain the initiation of localized corrosion in Cl- containing aqueous solutions involves the

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adsorption of Cl- on the passive film surface followed by the penetration of Clinto the oxide layer [1]. Density functional theory calculations of the adsorption and penetration of chlorides on passivated nickel surfaces have been performed to evaluate this mechanism [1]. The relaxation of the surface layers was allowed only in the direction perpendicular to the surface. Here, we report new results involving full optimization of the structures obtained after substitution of the surface hydroxide by chlorides and their insertion in the oxide lattice in order to evaluate the contribution of the in-plane relaxation of the atomic positions on the energetic of the different steps of the mechanism. The structural reorganization of the substituted and inserted structures will be analyzed in order to see if they can constitute an initial stage for the passivity breakdown. 2. Computational details Ab initio periodic calculations were performed using with the Vienna Ab initio Simulation Package code (VASP) [2-3]. The Kohn-Sham equations are solved with the generalized gradient approximation (GGA), using ultra-soft pseudopotentials [4-5] together with a plane wave basis set and a cutoff to 396 eV. The Brillouin Zone was mapped with a 4x4x1 grid of Monkhorst-Pack type [6]. The fully hydroxylated NiO(111) surface is modeled by a slab of three NiO layers covered by an outmost hydroxyl layer (Fig 1). The surface is unreconstructed according to experimental results on the passivation of Ni(111) surfaces in aqueous acid solution [7-9], confirmed by theoretical calculations [1]. Starting from the hydroxylated structure, we have investigated the successive substitution of the surface hydroxide by chlorides, followed by the insertion (penetration of a single Cl- ion in the first oxygen layer of the oxide lattice) resulting from the exchange between the Cl at the surface and an O from the O layer (Fig. 1). OH layer Relaxed layers

Ni layer O layer

Frozen layers Fully hydroxylated surface

Surface substitution

Bulk insertion

Fig. 1 : Model of the hydroxylated NiO(111) surface and of the steps of substitution and insertion by chlorides

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For all the calculations the three lower atomic layers (O-Ni-O) were frozen. Full optimization (full relaxation) of the four upper layers (Ni-O-Ni-OH/Cl) of the substituted and inserted structures was allowed. The surface energy after substitution (E sur/sub) and after insertion (E sur/insert ) was evaluated, yielding values for the substitution (DE sustitution) and insertion (DE insertion) energies, according to the following formulas:

DE substitution = (E surf/ Cl- + E H20 ) - (E surf/ OH- + E HCl) DE insertion = (E sur/ ins ) - (E sur/ sub ) Following this notation, a negative value of energy indicates an exothermic process. Values per chloride ion are reported. 3. Substitution of hydroxide ions by chloride ions The replacement of surface hydroxide ions by chloride ions were studied for increasing Cl- surface coverages (25%, 50%, 75%, 100%) (see Fig. 2). Cl

-

H O Ni

25%

50%

75%

100%

Fig. 2: Side view and top view of the substituted surface structure of increasing Cl- coverage (25%, 50%, 75%, 100%) before relaxation of the atomic positions

3.1 – Energetic of the substitution The calculated substitution energy increases with the Cl- coverage (see Fig. 3). This is due to unfavorable electrostatic repulsive interaction between the chloride ions, Cl- being larger than hydroxide ions (ionic radius of 1.81 Å instead of 1.40 Å). With the reference considered for the calculation of the substitution energy, this step is found to be energetically favorable at low Cl-

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coverage. The full relaxation of the structures yielded no significant energy gain for the substitution step. E n e rg y (e V /C l)

3 2 Insertion (partial relaxation )

1

Insertion (full relaxation)

0 10% 20% 30% 40% 50% 60% 70% 80% 90% 100%

-1

Substitution (full relaxation) Substitution (partial relaxation)

-2

Coverage Fig. 3: Energies of substitution and insertion vs. Cl- coverage after full and partial relaxation of the structures

3.2–Structural aspects Besides energetic aspects, the structural changes of the surface layers allowing to minimize the repulsive interactions between the ions are important. The examination of the optimized surface structure after substitution shows major restructuration. Depending on the coverage, the restructured surface adopts local atomic arrangements close to those found in compounds such as Ni(OH)Cl and NiCl2. The structures of these compounds have been reported in ref [10].

25 % Cl-

View through the xz plane

50 % Cl-

75 % ClView through the yz plane

Before optimization

After optimization

Two views of the structure of Ni(OH)Cl

Fig 4: Structural evolutions of substituted structures for Cl- coverages of 25%, 50% and 75%

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For Cl- coverages of 25%, 50% and 75% (see Fig. 4), the topmost layer containing chloride and hydroxide ions splits into two separated layers, one with only hydroxide ions and the other with chloride ions. The resulting three upper layers (Ni-OH-Cl) tend to form a nucleus having the structure of Ni(OH)Cl. For 100% Cl- coverage (see Fig. 5), the two upper layers (Ni-Cl) tend to form a nucleus of NiCl2 structure and there is only a minor restructuring of the surface layers.

Structure of NiCl2

Before optimization

After optimization

Fig. 5: Restructuring of substituted structures for a Cl- coverage of 100%

4. Insertion of chloride ions in the oxide lattice The insertion step involves a place exchange between a surface chloride ion and an oxygen of the first inner anionic layer. Depending on the Cl- surface coverage, some of the four oxygen atomic positions of this layer can be equivalent. For each coverage, we have investigated all the non equivalent insertion sites. The results concerning the most stable inserted structure after full optimization are reported here. The corresponding initial structures are shown in Fig. 6.

25% Cl-

50% Cl-

75% Cl-

100% Cl-

Fig.6: Side view of the inserted structure with increasing Cl- coverage before optimization

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4.1 – Energetic of the insertion The insertion energy decreases with Cl- coverage (see Fig. 3). This is attributed to the decrease of the repulsive interactions between surface chloride ions when one ion is inserted in the oxide lattice. For low coverage of Cl-, the insertion of Cl- is not favorable (positive energy difference). For high Cl- coverage, the full relaxation leads to a considerable energy gain for the final optimized inserted structures and the insertion process becomes energetically favorable for Clcoverages above 70%. 4.2 – Structural aspects The surface restructuring after insertion of Cl- leads to the formation of local structures close to the structure of compounds as Ni(OH)Cl and NiCl2 (Fig. 7). For a Cl- coverage of 25%, chloride and oxygen ions of the inner layer split into an oxygen layer and a chloride layer (Fig. 7). The three resulting upper layers (Cl-Ni-OH) tend to form a nucleus with a Ni(OH)Cl structure. For Cl- coverages of 50% and 75%, the mixed chloride-oxygen layers show a minor splitting, in contrast to the mixed chloride-hydroxide layer (see Fig. 7). The three resulting upper layers (Ni-OH-Cl) tend to form a nucleus of a Ni(OH)Cl structure.

25 % ClStructure of Ni (OH)Cl

50 % Cl-

Structure of Ni(OH)Cl

75 % ClBefore optimization

After optimization

Fig. 7: Restructuring after the insertion of Cl- for surface coverages of 25, 75 and 100%

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For a Cl- coverage of 100%, there is a major reorganization of the four surface layers (Fig. 8). The reconstructed layers (Ni-Cl) tend to form nuclei of NiCl structure with the formation of Cl-Ni-Cl sequences. Fig. 8: Restructuri ng after the insertion of Cl- for surface coverage of 100%

Structure of NiCl2 -

100 % Cl

Before optimization

After optimization

5. Conclusions DFT calculations of the substitution and insertion of chlorides on a hydroxylated NiO(111) surface have been performed to investigate the initiation of localized corrosion of passivated metal surfaces in the presence of Cl- . The penetration mechanism of localized corrosion was simulated using a two step mechanism with (i) substitution of surface hydroxide ions by chloride ions at increasing Cl- coverage, followed by (ii) the insertion of chloride ions in the first inner anionic layer of the oxide lattice. The results obtained after full optimization of the atomic positions show that : · the insertion step can be favorable above a 70% Cl- surface coverage, but increasing the surface coverage of Cl- costs energy due to repulsive interactions. · the reconstruction of the surface layers tends to form nuclei of compounds such as Ni(OH)Cl or NiCl2. The formation of such compounds disrupts the NiO lattice, which can be a pathway for the breakdown of passivity. References 1. N. Pineau, C. Minot, V. Maurice, P. Marcus Electrochemical and Solid State Letters 6 (2003) B47 2. G.Kress, J.Hafner, Phys.Rev.B, 47, C558 (1993). 3. G.Kress, J.Hafner, J.Furthmuller, Comp.Mat.sci, 6,15(1996). 4. D.Vanderbilt,Phys.Rev.B,41,7892(1990). 5. G.Kresse,J.Hafner,J.Phys.Condens.Matter,6,8245(1994). 6. H.J.MonKhorst ,J.D.Pack, Phys Rev. B,13,5188(1976). 7. V.Maurice, H.Tallah, P.Marcus, Surface Sci, 304,98(1994). 8. D.Zuili, V.Maurice, P.Marcus, J.Phys.Chem, 147,1393(2000). 9. V.Maurice, L.Klein, P.Marcus, ELectrochem.Sol.StateLett.4, B1 (2001). 10. S.Antoci, L.Mihich, Phys Rev. B, 18,5768 (1978).

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Simulation of corrosion processes with anodic and cathodic reactions separated in space Janusz Stafieja, Abdelhafed Talebb, Chrisitine Vautrin-Ulc, Annie Chausséc, Jean Pierre Badialib* a

Institute of Physical Chemistry, Polish Academy of Sciences, ul. Kasprzaka 44/5201224, Warsaw, Poland. acc [email protected] b Laboratoire d’Electrochimie et de Chimie Analytique, UMR 7575, ENSCP Université P. et M. Curie, 4 Place Jussieu, 75005 Paris, France. [email protected], [email protected] c Laboratoire Analyses et Environnement, UMR 8587, Université d’Evry Val d’Essonne Bd F. Mitterrand, , 91025 Evry, France

Abstract - Using a cellular automata model we investigate corrosion processes in which anodic and cathodic reaction can be spatially separated. We show that for certain diffusion conditions acidic or basic regions are generated. The simulation reveal also the existence of an incubation time characterized by a stochastic distribution. Keywords: corrosion, cellular automata, incubation time

1. Introduction Protection of the natural environment and storage of nuclear waste require that we have reliable predictions on the evolution of corroding systems over a very long period of time. Then the computer simulations are a natural complement of the usual experiments. In numerical simulations we deal with simple model systems and a small number of processes. We expect these models to predict general trends of a large class of corrosion processes. Here we focus on induction time preceding a pitting corrosion. The existence of this time is documented experimentally [1,2] and defined as the period of time elapsed in between chloride ions additions and the passive layer breakdown. It is of crucial importance for the long time security of the natural environment. We show that

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a simple model predicts behaviors in qualitative agreement with those observed in real experiments. The induction time is a result of the passivation/depassivation competition. In our model this competition is essential. However as it will be shown it is not sufficient to produce an induction time. Hereafter we introduce anodic and cathodic reactions that can be separated in space. As a result of this separation heterogeneities in the distribution of H+ and OH- ions can be formed [3]. Diffusion of these ions and their possible neutralization counteract the formation of these heterogeneities and diffusion determines to a large extent the value of the induction time. Since the pitting corrosion is a stochastic process in our model we use a stochastic cellular automata (CA) approach as we have already done in [4-6]. We present our model in Section 2 and the simulation results in Section 3 with a brief conclusion in Section 4. 2. The Model In ref. 5 we describe the model first in terms of electrochemical and chemical reactions. Here we recall the discrete lattice representation with the CA transformation rules. The state of the lattice is defined by the ensemble of states at each lattice site. It is customary to speak about states of a lattice site as about chemical species occupying this site. The metal sites (M sites) represent the bulk material. M sites in contact with the corrosive environment are put into two separate categories: R and P - reactive and passive sites respectively. The sites denoted A, C and E represent the environment. The site A (C) indicates that the environment at the position occupied by this site is more acidic (basic) than in the original solution E, which is assumed to be neutral. The memory size and run time limitations of the computers force us to work on a two dimensional lattice. We construct our CA model as a combination of several stochastic processes characterized by their probabilities. They represent electrochemical and chemical reactions of ref. [6]. The spatially separated electrochemical (SSE) reactions have an a priori probability, psse, to be realized. We assume that anodic dissolution of a piece of metal is accompanied by a simultaneous cathodic reaction. The sites of the two reactions belong to the surface of a single connected piece of material. The connectivity of two sites in our lattice model is defined such that there is a suite of nearest neighbor (nn) sites of the types M, P or R forming a path from one site to another [6]. The Von Neumann or 4 - connectivity is used except for the front. Two neighboring front sites are connected if they are both von Neumann nearest neighbors or they have an M site as a common von Neumann neighbor. This prevents the surface of a connected M piece from splitting into disconnected pieces. Since the transformation of an R site depends on the state of the solution in its environment we use the quantity, Nexcnn, that determines the excess of A over C

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sites nearest neighbors of the R site. This quantity roughly estimates pH in the vicinity of the reactive site. We mimic anodic reaction as R A and as R+C(nn) P+E(nn) for Nexcnn ≥ exc 0 and for N nn < 0 respectively. The “(nn)” is put to indicate the nearest neighbor of the considered site. Here we adopt the rule that any surface site S, where S=R or S=P, mediates the cathodic processes with an a priori equal probability. We represent cathodic reaction by S+A(nn) S+E(nn) and S+E(nn) S+C(nn) respectively. We also adopt a simple exclusion rule. If the nearest neighbors of the surface site are neither E nor A it follows that the S site is blocked by nearest neighbors of C type and such an S site cannot mediate the cathodic process. Therefore the a posteriori probability of the SSE reaction depends on the blocking of its cathodic part by the presence of C sites. We repeat the selection from among the other S sites until a free site is found, whatever its position provided it is connected to the site on which the anodic reaction takes place, or until the list of connected surface sites is exhausted. If the cathodic reaction is impossible, the spatially separated anodic and cathodic reactions cannot be realized. There are three processes: R E, R P and P E in our model occurring with probabilities PRE , PRP and PRP respectively. These probabilities depend on the value of Nexcnn . We express the dependences in terms of pcor1, pcor2 , and poxi according to : if Nexcnn > 0 ; PRE= pcor2 Nexcnn ; if Nexcnn = 0 ; PRE = 0 ; if Nexcnn < 0 ; PRE = 0 ;

PRP = 0 ; PPE=0.25 Nexcnn PRP= pcor2 ; PPE = poxi PRP = 1 ; PPE = 0

We adopt the rule that M sites in contact with the solution become R sites. To mimic neutralization the C and A species perform a random walk over E sites according to the following rules. At a given time step each C and A attempt to jump from their initial position to a nearest neighbor position. The result of the attempted random step depends on the nearest neighbor. If it is an E1+ A2 and C1 + E2 E site the walker is moved to it according to: A1 + E2 E1+ C2 where subscripts "1" and "2" denote the source and the target site for the random step. If the walker is an A site and it attempts to step on a C site both annihilate according to: A1 + C2 E1+ E2 and vice versa: C1 + A2 E1 + E2 In the other cases the configuration is unchanged: A1 + X2 A1+ X2 where X ≠ E nor C and C1 + X2 C1+ X2 where X ≠ E nor A. With the choice of our parameters we should like to mimic the fact that the surface is hard to depassivate poxi ~0 and readily repassivated pcor1~1. In contact with a basic environment a reactive site gets immediately passivated. To regulate the diffusion rate with respect to corrosion rate in a simple way we introduce an integer number Ndiff. It means that for each time step related to corrosion we perform Ndiff steps of random walk assuming that the diffusion processes are faster than the corrosion ones.

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We launch the simulations with initial conditions as follows. The configuration is of 998x998 M sites surrounded by a frame of neutral wall sites so that the size of the box is 1000 x 1000. Two M sites at the positions (999, 500) and (999, 501) are then turned into R sites mimicking local damage initiating the evolution of the corrosion process. The above schemes and the values that we use for the set of probabilities are only qualitative and conceived on what is generally known or accepted about corrosion processes in some class of materials. The generic features of the corrosion process should not depend, however, on the detailed form of our assumptions. 3. Results Consider the particular case in which there are no SSE reactions, psse=0. The initial spot of the surface is passivated and we have to wait a number of time steps Nt (of the order of 1/poxi) for the corrosion process to start effectively. The development of the front can be characterized by the time evolution of the equivalent radius, R, of the forming cavity defined according to R=(2Ncorroded/ヾ}½ where Ncorroded is the number of metal sites corroded after Nt time steps. For psse=0 this evolution is simply proportional to Nt with the slope dR/dNt is consistent with simple mean field predictions [6]. In this paper we concentrate on the effects of diffusion and spatially separated anodic and cathodic reactions on the incubation time. For all the simulations reported we select:, pcor1=0.9, pcor2 =0.02 and poxi =0.001 or 0.002. Concerning psse, we have retained two values psse=0.1 or 0.2. We take the rate of diffusion, Ndiff =1 or 10. For the case Ndiff = 1. in Figure 1a and 1b there are snapshots corresponding to psse=0.1 and 0.2 at the moment when the corrosion front reaches the boundary of the simulation box. Roughly the corrosion front appears as a half-circle decorated with cauliflower like structures. The interior of the circle looks like a patchwork of acidic and basic zones. In Figure 2 we present an example of an R vs Nt dependence for pss=0.1 and poxi=0.001. Initially we observe a rather slow front development characterized by a suite of plateaus corresponding to the full repassivation of the surface after isolated dissolution events. Then there is a sharp high rocketing of the front development after a time somewhat above 4000 steps for this particular simulation run. We refer to the moment of this abrupt change as the incubation time. This behavior is quite general. In Figure 3 we show this dependence for another simulation run with the same set of parameter values (case a). For this run the incubation time is about 2000 time steps. In the same figure we plot the number N+ of A sites as the function of Nt. We can see that there is no net production of A sites until the incubation where N+ starts high rocketing too. This means that there is a certain distribution of A and C sites in the solution creating heterogeneities observed in Figure 1. For Nt larger than the incubation time. For Ndiff=10 we have basically the same type of behavior (case (b) in

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Figure 3). However the incubation time is longer and we can observe relaunching of the SSE mechanism and its regression due to the neutralization by diffusion. This is marked by the peaks in the N+ dependence on Nt before the incubation. This suggests that the faster is diffusion compared to chemical reaction rates the longer is the incubation time on the average. This we indeed observe in our simulations. To characterize the statistical distribution of incubation time and its dependence on the passivation probability poxi in Figures 4a and 4b we present histograms for 100 runs with poxi=0.001 and poxi=0.002 respectively. The histograms indicate a roughly distorted Gaussion distribution of the incubation times as observed experimentally [2]. In our model the probability poxi is related to a mechanism of dissolution of a passive site due to the presence of aggressive anions in the environment. Thus to make a comparison with experimental results we may assume that poxi is related to the concentration of Cl- ions, in this case we predict, in agreement with experiments that a decrease in the Clconcentration increases the induction time [1]. 4. Conclusion A simple stochastic model reproduces the main features observed in real experiments. The results seem quite general. Before the incubation time slow passivation/depassivation processes determine the evolution. After that the corrosion speeds up. This simple scenario suggests existence of scaling laws from which a reliable prediction of the long time behavior is possible. Acknowledgment On the Polish side this work is financed by the Ministry of Science and Informatization for the years 2004-2005, grant no. 3 T08A 010 26. References 1. Hoar T. P., Jacob W. R. Nature, 216(1967) 299-1301. 2. Shibata T., Takeyama T. Nature, 263 (1976).315-316. 3. Saunier J., Chausse A., Stafiej J., Badiali J.P. Journal of Electroanalytical Chemistry, 563 (2004) 239-247. 4. Wang M., Pickering H. W., Xu Y., Journal of Electrochemical Society, 142 (1995) 29862995. 5. Taleb A., Chausse A., Dymitrowska M., Stafiej J., Badiali J.P., Journal of Physical Chemistry B..108 (2004) 952-958. 6. Vautrin-Ul C., Chausse A., Stafiej J., Badiali J.P., Polish Journal of Chemistry. 78 (2004) 1795-1810.

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(a)

(b)

Figure 1. Snapshots obtained when the corrosion front reaches the border of the simulation box. The simulation parameters are: poxi=0.001, pcor1=0.9, pcor2=0.02, Ndiff=1, a) psse=0.1 and Nt=3776 b) psse=0.2 and Nt=4693.

Figure 2. The equivalent radius of a half-circle R=(2 Ncorroded/ヾ)½ versus Nt for a typical simulation run with parameters as in Figure 1a.

Figure 3. (a) Nt dependence of the equivalent radius R and the number of A sites N+ for a simulation run different than the run of Figure 2 although corresponding to the same set of parameters. Case (b) is different in that diffusion is ten times faster, Ndiff=10, than in case (a) where Ndiff=1.

(a) (b) Figure 4 Histograms of incubation times obtained from 100 independent runs for (a) the set of parameters of Figure 4. (b) with depassivation probability poxi=0.002 - two times higher than in (a).

Section J Surface Modifications and Inhibitors (for Improved Corrosion Resistance and/or Adhesion)

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On the Interaction of Organic Molecules with Metal Oxide Surfaces H.Terryna,b*,O.Blajeva,S.Van Gilsa,J.van den Brandb, A. Ithurbided, E.Vandeweertc, J. Snauwaertc, A.Hubina, C.Van Haesendonckc a

Dep. Metallurgy, Electrochemistry and Material Science Vrije Universiteit Brussel, Pleinlaan 2 1050 Brussel, Belgium Tel 003226293537, Fax 003226293200 e-mail:[email protected] b Netherlands Institute for Metals Research [NIMR] , Rotterdamseweg 137 2628 AL Delft, The Netherlands c Laboratorium voor Vaste-Stoffysica en Magnetisme, Katholieke Universiteit Leuven Celestijnenlaan 200 D, B-3001 Leuven, Belgium d Ecole Nationale Supérieure de Chimie de Paris, 11 Rue Pierre and Marie Curie 75231 Paris, Cedex 05

Abstract - In this paper the interaction between organic molecules and aluminium oxide is considered. The overall study concentrates on the chemical interface which is created between an organic layer and aluminium oxide. The knowledge of the composition and structure of hydrated films on aluminium is important for the bonding between aluminium and polymers. Native surface layers on aluminium consist of an anhydrous inner layer of amorphous Al2O3 and an hydrated outer layer, where the outer layer determines the adhesion properties. The pseudoboehmite oxyhydroxide surface can be seen as a model for the chemistry found at the surface of the air-formed film after ambient exposure of the aluminium metal. XPS was used to determine the hydroxyl activity on prepared Al substrates. The studied surface treatments were vacuum evaporated, thermally oxidized, alkaline etched, acid etched and boiled in water. As roughness is important electropolished aluminium samples were used. FESEM, FEAUGER, AFM and XPS were used to see the distribution of the organic molecule over the surface. FTIR-RAIRS was used to study the nature of the interfacial bonding. In this contribution some data about phosphonic

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acids and carboxylic acids are reported. The bonding of the carboxylic organic molecules was correlated with the presence of OH activity on the Al surface. Keywords: adhesion, hydration, aluminium oxide, optical reflection 1. Introduction In practice, bonds are often made between metals and polymers. Examples of this are organic coatings that are applied onto metallic substrates for the purpose of corrosion protection and also structural adhesive bonds. When these bonds are exposed to aqueous environments, it is often observed that there is a deterioration of the bond between the metal and the polymer. A corrosion protective coating will then no longer be protective towards the underlying metallic substrate and for a structural adhesive bond this will result in failure of the structure. Despite investigations over the past decades, there is still a lack of knowledge of how metals bind to polymers and why the stability of this bond deteriorates in the presence of water. A significant part of this lack of knowledge is because the metal-polymer interface cannot be directly and nondestructively investigated using standard measuring techniques. The aim of this work is to obtain some basic knowledge on the type of bonds that are formed between typical polymers and aluminium oxide surfaces and to determine whether these bonds are stable in the presence of water. Moreover, we want to determine which factors of the aluminium oxide surface are of influence to this bonding. Ultimately, this knowledge is to be used to develop an aluminiumpolymer system which is more durable in the presence of water. In this contribution, it will be shown that the knowledge of the composition and structure of hydrated films on aluminium is important for the bonding between aluminium and organic molecules. Native surface layers on aluminium consist of an anhydrous inner layer of amorphous Al2O3 and an hydrated outer layer, where the outer layer determines the adhesion properties. The pseudoboehmite oxyhydroxide surface can be seen as a model for the chemistry found at the surface of the air-formed film after ambient exposure of the aluminium metal. Spectroscopic ellipsometry (SE) was studied in the past to quantify these aspects [1,2,3]. XPS was used to determine the hydroxyl activity on the prepared Al substrate [4,5]. Finally, the interaction of the aluminium oxide surfaces prepared in different ways, with different organics such as carboxylic acids esters (mono functional myristic acid and bifunctional succinic acid), phosphonic acids, benzohydroxamic acids and silane molecules is considered. To allow us to study the interaction of these polymers with the aluminium oxide surface, a model compound approach is followed. The polymers are simplified to so-called model adhesion compounds. The model compounds are adsorbed as a thin layer on the differently pretreated aluminium substrates and the presence of the molecule is studied using FESEM, FEAUGER, AFM and XPS. The nature of the bonding is studied with infrared reflection absorption spectroscopy

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(FTIR-RAS). By combining the study of the bonding behaviour with the study of the oxide layer chemistry it can be determined how bonding of organic functional groups is influenced by the chemistry of the aluminium oxide surface [6,7,8]. 2. Results and conclusions 2.1. Study of the deposition of the organic molecule on the aluminium surface In the first part of the study we want to show how we study the repartition of the molecules over the surface. AFM was used to study the flatness of the polished Al samples and to observe the result of the interaction of the organic molecules with the Al surface. In order to obtain very flat surfaces, Al was first electropolished followed by the formation a barrier anodised layer to control the thickness of the aluminium oxide film (Figure 1, left panel). The adsorption of an organic molecules (for example benzohydroxamic acids from water solutions) results in an increase of the surface roughness (Figure 1, right panel).

Figure 1 Left: AFM picture of a barrier anodized Al surface (200V) resulting in a relatively smooth surface with a rms roughness of (0,46 ± 0,05) nm measured on an area of 1x1 µm². Right: upon deposition of an organic molecular (benzohydroxamic acids) overlayer on an anodized surface (40V), the rms roughness determined from the same area increases to (3,84 ± 0,13) nm.

A second example which is shown is the deposition of phosphonic acid (R-CPO(OH)2 on different alumina surfaces (Figure 2) as observed in the FESEM. The substrates were prepared by alkaline treatment, electro-polishing, and anodizing at different voltages of 99,99% pure Al. Part of the samples was then immersed in a 10-3 M water solution of the phosphonic acid in deionized water.

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Figure 2: FE-SEM images of a barrier anodized aluminum (200V) after deposition of the organic compound; x10 000 (left), x23 000 (right).

XPS survey spectra were used for a qualitative chemical analysis of the surface before and after the deposition of the organic compound. The results relevant for the case of an alkaline pretreated specimen are shown on Figure. 3. It follows from them that, along with the components of the oxide and the carbon contamination, a certain amount of phosphorous is detected.

Figure 3: XPS spectra of an alkaline treated surface: blank (left); after a deposition of the phosphonic acid (right)

The deposition of the phosphonic acid increases the complexity of the studied system and brings significant difficulties in the interpretation of the XPS data. Indeed, the layer due to the organic deposit has to be considered as a further contribution to the carbon and oxygen lines. The intensities of the four previous components of the C 1s core level would increase, because they are all present in the organic molecule, and in addition a new line of the C – P interaction can be expected. The composition of the O 1s peak is complicated as well because it

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depends on the number and nature of the X elements as it can be seen in the scheme below. R1

O Xh+ Oh-

P

C

Xh+ Oh-

where X can be H, Al, or vacant

R2

R3

Figure 4 : XPS spectra of the compound deposited on a anodized (10V) sample: P2s (left), P2p (right).

Nevertheless, the XPS results can be used to estimate the concentration of the organic compound on the surface as for this purpose the intensity of the P2s and P2p signals can be used. An example of those peaks for an anodised sample is shown in Figure 4. An estimate of the surface concentration of the adsorbed compound jA is carried out by the following formula: jA

=

IA

S

EA N

IS

A

ES M

cos

where IP2s (IP2p) and IAl2p are the integrated adsorbate and substrate line intensities, A and S are the modified Scofield photo ionization cross sections, is the density of the substrate, EA and ES are the kinetic energies of photoelectrons, is the mean free path of electrons within the substrate, is the angle of collection with respect to the sample normal, M is the relative atomic mass of the substrate, and N is the Avogadro constant.

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The results show that for the investigated samples, namely, the alkaline pretreated and anodized (10V) ones, the surface concentration of the organic compound is about 1014-1015 molecules/cm². This outcome was crosschecked by assuming that the footprint of an individual molecule occupies 0.36 nm² which area was obtained from the molecular lateral dimensions. In that case, the surface concentration of the organic compound was found to be about 3.1014 molecules/cm². There is, therefore, a good agreement between those two evaluations, and thus, the XPS data suggests that there is a dense organic layer coverage. While the above results argue that the particular phosphonic acid is strongly interacting with the alumina surface, they as well contain an inherent uncertainty pertinent to the described method for XPS based surface coverage determination. It is invoked by the fact that the same result can be achieved from the same amount of organic compound, which is however distributed over the surface not as a monolayer, but as multilayer islands. In order to resolve this controversy a high-resolution FE-Auger mapping was performed. On Figure 5 a map of the PLMM transition on an anodized sample with a deposited layer is shown. It follows from it that indeed the coverage of the phosphonic acid is continuous over the surface and moreover there are not significant variations of the phosphorous concentration. It can be thus concluded on the base of the XPS and FE-Auger experiments that the investigated phosphonic acid tends to form uniform and likely monolayer deposits on the aluminium oxides.

Figure 5. FEG-Auger mapping of the PLMM transition on an anodized sample. The magnification is x23000. The length of the red marker corresponds to 1 mkm on the surface.

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2.2. Study of the bonding of organic molecules as function of the Al pretreatment An angle-resolved XPS investigation was performed to determine the amount of hydroxyls on five differently pretreated aluminum substrates [4], see figure 6. These surface treatments were an alkaline and an acidic pretreatment, oxidizing of aluminum in vacuum, dehydroxylation of aluminium at 275 oC and boiling in water, resulting in the formation of a pseudoboehmite oxyhydroxide layer. To determine the amount of hydroxyls, restricted curve-fitting of the O 1s core level peak was performed (figure 6).

Figure 6: Example of XPS fitting of the O 1s peak [4].

Type of oxide Dehydrated 24 hours 275°C Vacuum evaporated and oxidised Acid Pretreated Alkaline Pretreated Pseudoboehmite

OH fraction (15° angle) 0.09 0.21 0.35 0.43 0.47

Table 1: OH fractions as determined with XPS [4].

The differently pretreated aluminium substrates were found to be enriched in hydroxyls towards their surface region and showed clear variations in the amount of hydroxyls on their surfaces as shown in Table 1 [4]. The highest hydroxyl fraction 47% (of the total oxygen content present on the oxide layer surface), was obtained for pseudoboehmite, corresponding well to its composition, AlOOH. Next ended the alkaline pretreated aluminium, which had a hydroxyl fraction of 43% on its surface. The lowest hydroxyl fraction was found for dehydroxylated alumina, which had a hydroxyl fraction of only 9%.

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The same set of pretreated aluminium substrates was investigated with respect to their bonding with polymers. To allow for a detailed investigation, a model adhesion compound approach was followed. As a simplified model for a given polymer, a small molecule, containing the same, for bonding relevant, functional groups is selected. The compounds were adsorbed on the aluminium substrates as thin layers. Subsequently, bonding of the compounds was studied using FTIR-RAS [7-8]. The bonding of two general classes of organic functional groups was investigated: functional groups capable of chemisorption with the oxide surface and functional groups capable of physisorption. From a macroscopic adhesion point of view, chemisorptive bonding is preferred over physisorptive bonding because the bonding energy is roughly an order of magnitude larger. Model compounds based on carboxylic acids were chosen to represent functional groups capable of chemisorption and model compounds based on ester groups were chosen to represent functional groups capable of physisorption.

myristic acid

Figure 7: Intensity of IR peak responsible for binding of the myristic acid (carboxylate peak) as function of the OH fraction as determined with XPS [7].

Clear differences were found among the differently pretreated substrates with respect to the amount of organic functional groups that could bond to their surfaces. It was found that the substrates with more hydroxyls on their surfaces were capable of bonding both more carboxylic (as shown in Fig. 7) [7] and also more ester functional groups [8] (not shown in figure 7). It was as well clearly shown in ref [7] that interaction of the organic functional groups occurs through the hydroxyls on the oxide surface. In view of this and also of the fact the hydroxyl groups represent the most probable termination of alumina surface in general [9], it is believed that other than investigated here functional groups with enhanced reactivity towards them will also possess a good adsorption ability.

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References 1. S.VAN GILS, C.MELENDRES,E.STIJNS,H.TERRYN, “Use of in situ ellipsometry to study Al surfaces” Thin Solid Films,455-456 (2004) 742-746. 2. S. VAN GILS , C.A. MELENDRES, H. TERRYN, “Quantitative chemical composition of thin films with infrared spectroscopic ellipsometry : application ot hydrated oxide films on aluminium”, Surface and Interface Analysis, 35 (2003), 387-394. 3. C.A. MELENDRES, S. VAN GILS, H. TERRYN , “Toward a quantitative description of the anodic oxide films on aluminium”,Electrochem. Communications, 3 (2001) 737-741. 4. J.VAN DEN BRAND, H.TERRYN, W.SLOOF, H.DE WIT, “Quantitative determination of the amount of hydroxyls in aluminium oxide layers using XPS: two different approaches” Surface and Interface Analysis, 36, (2004),81-88. 5. J. VAN DEN BRAND, P.C. SNIJDERS, W.G. SLOOF, H. TERRYN, J.H.W. DE WIT “Acid-base characterisation of aluminium oxide surfaces using XPS” The Journal of Physical Chemistry B. B 108 (19) (2004) 6017-6024 6. J. VAN DEN BRAND, S. VAN GILS, P.C.J. BEENTJES, et al. ”Ageing of aluminium oxide surfaces and their subsequent reactivity towards bonding with organic functional groups“ Applied Surface Science 235 (4): 465-474 AUG 31 2004 7. J. VAN DEN BRAND, O. BLAIJEV, P.C.J. BEENTJES, et al. ”Interaction of anhydride and carboxylic acid compounds with aluminum oxide surfaces studied using infrared reflection absorption spectroscopy” Langmuir 20 (15): 6308-6317 JUL 20 2004 8. J. VAN DEN BRAND, O. BLAIJEV, P.C.J. BEENTJES, et al. ”Interaction of ester functional groups with aluminum oxide surfaces studied using infrared reflection absorption spectroscopy” Langmuir 20 (15): 6318-6326 JUL 20 2004 9. G.E. BROWN Jr. ”How minerals react with water” Science 294: 67-70 OCT 5 2001

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

685

Formation of Al-Si Composite Oxide Films on Aluminum by Electrophoretic Sol-Gel Coating / Anodizing H. Takahashia*, M. Sunadaa, T. Kikuchia, M. Sakairia, and S. Hiraib a

Graduate School of Engineering, Hokkaido University; [email protected] b Faculty of Engineering, Muroran Institute of Technology, Muroran, Japan

Abstract - Highly pure aluminum specimens (99.99 %) with electropolishing or DCetching were covered with SiO2 films by electrophoretic sol-gel coating, and were anodized in a neutral borate solution. Structure and dielectric properties of the anodic oxide films were examined by SEM, TEM, EDX, and electrochemical impedance spectroscopy. It was found that anodizing of both specimens after electrophoretic deposition lead to the formation of anodic oxide films consisting of an inner alumina layer and an outer Al-Si composite oxide layer. The anodic oxide films formed thus had slightly higher capacitances than anodic oxide films on aluminum without any coating. Higher heating temperatures after electrophoretic deposition caused the capacitance of anodic oxide films to increase. Keywords : Aluminum, Electrolytic capacitor, Electrophoretic deposition, Sol-gel coating, Composite oxide

1. Introduction Aluminum electrolytic capacitors, which use aluminum anodic oxide films as a dielectric layer, are used in many electric circuits. Recent development in mobile electronic devices, such as notebook computers, cellular phones, and so on, strongly requires electrolytic capacitors with high electric capacitance. The capacitance, C, of the electrolytic capacitor can be expressed by the following equation. C = e0eS / d (1)

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where ε0 is the vacuum permittivity, ε the relative dielectric constant of anodic oxide films, S the surface area, and δ the film thickness. Increase in S is attained by electrochemical etching of the aluminium substrate before anodizing. Decrease in δ can be achieved by forming anodic oxide films with low K-values. The K-value is defined by the ratio of film thickness against film formation potential ( δ / Ea ). Anodic oxide films containing crystalline oxides have low K-values, since strong … chemical bonds in crystalline oxides enable … higher electric fields to be sustained. Increase in ε of anodic oxide films can be achieved by incorporating valve metal oxides with high ε values. In a previous work, the authors developed new processes, which involved anodizing after sol-gel dip coating of valve metal oxides1-3). It was found that anodizing of Al specimens after SiO2-coating causes the formation of anodic oxide films including an Al-Si composite oxide layer, and that the films have high voltage sustainability and at most about 20% higher parallel capacitance than Al2O3 films formed on electropolished specimens. However, the sol-gel dip coating was insufficient for uniform coating of SiO2 films on DC-etched aluminum with micro-tunnel pits, which is used in aluminum electrolytic capacitors. In the present investigation, the authors attempt to form Al-Si composite oxide films by electrophoretic sol-gel coating4, 5) and anodizing on both electropolished and DC-etched aluminum foils to examine dielectric properties of anodic oxide films as functions of heating temperature in SiO2 coating. 2. Experimental 2.1. Specimens and SiO2 -film coating by electrophoretic sol-gel method Highly pure aluminum foil (99.99 %) was used as specimens after electropolishing (Specimen-I) or DC-etching (Specimen-II, Nippon Chemicon Co). Preliminary experiments showed that the surface area of specimen-II is twenty times as large as that of Specimen-I. In the present study, currents in electrophoretic sol-gel deposition and anodizing were determined so that the current density is in the same range on both types of specimens. A sol composed of tetra-ethoxysilane ( Si(OC2H5)4 ), NH3, distilled water and dehydrated ethanol (1 : 0.07 : 5.75 :17.6 in molar ratio) was prepared in a glove box at room temperature under N2-atmosphere. Specimens-I and -II as anodes and a platinized Pt electrode as a cathode were dipped into the sol, and constant currents of ic = 0.05 - 0.2 A m-2 were applied until cell voltage, Vc, reaches 10 V. After the electrophoretic deposition, the specimens were withdrawn from the sol at 0.3 mm s-1, dried at 298 K for 300 s, and heated at Th = 573, 673, and 773 K

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for th = 1.8 ks in O2 atmosphere. After coatings, mass gains of the specimens were measured by a microbalance. For comparison, Sol-gel dip coating was also carried out by dipping Specimens-I, and -II into the sol for 300 s on open circuit, and then drying / heating under the conditions described above. 2.2. Anodizing and Characterization The specimens coated with / without SiO2 were anodized with a constant current of ia = 10 A m-2 at Ta = 293 K in 0.5 kmol m-3-H3BO3 / 0.05 kmol m-2– Na2B4O7 solution. The change of anode potential, Ea (vs. Ag/AgCl), with time, ta, during anodizing was monitored by a digital multimeter connected to a PC system. Anodic oxide films formed after SiO2-coating were examined by transmission electron microscopy (TEM), energy dispersed X-ray analyzer (EDX), and scanning electron microscopy (SEM). Electrochemical impedance spectroscopy (EIS) was performed in 0.5 kmol m-3-H3BO3 / 0.05 kmol m-2 Na2B4O7 solution using sinusoidal wave of 10 mV amplitude in 0.1 Hz – 50 kHz and parallel capacitances of anodic oxide films, Cp, were estimated by analyzing Bode diagrams of EIS. 3. Results and discussion 3.1. SiO2 films coated by electrophoretic deposition Figure 1-a shows the time-variations in cell voltage, Vc, during electrophoretic deposition of SiO2 on Specimen-I with different current densities of ic = 0.05, 0.1, and 0.2 A m-2. On all specimens, Vc shows a jump at the initial stage of deposition, and then, increases linearly with coating time, tc. With increasing ic, the Vc-jump becomes more remarkable and the slope of Vc vs. tc curves becomes steeper. The Vc-jump is due to IR-drop of sol, and the increase in Vc with ta is due to the deposition of SiO2 sol and the formation of anodic oxide films. TEM image of the cross section of Specimen-I after electrophoretic deposition up to Ec = 10 V showed that there are two layers on the specimen surface: an inner Al2O3 layer with 15 nm thickness and an outer SiO2 layer with 80 nm thickness. Figure 1-b shows the time-variations in cell voltage, Vc, during electrophoretic sol-gel coating of SiO2 on Specimen-II with ic = 0.01. 0.02, 0.05, 0.1, and 0.15 A m-2. All the specimens show a jump in Vc at the initial stage and then a linear increase in Vc with tc.

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8 Cell voltage, Vc / V

10

ic = 0.2 Am-2

ic = 0.15 Am-2

8

ic = 0.1 Am-2

6 4 -2

i c = 0.05 Am

2 0

Cell voltage, Vc / V

10

ic = 0.10 Am-2

6

ic = 0.05 Am-2

4 ic = 0.02 Am-2

2

ic = 0.01 Am-2

0

1.0 2.0 Coating Time, Tc / ks

3.0 0 0

1.0 2.0 Coating Time, Tc / ks

3.0

Fig.1: Change in cell voltage, Vc with time, ta during electrophoretic sol-gel coating on a) specimen-I (left) and b) specimen-II (right) Comparison of Fig. 3 with Fig. 1 indicates that the voltage-jump on Specimen-II is higher than that on Specimen-I, and that the slope of Vc vs. tc curve on Specimen-I is as high as that of specimen-I at each ic. This suggests that IR drop of sol in micro-pits of Specimen-II is appreciably large, and that electrophoretic deposition on Specimen-II is similar to that on Specimen-I. Table 1 Mass gain of Specimens-I and -II by electrophoretic- and dip coating, Dip coating Specimen-I SpecimenII

0.03 g m-2

ic=0.05A/ tc=5440s 0.15 g m-2

0.05 g m-2

0.36 g m-2

Electrophoretic coating ic=0.2A/m2 ic=0.1A/m2 tc=1117s tc=2497s 0.13 g m-2 0.15 g m-2 0.35 g m-2

0.31 g m-2

(Th = 573 K, th = 1.8 ks)

Table 1 shows the mass gain of Specimens-I and -II by dip-coating and electrophoretic coating with ic = 0.05, 0.1, and 0.2 A m-2. After electrophoretic deposition, the specimens were heated at Th = 573 K for th = 1.8 ks in O2atmosphere. The mass gain of each specimen by electrophoretic deposition is much larger than that by dip coating, suggesting the effect of electric field is significant for uniform coating of SiO2. Larger mass gain after electrophoretic coating on Specimen-II than that on Specimen-I is due to the deposition in tunnel pits of Specimen-II.

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3.2. Anodic oxide film formation on SiO2-coated specimen

300

300

250

ic = 0.05 Am-2

250 i = 0.1 Am-2 c

200

200

150

ic = 0.2 Am-2

150

100 50

Anode potential, Ea / V (vs. Ag / AgCl)

Anode potential, Ea / V (vs. Ag / AgCl)

Figure 2-a shows Ea vs. ta curves for Specimen-I, which was coated with SiO2 layers at ic = 0.1 and 0.2 Am-2 before heat treatment at Th = 573 K for th = 1.8 ks. The Ea vs. ta curve for Specimen-I without any coating is also indicated as a solid line. Specimen-I coated with SiO2 shows a 10 V-jump in Ea at the initial stage, and then linear increase in Ea with ta. The slope of Ea vs. ta curves for Specimen-1 with SiO2-coating is slightly steeper than that without SiO2-coating, and does not depend on ic. Figure 2-b shows the Ea vs. ta curves for Specimen-II coated with SiO2 layers at ic = 0.05, 0.1 and 0.2. A/m2. Heat treatment conditions are the same as those in Fig. 2-a. The Ea vs. ta curve for each specimen shows a slightly up-curved line, and the slope of the curve is steeper at lower ic. The steeper slopes of the Ea vs. ta curves on SiO2-coated specimens suggest the formation of Al-Si composite oxide films. The effect of heat treatment temperature, Th, after electrophoretic deposition on the Ea vs. ta curve was examined, and it was found that higher Ths lead to higher slopes of the curve.

No-coating ic= 0.1 Am-2 ic= 0.2 Am-2

0 0 100 200 300 400 500 600 700 Anodizing time, ta / s

100

No-coating

50 0 0 100 200 300 400 500 600 700 Anodizing time, ta / s

Fig. 2 Change in anode potential, Ea, with time, ta, during anodizing of a) Specimen-I (left) and b) Specimen-II (right) after coating with SiO2 film

The TEM and EDX analysis of anodic oxide films formed on Specimen-I by SiO2-coating and anodizing showed the formation of anodic oxide films consisting of Si-Al composite oxide and pure alumina layers. These layers became thicker with Ea, while SiO2 layer became thinner. The anodic oxide films were estimated to have 1.1 nm / V of K-value.

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3.3. Dielectric properties of anodic oxide films formed on SiO2 coated specimen The EIS analysis showed that the reciprocal of parallel capacitance, 1 / Cp, of anodic oxide films is proportional to anode potential, Ea, on Specimens-I and -II with / without SiO2. The slope of 1 / Cp vs. Ea curves on the specimens with SiO2 coating was slightly flatter than those without coating, and that the tendency became more remarkable at higher heating temperatures. Consequently, electrophoretic sol-gel coating of SiO2 is effective for developing new types of aluminum electrolytic capacitors with high electric capacitance. 4. Conclusion Electrophoretic sol-gel coating of SiO2 was performed on DC-etched and electropolished aluminum, and then anodizing was carried out to examine the dielectric properties of anodic oxide films formed by the two successive steps. The following conclusions may be drawn. 1. Electrophoretic sol-gel coating can form SiO2 films uniformly on the surface of plain and DC-etched specimens. A thin oxide layer is formed by anodic oxidation during the formation of SiO2 film. 2. Anodizing of aluminum coated with SiO2 films by electrophoretic sol-gel coating leads to the formation of oxide films, which consist of an outer Al-Si composite oxide layer and an inner Al2O3 layer underneath SiO2 layer. The anodic oxide films formed thus has high electric field sustainability. 3. The anodic oxide films formed on specimens with SiO2 film have at most 10% higher parallel capacitance than those without SiO2 layer. Higher heating temperatures after electrophoretic deposition lead to the increase in capacitance of anodic oxide films. References [1] K. Watanabe, M. Sakairi, H. Takahashi, S. Hirai and S. Yamaguchi; J. Electroanal. Chem., 473, 255 (1999) [2] K. Watanabe, M. Sakairi, H. Takahashi, K. Takahiro, S. Nagata, and S. Hirai; Electrochemistry, 67,1243 (1999) [3] K. Watanabe, M. Sakairi, H. Takahashi, K. Takahiro, S. Nagata, and S.Hirai; J. Electrochem. Soc., 148, B473 (2001) [4] K. Hasegawa, M. Tatsumisago, and T. Inami: J. Ceramic Soc. Jpn., 105, 569 (1997) [5] S. Sakka ; “Science of sol-gel method”, p.8 (Agne-shouhusha, 1996)

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Published by Elsevier B.V.

691

In-situ Raman Spectroscopy and Spectroscopic Ellipsometry study of the iron/Polypyrrole interface. T. Van Schaftinghena, S. Joiretb, C. Deslouisb, H. Terryna a

Department of Metallurgy, Electrochemistry and Materials Science, Vrije Universiteit Brussel, Pleinlaan 2, 1050 Brussels, Belgium, [email protected] b LISE "Laboratoire Interface et Systèmes Electrochimiques", Université Pierre et Marie Curie, 4 place Jussieu, 75252 PARIS Cedex 05, France

Abstract - Different pre-treatment procedures of the iron surface have been developed with the intention to alter the iron/Ppy interfacial properties. These results were correlated with the impedance response of Ppy coated iron samples in a corrosive aqueous environment. Keywords : Iron, Magnetite, Iron oxalate, EIS, In-situ Raman Spectroscopy

1. Introduction. In the corrosion protection mechanism of Intrinsically Conducting Polymers (ICP), the metal/polymer interface plays a very important role [1-4]. It partly determines the adhesion of the ICP film on the metallic surface, it has intrinsic barrier properties and it affects the charge carrier transfer related to the electrochemical reactions that take place during the corrosive degradation of the system. Different pre-treatment procedures of the iron surface have been developed with the intention to alter the iron/Polypyrrole interfacial properties [5]. During the pre-treatments, different passive states of the iron surface in an aqueous oxalic acid environment are combined and the resulting surface properties were determined in-situ with the Raman Spectroscopy and the Spectroscopic Ellipsometry techniques.

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2. Experimental. A LABRAM® Raman spectrophotometer from Jobin-Yvon® was used with a Helium-Neon laser (l = 632.81 nm) power of 0.01 mW and an analysed surface of approximately 1 µm² (lens magnification = 80). The electrochemical experiments were carried out in a three-electrode set-up (in a 0.05 M oxalic acid-dihydrate electrolyte) using a Saturated Sulphate reference Electrode (SSE) and a platinum wire counterelectrode, both connected to a PGSTAT30Ò potentiostat from EcochemieÒ. The working electrode consisted of glass substrates covered with a 2 µm Radio Frequency (RF) plasma deposited iron layer with a characteristic roughness shown in Figure 1 a). 668

I / AU

Therm al Rough Bulk

B/ A = 25

B

538

A

500 nm a)

250 b)

500

750

Wavenum ber / cm

1000 -1

Figure 1: a)SEM pictures of RF plasma deposited iron. b)Ex-situ Raman spectra of a polished iron sample and the rough substrate of a) after curing at 250 °C during 1 h in an Ar atmosphere.

This roughness enhances Raman scattering and in order to determine the nature of this enhancement a polished iron substrate and the rough substrate of Figure 1 a) were cured at 250 °C during one hour in an argon atmosphere to grow identical magnetite layers. According to Oblonsky et al. [6] the classical Raman spectrum of magnetite contains a strong band at 668 cm-1 and a weak band at 538 cm-1, while the SERS spectrum contains a weak band at 680 cm-1 and a strong band at 578 cm-1. The Raman spectrum of the bulk substrate in Figure 1 b) shows that pure magnetite is formed during the thermal process. The enhanced spectrum measured on the rough sample shows the same characteristic bands, which excludes the SERS nature of the enhancement. One straightforward way to explain the signal amplification is the increase of the exposed surface due to the roughness. We can namely expect that a larger quantity of products accumulates on the surface, which generates more Raman scattering.

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The in-situ Visual Spectroscopic Ellipsometry (Vis-SE) measurements were done with a J.A. Woollam Co. VASEÒ (Variable Angle Spectroscopic Ellipsometer) working in the UV-visible-NIR (300-1700 nm) spectral range with a wavelength resolution of 10 nm at 70° angle of incidence. The electrochemical experiments (in a 0.05 M oxalic acid-dihydrate electrolyte) were carried out in a three-electrode set-up with a Saturated Calomel reference Electrode (SCE) and a platinum grid counter electrode, both connected to a PGSTAT30Ò potentiostat from EcochemieÒ. 3. Results. The following table gives an overview of the characteristic Raman bands of the species that are expected to grow on the analysed surfaces. Table 1: Characteristic Raman bands for magnetite and iron oxalate [7,8]. Wavenumber (cm-1) 668 705 519 582 913 1468 467 554 644 1621 1664 1688

Product Fe3O4

Fe(II)C2O4

j / A/ cm ²

Fe(III)C2O4

Associated vibrational mode Symmetric Fe-O stretching vibration C-C-O scissoring vibration O-C-O scissoring vibration C-C stretching vibration C-O stretching vibration O-C-O scissoring vibration O-C-O twisting vibration C-O stretching vibration C-O stretching vibration 554

1 V ( Ag/ AgCl)

668

1 st ep Spect rum 1 Spect rum 2

1688 1664 1621

705 I / AU

0.01 1 st ep Spect rum 1 Spect rum 2

467

x200

1E- 3 0 a)

50 s

500 45 0 s

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500 b)

750

1000

1250

Wavenum ber / cm

1500

1750

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Figure 2: a)Chronoamperometric curve of 1 step pre-treatment along with the points in time where the Raman spectra are recorded. b)Raman analysis of the sample surface during the 1 step pre-treatment. Spectrum 2 is amplified with a factor 200.

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The bands of spectrum 1 situated between 1000 cm-1 and 1500 cm-1 can be attributed to Fe(III)C2O4 since they are always observed in the presence of the characteristic bands for this product. The band at 705 cm-1 of spectrum 2 is associated to the oxidised form of magnetite. Spectrum 1 shows that Fe(III)C2O4 and Fe3O4 are present on the metallic surface at the initial stages of the process, while magnetite is detected along with its oxidised form at the end of the treatment when the system has reached a stable passive regime. This could indicate that Fe(III)C2O4 dissolves and/or oxidises to form Fe3O4 while the magnetite that is present from the beginning reaches a higher oxidation level as the pre-treatment goes on.

0 V ( Ag/ AgCl)

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2 st ep Spect rum 1 Spect rum 2 1E- 3

1688

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x5 1E- 4 0

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500 240 s 360 s

1000 Tim e / s

1500 12 00 s

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2000 18 00 s

b)

750

1000

1250

Wavenum ber / cm

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-1

Figure 3: a)Chronoamperometric curve of 2 step pre-treatment along with the points in time where the Raman spectra are recorded. b)Raman analysis of the sample surface during the 2 step pre-treatment. Spectrum 2 is amplified with a factor 5.

As can be observed in Figure 3 a), spectrum 1 has been measured during the first potential step while spectrum 2 was recorded during the second stage of the so-called 2 step pre-treatment. Fe3O4 and Fe(II)C2O4 are predominantly present in spectrum 1 although the weak bands between 1600 cm-1 and 1700 cm-1 reveal the presence of Fe(III)C2O4. The situation is analogue at the end of the process where Fe(III)C2O4 is mostly detected next to magnetite and small traces of Fe(II)C2O4 (weak bands at 519 cm-1, 582 cm-1 and 913 cm-1). This means that Fe(II)C2O4 is formed during the first step and oxidises at the higher potential of the second step to form Fe(III)C2O4. The chemical information obtained with Raman Spectroscopy is used to define the optical model for the Vis-SE data analysis. Figure 4 shows that the passive film in this model is composed of a magnetite and an iron oxalate layer where no distinction is made between Fe(II)C2O4 and Fe(III)C2O4. Since the measurements were carried out in-situ, the electrolyte must be included in the optical model. To determine the pre-treatment film thickness (sum of d1 and d2),

In-situ Raman Spectroscopy and Spectroscopic Ellipsometry study

695

the Mean Squared Error (MSE), which quantifies the difference between the measured value and the model value, is minimised in a non-linear least squares regression analysis. Wat er Opt ical Model

Physical Syst em

Oxalic acid

FeC O+ Fe 3 O 24 4 +FeC Fe 3O 2O44

FeC2 O4

d1

Fe 3 O4

d2

Fe

Fe

Figure 4: Representation of pre-treated iron with the corresponding optical model used for Vis-SE analysis.

The chronoamperometric graphs of Figure 5 a) and c) show that the in-situ VisSE measurements were carried out at the end of the potentiostatic pretreatments under stabilised conditions. 100

1 V ( Ag/ AgCl)

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0 c)

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6000 6140 s

400 d)

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Wavelengt h / nm

Figure 5: Chronoamperometric curve of a)1 step and b)2 step pre-treatment with the indications of the in-situ Vis-SE measurement timeframes. b) and d) Experimental and fit data for the respective Vis-SE measurements.

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The graphs in Figure 5 b) and d) show that the overall fit quality is rather poor for the two pre-treatments. Nevertheless, it is more satisfactory for the Dspectra, which relate more to the morphology, than for the Y-spectra, which are more closely related to the chemical properties. Table 2 displays the corresponding fit results. The high MSE value and the high error on the calculated value of d1 for the 2 step sample confirm these qualitative findings. This indicates that the used optical model isn’t fully suited for the interpretation of the ellipsometric data. Nevertheless, the values of Table 2 can be considered as good indicators for the surface morphology of the pre-treated samples. Table 2: Fit results and 90 % confidence limits for the in-situ Vis-SE measurements of 1 step and 2 step pre-treated samples. Fit parameter d1 (nm) d2 (nm)

1 step; MSE = 10.26 Value Error (%) 22.30 ± 16.54 20.99 ± 2.03

2 step; MSE = 24.58 Value Error (%) 19.08 ± 47.13 26.86 ± 3.62

According to the configuration of the optical model, the sum of d1 and d2 equals the thickness of the pre-treatment film. The values of Table 2 indicate that the thickness of the potentiostatically formed layers doesn’t exceed 50 nm. 4. Conclusions. The in-situ Raman Spectroscopy analysis shows that magnetite is predominantly present on the 1 step sample whereas the 2 step sample is mainly composed of Fe(III)C2O4. The thickness of the potentiostatic pre-treatment layers is estimated at about 50 nm. These interfacial properties have a strong impact on the electrochemical behaviour of ICP systems [5]. The optical model for the Vis-SE data interpretation has some shortcomings but the information input from the in-situ Raman spectra proved to be very useful. References. [1] Bernard M-C., Hugot-Le Goff A., Joiret S., Dinh N.N. and Toan N.N., J. Electrochem. Soc., 146 (1999) 995. [2] Bernard M-C., Joiret S., Hugot-Le Goff A. and Viet Phong P., J. Electrochem. Soc., 148 (2001) B12. [3] DeBerry D.W., J. Electrochem. Soc., 132 (1985) 1022. [4] Lacaze P.C., Camalet J.L., Lacroix J.C., Aeiyach S. and Chane-Ching K., Synth. Met., 93 (1998) 133. [5] Van Schaftinghen T., Deslouis C., Hubin A. and Terryn H., Accepted in Electrochim. Acta. [6] Oblonsky L.J., Virtanen S., Schroeder V., Devine T., J. Electrochem. Soc. , 144 (1997) 1604. [7] Shebanova O.N. and Lazor P., J. Solid State Chem., 174 (2003) 424. [8] Edwards H.G.M. and Russell N.C., J. Mol. Struct., 443 (1998) 223.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Spontaneous grafting of iron surfaces by reduction of aryldiazonium salts in acidic water. Applications to the inhibition of iron corrosion Fetah I. Podvoricaa, Catherine Combellasb, Michel Delamarc, Frédéric Kanoufib and Jean Pinsond a

Chemistry Department of Natural Sciences Faculty, University of Prishtina, rr. “Nëna Tereze” nr. 5, Prishtina, Kosovo. E-mail: [email protected] b Laboratoire Environnement et Chimie Analytique, CNRS-ESPCI, 10 rue Vauquelin, 75231 Paris Cedex 05, France. c Institut de Topologie et de Dynamique des Systèmes, Université Paris VII Denis Diderot, 1 rue Guy de la Brosse, 75 005 Paris, France. d Alchimer, Z. I. de la Bonde, 15 rue du Buisson aux Fraises, 91300 Massy.

Abstract - Spontaneous grafting of iron in acidic aqueous solution was characterized by several methods. The organic film can be used either for direct protection of the metal against corrosion or for further derivatization. Inhibition efficiencies up to 97 % can be obtained when the grafted metal is left in the corrosive medium in the presence of the diazonium salt. Keywords: iron; diazonium salts; spontaneous grafting; corrosion inhibition

1. Introduction Electrochemical reduction of aryldiazonium salts at various carbon electrodes in an organic medium leads to the covalent grafting of the organic groups to the electrode surface [1-6]. This has also been demonstrated at various other semiconductive or conductive surfaces, such as silicon [7], iron and mild steel [8], zinc, nickel, cobalt, copper, gold, platinum [9],... The same electro-grafting reaction can be performed in acidic aqueous solution in the case of carbon [2] , iron and mild steel [8]. The grafting can also take place spontaneously without

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any electrochemical assistance in an aprotic medium onto carbon, copper, iron and zinc [10,11], palladium and semi-conductors (Si and GaAs)[12].Moreover, McCreery has described the spontaneous grafting of copper and aluminium alloys in acidic aqueous solution [13]. The electro-grafted film protects the surface against corrosion in acidic media [14]. It acts as a corrosion inhibitor, whose efficiency can be compared to that of classical corrosion inhibitors described in the literature. Another advantage of covalently bonded organic films is that they can be further derivatized chemically or photochemically[15]. In the present paper, we will show that it is possible to graft chemically aryl groups derived from diazonium salts onto iron surfaces in acidic aqueous or neutral water solution and that these groups protect iron against corrosion in a 0.01 M H2SO4 aqueous solution. The following diazonium salts have been used: + N2Ar BF4-, with Ar = C6H5 (1), C6H4Br (2), C6H4(CH2)11CH3 (3), C6H4CO2H (4). 2. Experimental Section The iron samples (iron electrodes and iron plates) were first polished with 1 mm diamond paste and then with 0.02 mm alumina slurry in Nanopure water to avoid organic contaminants on the electrodes. They were finally rinsed in acetone. The synthesis of diazonium salts has been described in previous publications [2,14]. The iron samples (plaques or 1mm diameter electrodes) were immersed into a 10 mM H2SO4 solution containing 10 mM of diazonium salt or into a 10 mM NaBF4 aqueous solution containing 10 mM of diazonium salt for 60 min. Following grafting, all samples were ultrasonicated in acetone for 10 min to remove any species adsorbed onto the surface. Polarization curves and impedances measurements were obtained with a 5 mV amplitude at the open circuit potential in the 104 – 10-2 Hz interval with 6 points per decade, with a potentiostat/galvanostat (CH660A, CH Instruments, USA). 3. Results and discussion 3.1. Reaction and analysis of the surface Spontaneous reduction of aryldiazonium salts by iron in acidic aqueous solution is possible because of the difference between the oxidation potential of iron (E°Fe2+/Fe = -0.68 V/SCE) and the reduction potential of aryldiazonium salts (peak potentials measured at a C or Fe electrode between –0.45 and 0.2 V/SCE in an aprotic medium) [2,14]. The reduction mechanism involves a first dissociative

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electron transfer that gives dinitrogen and aryl radicals. The latter bind to the iron surface to give organic mono or multilayers. XPS analysis show that organic film is present onto iron surface after chemical grafting and that it is several layers thick. XPS measurements also show that spontaneous grafting in water is at least as efficient as the electrochemical grafting [16]. The characteristic peaks for diazonium at 403.8 and 405.1 (N[1s]) [17] were not observed and therefore there is not any adsorption of the diazonium salt onto the iron surface. IR spectroscopy confirms the presence of phenyl moieties on iron surfaces after chemical grafting by (1). The method allows conclude to the presence of a multilayer film. 3.2. Impedance measurements The inhibition effect of the spontaneous grafting in an acidic or a neutral aqueous medium has been assessed by determining the polarization resistance, Rp, by impedance measurements. Rp values were calculated from Nyquist impedance diagrams (see Figures 1 and 2). Polarization resistances for bare iron electrode and the same electrode grafted by (1)-(3) in acidic aqueous solution are shown respectively in Figure 1) a, b, c and d. The polarization resistance value increases upon grafting, which indicates that the grafted film has inhibitor properties. The highest value is observed for the dodecylphenyl film (3), for which the polarization resistance is about 1,5 times higher than with (1) and almost four times higher than for bare iron. The latter result strongly suggests that the hydrophobic alkyl chains are standing upright from the iron surface as we have already observed when iron was grafted electrochemically in an organic medium [14]. When the iron electrode has been immersed for one hour in a 10 mM neutral solution of (2) (figure 2), its Rp value increases about 50 %. 21000

12000 1 Hz 10 Hz

8000 - Zim ( W )

- Zim ( W )

14000 7000

1 Hz

0 -7000

b c

a

0 10000

20000

d

4000 a

0

30000

0 -4000

Zre ( W )

0.1 Hz

6000

12000

b 18000

24000

Zre ( W )

Fig. 1. Impedance spectra for an iron

Fig. 2. Impedance spectra for an iron

electrode in a 0.01 M H2SO4 solution a) bare. b- d) modified chemically by respectively (1)-(3) in a 10 mM in a 10 mM H2SO4 solution.

electrode in 0.01 M H2SO4 solution. (a) bare and b) spontaneously grafted by (2) in neutral aqueous solution.

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3.3. Polarization measurements In all the examples of grafting in acidic or neutral aqueous medium, polarization curves exhibit a decrease of iron oxidation upon grafting (compare Figures 3, curves a-d), and (Figure 4 curves a-b), which means that the polyphenylene grafted organic film acts as a physical barrier between the metal and the aqueous medium. The grafted film plays the role of a cathodic inhibitor for iron as we observed a negative shift of the corrosion potential when the modified electrode was immersed into a 0.01 M sulfuric acid solution. The inhibition efficiency, I. E., was calculated from: I .E. = 100(1 -

film I corr film ) , where I corr and I corr

I corr represent the current intensities at respectively the coated and bare iron electrodes. The results for the corrosion intensities and inhibition efficiencies are gathered in Table 1. The best results were obtained with salts (1), (2) and (3) after a one-hour grafting. We can observe a decrease of the corrosion current by ~ 70 % upon grafting by (1) - (3). This is in agreement with the above determination of Rp by impedance measurements and also with previous results obtained for electrochemical grafting [14]. Figure 3 (curve e) shows the electrode was highly pasivated when after grafting by (2) it had been maintained in the diazonium solution. A 97 % decrease of the cathodic current was then observed. One can compare this result with that of the literature using classical corrosion inhibitors in aqueous sulfuric acid solution where inhibition efficiencies between 95.2 to 100 percent were obtained [18-21]. -3

-3 log I (I in A)

-5

d

-6

log I (I in A)

-4

-4

a cb e

-7

a b

-5 -6 -7 -8

-8

-9

-9 -1.1

-1

-0.9

-0.8

-0.7

-0.6

V/Ag/AgCl (V)

-1.1

-1

-0.9

-0.8

-0.7

V / Ag/AgCl (V)

Fig. 3. Semilogarithmic polarization curves for

Fig. 4. Semilogarithmic polarization

an iron electrode in 0.01 M H2SO4. a) bare. b- d) modified chemically by respectively (4), (2) and (3) in 0,01 M H2SO4. (e) left in a 10 mM solution of (2) after grafting. Scan rate, v = 1 mV/s.

curves for an iron electrode in 0.01 M H2SO4. a) bare and b) spontaneously grafted by (2) in neutral solution. Scan rate, v = 1 mV/s.

-0.6

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Applications to the inhibition of iron corrosion

It is possible to chemically derivatize the covalently grafted film of an iron electrode modified with diazonium salt (4) with no risk of ungrafting the overlayer. For that purpose, the electrode was reacted with octyltriethoxysilane, C8H17Si(OC2H5)3 (C8TES), in order to overcoat the primary layer by a polymer film using the procedure described by Aramaki et al [22]. Polarization curves (not shown) obtained with an iron electrode chemically grafted by (4), reacted with C8TES and left in a 0.01 M H2SO4 solution for two hours showed an inhibition efficiency of 85%. Table 1. Corrosion currents and inhibition efficiencies of bare and modified iron electrodes by chemical grafting of diazonium salts. Measurements in a 0.01 M H2SO4 aqueous solution. Diazonium salt, +N2Ar BF4none none 1 C6H5 2 C6H5 3 C6H4Br 2 C6H4Br 2,4 C6H4C12H25 2 C6H4CO2H 2 C8H17Si(OC2H5)3 5 C6H4CO2H + C8H17Si(OC2H5)3 2,6

1 1 2 2 3 4 4

Rp kW

Ecorr V/Ag/AgCl

Icorr 10-6 A

I. E. %

10.8 16.0 28.2 24.2 28.7 54.2 36.6 20.6 11.5 41.3

- 0.742 -0.825 -0.830 -0.841 - 0.827 - 0.685 - 0.864 - 0.787 - 0.778 - 0.899

3.20 1.41 1.05 0.71 1.07 0.12 0.89 1.60 2.44 0.57

68 50 67 97 73 50 25 85

1

bare iron electrode maintained 1 h in a neutral aqueous solution. 2electrode grafted for 1 hour by a 10 mM diazonium salt in 0.01 M H2SO4 aqueous solution. 3 same as 1) in the presence of 10 mM diazonium salt. 4 polarization measurements in the presence of 10 mM diazonium salt. 5 iron electrode treated with octyltriethoxysilane but not by a diazonium salt. 6 treatment by octyltriethoxysilane after grafting by (4).

4. Conclusion Spontaneous grafting of iron in acidic and neutral aqueous solution was observed by several methods. The organic film is several layers thick and can be used either for direct protection of the metal against corrosion or for further derivatization. Inhibition efficiencies up to 97 % can be obtained when the grafted metal is left in the corrosive medium in the presence of the diazonium salt. The advantages of this method are that it operates under simple and

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environment friendly experimental conditions in water or a fortiori in any organic solvent and in the presence of air. References [1] M. Delamar, R. Hitmi, J. Pinson, J.-M. Savéant, J. Am. Chem. Soc. 114 (1992) 5883. [2] P. Allongue, M. Delamar, B. Desbat, O. Fagebaume, R. Hitmi, J. Pinson, J.-M. Savéant, J. Am. Chem. Soc. 119 (1997) 201. [3] K Ray III, R.L. McCreery, Anal. Chem. 69 (1997) 4680. [4] C. Saby, B. Ortiz, G.Y. Champagne, D. Bélanger, Langmuir 13 (1997) 6805. [5] A.J. Downard, A.D. Roddick, Electroanalysis 7 (1995) 376. [6] J. K. Kariuki, M.T. McDermott, Langmuir 17 (2001) 5947. [7] C. Henry de Villeneuve, J. Pinson, M.C. Bernard, P. Allongue, J. Phys. Chem.B 101 (1997) 2415. [8] A. Adenier, M. C. Bernard, M. M. Chehimi, E. Cabet-Deliry, B. Desbat, O. Fagebaume, J. Pinson, F. Podvorica, J. Am. Chem. Soc. 121, (2001) 4541. [9] M. C. Bernard, A. Chaussé, M. M. Chehimi, E. Cabet-Deliry, J. Pinson, F. Podvorica, C. Vautrin-Ul, Chem. Mater. 15 (2003) 3450. [10] J.A. Belmont, R.M. Amici, C.P. Galloway, (Cabot Corp.) PCT WO 96 18, 688 [Chem. Abstr. 125 (1996) 144212t]. [11] A. Adenier, E. Cabet-Deliry, A. Chaussé, S. Griveau, F. Mercier, J. Pinson, C. Vautrin-Ul, Chem. Mat., 17 (2005) 491. [12] M. Stewart, F. Maya, D. Kosynkin, S. Dirk, J. Stapleton, C. McGuiness, D. J. Allara, J. Tour, J. Am. Chem. Soc. 126 (2004) 370. [13] B. L . Hurley, R. L. McCreery, J. Electrochem. Soc. 151 (2004) B252. [14] A. Chaussé, M. M. Chehimi, N. Karsi, J. Pinson, F. Podvorica, C. Vautrin-Ul, Chem. Mater. 14 (2002) 392. [15] A. Adenier, E. Cabet-Deliry, Th. Lalot, J. Pinson, F. Podvorica, Chem. Mat. 14 (2002) 4576. [16] C. Combellas, M. Delamar, F. Kanoufi, J. Pinson, F. I. Podvorica Chem. Mat. , 17 (2005) 3968. [17] P. Finn, W. L. Jolly, Inorg. Chem. 11 (1972) 1434. [18] A. Chetouani, K. Medjahed, K. Benabadji, B. Hammouti, S. Kertit, A. Mansri, Prog. in Org. 46 (2003) 312. [19] A. Chetouani, K. Medjahed, K. E. Sid-Lakhadar, B. Hammouti, M. Benkaddour, A. Mansri, Corr. Sci. 46 (2004) 2421. [20] T. Tüken, G. Arslan, B. Yazici, M. Erbil Corr. Sci. 46 (2004) 2743. [21] F. Bentis, M. Traisnel, H. Vezin, H. F. Hildebrand, M. Lagrenée, Corr. Sci. 46 (2004) 2781. [22] K. Aramaki, T. Shimura, Corr. Sci. 46 (2004) 2533.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

703

Conversion coating of zinc heptanoate in aqueous media on electrogalvanized steel S. JACQUESa, E. ROCCAa, M-J. STÉBÉb, H. DERULEc, N. GENETd, J. STEINMETZa a

Laboratoire de Chimie du Solide Minéral (UMR 7555) and bEquipe Physico-Chimie des Colloïdes (UMR 7565), Université Henri Poincaré, Nancy 1/CNRS, BP 239, F-54506 VANDOEUVRE LES NANCY, [email protected] c Arcelor Research, Voie Romaine, F-57280 MAIZIERES LES METZ d Total France, CReS, Chemin du canal, BP 22, F-69360 SOLAIZE

Abstract - The aim of this investigation is to develop a protective coating of zinc heptanoate in aqueous media on electrogalvanized steel. Fatty acids are insoluble in water so it requires adding surfactant. Firstly, heptanoic acid/water/surfactant diagram was established at 20°C to determine the different phase domains. In order to study these domains, further techniques were employed, such as optical microscopy and quasi-elastic light scattering. Then, coatings were performed and its morphology and composition were characterized by SEM and XRD. Afterwards, corrosion tests of coatings led the characterization of the protective properties. Finally, this ternary system easily allows deposition of both hydrophobic and corrosion resistant zinc carboxylate coatings. Keywords: zinc, carboxylate, surfactant, passive film, environment

1. Introduction The present study proposes an environmentally friendly alternative to the zinc phosphating conversion process. In previous work, a protective coating based on metallic soaps has been developed using heptanoic acid (HC7) solubilized in a water-ethanol solution [1]. In this study, to avoid addition of solvent, non ionic surfactant has to be employed for the solubilization of fatty acid in aqueous media.

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2. Experimental details Phase diagram Brij®58/HC7/water at 20°C: Brij®58 (C16H33(OCH2CH2)20OH) is an ethoxylated fatty alcohol surfactant. Organized molecular systems were identified by visual observation and by optical microscope equipped with cross polarizer. Droplet sizes were evaluated by optical microscopy (OM, Olympus BX50) and quasi-elastic light scattering (QLS, Malvern Zetasizer 300HsA). Coating elaboration on electrogalvanized steel: The process requires five steps: degreasing, rinsing, treatment, rinsing and drying. The treatment bath contains heptanoic acid (HC7), surfactant, water and an oxidizing agent. Surface analyses: Morphology and covering were determined by SEM and the crystallographic structure of the deposit was characterized through XRD. Aqueous corrosion test: The corrosion behaviour (Ecorr, Rp) of coatings was tested in ASTM D1384-87 solution (148 mg L-1 Na2SO4, 138 mg L-1 NaHCO3, 165 mg L-1 NaCl) [2] using a three-electrode electrochemical cell. Atmospheric corrosion test: Test cycles in a climatic chamber (KBEA 300, LIEBISH) were carried out to simulate atmospheric corrosion [1,3]. 3. Experimental results 3.1. Phase behaviour Solubilization is the dissolving of a predominantly lipophilic substance by the mediation of the micelles or other surfactant systems [4,5].The phase diagram Brij®58/HC7/water at 20°C is reported in Figure 1. HC7

90

HC7 90

80

10 80

70

8

70

60

60

6 50

50

40

40

4 30

30

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20

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Solid cr yst als

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20

Lm 20

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40

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80

90

E

2

L1

10

BRI J 58

WATER

2

4

6

(% W/ W)

(% W/ W)

a)

b) zoom ®

Figure 1 - Phase diagrams of Brij 58/HC7/water system at 20°C

Different self- assemblies (Figure 2) have been observed:

8

10

BRI J 58

Conversion coating of zinc heptanoate in aqueous media on electrogalvanized steel 705

surfactant

surfactant

oil

water oil

water

a) L1

b) L2

c) E

d) Lg

e) H

Figure 2 – Structures of different surfactant systems

Microemulsion is an isotropic solution, thermodynamically stable (1nm<Ødroplets<100nm). Direct micelle (L1) is constituted of droplets of oil in water (O/W) and inverse micelle (L2), of droplets of water in oil (W/O), Emulsion (E) is a several-phase milky system, kinetically stable (0.1µm<Ødroplets<100µm), Liquid crystals (LC) are highly viscous, have specific properties such as optical properties and are organized at long distances. Three structures have been observed: lamellar (Lg), hexagonal (H) and cubic (C) phases. 3.2. Droplet size of microemulsion and emulsions In order to characterize the different studied systems, the droplet size was determined by OM for emulsions and by QLS for microemulsions (Table 1). Solution prepared with ultra-turrax homogenizer was mixed during 1 minute at 20 000 rpm. Table 1 - Studied systems and their droplet size

Systema 1B1MB 3B1MB 5B1MB 3B3MB 3B3UT 3B5MB 3B7MB

%wt HC7 1 3 5 3 3 3 3

%wt Brij®58 1 1 1 3 3 5 7

Structure E E E E E Lg L1

Ødroplets (nm) 1 000 – 3 500 1 000 – 5 000 1 000 – 3 000 1 500 – 20 000 1 500 – 5 000 10 - 50

a

Criterion for codes xBy(MB/UT) is: x for % HC7, B for Brij®58 and y for % Brij®58, MB for magnetic stirring bar and UT for ultra-turrax.

According to the phase diagram, emulsion 1B1MB, located on the limit of microemulsion, is probably constituted of big droplets of HC7 dispersed in the L1 phase. System 5B1MB is put into emulsion by the effect of the agitation process. When agitation was stopped, this system separated into two phases and in this case, water seemed to be in equilibrium with the inverse micelle phase L2. Emulsion 3B1MB is supposed to be of the same type. No difference in droplet size was observed for these emulsions. On the other hand, the droplet size is sensitive to surfactant concentration. Indeed, emulsion 3B3MB prepared in the same conditions than 3B1MB, is coarser. The agitation process has an

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effect on droplet size; the use of high speed dispersing (ultra-turrax) induced finer droplet dispersion. 3.3. Influence of different parameters on the coating process Figure 3 shows that crystal size, covering rate and weight gain are closely dependent on the phase domain and on HC7 concentration. The best coating is obtained with the thermodynamically unstable solution: emulsion 3B3MB. The two other solutions are thermodynamically stable and the treatment is less efficient because HC7 is less available. Moreover, the higher the HC7 content, the smaller the crystals, the greater the covering and the denser the coating.

a) 3B3MB E (6s, 1.11g/m2)

b) 3B5MB Lg (10s, 0.61g/m2)

c) 3B7MB L1 (10s, 1.27g/m2)

d) 1B1MB E (10s, 0.95g/m2)

e) 3B1MB E (10s, 0.85g/m2)

f) 5B1MB E (5s, 0.87g/m2)

g) 3B3UT E (10s, 0.73g/m2)

Figure 3 - SEM micrographs of electrogalvanized steel sheets treated for few seconds in different phase domains

At constant HC7 concentration and increasing surfactant concentration, crystal size and covering rate are roughly the same (Figure 3 a, e). The deposition rate is higher with the less stable emulsion, 3B3MB; indeed, coating weight is higher with less dipping time. For emulsions, two mixing processes using respectively a magnetic stirring bar or an ultra-turrax homogenizer, were implemented (Figure 3 a, g). The deposition rate is higher with the coarsest emulsion, 3B3MB; indeed, coating weight is higher with less dipping time. 3.4. Corrosion behaviour Coatings were elaborated in order to have around 1 g.m-2 in less than 10 seconds. XRD analyses have revealed that whatever the self-assembly of the Brij®58/HC7/water system used, zinc heptanoate Zn(C7)2 [6] is formed.

Conversion coating of zinc heptanoate in aqueous media on electrogalvanized steel 707

The corrosion potential (Ecorr) and polarization resistance (Rp) of some coatings are shown in Figure 4a. The results can be compared to those obtained for non treated electrogalvanized steel EG. Values of treated samples are close to those of EG. For example, emulsion 3B3MB, has a Rp value higher by 1 kW.cm2 than EG during the first minutes of immersion. These coatings were then tested in a climatic chamber to determine their efficiency against atmospheric corrosion. The percentage of the zinc surface covered by white rust as a function of the number of cycles is presented in Figure 4b. The non treated electrogalvanized steel in the climatic chamber is completely covered with white rust after 2 cycles. 3B3UT

3B7MB

EG

3B5MB

100

4

2

0

80 60 40 20 c d EG

3B3MB

% of corroded surface

0 0

50

100 150 Immersion time (min)

200

a)

250

a b

Rp(kW.cm²)

6

5

10

15

Number of cycles

20

b)

Figure 4 – a) Rp value of treated samples versus immersion time in the ASTM corrosive water and b) percentage of corroded surface (white rust) versus the number of cycles in the climatic chamber for non treated electrogalvanized steel (EG) and treated in different solutions: a)3B3MB (E), b)3B3UT (E), c) 3B5MB (Lg), d)3B7MB (L1)

The products on the corroded surface have the same spectra as those analyzed on non treated steel after 2 cycles: white rust is mainly composed of Zn(OH)2 and Zn5(CO3)2(OH)6 as reported in the literature [7]. 4. Discussion and conclusions From the results obtained in this study, it is possible to synthesize a coating with heptanoic acid solubilized in aqueous media with a non ionic surfactant. This treatment allows the reduction of zinc corrosion and so appears to be an effective and suitable replacement for some phosphating or phosphating/chromating processes. Whatever the surfactant system (emulsion, microemulsion or liquid crystal) in which the deposition was realized, coatings have always the same composition: zinc heptanoate Zn(C7)2. But crystal size, covering rate and weight gain are closely dependent on the surfactant system. Microemulsion and liquid crystal are two thermodynamically stable solutions, so HC7 is less available and the reaction is delayed. Surfaces are then not well covered and corrosion resistance is reduced. In comparison, emulsions are thermodynamically unstable solutions and allow an easier release of HC7. Surfaces are then well covered and this

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observation is confirmed by a better resistance against corrosion, according to the climatic chamber tests. When this emulsion is prepared with a magnetic stirring bar, droplet size is larger than with ultra-turrax. By the reduction of the droplet size, the stability of the solution increased and the reaction rate decreased. Indeed, coating weight is lower with greater dipping time. Protective properties are damaged; white rust appeared more quickly in atmospheric corrosion. In the emulsion domain, coating properties are linked to emulsion quality. The coarsest emulsion is more easily destabilized and induces a better covering of the surface. Covering rate is increased with the increase of HC7 content; but at too high HC7 concentration, visual observation indicated the presence of heptanoic acid which had not reacted with the surface. To conclude, morphology, covering and corrosion resistance of the passive layer are mainly dependent on HC7 and surfactant concentrations, and on the phase domain. The best coating is obtained with a coarse emulsion. In all cases the formation of a passive layer in HC7 solutions requires a preliminary step corresponding to the oxidation of the metal and then precipitation of zinc carboxylate. These results are very promising and this process appears to be a versatile “green” conversion treatment of zinc. It allows depositing easily both hydrophobic and corrosion resistant zinc carboxylate coatings. Finally, the coating properties could be improved by increasing the time of treatment for a greater coating weight, as well as by increasing carbon chain length for having a less insoluble protective compound. Acknowledgements This study was financially supported in part by Arcelor Research, Total France and ADEME-AGRICE (convention n° 0201024). References [1] J. Peultier, E. Rocca, J. Steinmetz, Corrosion Science, 45 (2003) 1703 [2] ASTM standard D1384-87 (1988) [3] AFNOR standard NFT60-174 (1994) [4] P. Ekwall, L. Mandell and K. Fontell, Mol. Cryst. and Liq. Cryst., 8 (1969) 157 [5] A. Berthod, Journal de chimie physique, 80 (1983) 409 [6] J. Peultier, M. François, J. Steinmetz, Acta Cryst. C55 (1999) 635 [7] T.E. Graedel, J. Electrochem. Soc., 136 (1989)193

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

709

Study of carbon steel corrosion inhibition in alkaline solution by means of EIS Gianni Rondelli a, Luciano Lazzari b, Marco Ormellese b, Ramon Novoa c, Elmer Pérez b a

CNR - IENI, Via Cozzi 53, Milan 20100, Italy Politecnico di Milano, Via Mancinelli 7,20131 Milan,Italy. c Universidad de Vigo, E,T.S.E.I., Campus Universitario, 36319 Vigo, Spain [email protected] b

Abstract - Corrosion inhibitors are one of the additional prevention methods used to delay or stop chloride induced corrosion in concrete structures. Nowadays, both inorganic compounds (mainly based on nitrites) and organic substances (based on amines) are proposed world-wide. However, there are conflicting opinions about the effectiveness of these products: there is a lot of interest in studying new mixtures and in understanding the mechanism of inhibition. This paper reports results on the effectiveness of some organic substances in preventing chlorides induced corrosion. Tests were performed on carbon steel rebar samples in a simulated concrete pore solution. Potential and corrosion rate measurements, as well as EIS spectra were performed. Results have been discussed in term of influence of organic compounds on time-to-corrosion and in understanding the mechanism of inhibition of the analysed substances. Keywords: Concrete, pitting corrosion, corrosion inhibitors, EIS tecnhiques

1. Introduction Reinforcements steel in concrete are normally passivated due to concrete’s high alkalinity. However, this film can be destroyed mainly by two specific conditions: concrete carbonation and chlorides presence at the metal surface. The former causes a decrease in concrete’s pH, changing the steel condition from passive to active, whereas the latter provokes local breakdown of the passive film, originating pits on the metal surface. In most of the cases, the use of high quality concrete and adequate cover are enough to prevent corrosion. Additional prevention methods can be used in severe environmental conditions,

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710

or when the structures require a very long service life [1]. Among these methods, corrosion inhibitors are very attractive due to their low cost and easy handling [2]. The inhibition properties of different substances have been widely studied: mechanism of adsorption, efficiency, how they affect the chloride threshold; but controversial results were obtained [3-8]. This paper deals specifically with three different organic substances, and their mixtures, in order to analyze if they could minimize the effects of chloride-induced corrosion in simulated concrete pore solution, and how they interact with the passive layer. These substances were two carboxylic acids, a bicarboxylate (CARB-1) and a monocarboxilate (CARB-2), and a ferrous complexant amine (AM-1). The inhibitors mixtures were considered in order to study whether or not there was a synergic effect between them. Electrochemical impedance spectroscopy (EIS) and linear polarization (LPR) were performed in order to monitor the corrosion initiation on the carbon steel samples, and assess the inhibition mechanism of the analysed substances. 2. Experimental Simulated concrete pore solutions were made by mixing the organic compounds with NaOH, and saturated with Ca(OH)2 (pH 12,6). NaOH concentration was kept at 0.01 M, whilst the inhibitors proportions are summarised in Table 1. Tests with sodium nitrite 0.1 M and without any inhibitor (reference) were performed for comparison purposes. Table 1. Inhibitor concentration (mol/L) in solution Mixture

CARB-1

CARB-2

AM-1

S1

0.017

-

-

S2

-

0.207

-

S3

-

-

0.125

Mix1

0.017

0.207

-

Mix2

0.012

-

0.125

Mix3

-

0.075

0.062

Steel samples were machined from commercial reinforcing steel (10 mm diameter and 60 mm long), sandblasted and coated with self adhesive tape on both ends. The area exposed to the solution was approximately of 20 cm2. Five samples were immersed simultaneously in every solution for fifteen days to allow the passive film to grow. Then, 0.1 M of NaCl was added every seven days until all the samples presented pits. Free corrosion potential (Ecorr) was monitored everyday with respect to SCE reference electrode. Linear Polarization Resistance measurements (LPR) were made after the addition of

Study of carbon steel corrosion inhibition in alkaline solution by means of EIS

711

chlorides, while EIS ones were done the day before. Both tests were performed using an Autolab Potentiostat (PGSTAT30). EIS spectra were obtained at Ecorr with a perturbing signal of ±10 mV. The frequency span was from 10 kHz to 1 mHz. The experiment and registration of data were conducted under computer control using the software FRA v4.8 [9]. Fitting of the data was done using the software ZsimpWin v3.10 [10]. LPR values were determined by potentiodynamic tests with a scan rate of 10 mV/min, from -10 mV to +10 mV with respect to Ecorr. The test was carried out using the software GPES v4.8. Corrosion rate was calculated by the formula: icorr = B/LPR, with B = 26 mV [11, 12]. Results of these tests are discussed in [13, 14]. 3. Results and Discussion Fig. 1 reports the chloride content in correspondence of carbon steel corrosion initiation for every analysed organic substance. Values were determined by LPR and potential measurements [13, 14]: when a sudden decreases both on potential and LRP was measured the sample was considered in corrosion. Localised corrosion was also confirmed by visual inspection. These tests proved the analysed organic substances increased the maximum chloride threshold with respect to the reference condition. Organic compound S3 showed the best performance among the monocomponent solutions, with a medium tolerance 0.25 M of chlorides. Mix1 and Mix3 showed chloride threshold values higher to the one observed for their parent components, as shown in Fig. 1, with a mean value of 0.35 M for the former and 0.45 M for the latter. No significant effect was observed in the presence of Mix2. Fig. 2 reports the Nyquist plot and Bodephase diagram obtained in solution without inhibitors addition, in the presence of pure organic substances and of sodium nitrite, after 2 weeks of passivation, before chlorides addition. It can be seen that the addition of the inhibitor to the solution resulted in an increment in the value of the general impedance of the system, with S2 being the organic substance with the highest impedance. Analysis of the experimental data was done by fitting the results to an equivalent circuit. Two models were considered to fit the data: 1) a hierarchical model, where the relationship between the passive layer and the metal is similar to the one observed in the paint-metal system [15, 16]; 2) a series model, with the circuit form Rsol(Q1R1)(Q2R2), where the elements in brackets are in parallel [17, 18]. The data have been fitted to the latter, where: Rsol represents the solution resistance; Q1 and R1 represent the double layer capacitance and the charge transfer resistance; Q2 and R2 are related to redox processes in the oxide layer [17, 18]. In Table 2, the calculated values are reported for the following conditions: after two weeks passivation (before chlorides addition), just before passivity breakdown and after the pit formation. Also the measured free corrosion potential values are reported: it is evidente that potential is close to – 0.2 V SCE in passive condition, while potential dropped to –0.4 V SCE when corrosion occurred.

G. Rondelli et al.

712 0.6

Chlorides (mol/L)

0.5 0.4 0.3 0.2 0.1 0.0 ref

Nit

S1

S2

S3

Mix1 Mix2 Mix3

400000

90

320000

70

240000

Ref Nit S1 S2 S3

160000 80000 0 0

60000

120000 2

Z' (Ohm cm )

180000

- Phase angle

-Z'' (Ohm cm2)

Figure 1: Chloride threshold value for all the substances and their mixtures determined by LPR and potential measurements. The minimum, mean and maximum values are plotted. Ref Nit S1 S2 S3

50 30 10 -10 0.0001

0.01

1

100

10000

Log | f |

(a) (b) Figure 2: Nyquist plot (a) and Bode-phase diagram (b) for the inhibitors and the reference solution after 2 weeks of passivation. Some authors proposed that organic inhibitors adsorb on the passive layer defects, screening the active sites present in the film [19-21]: the values reported in Table 2 show how the inhibitor presence is related with an increment in R1 with respect to the reference solution, increasing its value until the film breakdown, when it drops violently. The double layer capacitance, Q1, has similar values for all the solutions, and it remains almost constant until the

Study of carbon steel corrosion inhibition in alkaline solution by means of EIS

713

pit is formed. The n1 values, for all the solutions, are very close to one, meaning that the double layer behaves very similar to a pure capacitor, reflecting the homogeneity of the solid/liquid interface. The relative low values observed for R2, may reflect fast redox transformations in the passive layer film, which is not homogeneous in reactivity as suggested by the low n2 values. Table 2. Values obtained from the EIS fitting: after two weeks of passivation, before breakdown of the passivity, after the pit formation. Resistances values are in [W cm2]. Capacitances values are in [S sn cm-2]. Chloride concentration is in [mol/L]. Nit stands for NaNO2. Sol. Ref

Nit

S1

S2

S3

Mix1

Mix2

Mix3

[Cl-]

Pot (V)

Rsol

R1 (104)

Q1 (10-4)

n1

R2 (102)

Q2 (10-4)

n2

0

- 0.14

133.0

174

2.56

0.90

1.29

1.58

0.63

0.1*

- 0.48

92.8

1.13

0.23

0.76

87.40

66.9

0.75

0

- 0.10

96.5

82.3

3.79

1.00

19.30

7.08

0.72

0.5

- 0.11

75.1

90.3

3.21

0.82

0.23

110

0.82

0.6

- 0.45

28.5

2.05

3.01

0.70

27.40

6.35

0.75

0

- 0.20

181.0

315

2.20

0.90

3.30

14.4

0.56

0.1*

- 0.41

84.8

5.59

3.80

0.79

0.53

6.40

0.66

0

- 0.19

84.2

231

2.56

0.89

7.86

8.21

0.58

0.2

- 0.17

70.7

984

2.15

0.84

50.20

6.19

0.53

0.3

- 0.42

6.3

13.3

1.83

0.56

0.64

7.32

0.42

0

- 0.15

220.0

127

3.14

0.90

1.34

22.2

0.45

0.2

- 0.14

65.6

379

2.75

0.90

4.85

8.11

0.52

0.3

- 0.44

57.8

3.39

7.08

0.88

0.01

9.11

0.50

0

- 0.10

77.0

283

2.11

0.83

3.07

12.4

0.59

0.3

- 0.11

47.5

74.7

6.63

0.79

310.00

0.14

0.45

0.4

- 0.41

50.3

2.98

7.29

0.72

6.31

11.3

0.44

0

- 0.15

190.0

131

2.06

0.86

29.60

6.93

0.56

0.1

- 0.17

108.0

115

2.89

0.89

3.00

2.89

0.89

0.2

- 0.43

67.0

7.72

4.42

0.84

0.18

8.80

0.73

0

- 0.11

134.0

182

2.58

0.89

6.10

8.30

0.46

0.4

- 0.13

41.5

947

2.28

0.84

13.30

1.55

0.55

0.5

- 0.40

42.5

18.7

2.60

0.72

17.60

1.08

0.58

* Breakdown of passivity occurred before EIS measurement

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G. Rondelli et al.

The chloride concentration values reported in Table 2, in combination with the results illustrated in Fig. 1, show that there is synergy between the studied substances: Mix1 and Mix3 have a critical chloride threshold higher than the one observed for their parent substances. Only Mix2 does not exhibit this effect. Nevertheless, neither of the mixtures perform better than the sodium nitrite. 4. Conclusions The tested substances proved efficient as corrosion inhibitors for chloride induced corrosion in concrete; their mixtures showed a synergic effect, i.e. they behave better than their parent compounds. Characterization of this effect was done by means of EIS, LPR and potential measurements. References [1] L. Bertolini, B. Elsener, P. Pedeferri and R. Polder (Wiley), Corrosion of steel in concrete: prevention, diagnosis, repair, Weinheim, 2003. [2] B. Elsener (Maney Publishing), Corrosion inhibitors for steel in concrete - an EFC state of the art report, EFC no. 35, Institute of Material, London, 2001. [3] C. Monticelli, A. Frignani and G. Trabanelli, Cement and Concrete Research, 30 (2000) 635. [4] L. Dhoubi, E. Triki and A. Raharinaivo, Cement and Concrete Composites, 24 (2002) 35. [5] F. Bolzoni, F. Fontana, G. Fumagalli, L. Lazzari, M. Ormellese and MP. Pedeferri, Eurocorr 2001, Riva del Garda, Italy, (2001). [6] J.M. Gaidis, Cement and Concrete Composites, 26 (2004) 181. [7] F. Wombacher, U. Maeder and B. Marazzani, Cement and Concrete Composites, 26, (2004) 209. [8] C.K. Nmai, Cement and Concrete Composites, 26 (2004) 199. [9] FRA v4.8, Eco Chimie B.V. (2001). [10] B. Yeum, ZsimpWin v3.10, Echem Software (2002). [11] M. Stern and A.L. Geary, Journal of the Electrochemical Society, 105 (1957) 56. [12] C. Andrade and J.A. Gonzalez, Werkstoffe und Korrosion, 29 (1978) 515. [13] V. Sesia, Politecnico di Milano, Graduate Thesis (2003). [14] A. Gambarella and A. Guidi, Politecnico di Milano, Gradaute Thesis (2004). [15] C. Andrade, M. Keddam, X.R. Novoa, M.C. Perez, C. M. Rangel and H. Takenouti, Electrochimica Acta, 46 (2001) 3905. [16] S. Joiret, M. Keddam, X.R. Novoa, M.C. Perez, C. Rangel and H. Takenouti, Cement and Concrete Composites, 24 (2002) 7. [17] Q. Yin, G.H. Kelsall, D.J. Vaighan and N.P. Brandon, Journal of the Electrochemical Society, 148 (2001) A200. [18] F. Bolzoni, L. Lazzari, M. Ormellese, M.P. Pedeferri and G. Rondelli, Eurocorr 2003, Budapest, Hungary (2003). [19] Y.I. Kusnetsov (Plenum Press), Organic inhibitors of corrosion of metals, New York, 1996. [20] S. Sastri (Wiley), Corrosion Inhibitors: principles and applications, Chichester, 1998. [21] I. Felhosi, J. Telegdi, G. Palinkas and E. Kalman, Electrochimia Acta, 47 (2002) 2335.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

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Nitrided/nitrocarburized and oxidized steel: Corrosion data in dependence on the N- and Ccontent of the e-phase under the oxide layer U. Ebersbach Leipzig University of Applied Sciences, Postfach 301166, D-04251 Leipzig ,Germany, e-mail: [email protected]

Abstract - The corrosion behaviour of nitrided/nitrocarburized and oxidized steel 20MnCr5 was characterized by polarisation curves in 0.9 M NaCl solution and discussed in relation to the structure of the compound layer. Oxidized e-carbonitride layers on steel effectively improve the corrosion resistance, if the interstitials (N, C) in the e-phase under the oxide layer are within specific concentration ranges. Keywords:

nitriding, oxidation, structure, uniform corrosion, pitting corrosion

1. Introduction The oxidation of nitrided and nitrocarburized steel in water vapour at temperatures of 450 ... 550 °C leads to effective enhancement of corrosion resistance in neutral media. During the oxidation, the outer range of the compound layer consisting of iron nitrides and iron carbonitrides, is converted into a magnetite layer of about 0.2 to 2 µm thickness. In addition the released nitrogen effects a N-enrichment of the compound layer directly under the oxide layer. The best corrosion resistance is reached by the existence of an e-carbonitride compound layer [1], but not much is known however about the optimal composition of the e-carbonitride e-[Fe2(N,C)1-x] as an interstitial compound. Therefore in this paper the influence of the N- and C-concentration on the corrosion behaviour of e-carbonitride is explored.

U. Ebersbach

716

2. Experimental Specimens of the steel 20MnCr5 were treated by gas nitrocarburizing, salt bath nitrocarburizing and gas oxinitriding under conditions relevant to the practise. The produced e-compound layers of about 20 µm thickness are processed by grinding with 3 µm surface abrasion to obtain an equal surface quality. The following oxidation was applied in the sensor-controlled water vapour atmosphere with optimized parameters [2]: J = 450; 500; 550 °C; t = 20; 40; 60 min; pO2 = 10-30, ..., 10-23 bar. Glow Discharge Optical Spectroscopy was used to determine the concentration-depth profiles in the compound layers of the treated specimens. These data were used to determine the N- and C-concentrations of the e-phase under the oxide layer. The corrosion behaviour of the treated specimens was investigated by the current density-potential curves in 0.9 M NaCl solution. 3. Results and discussion 3.1. N- and C-concentrations in the e-carbonitride Fig. 1 shows typical concentration profiles of two different oxidized e-carbonitride compound layers. In view of the corrosion behaviour, the N- and Cconcentration in the interface oxide/ e-phase were determined. A gas nitrocarburized and at 550 °C oxidized specimen contains 7 % N and 1.6 % C in the ephase directly under the oxide layer (dO = 2,1 µm), whereas a salt bath nitrocarburized specimen which was oxidized at 450 °C indicates 9 % N and 0.6 % C directly under the oxide layer (dO = 0,4 µm). Depending conditions on the treatment the concentrations of the nitrogen and carbon in the e-phase directly under the oxide layer result in wide differences. The ranges of the values are given in the table 1; the detailled connection with the treatment parameters can be found in an earlier paper [2]. Table 1: N-and C-concentration of the e-phase e-[Fe2(N,C)1-x] on the "surface" of compound layers produced by different treatment variants of the steel (GNC gas nitrocarburizing, SNC salt bath nitrocarburizing, GON gas oxinitriding, PR processing, OX oxidation)

treatment GNC GNC + PR + OX SNC SNC + PR + OX GON GON + PR + OX

N, % 6.4 5.4 ... 11.2 7.7 5.9 ... 11.4 7.5 7.6 ... 9.6

C, % mean value 1.4 1.4 0.6 0.6 0.1 0.1

N + C, % 7.8 7 ... 12.1 8.3 6.9 ... 12.1 7.6 7.7 ... 9.7

C N+C 0.18 0.07 ... 0.25 0.07 0.03 ... 0.15 0.01 0.01

Nitrided/nitrocarburized and oxidized steel

717

16 dO

14

concentration (%)

12 N+C

10

8,6 % N+C

8 N 6 4 O C

2 0 0

(a)

5

10

15

20

25

30

35

30

35

distance to outer surface (µm) 16 dO 14

concentration (%)

12 N+C

10

8,6 % N+C

8 N

6 4

O C

2 0

(b)

0

5

10

15

20

25

distance to outer surface (µm)

Fig. 1: GDOS concentration profiles of the elements N, C and O in the compound layer of the nitrocarburized + processed + oxidized steel 20MnCr5 comparison between the calculated (N+C) profiles and the characteristical (N+C) concentration of 8,6 % with respect to an improved pitting corrosion resistance (a) gas nitrocarburized + processed + oxidized (550 °C, 60 min, pO2 = 10-23,7 bar) (b) salt bath nitrocarb. + processed + oxidized (450 °C, 60 min, pO2 = 10-27 bar)

3.2. Corrosion behaviour Uniform corrosion: Analogous to the mixed crystals the active dissolution of the e-carbonitride was investigated in dependence of the mole fraction of the components. Here the concentration ratio [C/(N+C)] of the e-phase is more suitable.

U. Ebersbach

718

The corrosion data (UR, Rp-1) are shown in dependence of the [C/(N+C)]-ratio of the e-phase (Fig. 2 and 3). The curves characterise nitrided/nitrocarburized surfaces and areas after defined surface abrasion of the earlier charge [3]. The areas, which consisted of e-carbonitride with an [C/(N+C)]-ratio of 0.02 to 0.2, show a high corrosion resistance. The rest potential becomes nobler up to 300 mV (as result of the favoured formation of a passive layer) and the reciprocal polarisation resistance decreases by more than one order of magnitude. The resistant areas exist in the non porous part of the compound layer (depth ³ 8 µm under the surface). The here tested surfaces of the 2.charge (Tab. 1) fit with Rp-1-curves, but the UR-values are differing. The oxidation of the specimens effects the formation of the thin magnetite layer and an N-enrichment of the 4 8

-100

18 30

16

rest potential UR (mV [SCE])

12

12

-200 16 0

8 4

0

0

-300 0 0 4

-400

-550 -600 0,0

0,1

0,2

0,3

0,4

0,5

concentration ratio C/(N+C) gas nitrocarburized + processed + oxidized specimens salt bath nitrocarb. + processed + oxidized specimens gas oxinitrided + processed + oxidized specimens mean curve of nitrided/nitrocarburized specimens of the earlier charge: mean value of unoxidized materials (with scattering range) gas nitrocarburized 20MnCr5 resp. 45Cr2 salt bath nitrocarburized 20MnCr5 gas oxinitrided 20MnCr5 resp. 45Cr2 after surface abrasion: 0; 4; 8; 12; 16 ;18; 30 µm (s. the number next to a sign) steel, untreated

Fig. 2: Rest potential UR of treated specimens of the steel 20MnCr5 in 0.9 M NaCl solution in dependence on the concentration ratio [C/(N+C)] of the e-phase directly under the oxide layer in comparison with the unoxidized materials after different surface abrasion

Nitrided/nitrocarburized and oxidized steel

719

1

-1

-1

-2

reciprocal polarisation resistance Rp (mA V cm )

e-phase under the oxide layer. Numerous oxidized surfaces with a [C/(N+C)]ratio of 0.02 to 0.2 reach the noble rest potentials of the curve in Fig. 2, divergences are caused by the surface defects. The Rp-1-values of the oxidized surfaces (Fig. 3) demonstrate the best corrosion resistance, if the e-carbonitride directly under the oxide layer provides a [C/(N+C)]-ratio of 0.02 to 0.2. Pitting corrosion: The pitting potentials of the treated specimens depend significantly on the content of the interstitials (N, C) directly under the oxide layer (Fig. 4). For the gas nitrocarburized and salt bath nitrocarburized, respectively, and oxidized specimens, showed [N+C] < 8.6 % in the e-phase, are measured pitting potentials of 630 to 850 mV, but the pitting potentials of the specimens with [N+C] > 8.6 % rise to values of 1000 to 1200 mV. Thereby the e-carbonitride obtains about 0.6 % to 1.4 % carbon (Tab. 1). The pitting potentials of the gas oxinitrided and oxidized specimens however showed small values of 630 ... 850 mV for the all (N+C)-contents of 7.0 to 9.7 %. The content of the carbon in the e-carbonitride amounts to only 0.1 %. Therefore the ecarbonitride, which shows an interstitials concentration > 8.6 % and a carbon content of at least 0.2 %, leads to a better adaption of the outer magnetite layer [2] and to a decrease of the pitting corrosion susceptibility.

0 0 4 0 4

0,1

0

8

16 12

12 30 4

0

8

16 18

0,01

0,0

0,1

0,2

0,3

0,4

0,5

concentration ratio C/(N+C) Fig. 3: Reciprocal polarisation resistance Rp-1 of treated specimens of the steel 20MnCr5 in 0.9 M NaCl solution in dependence on the concentration ratio [C/(N+C)] of the e-phase directly under the oxide layer (farther illustrations see Fig. 2). Rp-1 was corrected with the roughness factor due to pore formation on the surface.

U. Ebersbach

720

1300 1200

U L (mV [SCE])

1100 gnc+pr+ox

1000 snc+pr+ox

900 800 gon+pr+ox

700 600 500 6

7

8

9

10 N+C (%)

11

12

13

Fig. 4: Pitting potential UL of treated specimens of the steel 20MnCr5 in 0.9 M NaCl solution in dependence on the concentration [N+C] of the e-phase directly under the oxide layer ̊

D

»

̈u

gas nitrocarburized + processed + oxidized specimens salt bath nitrocarb. + processed + oxidized specimens gas oxinitrided + processed + oxidized specimens mean values of unoxidized materials

4. Conclusions The corrosion behaviour of nitrided/nitrocarburized and oxidized steel specimens in 0.9 M NaCl solution shows a significant dependence on the N- and Ccontent of the e-carbonitride directly under the oxide layer. The best resistance against uniform corrosion was achieved for the concentration ratios [C/(N+C)] in the e-carbonitride of 0.02 up to 0.2. The pitting corrosion susceptibility is smallest for concentrations of the interstitiales [N+C] > 8.6 % and carbon content > 0.2 %. 5. References 1. H.-J. Spies, H. Winkler and B. Langenhan, Härterei-Tech. Mitt. 44 (1989) 75. 2. U. Ebersbach, F. Vogt et al., Härterei-Tech. Mitt. 54 (1999) 241. 3. U. Ebersbach, S. Friedrich and T. Nghia, International Conference on Surface Engineering, Bremen (1993), Proceedings, p.299.

Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

721

Passivation treatments on Zn and Zn alloys substrates involving Mo and W compounds L.Anicaia, A.Peticab, M.Budac, T.Visanc a

Petromservice S.A., Div.of Ecological Technologies Implementation, Gral Budisteanu 11-bis, 010772, Bucharest, Romania, e-mail:[email protected] b ICPE Advanced Research, Non-Conventional Engineering Laboratory, Splaiul Unirii 313, sector 3, Bucharest, Romania c University POLITEHNICA, Dept.Applied Physical Chemistry and Electrochemistry,Calea Grivitei 132, sector 6, 010737, Bucharest, Romania

Abstract - Formation and characterization of chemical conversion layers on Zn and Zn-Fe alloy (0.3-0.6% Fe) coatings involving Mo (VI) and W (VI) based iso- or heteropolycombinations are discussed, as ecological alternatives to replace Cr(VI) ones. They offer a large diversification of decorative aspect due to an extended range of colors and tones, as a function of operating parameters. The passive films have been analysed from color view point using CIELAB system in trichromatic coordinates L*, a*, b*. The coexistence of Mo(VI) and Mo(V), respectively of W(VI) and W(V) species were suggested to be responsible for color appearance, by means of recorded electronic diffuse reflectance spectra. Some aspects regarding their corrosion resistance as compared with classical Cr(VI) based conversion coatings are also discussed. Keywords:Conversion coatings, Zinc; Zn alloy; Molybdate; Tungstate

1. Introduction With the actual trend to use environmentally friendly electrolytes to process metallic surfaces, an important role is assigned to Cr(VI) replacement, especially for Zn and Zn alloy surface treatments, to improve the decorative aspect and to offer at least the same corrosion resistance. As suitable candidates, compounds of either transition metals or rare earth metals with at least two stable oxidation states may be considered. Thus, to minimize or eliminate Cr(VI) use, chemical conversion based on Mo(VI) and W(VI) compounds may

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be proposed, even though the formation and protection mechanisms are poorly understood. Several investigations on Mo based conversion coatings involving ESCA and XPS analysis showed the presence of Mo in various oxidation states, including Mo(V) and Mo(IV) as well as Mo(VI) as MoO3 or MoO42- [1,2,3]. Recently, based on EXAFS investigations Wharton et al.[2] proposed the formation of a Zn(II) and Mo(IV) mixed oxides containing a significant amount of Mo(VI). Mo(VI) species have an octahedral distorted arrangement suggesting the presence of heptamolybdate species rather than MoO42- anions which are characterized by a tetrahedral symmetry. The proposed formation mechanism included an initial Zn dissolution and Mo6+ reduction to Mo4+ at Me/solution interface. Due to H2 evolution a local pH increase takes place near cathodic positions and thus precipitation of a Zn and Mo hydrated mixed oxide occurs with a cracked structure that may allow incorporation of a Mo(VI) residual amount. In this work some investigations regarding formation and characterization of conversion layers onto Zn and Zn-Fe alloy (0,3-0,6% Fe) coatings involving Mo (VI) and W (VI) based iso- or heteropolycombinations are presented, as ecological alternatives. 2. Experimental Chemical conversion treatment was performed on mild steel specimens of 70x140x1 mm, previously electrochemically coated with Zn and Zn-Fe alloy (0,3-0,6% Fe), using an acidic electrolyte of chloride type; the detailed procedure was described elsewhere [4]. The passive films have been formed using solutions based on molybdic acid, phosphomolybdic acid, ammonium heptamolybdate and tungstic acid, ammonium paratungstate, respectively, with different compositions and concentrations, at temperatures between 25-40oC. Various immersion times were applied, between 15-120 s. Open circuit potentials during conversion film formation were measured using Zahner IM6e electrochemical system, using a saturated calomel (SCE) reference electrode. The characterization of the formed layer was carried out by analyzing adherence to the metallic support, uniformity, and colour. The developed colour was examined both visually and by means of CIELAB system in trichromatic coordinates - L*,a*,b*, using a MICROFLASH 200D spectrophotometer with two pencils and integrating sphere. To investigate the passive films from existent chromophoric species view point, diffuse reflectance electronic spectra were recorded in the visible range (360-900 nm), using the above mentioned equipment. To estimate corrosion resistance of the passive films, laboratory tests were performed involving the continuous immersion in a NaCl 3% solution at 20oC for 168 h. The appearance of coating metal corrosion was observed and compared with those formed in Cr(VI) based solutions.

Passivation treatments on Zn and Zn alloys substrates involving Mo and W

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3. Results and Discussion To form chemical passive films on Zn-Fe alloy surfaces, the following Mo(VI) compounds based solutions have been tested: (A) 2% H2MoO4 H2O + 1% NH4Cl; (B) 2% (NH4)2[Mo3O10] 4H2O + 1% H2O2 (30% H2O2 solution); (C) 2-5% H3[PMo12O40] . For H2MoO4 based solution (solution A) the passive films obtained develop yellow tones for 15-30 s. When immersion time increased up to 1-2 min., blue and then reddish iridescence appear, over the dominant yellow colour. In solution B, containing ammonium heptamolybdate, the addition of H2O2 facilitates peroxomolybdate formation. Here, the first conversion layers with blue-greenish tones appeared for 15-30 s; then the predominant colour was yellow. We notice that the temperature rise from 25oC to 30oC leads to a slight intensification of colour. Fig.1 shows an example of the variation of trichromatic coordinates against conversion time for two working temperatures. In the case of conversion layers formed in more acid solutions (pH=1-1.5) based on phosphomolybdic acid (solution C) the range of colours varies in the same way, respectively between blue-greenish to golden-yellowish tones. Light tones are facilitated by a relatively low temperature. An increase of phosphomolybdic acid concentration, from 2% to 5%, determines a displacement of chromatic space from blue to yellow. Also, an increased temperature, e.g. at 40oC, leads to a deeper yellow tone with reddish iridescence. Ammonium heptamolybdate 2% +H O 1% 2

2

100

yellow

25 oC

30

2

30 oC 80

1

20

1 60

b*

L*

10

40

0.5

0

green

red 0.25

-10

20

25oC 30oC

0.5 0.25 blue

0 0.0

-10

0

10

20

30

0.5

1.0

1.5

2.0

time, min.

Figure 1. Chromatic coordinates (CIELAB system) vs. conversion time at 25oC and 30oC temperatures, for layers formed in solution (B) based on ammonium heptamolybdate and hydrogen peroxide; the figures placed near experimental points show immersion duration, min.

We may conclude that, generally, the use of Mo(VI) based solutions allows chemical conversion films to form on Zn-Fe alloys coatings with colours varying in the range of yellow - bluish tones, as a function of immersion time

L. Anicai et al.

724

and solution composition. It must also be mentioned that all solutions lead to very adherent passive films on the metallic substrate. To characterize the passive films composition, diffuse reflectance electronic spectra were recorded, taking into account the possibility of incorporation of Mo(VI) species, as well as the formation of Mo(V) or even Mo(IV) species through a reduction process favourized by metallic Zn and Fe presence. The presence of Mo(V) species is proved by the existence of a charge transfer band between 380-450 nm in the visible spectra. When Mo(V) is covalently bonded as (MoO)3+ ion, a characteristic band in the range 666-770 nm occurs, specific to a d-d transition, as well as in a band in the domain 454-555 nm [5,6]. Mo4+ ion is characterised by the appearance of two low energy bands for d-d transitions, followed by intense charge transfer transitions, with a prominent shoulder around 700 nm [5]. Fig. 2 shows an example of electronic absorption spectra in the case of chemical conversion layers formed in solution A, based on H2MoO4, for different values of immersion time. o

100-R, %

H2MoO4 2% + NH4Cl 1%, 25 C 360 380 400 420 440 460 480 500 520 540 560 580 600 620 640 660 680 700 720 100

100

90

90

80

80

70

70

60

15 s 30 s 60 s 120 s

60

50

50 360 380 400 420 440 460 480 500 520 540 560 580 600 620 640 660 680 700 720

l, nm Figure 2 - Electronic absorption spectra for passive films formed onto Zn-Fe alloy in 2% H2MoO4 + 1% NH4Cl solution, for various durations of immersion

In this case the ammonium molybdate (VI) may be formed, that might be adsorbed on the metallic surface. According to Fig. 2, a maximum absorption in the range 380-400 nm, recorded for short immersion times, of 15-30 s, was attributed to the charge transfer (CT) existence, a phenomenon characteristic to Mo5+ species [5]. Absorption becomes more intense with increasing conversion time, as indicated by an intensification of yellow tone of the specimen. For 60 s, the colour changes and an intense band at 700-720 nm is observed, assigned to d-d transitions for Mo5+ species or even (MoO)3+ (molybdenil) that may coexist with Mo6+ species. For films obtained at 120 s, the absorption maximum is

Passivation treatments on Zn and Zn alloys substrates involving Mo and W

725

positioned towards 440 nm, also characteristic to d-d transitions of Mo5+ species, which is slightly shifted at higher frequencies due to possibility of reducing Mo6+ to Mo5+ through metallic Zn/Fe [5-7]. Thus, the coexistence of MoO42- anion and Mo(V) oxy-hydroxide is proposed within the passive film. The other electronic absorption spectra also suggested the presence of the same species in different proportions that contribute to the final color of the coating. Thus, it suggests the occurence of molybdenum anions with mixed valence having Keggin type structure, (PMo12O40n-1) and the intervalence transitions Mo5+-O-Mo6+ between the corner-sharing of octahedra [8-10]. The colour of chemical conversion layers formed onto Zn-Fe alloys is mainly determined by the coexistence in various ratios of Mo6+ and Mo5+ species, depending on immersion time. The general tendency is for Mo(VI) to reduce to Mo(V), with possible formation of Mo5+-O-Mo6+ [10]. The experiments carried out to form conversion layers using W based solutions have shown that these have always a weaker acidity than solutions containing Mo species. As Table 1 shows, the conversion procedure is performed with HNO3 additions for both tungstic acid and ammonium paratungstate solutions to allow formation of a passive film. It is important to notice that W species favour formation of blue films the tone of which depends on W based compounds as well as of HNO3 concentration and treatment duration. For longer immersion times dark tones and color interferences are usually obtained. The color of the films using tungstic acid solutions are more reproducible than for the case of paratungstate ones. Table 1. W based chemical conversion solutions and passive film color evolution on Zn and ZnFe coatings Electrolyte composition, g/L

Operation parameters

Obtained colour of passive film

Tungstic acid

pH

Blue with goldish irizations (30-60s) - dark blue (300 s) for less concentrated solution;

Nitric acid

20-50 10

Ammonium paratungstate 20 Nitric acid

10

6-6.5 o

T

30-40 C

t

30-300 s

pH

7 o

T

25-60 C

t

30-300 s

Light blue (30 s) – blue (60 s) – blue iridescence (180 s) for concentrated solution Light blue (30-60 s) – reddish tones with iridescence (120 s)

Fig.3 shows an example of open circuit potential (Uoc) evolution against conversion time for an ammonium paratungstate solution, for various temperatures. It is clear that the formation process depends on the applied temperature. Thus, at a temperature of 25oC when the electrode is immersed in the conversion solution an initial open-circuit potential of around -1.03 V/SCE is established, that moves towards electronegative values in the first ~60 s,

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726

corresponding to Zn chemical dissolution with hydrogen evolution. After that Uoc follows the same trend but with a lower slope suggesting film formation. For durations longer than 240 s Uoc attains a quasi-constant value of about 1.045 V/SCE when the passive film entirely covers the metallic surface. The layer color varies from light blue at ~30 s to iridescent red at ~130 s, suggesting some W(IV)/W(V) compound formation. At higher temperatures, e.g. of 60oC (Fig.3), the curve shape is different. Zn dissolution is temperature activated; after ~30 s Uoc moves in electropositive direction, suggesting formation of a Zn and W mixed oxide. Then, it attains a potential plateau around -1.003 V/SCE suggesting a formation equilibrium. If conversion is continued, another displacement in electronegative direction is noticed assigned to a further dissolution process. It should be considered that for relatively long conversion times, over 200 s, dissolution-precipitation-redissolution steady state is modified due to local pH increase at high temperatures, leading to a diminishing of the passive layer thickness. o

25 C o 60 C

-1.005 -1.010

Uoc, V/SCE

-1.015 -1.020 -1.025 -1.030 -1.035 -1.040 -1.045 -1.050 0

100

200

300

400

immersion time, s

500

600

Figure 3. Open circuit potential against immersion time during conversion process in ammonium paratungstate 20 g/l(pH7) at 25o C and 60o C

To obtain some information on corrosion resistance, some laboratory tests have been performed by continuous immersion for 168 hours in 3% NaCl solution at 20oC on minimum 3 specimens, with intermediary examinations. The thickness of Zn-Fe metallic coating was 10 µm. The percentage of degraded surface was noted, showing coating-metal corrosion (CMC), other passive films modifications. The following procedures were investigated: (i) classical Cr(VI) based chemical conversion; (ii) chemical conversion in 5% H3[PMo12O40] solution, t=25oC, for 30 s; (iii) chemical conversion in 5% H3[PMo12O40] solution, t=25oC, for 120 s. As it may be seen in Fig.4, the best corrosion resistance is still offered by chromated Zn-Fe coating. The passive film formed in H3[PMo12O40] solution assures a medium protection, slightly improved by a longer conversion time, of about 120 s. However, the very good adherence of the passive layer onto Zn-Fe surface should be noted, the film not being removed even after 168 h of conditioning.

Passivation treatments on Zn and Zn alloys substrates involving Mo and W

727

35 30

Zn-Fe/Cr Zn-Fe/PMo30s Zn-Fe/PMo120s

CMC, %

25 20 15 10 5 0 4

8

12

24

36

48

60

hours

120 132 144 156 168

Figure 4. Corrosion evolution for Zn-Fe coatings subjected to chemical conversion in Cr(VI) and Mo(VI) compounds based solutions

4. Conclusion The experiments performed demonstrated the possibility of passive film formation on Zn and Zn-Fe alloys metallic surfaces involving Mo(VI) and W(VI) compound based solutions. They may offer a large diversification of decorative aspect, through the tones of the obtained colors, as a function of the involved solution composition. Based on electronic diffuse reflectance spectra, the presence of Mo(VI) and Mo(V) species was suggested, as well as of molybdenil (V) anion that coexist within the passive film in variable ratios, depending on immersion time and temperature. Associating the color evolution and open circuit potential variation, W(IV) and W(V) compounds might be present within the film. From corrosion resistance view point, the experiments showed that Mo(VI) based passive films may assure a medium protection, weaker than that noticed for Cr(VI) based ones. Still, the attractivity and interest for these treatments is due to their special decorative effect, as well as their very good adherence on the metallic support. References 1. E.Almeida, T.C.Diamantino, M.O.Figueiredo, C.Sa, Surf. and Coat. Technol.106 (1998) 8 2. J.A.Wharton, D.H.Ross, G.M.Treacy, G.D.Wilcox, K.R.Baldwin, J.Appl.Electrochem. 33 (2003) 553 3. A.A.O.Magalhaes, I.C.P. Margarit, O.R.Matos, J.Electroanal.Chem. 572 (2004) 433 4. L. Anicai, A. Petica, Romanian research project, 57/A328, 1998 5. A.B.P. Lever, Inorganic Electronic Spectroscopy (2nd ed.), Elsevier, NewYork, 1984 6. S.D. Russell, Physical Methods for Chemists, Saunders College Publ.,1992 7. G. Centi, V. Lena, F. Trifio et al, J. Chem. Soc. Faraday Trans. 86(15)(1990) 2775 8. M.B. Robin, P. Day, in: H.J. Eneleus, A.G. Sharpe (Eds.), Advanced Inorganic Chemistry and Radiochemistry, Vol.10, Academic Press, NewYork, 1967, p.340 9. J.J. Altenau, M. Pope, R.A. Prados, H. So, Inorg.Chem. 14 (1975) 417 10. C. Sanchez, J. Livage, J.P. Launay, M. Fournier, Y. Jeannin, J. Am. Chem. Soc., 104 (1982) 3194

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

729

Electrochemical Properties and Applications of Sputtered Iridium Oxide Thin Films Evelina P. Slavchevaa,, Uwe Schnakenbergb, Wilfried Mokwab a

Institute of Electrochemistry and Energy Systems - Bulgarian Academy of Sciences, Sofia, Bulgaria b

Institute of Materials in Electrical Engineering I, Aachen University, Aachen, Germany

Abstract - The study reports on deposition of iridium oxide thin films by DC reactive magnetron sputtering. The sputtering process has been investigated applying the method of the generic curves. The composition, surface structure and morphology of the sputtered iridium oxide films (SIROFs) have been characterised using X-ray diffraction, SEM, and AFM analysis. An optimal combination of sputtering parameters has been established (40 sccm argon –8/18 sccm oxygen – 40% effective pump power), which yields stable microporous amorphous films with highly extended fractal surface. The electrochemical properties of these films have been investigated in neutral and acid aqueous solutions by methods of cyclovoltyammetry (CV), electrochemical impedance spectroscopy (EIS), and steady state polarization curves (PC). The SIROFs have shown an excellent electrochemical reversibility in physiological saline solution and have demonstrated high catalytic activity towards oxygen evolution reaction (OER) in 0.5M H2SO4. A current density of 100 mA/cm2 at potential of 1.6V (vs.SSE) was obtained at catalyst load of only 10µg/cm2. Keywords : iridium oxide, functional electro-stimulation, water splitting

1. Introduction Iridium oxide is well known electrode material with interesting properties and variety of technical applications. It possesses excellent electrical conductivity, enhanced electrochemical activity and is biocompatible. Iridium oxide has a low overvoltage for oxygen evolution reaction (OER) and is stable at high potentials

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E. P. Slavcheva et al.

where other materials corrode with a significant rate. Therefore, regardless its comparatively high price, the material is intensively studied as electrocatalyst for water splitting [1-6]. Due to its biocompatibility most recently it has also become an object of increasing interest as signal recording and stimulating material in field of functional electrostimulation (FES) [7,8]. Iridium oxide can be produced by variety of methods such as thermal decomposition of iridium salts, anodic oxidation of bulk iridium, electroplating, and reactive magnetron sputtering. The method of preparation and the particular experimental conditions can influence both the activity and the stability of the material. For instance, the properties of sputtered iridium oxide films (SIROFs) depend in a great extent on the parameters of the processes as the partial pressure of the oxygen species, pox in the sputtering plasma is one of the most important factors [8]. The present paper demonstrates the decisive influence of oxygen flow into the sputtering chamber on the surface structure and morphology of the SIROFs. It reports on the electrochemical behaviour of these films in neutral and in acid aqueous solution and their potential application as electrode material for functional electrostimulation and as catalytic material for oxygen evolution in water splitting. 2. Experimental The iridium oxide films were fabricated by reactive DC-magnetron sputtering from an Ir target in Ar/O2 plasma using commercial equipment Nordiko NS 2550. The sputtering process was performed in a DC mode using a previously defined optimal set of process parameters (40 sccm argon / 40% effective pump power / 100W DC power / ”cold” sputtering) [8]. All films were deposited on the oxidized Si-wafers upon 50 nm thick sublayer of Ti. The structure, morphology and roughness were analysed by X-ray diffraction (XRD), SEM, and AFM techniques. The electrochemical properties were studied by methods of cyclovoltammetry (CV), electrochemical impedance spectroscopy (EIS), and steady state polarization curves (PC) using EG&G PAR 283 Potentiostat and SI 1260 Impedance Analyser. An aqueous 0.9 % NaCl solution, simulating the physiological environment in the human body and a 0.5M H2SO4, which is one of the most frequently used electrolytes for water splitting, were chosen as test solutions. The experiments were performed at room temperature on test samples with geometric area of 0.5 cm2 in a standard three electrode glass cell with a high area Pt-wire counter electrode and an Ag/AgCl reference electrode.

Electrochemical Properties and Applications...

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3. Results and Discussion 3.1. Sputtering A series of SIROFs test samples were deposited at varying oxygen flow rate, Fox, of 0-40 sccm (standard cubic centimetre per minute). The observed changes in the oxygen partial pressure, pox , and the deposition rate, Rd, are illustrated by the so called generic curve of the process shown in Fig.1. The increase of Fox up to 10 sccm does not influence the partial pressure in the chamber, while at Fox = 10 sccm the value of pox increases abruptly indicating deviation of the system from a steady state, known as a “transition point” of the sputtering process. The further increase in Fox influences slightly the partial pressure in the chamber. According to the literature, the transition point of the reactive sputtering usually is accompanied by a considerable increase of pox and a corresponding decrease of Rd [9]. However, in the case of iridium the latter effect has not been observed. After the initial rapid increase in the vicinity of the transition point, Rd reaches a maximal value of 9.5 nm/min and further on does not change essentially. At the combination of process parameters used, the pressure in the reactor is comparatively low (9.0-9.5 mPa) and although the partial pressure of oxygen reaches transition point, the target stays in metallic mode producing mechanically stable iridium oxide films. 0.012

250

Si

Si

Si

0.008

pox [Pa]

8

0.006 6

0.004 pox

0.002

4

Intensity [a.u.]

0.010

Rd [nm/min]

10

200

Ir

150

Ir Ir

Rd

Ir

0.000 0

5

10

15

20

Fox [sccm]

Fig.1: Generic curves of SIROFs

2

100 20

30

40

50

60

70

80

90

2q [degrees]

Fig.2: XRD spectra of SIROFs

3.2. Surface analysis The X-Ray diffraction spectra of all SIROFs showed XRD iridium reflections with very low intensity, indicating an amorphous state of the deposited oxide layers. This is illustrated in Fig.2 presenting the spectrum of a sample sputtered at Fox = 12 sccm. The SEM images in Fig.3 demonstrate that the film of pure Ir is composed of uniform, densely packed granular particles with a feature size of about 70 nm (Fig.3a), while the oxide films obtained at higher oxygen flows

732

E. P. Slavcheva et al.

have a microporous structure, where the grains can not be resolved clearly (Fig.3b).The dependence of the average roughness factor, Ra, calculated from the AFM on the oxygen flow rate (Fig.4) shows that the surface roughness increases rapidly with the increase in Fox up to 12 sccm, does not change much in the range 12-20 sccm, while at higher oxygen flows the trend is back to smoother surfaces. 150

a) Fox = 0 sccm

b) Fox = 12 sccm

Ra [nm]

100

50

200 nm 0 0

10

20

30

40

Fox [sccm]

Fig.3: SEM images of selected SIROFs

Fig.4: Influence of Fox on film roughness, Ra

3.2. Electrochemical tests The performed surface analysis have shown that the SIROFs deposited in the vicinity of the transition point have the most developed surface with fine homogeneous microporous structure, excellent adhesion to the substrate and good mechanical stability. For this reason, all electrochemical tests were carried out with SIROFs deposited at Fox = 12 sccm. The as-prepared samples were activated to reach a stable charge delivery capacity Qc, using the activation procedure described elsewhere [8] and this value was used as a measure for the electrochemical activity. In both solutions the voltammograms gradually expanded with cyclic due to hydration of the film and activation of more surface regions of difficult accessibility. The activation was completed after about 100 potential cycles. The CV curves of the activated samples are shown in Fig.5. The shape of the curves in acid and neutral media is rather similar, while the “water window” potential ranges ( -1.2V to 1.3V for neutral and -0.4 to 1.4V for acid media) as well as the values of Qc (95 mC/cm2 and 150 mC/cm2, respectively) depend essentially on the electrolyte type. The main feature of CV curves after activation is their symmetry along the potential axis and the equality of the integral anodic and cathodic charges passing through the phase boundary electrode/electrolyte which, as it has been reported by other authors [1,2], is typical for a condenser-like interfaces. In such an aspect it can be stated that the activated SIROFs show a reversible electrochemical behaviour. The

Electrochemical Properties and Applications...

733

high Qc value means a high ratio active to geometric electrode surface, which is advantageous for electrode materials used in field of neural stimulation and electrocatalysis as it allows to reduce the size of the stimulating electrodes and the catalyst loading respectively, without any sacrifice of efficiency [2,4,6]. The impedance spectra of the SIROFs were recorded in a broad frequency range (50mHz to100kHz) at open circuit potential using ac perturbing voltage signal with amplitude of 10 mV. They were used to calculate the double layer capacity of the electrodes under study. The obtained values of 11.5 mF/cm2 and 17.2 mF/cm2 for the activated films in neutral and acid solutions, respectively are more than 2 orders of magnitude higher in comparison to the values of 60 µF/cm2 typical for smooth oxide electrodes [12]. 20

NaCl

-2

0

E [V]

Q [mC.cm ]

experimental curve IRW - correction

1.4

10

1.3

-10 H2SO4

-20 v = 0.1667 mV.s

1.2

-1.0

-0.5

0.0

0.5

1.0

1.5

E [V]

Fig.5: Fig.5: Cyclovoltammetric curve of SIROFs; scan rate v = 100 mV/s

10

-4

10

-3

-2

10

-1

-2

j [A.cm ]

Fig.6: Steady state polarization curve of SIROFs in H2SO4

The quasi-stationary polarization curve of OER in acid solution is presented in Fig.6. In the high overpotential domain the experimental data deviate from linearity, thus requiring correction for ohmic drop, IRW. After correction of the raw data according the procedure poposed in [13], an ohmic resistance 0.23W was determined from the DE/DI 伊vs. 1/DI 伊plot. This value is in accordance with the literature data for resistance of IrO2 and other conductive metallic oxides [2,5,6]. The corrected polarization curve (the compact line) deviates slightly from the raw data only at high current densities and revels two linear segments with Tafel slopes of 48 and 130 mV.decade-1 for the low and high current density domains, respectively. The results suggest that the registered deviation from linearity is due mainly to the changes in the kinetics of the OER at higher overpotentials. The low Tafel slope of ca. 48 mv.dec-1 is consistent with a primary water discharge, while the second Tafel slope of ca. 130 mV.dec-1 can be tentatively prescribed to a second charge transfer step of deprotonation of the adsorbent, followed by formation of oxygen molecule. Although our data correlate well with the existing literature [3,4,5], further experiments including

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a study on the temperature effects, are required in order to offer more precise and unambiguous reaction path of the oxygen evolution for our SIROFs. From theoretical point of view a good catalyst is characterised with low Tafel slope and high exchange current. For practical applications it is relevant to consider the potential required to provide a given amount of current as well as the stability of the catalysts with time, e.g. the potential variation with time as the electrode is subjected to continuous oxygen evolution. In this work the stability of the catalysts has not been tested in long term galvanostatic tests. Instead, the development of the electrode potential with time for a period of 18 hours under galvanostatic conditions (j=100 A.cm-2) was followed. A potential value of about 1.7V was observed during the whole test duration, indicating a stable electrochemical behaviour. On completion of polarisation experiments the cyclic voltammograms were re-measured. They were identical in shape to that obtained on the freshly activated electrodes as the registered decrease in the integral charge (the area under the CV curve) was negligible. This implies a constancy of the catalytic properties, at least under the experimental conditions in this work. 4. Conclusions DC magnetron sputtering in argon/oxygen plasma is a reliable method for production of SIROFs yielding reproducible thin iridium oxide layers with extended micro porous surface. The SIROFs obtained in the vicinity of the transition point of the sputtering process possess a broad “safe” potential range, high charge delivery capacity and constant low impedance in the frequency range used for neural stimulation and signal recording. They possess excellent electrical conductivity and have demonstrated very good catalytic properties towards OER and corrosion resistance under intensive oxygen evolution. The results obtained give credence to regard the SIROFs as an ideal electrode material for neural stimulating electrodes as well as a potential electrocatalyst for water splitting. 5. Acknowledgment The research has been carried out in the frames of the projects EPI-RET-II, (grant 01KP001), supported by the German Federal Ministry of Education and Research and the EU project Prometheas, Contract No. ICA-CT-2002-10001 6. References 1. C. P. De Pauli, S. Trasatti, J. Electroanal. Chem., 369 (1995) 161 2. L. A. Da Silva, V. A. Alves, M. A. P. Da Silva, S. Trasatti, J. F. C. Boodts, Electrochimica Acta, 42 (1997) 272

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3. T. Pauporte, F. Andolfatto, R. Durand, Electrochimica Acta, 45 (1999) 431 De Lima-Neto, Electrochimica Acta, 45 (2000) 4467 4. E. Rasten, G. Hagen, R. Tundol, Electrochimica Acta, 48 (2003) 3945 5. M. H. P. Santana, L. A. De Faria, J. F. C. Boodts, Electrchimica Acta, 49 (2004) 1925 6. L.M. Da Silva, D. V. Franco, L. A. De Faria, J. F. C. Boodts, Electrchimica Acta, 49 (2004) 3977 7. I. S. Lee, S.N. Whang, J.C. Park, D.H. Lee, W.S. Seo, Biomaterials, 24 (2003) 2225 8 E. Slavcheva, R. Vitushinsky, W. Mokwa, and U. Schnakenberg, J. Electrochem. Soc., 151 (7), (2004) E226 9. S. Zhu, F. Wang, J. of Appl. Phys., 84 (1998) 6399 10. J. D. Klein, S. L. Clauson, S. F. Cogan, J. Mat. Res., 4 (1989) 1505 11. T. M. Silva, J. E. Rito, A. M. P. Simoes, M. G. S. Ferreira, M. da Cunha Belo, and K. G. Watkins, Electrochem. Acta, 43 (1998) 203 12. S. Levine, A.L. Smith, Discuss. Faraday Soc., 52 (1967) 1290 13. N. Krastajic, S. Trasatti, J Appl. Electrochem., 28 (1998) 1291

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Passivation of Metals and Semiconductors, and Properties of Thin Oxide Layers P. Marcus and V. Maurice (Editors) © 2006 Elsevier B.V. All rights reserved.

737

Study of DSA® deactivation E. Herrera Calderon a,b *, R. Wüthricha, H. Bleulera, Ch. Comninellisb a

Ecole Polytechnique Fédérale de Lausanne, Laboratoire de systèmes robotiques, CH1015 Lausanne, Switzerland ([email protected]) b Ecole Polytechnique Fédérale de Lausanne, Groupe de génie électrochimique, CH1015 Lausanne, Switzerland

Abstract - Ti-DSA®-IrO2 electrodes were prepared and characterized electrochemically. Two time constants were identified with chronoamperometry measurements and confirmed by electrochemical impedance spectroscopy. With the help of a simple model, they are attributed to the contribution of the inner and outer surface of the anode. Keywords: DSA®, inner/outer surface, chronoamperometry, electrochemical impedance spectroscopy, lifetime, deactivation

1. Introduction The dimensionally stable anodes (DSA®) are one of the greatest technological breakthroughs of the past 50 years of electrochemistry. Metallic oxide coatings have been studied and successfully developed since 1969. The invention of these electrodes by Beer, at the end of the 60’s, has caused researchers to look for the perfection of such electrodes and new possibilities had been created. One of the best electrodes today and purchased easily is the one on titanium support covered with a layer of IrO2 and Ta2O5 for the O2 evolution (DSA®-O2), as well as those covered with RuO2 and TiO2 for the Cl2 evolution reaction (DSA®-Cl2) [1]. Nowadays, the DSA®-Cl2 electrodes have replaced almost the totality of the anodes made of graphite used for the preparation of chlorine. A new DSA type electrode is the Boron-Doped-Diamond (BDD) electrode which is particularly suited for usage in the removal of harmful contaminants (i.e. organic or inorganic contaminants) from waste streams even when present in relatively low concentrations. It is of extremely importance that these electrodes have a long service life. Therefore various studies about their deactivation process were undertaken in the past years [2]. Despite of these various investigations the different mechanism in DSA deactivation are still under discussion. Further

E. Herrera Calderon et al.

738

experimental investigations which could monitor in-situ the change of morphology on the surface and inside the anode would be helpful. A possible way to differentiate between inner and outer surface was first presented by S. Ardizzone et al. [3] based on cyclic voltammetry. In this contribution an extension of this method by chronoamperometry and impedance spectroscopy is presented for the investigation of the DSA-O2 electrode (Ti/IrO2). 2. Experimental 2.1. Electrode Preparation Ti/IrO2 electrodes were prepared by the thermal decomposition technique from H2IrCl6 (0,2 g) in isopropanol (10%) + water (90%) onto a surface of Ti support (2 x 2 cm) previously sandblasted and etched in hot oxalic acid (1M) for one hour. The mixture of precursor (H2IrCl6 in isopropanol and water) was put onto the etched Ti surface by painting. Afterwards the solvent was evaporated at approximately 80°C for 10 minutes and the electrode was transferred to a furnace pre-heated at 500°C where it stayed for 10 minutes. The above operations were repeated until the desired oxide loading was attained. A final one-hour calcination at the same temperature completed the procedure. Each electrode was kept in plastic cubes and sealed with tape. Figure 1a, shows a typical voltammogram of the prepared anode in HClO4(1M). A scanning electron micrograph illustrates the IrO2 coating (loading of 5.69 mg/cm2) after electrochemical characterization, (figure 1b).

a)

b) b)

j(mA/cm2)

4.00 2.00 0.00 -1

-0.5

-2.00

0

0.5

1

-4.00 E(V /MSE)

Fig. 1: a) Typical voltammogram of DSA®-IrO2 with a loading of 5.69mg/cm2, at 50 mV/s with N2 deaeration in HClO4 (1M). b) Scanning electron micrograph after IrO2 coating is characterized electrochemically in 1M HClO4.

2.2. Cell A three compartment all teflon cell 30ml was used throughout experiments and a computer-controlled potentiostat (EG&G, Model 273). The DSA®-IrO2

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was used as working electrode (0.78cm2 exposed area). Ohmic drop was minimized using a lugging capillary approaching the electrode from below while the platinum counter electrode ensured uniformity of the current on the sample. The electrode potential was read against a mercurous sulphate electrode (MSE) saturated with potassium sulphate that was submerged directly in the 1M HClO4 solution by means of a lugging capillary. Solutions were prepared with ultrapure water. 3. Results and Discussion 3.1. Chronoamperometry Chronoamperometry is done for various potential steps starting from 0 V. Before each step the electrode is kept for 10s at 0V in order to reach equilibrium.

j(mA/cm 2)

100 10 1 0 0

0.2

0.4

0.6

0.8

1

t(s)

Fig. 2. Typical chronoamperometry curve, for an anode with a loading of 2.35 mg/cm2, at 100mV in HClO4 (1M). Dashed lines show the two time constants (t1 and t2).

Figure 2 shows a typical result of chronoamperometry (note the logarithmic scale). From this figure it is seen that the response is the addition of two exponential decays. All measurements were fitted using the least square method by following equation : j(t) = j1 exp (-t/t1) + j2 exp (-t/t2)

(1)

The evolution of the parameters j1, t1 and j2 , t2 are depicted on figure 3. The slower time constant is typically around 50ms and the faster one is about 30 ms. 3.2. Electrochemical Impedance Spectroscopy (EIS) In order to confirm the results of section 3.1, electrochemical impedance spectroscopy studies have been done. Impedance measurements were taken at fixed potentials in the region of water stability: -0.7V to 0.7V (MSE) in HClO4

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740

-0.4

b)

-0.2

0 -50 0

150

(ms)

100 50

100 50

1

j1(mA/cm2)

a)

0.2

0.4

0

-100 -150

-0.4

-0.2

0

0.2

0.4

E(V/MSE)

E(V/MSE)

-0.4

-0.2

d)

30 20

(ms)

60 40 20 0 -20 0 -40 -60

10

2

j2(mA/cm2)

c)

0.2

0.4

0 -0.4

-0.2

0

0.2

0.4

E(V/MSE)

E(V/MSE)

Fig. 3. Evolution of the parameters of equation (1) in function of step potential E for an anode with a loading of 2.35 mg/cm2. Freely drawn lines show the tendency of the values j1, j2, t1, t2.

20 1 0 1.E-02 1.E+00 1.E+02 1.E+04 1.E+06 f(Hz)

-1

60 40

0.10

Y''(

10

100 80 °)

|Z|( cm2)

100

cm-2)

(1M). The frequency of the ac signal was varied between 50mHz and 100kHz and had an amplitude of 5mV. b) a) 0.20

0.00 0.00

0.10

0.20 0.30 Y'( W-1cm-2)

0.40

Fig. 4. Nyquist Diagram of DSA®-IrO2 at 100mV for an anode with a loading of 2.35 mg/cm2, in 1M HClO4 deaerated with N2. a) Bode diagram: ◊: |Z|, filled ฀: F). b) admittance plane (+: experimental, ฀: fitting). Fitting parameters: R1= 5.067W , R2 = 6.54W, C1= 0.016F, C2= 0.0041F.

Figure 4 shows a typical Nyquist diagram in admittance plane, where the fit is done. Starting from equation (1) an equivalent circuit of the system is proposed on figure 5. The components are related to the parameters of equation

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741

(1), by the following expressions (by simply writing the impedance corresponding to the equivalent circuit): j1= E/R1, j2= E/R2, t1= R1C1, t2= R2C2.

Fig. 5. Equivalent circuit proposed for the system.

The parameters R1, C1, R2, C2 obtained by fitting the EIS measurements, as shown on figure 4b, agree well with the results from figure 3 (agreement better than 10%). 3.3. Discussion A possible interpretation of the done measurements is presented now According to S. Ardizzone et al. [3], the following electrochemical reaction is proposed: ka – + F IrOx+y + 2yH + 2ye IrOx + yH2O (2) kc Based on equation (2) and if p denotes the total fractional coverage of the surface of IrOx+y , one can write, assuming y » 1: dp = k a c H 2O (1 - p ) - k c c H2 + p (3) dt with CH2O the water concentration and CH+ the proton concentration. The rate

constants are expressed as ka = ka° exp (ßh) and kc = kc° exp (-ßh) where the over potential h = E- Eth with E is the measured potential, Eth the equilibrium potential, b the transfer coefficient and ka°, kc° the standard rate constants (anodic and cathodic respectively). Equation (3), gives the evolution of the electrochemical system. The current density j is directly proportional to dp/dt. Let us now assume that there are two different kinds of sites in the DSA® anode. The first one are the easily accessible (under kinetic control) while the second one (under diffusion control) are the more difficult to access, because of the H+ diffusion [3]. The evolution equation for the easily accessible sites, pout is written as follows:

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dp out

= k a c H 2O (1 - p out ) - k c c H2 + p out (4) dt where cH2O and cH+ are supposed to be almost constant. It follows dp

out

dt

= A exp (-t/k) with A =

k a cH O 2

k a cH O + k c c 2

2 H+

and k =

1 . ka cH2O + kc cH+ 2

A

similar equation can be written for the less accessible sites pin with however different reactants concentrations (because of the diffusion). The total contribution to the current density is proportional to dp/dt as calculated by:

dp

=

dpout

+

dpin

(5) dt dt The solution of the equations (3)-(5) is identical to the experimental relation (1). According to this simple model, the two observed time constants can be related to the inner and outer surface of the DSA® anode. Further investigations are needed in order to clarify which time constant is linked to the inner and which one to the outer surface.

dt

4. Conclusions

Ti/IrO2 anodes were prepared. From the electrochemical characterization of the anodes, using choronoamperometry and impedance spectroscopy, two time constants are identified in the system. From a simple model, these two different time constants could be interpreted as the contribution from the inner and outer surface of the anode. These results confirm that reaction (2) is limited by diffusion as proposed by S. Ardizzone et al. [3]. Acknowledgements

Financial support from the Fonds National Suisse de la Recherche Scientifique is gratefully acknowledged. The authors are thankful to the Interdisciplinary Center of Electron Microscopy (CIME) from the EPFL. References 1. 2. 3.

J. Rolewicz, PhD thesis, Ecole Polytechnique Fédérale de Lausanne, EPF Lausanne, (1987). L.M. Da Silva, K. Fernandes, L.A. de Faria, and J. F. C. Boodts, Electrochimica Acta 49 (2004) 4893. S. Ardizzone, G. Fregonara and S. Trasatti, Electrochimica Acta 35 (1990) 263.

743

Index A Acidic, 35–40, 189–90, 378–80, 697–702 Adhesion, 145, 383–8, 480, 732 Adsorption, 13, 102, 112, 318, 320, 321–2, 351–6, 357–63, 365–70, 521, 677 AES, 10–11, 63, 267, 386, 510, 587 AFM, 372, 555–60, 570, 677, 732 Ageing, 11, 81, 152, 153, 596 Air-formed films, 29, 69, 220, 676 Alkaline, 55, 71–6, 107–12, 137–42, 251, 591–6, 678, 709–14 Alloy 600, 303, 403–408, 425–30, 507–12 Alloy 625, 299, 302, 303, 509 Alloy 690, 303, 403–408 Alumin(i)um, 125, 212, 331–6, 338–9, 469–73, 538, 609–14, 633–8, 677, 681, 685–90 Alumin(i)um alloys, 538, 621–6, 627–32 Alumin(i)um oxide, 263, 675–7, 479, 423 Aluminates, 107, 112 Amorphous, 65–70, 211–16, 343–6 Anodic oxidation, 233, 236, 271–6 Anodic oxide, 75, 263–9, 447, 457–62, 689–90 Anodization, 179–83, 272, 273, 275 Artificial, 3–8, 317–23, 352–3, 355–6, 561–6, 605–607 B Barium, 149–54 Barrier layer, 35, 400, 414–15, 432 Bilayer, 121, 652 Biocompatibility, 371, 383, 384, 730

Biomaterial, 371, 372, 384 Biomolecule, 360, 366, 390 Borate, 9–14, 83–7, 89–90, 91, 286, 288, 302, 303, 397–402, 441, 442, 443, 446, 452, 454, 455, 557, 558, 685 Boride, 507–12 Bovine serum albumin (BSA), 352, 353, 357, 366 Bromide, 131–6 C Carbon Steels, 83–87, 597–602, 709–14 Carbonate, 54, 55, 60, 125–9, 600, 601 Carboxylate, 682, 708 Cellular automata, 667–8 Cerium, 47–52, 225–30 Cerium oxide, 47, 50, 228, 230 Chloride, 107–12, 117, 149–54, 245–50, 354, 513–18, 555–60, 573–8, 579–84, 591–6, 633–8, 659–65, 712, 713 Chromium, 9–10, 50, 65, 143–8, 400, 405, 406, 417–23, 442, 491, 651–7 Chromium oxide, 13, 405, 406 Cold rolling, 373, 374, 375 Composite oxides, 685–90 Concrete, 305–10, 709–10 Conduction mechanism, 59, 64 Conversion coatings, 226, 228, 703–708 Cooperative interactions, 641–8 Copper, 125–129, 131–6, 137–42, 167–72, 243, 470, 525–30, 531–6 Copper alloys, 137–42, 167–72 Copper oxides, 140–1, 605 Corrosion fatigue, 537 Corrosion inhibition, 107, 359, 709–14

Index

744 Critical threshold, 641–48 Crystalline, 66, 113, 196, 211–16, 259, 558, 686 Crystallization, 65, 66, 200, 212, 213, 214, 215–16 Current transients, 37, 237, 451–6, 480, 485, 538 D Deactivation, 607–608, 737–42 Dealloying, 621 Defects, 10, 35, 37, 213, 311–16, 325–6, 335, 411 Depassivation, 5, 482–3, 672 Deposition, 119, 121, 239, 242, 373, 383–8, 677–80, 687–8 Depth profiling, 9, 10, 11, 622, 627 DFT, 326, 616, 659–65 Diffusion, 12, 20, 69, 227, 402, 433, 529, 641, 644, 650, 651–7, 672 Diffusivity, 311–16 Dislocations, 374, 519–24 Disruption, 161–6 Dissolution, 4, 7–8, 35, 38, 79–80, 119–24, 251–4, 431, 515, 531–2, 535, 567–72, 652–5, 726 Duplex steels, 463, 585–90 E EIS, 60, 196, 199, 200, 202, 258, 301, 308, 332, 345, 687, 709–14, 740–1 Electrochemical noise, 609–14 Electron Stimulated Oxidation, 9, 10, 12, 14 Electronic defects, 59, 397 Ellipsometry, 4, 218, 258, 273, 676, 691–6 EQCM, 119–20, 479

Etching, 11, 114, 233, 234, 245–6, 251, 419, 428, 686 Exfoliation, 546, 550, 552, 553 F Fe-Cr alloy, 9–14, 15, 285, 287, 288, 417, 652 Fe-Cr-Mo alloy, 3–8, 38 Fe-Cr-Ni alloy, 279–84, 418, 423 Fe-Mn-Al steel, 77–82 Fermi level pinning, 257–61 Ferrite, 79, 83–4, 85, 87, 466 Fe-V alloy, 71 Fe-V-Cr alloy, 71 Film thickness, 8, 11, 214, 221, 222, 401, 478, 686 Fluoride, 149–54, 155–60, 188–9, 191, 192, 246 Friction coefficient, 439–40, 444, 445, 446, 448, 457–8, 460, 461 G Gate oxides, 271–6 Gloss, 206, 209, 210 Grafting, 697–702 Grain size, 78, 464, 472, 521 Graining, 173–8 Growth kinetics, 545–53 Growth mechanism, 10, 14, 81, 116, 478, 656 H Hafnium oxide, 271–2, 274–5 Hardening, 373, 469, 533, 559 Hardness, 372, 373, 376, 441–3, 458–9, 482, 491 Heat treatment, 214, 464, 533, 621, 624, 627–8, 630, 631, 689

Index High-field model, 37 High pressure water, 403, 404, 408 High-temperature water, 431 Hydration, 11, 110, 326, 732 Hydrogen, 19, 175, 246, 326, 413, 723 I III-V semiconductors, 264 Implant alloys, 193, 377–81 Implantation, 47–52 Inclusions, 555–6, 567–72, 643 Incorporation, 94, 203, 423, 433, 596, 617, 619, 724 Incubation time, 503, 670–1, 672 Inhibitors, 54, 712 Intergranular corrosion, 545–6, 581 Intermetallics, 203, 379, 538, 615 Iridium oxide, 729–34 Iron, 89–94, 451–6, 579–84, 691–6 Iron oxalate, 693, 694 Iron oxides, 75, 84, 589, 594 K Kinetic model, 59, 397, 405 KMC, 651, 655, 656–7 L Lead, 425–30, 501–506, 621–6 Lead-induced corrosion, 503 Lifetime, 532 Light water reactor coolant, 397, 431–6 Lithium, 131–6, 411–16 Lithography, 173 Local probes, 463 Localiz(s)ed corrosion, 45, 54, 56, 300, 545–53, 561–6, 615–20 Long-term test, 591

745 M Magnesium, 217–22, 225–30, 618 Magnesium alloys, 225–30 Magnesium oxide, 220 Magnetite, 86, 92, 94, 491, 594, 693 MEAM, 653, 654 Membrane potential, 319, 320 Microhardness, 472 Microindentation, 451–6 Microstructure, 78, 84–5, 378, 464, 466 Mixed oxides, 337, 340 Modeling, 478, 483, 651 Molten electrolyte, 59–64 Molybdate, 392, 609–14, 723 Mössbauer spectroscopy, 149, 150, 156, 158 Mott-Schottky, 259, 282–4, 288, 296, 299–304, 309, 315, 505 N Nanoindentation, 440, 458 Nanoporous, 187–90, 191, 192 Nanoscratching, 445, 446, 448 Nanostructure, 179 Nanotubes, 179, 189 Nickel, 23–8, 29–34, 41–6, 119–24, 325–30, 403–408, 513–18, 519–24, 659–65 Nickel alloys, 513 Nickel oxide, 62 Ni-Cr alloy, 59, 299 Ni-Cr-Mo alloy, 53 Niobium, 69, 211–12, 213, 214, 215, 229 Nitrogen, 35–46, 266, 267, 268, 386 Noble elements, 603–608 Nonlinear phenomena, 525

Index

746 O Order-disorder transitions, 155 Organic, 245–50, 271–6, 675–82 Organic field effect, 271–6 Oscillations, 239–44, 525–30 Oxidation kinetics, 11, 13, 265, 269 Oxide structure, 90, 622 Oxygen reduction, 80, 94, 305–10 P Palladium, 390–1, 698 Passivation kinetics, 3–8, 65–70, 91, 92, 119–24, 143–8, 150, 233–7, 245–50, 351–6, 403–408, 417–23, 425–30, 519–24, 651–7, 721–7 Passivity breakdown, 53–4, 463–8, 603–608, 621–6, 627–32 Pearlite, 84–7 Perchlorate, 101–106 Permeability, 317–18, 320, 323 Photocurrent, 281–2, 295, 331–6, 342 Photon rupture, 561–6 Pickling, 579–84 Pitting corrosion, 45, 422, 556, 582, 585–90, 719 Point defect model, 10, 37, 311 Precipitation, 67, 97, 152, 173, 221, 318 Protein, 351–6, 357–8, 366, 367 PVD, 371, 372, 469 PWR, 403–404, 413, 489 Q Quantum chemistry, 652 R Radiation, 102, 104, 299–304, 592, 598 Radioactive waste, 53–7

RBS, 168, 338 Repassivation, 43–4, 56, 173–8, 301, 302, 419, 420, 445–6, 478–80, 513–18, 537–42, 607 Restructuring, 397–402, 431–6, 663, 664 Risk assessment, 53 Rust film, 317–23 S Segregation, 16, 36, 623, 632 Self-assembly, 389, 393 Self-organization, 179, 187 SEM, 78, 81, 84, 108, 112, 131, 132, 133, 134, 176, 185, 196, 243, 253, 387, 511, 575, 582, 583, 692, 706, 732 Semiconducting properties, 279–82 Semiconductors, 33, 245, 263 SERS, 89, 90, 92, 692 SHG, 257, 258, 261 Silicon, 233–7, 239–44, 245–50, 251–4 SIMS, 357, 568, 569 Single-crystal, 247, 457–2 Sol-gel coatings, 685, 688, 690 Solid-state transport, 398, 433, 436 Spinel structure, 89, 90, 93, 296, 569 Sputtering, 7, 267, 387, 414, 429, 568, 605, 731 Stainless steels, 23–8, 29–34, 35–40, 41–6, 47–52, 417–23, 513–18, 567–72 Steam generator, 403, 507–508 Steels, 77–82, 561–6, 585–90 STM, 90, 114, 115, 326, 652, 656 Strain, 523, 524, 527, 529, 531–6 Stress, 465–6, 555–60 Stress corrosion cracking, 404, 411, 501, 537 Structure, 95–100, 664, 665 Sulf(ph)ate, 171, 419, 455–6, 609–14, 619

Index Sulf(ph)ur, 568, 589, 700 Surface analysis, 228, 630–1, 731–2 Surfactant, 705, 707 Synchrotron radiation, 95

747 V Vacancies, 38, 90, 311, 326, 329, 532 Valve metal, 187–90 Vanadium, 72, 73, 75, 384

T W Tantalum oxide, 68, 211, 446–8, 457–62 TEM, 168, 265, 266, 267, 471, 535, 618 Texturing, 251–4 TGSCC, 525–30, 535 Thermal oxidation, 264, 265 Thickness, 4–5, 8, 11, 38, 139, 162, 169, 180, 182–3, 188, 189, 208, 219, 221, 222, 265, 273, 361, 405, 406, 407, 434 Tin, 137–42, 155–60 Tin oxide, 138, 141, 144 Tinplate, 143–8 Titanium, 179–83, 193–8, 199–204, 205–10, 325–20, 371–6, 377–81 Titanium alloy, 199, 377, 378 Titanium oxides, 185, 325–30 Tribocorrosion, 477–86, 495–500 Tungstate, 264, 726 Tungsten, 216, 513–18

Wear, 480, 481, 489–93, 498–500 X XPS, 5, 11, 17, 20, 27, 49, 50, 51, 60, 62, 73–6, 137–42, 144, 219, 228, 247, 249, 294, 359–63, 365–70, 404, 414, 511, 517, 571, 587, 594, 605, 606, 678, 679, 681, 682 XRD, 66, 81, 132, 730, 731 Z Zinc, 561–6, 703–708 Zinc alloy, 721–7 Zinc coated steel, 561–6 Zinc oxide, 436 Zirconium, 225, 229, 496 Zirconium alloy, 225–30 Zirconium oxide, 189, 500

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