EUROPEAN COMMISSION 5th EURATOM FRAMEWORK PROGRAMME 1998-2002 KEY ACTION : NUCLEAR FISSION
COBECOMA
State-of-the-art document on the COrrosion BEhaviour of COntainer MAterials
CONTRACT N° FIKW-CT-20014-20138
FINAL REPORT
B. Kursten1, E. Smailos2, I. Azkarate3, L. Werme4, N.R. Smart5, G. Santarini6
1 SCK•CEN, Mol, Belgium (co-ordinator) 2 FZK.INE, Karlsruhe, Germany 3 INASMET, San Sebastian, Spain 4 SKB, Stockholm, Sweden 5 SERCO ASSURANCE, Abingdon, United Kingdom 6 CEA/SACLAY, Gif Sur Yvette, France
Reporting Period: September 1, 2001 to December 31, 2003
Dissemination level : CO: confidential, only for partners of the COBECOMA project MT: Mid-term
October 2004 W&SF – COBECOMA
TABLE OF CONTENTS
TABLE OF CONTENTS...... I
PREFACE ...... 1
OBJECTIVES...... 2
EXECUTIVE SUMMARY ...... 4
KEYWORDS ...... 15
GLOSSARY OF TERMS...... 16
LIST OF ABBREVIATIONS AND ACRONYMS...... 22
PART A
A.1 INTRODUCTION ...... 26
A.2 HOST ROCK FORMATIONS FOR HLW/SPENT FUEL DISPOSAL...... 34 A.2.1 DISPOSAL IN SALT ...... 34 A.2.1.1 Characterisation of the host rock formation considered in Germany...... 34 A.2.2 DISPOSAL IN CLAY...... 35 A.2.2.1 Characterisation of the argillaceous host rock formation and bentonitic backfill material considered in Belgium ...... 35 A.2.2.2 Characterisation of the host rock formation considered in France ...... 38 A.2.2.3 Characterisation of the argillaceous host rock formation and bentonitic backfill material considered in Spain...... 39 A.2.2.3.1 Characterisation of the argillaceous host rock formation...... 40 A.2.2.3.2 Characterisation of the bentonitic backfill material ...... 40 A.2.3 DISPOSAL IN GRANITE ...... 43 A.2.3.1 Characterisation of the host rock formation considered in Sweden/Finland...... 43 A.2.3.2 Characterisation of the host rock formation considered in Spain ...... 44 A.2.4 DISPOSAL IN CEMENTITIOUS BACKFILL ...... 46
A.3 SCIENTIFIC BACKGROUND OF CORROSION DURING THE AEROBIC AND ANAEROBIC PERIOD ...... 49 A.3.1 CORROSION OF FE-BASED MATERIALS...... 49 A.3.1.1 Effects of environmental factors...... 51 A.3.1.1.1 Role of the chemistry of the aqueous environment in contact with the metallic materials ...... 51 A.3.1.1.2 Role of temperature...... 53 A.3.1.2 Role of pH ...... 53 A.3.1.3 Role of radiation ...... 54 A.3.2 CORROSION OF NI-BASED MATERIALS ...... 56 A.3.3 CORROSION OF CU-BASED MATERIALS ...... 59 A.3.3.1 Copper corrosion in the presence of oxygen ...... 59 A.3.3.1.1 Influence of chloride in the presence of oxygen ...... 60 2- A.3.3.1.2 Influence of other anions (SO4 , ammonia) ...... 61 A.3.3.1.3 Mechanism for the corrosion of copper in compacted buffer material saturated with saline, O2-containing groundwater...... 61 A.3.3.2 Copper corrosion in the absence of oxygen ...... 62 A.3.3.2.1 Influence of chloride ...... 62
i A.3.3.2.2 Influence of sulphide ...... 62 A.3.4 CORROSION OF TI-BASED MATERIALS...... 64
A.4 EXPERIMENTAL RESULTS...... 67 A.4.1 CORROSION STUDIES RELEVANT TO THE GERMAN DISPOSAL CONCEPT IN ROCK SALT ...... 67 A.4.1.1 Laboratory-scale corrosion experiments...... 69 A.4.1.1.1 Screening experiments...... 69 A.4.1.1.2 Experiments to evaluate the candidate metallic materials ...... 70 A.4.1.1.2.1 Background...... 70 A.4.1.1.2.2 Corrosion of unalloyed and low-alloyed steels in brines ...... 72 A.4.1.1.2.2.1 General and localised corrosion ...... 72 A.4.1.1.2.2.1.1 Influence of brine composition and temperature on steel corrosion in brines...... 72 A.4.1.1.2.2.1.2 Influence of pH on steel corrosion in brines ...... 77 A.4.1.1.2.2.1.3 Influence of chemical species on steel corrosion in brines...... 79 A.4.1.1.2.2.1.4 Influence of gamma irradiation on steel corrosion in brines ...... 83 A.4.1.1.2.2.1.4.1 Investigations at 90°C (in the MgCl2-rich Q-brine) ...... 83 A.4.1.1.2.2.1.4.2 Investigations at 150°C (in the three test brines 1, 2, and 3) .....84 A.4.1.1.2.2.1.5 Influence of welding on steel corrosion in brines...... 88 A.4.1.1.2.2.2 Stress corrosion cracking (SCC)...... 91 A.4.1.1.2.2.2.1 MgCl2-rich ‘Q’-brine ...... 91 A.4.1.1.2.2.2.2 NaCl-rich brine ‘3’ ...... 93 A.4.1.1.2.3 Corrosion of Hastelloy C-4 in brines ...... 95 A.4.1.1.2.4 Corrosion of Ti 99.8-Pd in brines ...... 99 A.4.1.1.3 Experiments on real specimens originating from the POLLUX-disposal container ....102 A.4.1.1.3.1 Background...... 102 A.4.1.1.3.2 Results and discussion ...... 103 A.4.1.1.3.3 Conclusions ...... 106 A.4.1.2 In-situ corrosion experiments ...... 107 A.4.1.2.1 Investigations in rock salt at rock temperature ...... 107 A.4.1.2.2 Brine Migration Test (investigations in rock salt / limited amounts of brine at HLW design temperature)...... 108 A.4.1.2.3 Testing of welded tubes (model containers) in rock salt / brine at HLW design temperature...... 109 A.4.1.2.4 Long-term testing under normal operating conditions at high temperature ...... 110 A.4.1.2.4.1 Results from the DEBORA-Experiment...... 111 A.4.1.2.4.2 Results from the BAMBUS II-Experiment...... 112 A.4.1.2.5 Conclusions...... 114 A.4.2 CORROSION STUDIES RELEVANT TO THE BELGIAN DISPOSAL CONCEPT IN CLAY ...... 116 A.4.2.1 Background information ...... 116 A.4.2.1.1 In situ corrosion experiments ...... 116 A.4.2.1.2 Laboratory-scale corrosion experiments ...... 119 A.4.2.1.2.1 Electrochemical experiments...... 120 A.4.2.1.2.2 Long-term immersion tests ...... 122 A.4.2.2 Experimental results ...... 123 A.4.2.2.1 In situ corrosion experiments ...... 123 A.4.2.2.1.1 Carbon steel ...... 123 A.4.2.2.1.2 Stainless steels ...... 129 A.4.2.2.1.3 Ni- and Ti-based alloys...... 130 A.4.2.2.1.4 Conclusions...... 130 A.4.2.2.2 Laboratory-scale corrosion experiments ...... 133 A.4.2.2.2.1 Determination of ECORR (free corrosion potential) based on experimental results ...... 133 A.4.2.2.2.2 Pitting corrosion studies ...... 135 A.4.2.2.2.2.1 Experimental results ...... 135 A.4.2.2.2.2.2 Discussion...... 141 A.4.2.2.2.2.2.1 Carbon steel ...... 141 A.4.2.2.2.2.2.2 Stainless steels...... 141 A.4.2.2.2.2.2.2.1 Influence of alloy composition...... 141 A.4.2.2.2.2.2.2.2 Influence of temperature ...... 141 A.4.2.2.2.2.2.2.3 Influence of chloride...... 143 A.4.2.2.2.2.2.2.4 Influence of sulphate...... 147 A.4.2.2.2.2.2.2.5 Influence of thiosulphate ...... 148 A.4.2.2.2.2.2.2.6 Influence of hydrogen peroxide (gamma radiolytic product) ....150 A.4.2.2.2.2.2.3 Ni-alloys...... 151
ii A.4.2.2.2.2.2.4 Ti-alloys ...... 153 A.4.3. GRANITIC ENVIRONMENT ...... 154 A.4.3.1 Corrosion studies relevant to the Swedish and Finnish disposal systems in granite ...... 154 A.4.3.1.1 Corrosion prior to water saturation ...... 154 A.4.3.1.2 Corrosion during water saturation ...... 155 A.4.3.1.3 Corrosion after water saturation ...... 156 A.4.3.1.3.1 General corrosion...... 157 A.4.3.1.3.1.1 Determination of ECORR (corrosion potential) and iCORR (corrosion rate) based on modelling studies ...... 157 A.4.3.1.3.1.2 Determination of ECORR (corrosion potential) and iCORR (corrosion rate) based on experimental results ...... 160 A.4.3.1.3.1.3 Influence of welding on the corrosion properties ...... 163 A.4.3.1.3.1.4 Influence of increased pressure at repository depth ...... 163 A.4.3.1.3.1.5 Influence of methane ...... 163 A.4.3.1.3.2 Localised corrosion ...... 163 A.4.3.1.3.2.1 Pitting corrosion ...... 164 A.4.3.1.3.2.2 Crevice corrosion ...... 165 A.4.3.1.3.2.3 Ants-nest corrosion ...... 165 A.4.3.1.3.3 Stress corrosion cracking (SCC) ...... 166 A.4.3.1.3.3.1 Influence of ammonia...... 166 A.4.3.1.3.3.2 Influence of acetate...... 166 A.4.3.1.3.3.3 Influence of chloride...... 166 A.4.3.1.3.3.4 Influence of temperature ...... 167 A.4.3.1.3.3.5 General conclusions concerning SCC ...... 167 A.4.3.1.3.4 Microbially influenced corrosion (MIC) ...... 168 A.4.3.1.3.5 Corrosion induced by radiation effects ...... 168 A.4.3.1.3.5.1 Corrosion in air in the presence of γ-radiation ...... 168 A.4.3.1.3.5.2 Corrosion in water in the presence of γ-radiation ...... 169 A.4.3.2 Corrosion studies relevant to the Spanish disposal concept in granite (investigations performed at ENRESA/INASMET)...... 170 A.4.3.2.1 Experimental procedure ...... 170 A.4.3.2.2 Carbon steels ...... 173 A.4.3.2.2.1 General and localised corrosion...... 173 A.4.3.2.2.2 Stress Corrosion Cracking (SCC)...... 177 A.4.3.2.3 Ni-Cr steels...... 179 A.4.3.2.3.1 General and crevice corrosion of 316L stainless steel...... 179 A.4.3.2.3.2 Stress corrosion cracking (SCC) of 316L stainless steel...... 180 A.4.3.2.4 Cu-base alloys...... 182 A.4.3.2.5.1 General and localised corrosion of Cu-OF and Cu30Ni ...... 182 A.4.3.2.5.2 Stress Corrosion Cracking (SCC) of Cu-OF and Cu30Ni...... 183 A.4.4 CORROSION STUDIES PERFORMED IN CEMENTITIOUS BACKFILL MATERIAL...... 187 A.4.4.1 Carbon steel...... 188 A.4.4.1.1 Uniform corrosion ...... 188 A.4.4.1.1.1 Aerobic corrosion behaviour ...... 188 A.4.4.1.1.2 Anaerobic corrosion behaviour ...... 188 A.4.4.1.2 Localised corrosion ...... 192 A.4.4.2 Stainless steel ...... 194 A.4.4.2.1 Uniform corrosion ...... 194 A.4.4.2.1.1 Aerobic corrosion behaviour ...... 194 A.4.4.2.1.2 Anaerobic corrosion behaviour ...... 196 A.4.4.2.2 Localised corrosion ...... 197 A.4.4.2.2.1 Pitting corrosion...... 197 A.4.4.2.2.2 Crevice corrosion ...... 199 A.4.4.2.2.3 Stress Corrosion Cracking (SCC) ...... 199 A.4.4.2.3 Galvanic corrosion...... 202 A.4.4.3 Passive film formation on carbon steel and stainless steel...... 203
A.5 MODELLING ...... 206 A.5.1 MODELLING STUDIES OF CARBON STEEL AND LOW-ALLOYED STEEL CONTAINERS ...... 206 A.5.2 MODELLING STUDIES OF COPPER CANISTERS...... 207 A.5.2.1 Modelling studies to predict the long-term general corrosion behaviour of Cu canisters in granitic environments...... 207 A.5.2.1.1 Mass-transport limited approach ...... 207 A.5.2.1.2 Kinetic approach...... 210
iii A.5.2.2 Modelling studies to predict the long-term pitting behaviour of Cu canisters ...... 213 A.5.2.3 Modelling studies to predict the long-term SCC behaviour of Cu canisters ...... 217 A.5.2.4 Lifetime predictions of Cu canisters ...... 218
APPENDIX A - CONTRIBUTION (%) OF NUCLEAR POWER PLANTS TO THE WORLD'S ELECTRICITY PRODUCTION [3] ...... 220
APPENDIX B - THE DIFFERENT DISPOSAL CONCEPTS ENVISAGED IN VARIOUS COUNTRIES OPERATING NUCLEAR POWER PLANTS (E.G. BELGIUM, FINLAND, FRANCE, GERMANY, SPAIN, SWEDEN, UK)...... 221 B.1 CURRENT DISPOSAL CONCEPT IN BELGIUM...... 221 B.2 CURRENT DISPOSAL CONCEPT IN FRANCE...... 223 B.2.1 Concept for the disposal of vitrified HLW ...... 223 B.2.2 Concept for the disposal of spent fuel ...... 225 B.3 CURRENT DISPOSAL CONCEPT IN SWEDEN/FINLAND ...... 228 B.4 CURRENT DISPOSAL CONCEPT IN GERMANY ...... 230 B.5 CURRENT DISPOSAL CONCEPT IN SPAIN ...... 234 B.6 CURRENT DISPOSAL CONCEPT IN THE UNITED KINGDOM ...... 237
PART B
B.1 COMMON SCIENTIFIC APPROACH ...... 240
B.2 POSSIBLE MODES OF CORROSION AFFECTING THE CONTAINER ...... 241
B.3 PARAMETERS AFFECTING THE CORROSION BEHAVIOUR OF CANDIDATE CONTAINER MATERIALS ...... 244
B.4 INTERNAL CORROSION ...... 248
B.5 INFORMATION FROM ARCHAEOLOGICAL ANALOGUES...... 249
B.6 DETERMINING AN EVOLUTIANARY PATH: CHARACTERISATION OF THE NEAR- FIELD ENVIRONMENT AS A FUNCTION OF TIME...... 253
B.7 APPROACH FOR MODELLING...... 255 B.7.1 LIFETIME PREDICTION FOR CORROSION-ALLOWANCE MATERIALS (CARBON STEEL AND LOW- ALLOY STEEL)...... 255 B.7.1.1 General corrosion ...... 255 B.7.1.2 Localised corrosion ...... 258 B.7.2 LIFETIME PREDICTION FOR PASSIVE MATERIALS (STAINLESS STEEL, NI-CR-MO ALLOYS)..... 261
B.8 CEMENTITIOUS BACKFILL MATERIALS...... 264 B.8.1 INTRODUCTION OF CRACKING DUE TO VARIATION IN THERMAL CONDUCTIVITY OF THE VARIOUS METALLIC AND CEMENTITIOUS COMPONENTS...... 264 B.8.2 CONCRETE SPALLING ...... 264
ACKNOWLEDGEMENTS...... 265
REFERENCES...... 266
iv PREFACE
This report is the outcome of a Thematic Network concerning the 'COrrosion BEhaviour of COntainer MAterials (COBECOMA)' considered for the underground disposal of radioactive waste in deep geological host rock formations. This Thematic Network is composed of five research organisations, i.e. SCK•CEN (Belgium), SKB (Sweden), SERCO ASSURANCE (United Kingdom), INASMET (Spain), and FZK.INE (Germany), originating from five different member states of the European Union. Also, CEA/SACLAY (France) was appointed as an expert. The consortium is a complementary group of research organisations that have a vast amount of specific expertise and scientific know-how in the field of container corrosion in the different envisaged geological formations and backfill materials, e.g.. salt, clay, granite, and cement.
This Thematic Network was co-ordinated by SCK•CEN and was entirely financed by the Commission of the European Communities (CEC) (Contract No. FIKW-CT- 2001-20138).
1 OBJECTIVES
All the spent fuel (SF) and conditioned high-level radioactive waste (HLW) that has been produced so far is currently being stored under water in special ponds, or in dry concrete structures or casks. However, these options were not designed to store the waste for an indefinite period of time. Although there exists no permanent method of disposal for HLW and SF yet, worldwide, deep underground burial in stable geological formations is believed to be the favoured approach to deal with long-lived radioactive waste (spent fuel and HLW derived from reprocessing spent fuel). This radioactive waste would be disposed of after a period of storage on the surface in order for both the temperature and the radioactivity of the wastes to gradually decrease and thereby simplifying their handling and disposal considerably.
Various repository design options are currently being pursued internationally of which all are based on the multibarrier concept. This multibarrier concept relies on the combined effectiveness of various natural and engineered barriers to ensure an effective long-term isolation and containment of radionuclides by preventing or retarding their release from the repository site to the human environment (the biosphere). In this concept, the waste packaging (container, overpack), which encapsulates the radioactive waste, plays an important role in separating the buried waste from the geosphere.
Over the years, a lot of investigations have been performed to choose potential container materials and to characterise their corrosion behaviour, thereby creating a lot of scientific data and know-how through the separate national HLW/SF management programmes. The European Commission (EC) has participated, within the framework of its R&D programme on Management and Storage of Radioactive Waste, in several research projects aimed at the selection of suitable container materials for the geological disposal of conditioned HLW and spent fuel. This research has been performed in the framework of cost-shared contracts with research centres in various member states of the European Communities and mainly focused on assessing the long-term interactions of potential container materials in contact with their potential disposal environments, i.e. salt, clay, and granite.
However, so far, the results from the individual research programmes have only been presented and discussed by the participating laboratories in six-monthly reports, final reports, papers, peer-reviewed literature, etc. Therefore, except in overview papers given at CEC conferences, very little of the research from the separate national programmes is integrated. As many national HLW/SF disposal programmes are heading in a similar direction towards final choices of materials and are faced with similar challenges and issues, the time seemed opportune for a more thorough and critical review of existing research in the field of the corrosion behaviour of container materials to help define the areas where future research is required.
The primary objective of the Thematic Network was to prepare a state-of-the-art document, concerning the various aspects of corrosion affecting containers used for the final disposal of HLW/SF, with the aim of formulating the major needs for future R&D on container corrosion-related issues.
2 The major topics that are addressed in this report are: • collecting results on container corrosion studies in the field of final disposal of HLW and spent fuel in various geological formations (e.g. salt, granite, clay) under investigation in various European Countries, gathered over the last 15-20 years (PART A), • critically evaluating the available results in order to put forward more profound interpretations, conclusions, and/or decisions, and establishing the needs for future R&D on container corrosion-related issues (PART B).
An important aspect of this state-of-the-art document was the integration of the various different approaches that exist in container corrosion studies within the various European Countries, such as: - corrosion-allowance and/or corrosion-resistant materials are considered as candidate container materials, - different disposal concepts are being studied, - different views are taken about to the role of the canister/overpack as an engineered barrier, - different constraints and/or boundary conditions exist, depending on the type of waste for disposal (e.g.vitrified HLW or spent fuel).
An important additional objective of the Thematic Network is to identify ways of integrating R&D amongst the various EU member states.
The Thematic Network further contributes to different sub-objectives of the EC programme under the Safety of the Fuel Cycle: • Performance assessment of repository systems. This report could contribute towards creating a quality assurance system for computer models used in the performance assessment of waste disposal systems. The possibility of developing a materials degradation model, based on the currently available data, is investigated. In case the current data turn out to be insufficient. A list of areas where further corrosion data are required has been formulated and experiments that could provide this data are proposed. • Long-term behaviour of repository systems. The Thematic Network is situated within the framework of safety assessment of the engineered barriers. An additional objective of this state-of-the-art document was to identify the topics that need further investigation to assess the long-term corrosion behaviour of waste containers in a repository system.
3 EXECUTIVE SUMMARY
The intensive industrialisation worldwide requires an ever-increasing production of energy. Currently, almost one fifth (16%) of the world’s total electricity production arises from nuclear power plants. Besides producing electricity for the benefit of society, nuclear power, unfortunately, is also leaving a legacy of dangerous radioactive residues for which a safe long-term solution has to be found to protect the many generations to come.
At present, deep underground disposal in stable geological formations, such as salt, clay, granite and tuff is the favoured option that is being pursued worldwide to deal with long-lived radioactive waste in a feasible and safe manner. The repository designs and the disposal concepts considered within the various European countries vary from country to country and according to the type of waste. However, the multibarrier principle forms the backbone of each of the deep geological disposal concepts. This multibarrier principle relies on a series of natural and engineered barriers that act in concert to isolate the radionuclides, on the one hand, and to retard any radionuclide release from the waste to the biosphere, on the other hand. The engineered, or man-made, barriers are the waste matrix, the metallic container and the buffer material. From all these engineered barriers, the metallic container is one of the principal ones.
Over the years, two approaches have evolved: the corrosion-allowance concept and the corrosion-resistant concept. In the corrosion-allowance concept, materials are considered of lower corrosion resistance that can be used economically in sufficient thickness for corrosion penetration for a desired lifetime. These materials corrode actively at a significant, but predictable general corrosion rate. No localised corrosion such as pitting or crevice corrosion is expected. In the corrosion-resistant concept, materials are considered that exhibit a very high corrosion resistance in the disposal environment. These materials corrode passively at a very low uniform corrosion rate. They owe this high durability to the existence of a protective adherent oxide layer, which is called ‘the passive surface film’. Therefore, they can be used with a relatively small thickness. However, for these materials, the risk of localised corrosion, such as pitting and crevice corrosion, has to be taken into account because the protective film may break down locally.
Table 1 gives a comparison of the geochemical composition of different underground host rock formations (together with the buffer material) considered as potential disposal environments within various EU-countries. The compositional factors that vary considerably between the various potential host rock formations and that can significantly influence the corrosion behaviour of the candidate container materials are the pH and the concentrations of hydroxide, chloride and sulphate ions:
• the three salt brines (brines 1 and 2 are MgCl2-rich brines and brine 3 is a NaCl- rich brine) – all of which are representative for the Gorleben rock salt repository environment in Germany – are slightly acidic (the pH ranges from 4.1 to 6.5). The various argillaceous and granitic host rock formations, studied in Belgium, France,
4 Spain, Sweden and UK, have a pH in the range from 6 to 9.7. The bentonite-based buffer materials have a near-neutral to slightly alkaline pH (the pH ranges from 6 to ~8). The cement-based buffer material considered in the UK is highly alkaline (pH = 13.1). • the high alkalinity of the cement-based buffer material considered in the UK, is also reflected in its high hydroxide concentration (3.9 × 10³ mg/L). • the salt brines (Germany) contain extremely high chloride concentrations (between 190,000 and 350,000 mg/L Cl-). The disposal environment considered in Belgium contains the lowest amount of chloride (less than 70 mg/L). The disposal environments studied in France, Spain, Sweden and UK contain between 300 and 22,000 mg/L Cl-. • the salt brines (Germany) contain the highest concentration of sulphate (up to 2- 15,500 mg/L SO4 ). The sulphate content of the argillaceous and granitic disposal environments (Belgium, France, Spain, Sweden and UK) is comparable and depending on the degree of oxidation varies in the range between 0 and 3,000 mg/L.
TABLE 1. Geochemical composition of potential disposal environments (together with the buffer material) considered within various EU-countries.
A wide variety of metallic materials has been investigated as candidate container material within the various EU-countries, as illustrated in Table 2. Although most of the individual national corrosion programmes were developed separately, a lot of identical metallic materials have been considered (highlighted in red in Table 2). In Belgium, Germany and Spain, a lot of types of materials (carbon steels, stainless
5 steels, Ni-, Ti and Cu-alloys) have been investigated, while in Sweden and Finland, the main effort was directed on copper.
TABLE 2. Materials selection of candidate container materials investigated within various Eu- countries.
A consensus exists among scientists on which parameters can influence the lifetime of the metallic container, which techniques are most appropriate to gain the necessary information and which modes of corrosion are most significant in relation to degradation of the integrity of the metallic container. Table 3 summarizes the parameters investigated, the techniques applied and the modes of corrosion evaluated during the various individual national research programmes within the EU- countries.
6 TABLE 3. Investigated parameters, applied techniques and evaluated modes of corrosion from the various individual national research programmes within the EU-countries.
The scientific approach adopted in the individual national corrosion programmes was found to be somewhat similar:
• in a first stage, screening studies were carried out. These studies usually involved some preliminary experiments on a very large number of metallic materials in order to make a classification of their corrosion behaviour, thereby enabling a first selection of the most appropriate candidate container materials to be made. • afterwards more detailed studies were performed on this limited selection of metallic materials. These detailed studies involved in situ and/or laboratory experiments depending on the availability of an underground research facility. The laboratory experiments that were used were immersion tests, electrochemical experiments, radiochemical experiments and specific experiments investigating SCC-susceptibility. These detailed studies aimed at investigating the influence of several parameters such as temperature, pH, environmental composition (concentration of aggressive ions), presence of a γ-radiation field (including the influence of radiolysis products), etc. • at the same time full scale demonstration tests were started. These experiments usually had the following objectives: (i) to demonstrate the feasibility of constructing an underground disposal facility or (ii) to demonstrate the feasibility of welding large scale thick-walled containers. Often several metallic coupons and instrumentation for monitoring were installed. Some of these experiments are ongoing or have only recently been dismantled and therefore not all experimental data are yet available. • over the last decade, an increasing effort has been invested in modelling studies and also in determining whether valuable information can be gained by studying natural analogues such as archaeological artefacts. This approach could become
7 an important tool in the future to predict the long-term corrosion behaviour of the candidate container materials because we are dealing with one of the biggest scientific challenges – a challenge which does not apply to any other industrial application – namely the prediction of the corrosion behaviour of the metallic container over a period up to hundreds of thousands of years.
This report presents a review of the current knowledge of the corrosion behaviour of potential candidate container materials for the final disposal of HLW/spent fuel in deep underground formations (salt, clay, granite), which has been gathered during the last 15 to 20 years within various European countries and provides recommendations for future R&D.
The main results for salt disposal environments (Germany), obtained for carbon steel TStE 355 (corrosion-allowance concept) and Ti99.9-Pd (corrosion-resistant concept), are summarized below:
• carbon steel TStE 355: - carbon steel corrodes actively in MgCl2- and NaCl-rich brines. The average general corrosion rates are given in Table 4. The higher corrosivity of the MgCl2-rich brines compared to the NaCl-rich brine is attributed to its higher chloride concentration and to the presence of Mg2+. It appears that Mg2+, which is incorporated into the ferrous hydroxide layer replacing Fe2+, interferes with the normally expected conversion of Fe(OH)2 to Fe3O4. The corrosion product formed (Fe, Mg)(OH)2 appears to have little or no ability to protect the carbon steel from corrosion in contrast to Fe3O4 which is formed in the NaCl-rich brine. The large difference in general corrosion rate derived from the laboratory and the in situ experiments is attributed to the differences in the available amount of brine: in the in situ experiments, the samples were exposed to a limited amount of brine, whereas in the laboratory-scale experiments, the samples were exposed to an excess of brine. - the initial pH has no significant influence on the general corrosion rate neither in MgCl2-rich brines in the pH-range between 3 and 7 nor in the NaCl-rich brine in the pH-range between 1 and 5. A decrease of the general corrosion rate was observed in the NaCl-rich brine from 50 to 26 µm/year when the pH was increased from 1 to 10. - - - the addition of salt impurities (B(OH)4 ), radiolytic products (H2O2, ClO ) and corrosion products (Fe3+) caused an increase in the general corrosion rate from 5 to 236 µm/year in the NaCl-rich brine at 90°C and from 17 to 120 µm/year in MgCl2-rich brines at 170°C. - welding caused a reduction of the resistance to localised corrosion in MgCl2-rich brines (150°C). The TIG- and EB-welded specimens suffered from severe localised corrosion attack in the weld region and the heat affected zone (the maximum depth of local attack reached values up to 500-700 µm after ~300 days). The application of a thermal heat treatment (2 hours at 600°C) improved the resistance to localised corrosion. - carbon steel exhibited a slight sensitivity to stress corrosion cracking and loss of ductility in the NaCl-rich brine. However, this was only observed at very slow strain rates – strain rates that are not expected to occur under realistic repository conditions. - the imposition of a γ-radiation field (10 Gy/h) did not affect the general corrosion rate in the NaCl-rich brine up to 150°C. Also, in the MgCl2-rich brines, a γ-dose
8 rate of 1-10² Gy/h had no influence of the general corrosion rate. At high dose rates of 10³ Gy/h, on the contrary, the general corrosion rate increased by a factor of about 10.
TABLE 4. General corrosion rates (µm/year) of carbon steel TStE 355 in salt disposal environments.
T MgCl2 NaCl NaCl (°C) (lab. exp.) (lab. exp.) (in situ exp.)
90 70 5 < 0.1 170 199 46 < 0.1
• Ti-alloy Ti99.8-Pd: - Ti99.8-Pd corrodes passively in MgCl2-rich brines with an average general corrosion rate of the order of 0.1-0.2 µm/year, which was not influenced noticeably by the temperature up to 200°C. - Ti99.8-Pd is not susceptible neither to pitting corrosion, nor to stress corrosion cracking both in MgCl2-rich and NaCl-rich brines. - neither the presence of a γ-radiation field, nor the addition of radiolytic products - H2O2 and ClO exerted a significant influence.
In argillaceous disposal environments, two approaches are being considered: the use of low-alloy or unalloyed steels (e.g. France) or the use of passive Fe-Ni-Cr-Mo alloys (e.g. Belgium). This difference in approach is mainly based on one of two factors: (i) the possible occurrence of localised corrosion during the oxic period, which depends on the chloride content of the disposal environment (relevant to passive materials) and (ii) the amount of hydrogen gas produced during the anoxic period (relevant to low-alloy and unalloyed steels). The main conclusions of the experiments performed in an argillaceous disposal environment (Belgium) are summarized below:
• carbon steel: - the average general corrosion rates are given in Table 5. The corrosion rate increased with increasing temperature from 1.8 µm/year at 16°C to 8.6 µm/year at 170°C. The corrosion rate decreased with increasing exposure time. - carbon steel was susceptible to pitting corrosion, with maximum observed pit depths up to 240 µm (after a 2 year exposure period at 90°C). - the weld region and the heat affected zone were slightly more susceptible to pitting corrosion than the parent material: pits up to 90 and 130 µm deep were measured in the parent material and the weld region respectively (after 7 years of exposure at 90°C). - the susceptiblity to pitting of the weld became even more pronounced in the presence of a γ-radiation field (400 Gy/h): pits up to 2.2 mm deep were observed after a 5 year exposure period at 80°C.
9 TABLE 5. Average general corrosion rates (µm/year) of carbon steel in an argillaceous disposal environment (Boom clay formation).
T t γ vCORR Dmax (°C) (year) (Gy/h) (µm/year) (µm)
170 4.7 0 8.59 - 90 1.7 0 7.68 240 (parent mat.) 90 7.0 0 4.65 90 (parent mat.) 130 (weld region) 16 4.7 0 1.81 100-120 (parent mat.)
80 5.0 400 3.8 150 (parent mat.) 20.3 2,170 (weld region)
• stainless steels: - the average general corrosion rate of stainless steel is < 0.1 µm/year and independent of temperature (up to 170°C). - stainless steel is resistant to pitting corrosion under ‘normal’ repository - 2- conditions, i.e. [Cl ] < 100 mg/L and [S2O3 ] < 17 mg/L. - under oxic conditions (140°C), pitting is expected to occur at [Cl-] > 10,000 mg/L. - under anoxic conditions (140°C), stainless steel is resistant to pitting corrosion in solutions containing chloride concentrations up to 50,000 mg/L. - temperature had a distinct influence on the pitting behaviour: with increasing temperature, the pit depth and pit density increased significantly and the critical potential for pit nucleation (ENP) shifted drastically in the active direction.
• Ni- and Ti-alloys: - the average general corrosion rate of Ni- and Ti-alloys is < 0.1 µm/year and also independent of temperature (up to 170°C). - Ni- and Ti-alloys are resistant to pitting corrosion under all tested conditions, i.e. - 2- temperature up to 140°C, [Cl ] up to 50,000 mg/L and [S2O3 ] up to 200 mg/L. - the Ni-alloys are however susceptible to crevice corrosion under oxic conditions at 140°C at [Cl-] > 20,000 mg/L.
The most detailed assessments of the lifetimes of copper containers have been performed in Sweden/Finland and Canada (Table 6). The corrosion processes considered in these studies include (i) general corrosion under oxic and anoxic conditions, (ii) localised corrosion (pitting), (iii) microbially influenced corrosion (MIC), and (iv) stress corrosion cracking (SCC). In both the Scandinavian and Canadian assessments, it is believed that general corrosion and pitting corrosion are the processes most likely to cause corrosion of the container. Both mass-balance and kinetic modelling have been used to predict general corrosion. For pitting, empirical pitting factors and statistical analyses have been used. The impact of microbially 2- influenced corrosion is expected to be limited, due to the limited quantities of SO4 that can be reduced to HS-. Stress corrosion cracking is believed to be highly unlikely to occur, either because the maximum concentration of SCC agents and the corrosion potential, lie below the respective threshold values for SCC, or because the creep rate will exceed the crack growth rate. Despite differences in the repository
10 conditions and in the approaches taken to make long-term predictions, the outcome of the Scandinavian and Canadian assessments is that the copper container lifetimes will exceed 106 years. It was apparent from these assessments that corrosion will not be the limiting factor when designing the container thickness.
TABLE 6. Comparison of various predictions of the service life of copper containers.
The main conclusions for a granitic disposal environment (Spain) are summarized below:
• carbon steel TStE 355: - the general corrosion rate diminished asymptotically with exposure time; the corrosion rates in granitic-bentonite water at 90°C and 120°C under anoxic conditions were 6 and 14 µm/year respectively. - carbon steel was resistant to pitting corrosion in aqueous bentonite-buffered granitic groundwater, but was found to be susceptible to non-uniform corrosion, with maximum pit depths up to ~100 µm, in compacted and saturated bentonite. This is attributed to the difference in test media (inhomogeneity of the saturated bentonite versus the homogeneity of an aqueous solution such as the granitic groundwater): the rehydration of the compacted bentonite surrounding the specimens, can lead to the occurrence of local cells, which result in the
11 formation of zones with a different corrosion potential on the surface of the specimen, resulting in local attack. - the weld region and the heat affected zone (HAZ) were more susceptible to pitting corrosion than the parent material: pits up to 280 µm deep were detected at the weld/HAZ interface after an exposure period of 12 months at 120°C in bentonite buffered granitic groundwater. - the parent and weld material were resistant to SCC in bentonite buffered granitic groundwater at 90°C.
• stainless steel AISI 316L: - the general corrosion rate was found to be less than 0.1 µm/year. - the parent and weld material were resistant to any type of localised corrosion, including stress corrosion cracking, up to 100°C. There was no loss of ductility, but isolated pits could be observed near the fracture zone. The fracture of the welded specimens was always located in the base material. The fracture surface showed ductile fracture (formation of dimples). - the imposition of a γ-radiation field of 10 Gy/h had no effect on the mechanical properties, nor on the resistance to SCC.
Relevant data with respect to the general and localised corrosion behaviour of carbon steel BS4360 and AISI 304L and 316L type stainless steel in cementitious environments (UK) are also given, because large amounts of concrete will be present in repositories as a structural material:
• carbon steel BS4360: - the general corrosion rates in alkaline media at 30 and 80°C under oxic and anoxic conditions are given in Table 7. - hydrogen overpressure, in the range 100 kPa to 10 Mpa, had no discernible effect on the anaerobic corrosion rate in anoxic alkaline solutions. - the corrosion rate in anoxic solutions, increased with decreasing pH: the corrosion rates were about 5 (pH 7) to 15 (pH 4) times higher than the rates recorded in pH 13 solutions. - carbon steel was susceptible to pitting corrosion in cementitious porewater containing 6,000-28,000 mg/L Cl- at room temperature. The pitting susceptibility increased with increasing temperature: the critical chloride concentration decreased to <4,000 mg/L and <300 mg/L at 50 and 80°C respectively. - aeration period calculations and measurements of pit growth rates indicated that, for a fully saturated backfill, the rate of oxygen diffusion would be too slow to support corrosion for more than a few years, during which time pits may propagate only a few millimeters into the steel.
TABLE 7. General corrosion rates (µm/year) of carbon steel BS4360 in alkaline media.
T anoxic oxic (°C)
30 <0.12 0.08 80 <1 1.6
12 • stainless steels (AISI 316L): - the general corrosion rates in alkaline media at 30 and 80°C under oxic and anoxic conditions are given in Table 8. - AISI 316L type stainless steel was resistant to pitting corrosion in cementitious backfill (NRVB, Nirex Reference Vault Backfill) up to 100,000 mg/L Cl- at room temperature and up to ~50,000 mg/L Cl- at 45°C and 70°C. - a strong synergistic effect of chloride and thiosulphate on the susceptibility to 2- stress corrosion cracking was observed: the addition of 3,360 mg/L S2O3 to solutions containing 17,750 mg/L Cl- led to SCC at 80°C in the pH-range 7-13, while AISI 316L did not suffer from SCC in solutions containing these concentrations separately.
TABLE 8. General corrosion rates (µm/year) of stainless steel in alkaline media.
T anoxic oxic (°C)
30 0.001-0.1 0.03 80 0.001-0.1 0.5
From the available data, the following general conclusions can be drawn: the passively corroding Ti99.8-Pd is the primary choice for the thin-walled corrosion- resistant concept in rock salt formations, since its general corrosion rate is negligible low and it is highly resistant to localised corrosion and stress corrosion cracking in salt brines. The TStE 355 carbon steel is the prime candidate for the corrosion- allowance concept because it is resistant to pitting corrosion and SCC and its general corrosion rates are sufficiently low to provide corrosion allowance acceptable for thick-walled containers. Stainless steels, Ni-based alloys and Ti-based alloys are the most important candidate container materials in clay for the thin-walled concept, while carbon steel is considered the main choice for the thick-walled corrosion- allowance concept. Studies performed in granite indicated that copper containers provide an excellent corrosion barrier with an estimated lifetime exceeding 100,000 years. The TStE 355 carbon steel also seems to be a valid option for a thick- walled container in granite.
Various areas for which more in-depth R&D is required were identified and some are listed below (the level of scientific knowledge already available and priority differ according to the type of candidate rock formation):
• microbially influenced corrosion (MIC). To date, MIC in an underground repository environment has not been firmly identified from studies but the potential for bacterial activity in the oxic and/or anoxic phases of the disposal have been established. It will therefore be necessary to check the performance of candidate container materials for resistance to MIC through literature references, or potentially by experimental work. Already, some preliminary studies have started (and are still ongoing) to determine for example whether SRB can maintain their activity in compacted bentonite.
13 • atmospheric corrosion. This mode of attack applies to the period of interim storage and to the disposal period before resaturatioon of the backfill. • effect of fabrication aspects and container design on corrosion, such as welding, thermal treatment, metal surface condition, residual stresses, etc. • effect of γ-radiation. This is one of the most important parameters in convincing the public of the feasibility of disposing of radioactive waste in a safe manner. • effect of long-term metallurgical modifications, such as dealloying, segregation, creep, growth of grains, etc. • effect of nitric acid on the container integrity. Once the underground repository is sealed, the gaseous atmospheres within the containers may contain air and water vapour, which can, under irradiation, result in nitric acid production (HNO3). HNO3 could become concentrated at local sites, such as crevices, even if only small amounts of nitric acid are produced. • valuable information from archaeological analogues. Since it is difficult to simulate the time scales involved in the repositories during laboratory experiments, studying similar natural systems that have been exposed for a long period of time could provide useful information about corrosion kinetics and corrosion products and could contribute to an improved confidence in lifetime predictions. • modelling. In HLW/spent fuel disposal, the big challenge is to predict the corrosion behaviour of candidate container materials for periods well exceeding the normal service life of constructional materials (30-50 years). The prediction of this corrosion behaviour over a period of one thousand to hundreds of thousands of years, based only on relatively short-term experimental observations up to a maximum of 10 years (in situ experiments, immersion tests, etc.) and theoretical calculations, may be uncertain. The development of models to predict the lifetimes of nuclear waste containers is still incomplete. To be able to extrapolate corrosion phenomena up to thousands of years, a thorough understanding of the mechanisms and processes is necessary.
14 KEYWORDS
radioactive waste, disposal, underground repository, salt, clay, granite, cement, canister, container, overpack, carbon steel, stainless steel, nickel alloys, titanium alloys, copper, cast iron, corrosion, uniform corrosion, pitting corrosion, crevice corrosion, stress corrosion cracking (SCC), microbially influenced corrosion (MIC), electrochemical testing, immersion testing, modelling, review
15 GLOSSARY OF TERMS
The terms described in this Glossary have been drawn up with the objective to provide common definitions that can be understood to have the same meaning by all concerned and are therefore intended to increase the readability and comprehensibility of the state-of-the-art document. Some terms are, however, often defined in a different manner in various countries. The authors, therefore, advise the reader to consult the Glossary in order to understand the exact meaning of the terms in the context of this state-of-the-art document. The terms relating to radioactive waste management were defined based on the IAEA Radioactive Waste Management Glossary1. The terms relating to corrosion were defined based on the ISO 8044 European Standard2.
AISI. American Iron and Steel Institute numbering system for stainless steels. The AISI designates the wrought standard grades of stainless steels by three-digit numbers. The austenitic grades are designated by numbers in the 200 and 300 series. The 200 series numbers are used for high-manganese (>2%) austenitic stainless steels. The ferritic and martensitic grades are designated by numbers in the 400 series.
Aerobic. Used to describe a process taking place in the presence of oxygen.
Anaerobic. Used to describe a process taking place in the absence of oxygen (e.g. anaerobic corrosion proceeds when there is insufficient oxygen available at a metal surface to support corrosion. In this case, corrosion occurs by reduction of water to form hydrogen).
Anoxic. Used to describe an environment where there is a lack or complete absence of oxygen.
Argillaceous. Term applied to all rocks and substances composed of clay or having a notable proportion of clay in their composition.
Backfill. Material used to refill excavated portions of a repository (drifts, shafts, disposal rooms or boreholes) after the waste has been emplaced.
Barrier. Component of the repository system that reduces the rate at which radionuclides from the waste can be mobilised and reach the human environment.
1 IAEA, 1993. “Radioactive Waste Management Glossary,” IAEA (International Atomic Energy Agency, Vienna, Austria). 2 ISO 8044, 1998. “Corrosion of Metals and Alloys - Basic terms and definitions,” CEN (European Committee for Standardization, Brussels, Belgium).
16 Bentonite. A commercial term applied to numerous variously coloured (plastic) clay deposits essentially composed of the mineral montmorillonite (a dioctahedral smectite), formed by the weathering or hydrothermal alteration of volcanic ash or rock. Bentonites contain varying proportions of quartz, mica, feldspars, and other minerals. It is frequently specified as a backfill owing to its beneficial sorption and swelling characteristics.
Biosphere. Regions of the Earth able to support life. This includes the land surface, the oceans (hydrosphere), and the atmosphere. In repository performance assessments, the biosphere is generally taken to be those parts of the environment accessible to human beings.
Breakdown potential, ENP. See critical potential for pit nucleation.
Buffer. Material placed around a waste package in a repository to serve as an additional barrier to (i) stabilize (mechanically) the surrounding environment, (ii) restrict the access of groundwater tot the waste package, and (iii) provide chemical conditioning (e.g. reduce, by sorption, the rate of eventual radionuclide migration from the waste into the repository).
Canister. The primary closed or sealed metallic vessel in which the waste form is initially packaged for handling, storage, and transport to a repository. It may then be placed in a final disposal container, chosen for its long-term corrosion properties.
Clay. Minerals that are essentially hydrated aluminium silicates or occasionally hydrated magnesium silicates, with sodium, calcium, potassium, and magnesium cations. Also denotes a natural material with plastic properties, essentially composed of fine to very fine clay particles. Clays differ greatly mineralogically and chemically, and consequently in their physical properties.
Container. A secondary (or additional) metallic vessel, surrounding the canister, serving as the outer containment of the waste package intended for emplacement in a repository. See also overpack.
Corrosion. Physicochemical interaction between a metal and its environment that results in changes in the properties of the metal, and which may lead to significant impairment of the function of the metal, the environment, or the technical system, of which these form a part.
Corrosion potential, ECORR. The electrode potential of a corroding surface in an electrolyte, relative to a reference electrode. Also called rest potential, open circuit potential, or free corrosion potential.
Crevice corrosion. Localised attack of a metal surface associated with, and taking place in, or immediately around, a narrow aperture or clearance formed between the metal surface and another surface (metallic or non-metallic).
17 Critical potential for pit nucleation, ENP. The lowest value of the corrosion potential at which pit initiation is possible on a passive surface in a given aggressive environment. Also called breakdown potential.
Cyclic potentiodynamic polarisation (CPP) curve. A diagram relating electrode potential to current density for a specific electrode/electrolyte combination, in which the potential of the electrode is varied in a continuous manner at a preset rate. The potential is first scanned in the noble (positive) direction until a specified anodic current density or potential is reached (forward scan). The direction of the scan is then reversed and the potential is scanned in the active (negative) direction until the original starting potential is reached (reverse scan). The potential at which the scan is started is usually the corrosion potential.
Disposal. The emplacement of waste in an approved, specified facility (e.g. near surface or geological repository) without the intention of retrieval.
Electrode potential. The voltage measured in the external circuit between an electrode and a reference electrode in contact with the same specified electrolyte.
Far-field. Formations outside of the repository, including the surrounding strata, at a distance from the waste disposal site such that, for modelling purposes, the repository may be considered as a single entity, and the effects of individual waste packages are not distinguished.
Free corrosion potential. See corrosion potential, ECORR.
Fuel cycle (nuclear). Processes connected with nuclear power generation, including the mining and milling of fissile materials, enrichment, fabrication, utilisation, and storage of nuclear fuel, optional reprocessing of spent fuel, and processing and disposal of resulting wastes.
Geological disposal. Isolation of waste, using a system of engineered and natural barriers at depths up to several hundred metres in a geologically stable formation.
Geosphere. In general usage this is synonymous with the lithosphere; the outer layer of the earth. In repository performance assessment it is used to refer to the rocks between the repository and the biosphere.
Granite. Broadly applied, any holocrystalline quartz-bearing plutonic rock. The main components of granite are feldspar, quartz, and as a minor essential mineral, mica.
Grout. A fluid cement-aggregate-water mixture.
Half-life. The time required for the activity of a radioactive nuclide to decrease, by radioactive decay, to one half of its initial value.
18 Heat-affected zone (HAZ). That portion of the base metal that was not melted during brazing, cutting, or welding, but whose microstructure and mechanical properties were altered by the heat. High-level waste (HLW). The highly radioactive waste which is separated during the first solvent extraction cycle of spent fuel reprocessing and, in the UK, subsequently converted to a solid glass waste form by the process of vitrification. It has a high thermal power. It may take many thousands or millions of years for the radioactivity of HLW to decay to background levels. A high level of shielding and heat dissipation is required during storage and disposal.
Intergranular attack (IGA). Localised corrosion (dissolution) in or adjacent to the grain boundaries of a metal which otherwise exhibits corrosion resistance (e.g., austenitic stainless steels are attacked due to chromium depletion in the grain boundaries during heat treatment).
Leaching. All processes that, as a consequence of attacks by aqueous solutions, can lead to the dissolution of a radionuclide from the emplaced solid waste package, irrespective of whether the waste product is preserved or not.
Localised corrosion. Corrosion preferably concentrated on discrete sites of the metal surface exposed to an agressive environment.
Multibarrier concept. Repository concept whereby a number of barriers (typically, solid waste form, container, buffer, and surrounding rock) act in concert to contain the wastes and ensure that any radionuclides released from the waste return to the biosphere in concentrations which do not pose an environmental hazard.
Near-field. The repository and its contents, and the volume of immediately surrounding rock which is significantly affected by thermal or chemical interactions with the wastes and engineered barriers.
Open circuit potential (OCP). See corrosion potential, ECORR.
Overpack. A secondary (or additional) metallic vessel, surrounding the canister, serving as the outer containment of the waste package intended for emplacement in a repository. See also container.
Oxic. Used to describe an environment where there is sufficient oxygen present.
Passive. The state of a metal surface such that a stable protective film, usually an oxide film, is formed. pH. A measure of the acidity/alkalinity of a solution. pH values around 7 indicate chemical neutrality; higher values are alkaline; lower values are acidic.
Pitting corrosion. Localised attack of a metal surface resulting in pits, i.e. cavities extending from the surface into the metal.
Pourbaix diagram. A plot of the redox potential of a corroding system versus the pH of the system, showing the regions within which the metal itself or some of its
19 compounds are thermodynamically stable. Also called potential-pH (E-pH) diagram.
Protection potential, EPP. Threshold value of the corrosion potential that has to be reached to enter a range of corrosion potential values in which an acceptable corrosion resistance is achieved for a particular purpose. Also called repassivation potential.
Reference electrode. An electrode, having a stable and reproducible potential that is used as a reference in the measurement of electrode potentials.
Repassivation potential, EPP. See protection potential.
Repository. A nuclear facility for waste disposal located underground (usually more than several hundred metres below surface) in a stable geological formation to provide long-term isolation of radionuclides from the biosphere.
Rest potential. See corrosion potential, ECORR.
Salt. A geological formation resulting from the evaporation of seawater containing mainly halite (NaCl) with smaller inclusions or layers of other minerals, usually the chloride or sulphate derivates of the alkali or alkali earth elements and clay. Salt formations occur either as bedded or domal deposits.
Spent fuel. Uranium or uranium oxide fuel elements that have been removed from a nuclear reactor at the end of their useful lives. Spent fuel builds up a substantial content of fission products and actinide elements as it is 'burned' in the reactor. These progressively reduce the efficiency of the fuel, until it has to be replaced.
Storage. The placement of waste in a nuclear facility where isolation, environmental protection, and human control (e.g. monitoring) are provided with the intent that the waste will be retrieved for processing and/or disposal at a later time.
Stress corrosion cracking (SCC). Cracking of a metal caused by the simultaneous action of corrosion and sustained straining of the metal (due to applied or residual stress).
Strip mining. A technique used for the extraction of coal, which involves removing up to 60 metres of covering soil.
Tarnishing. Dulling, staining, or discoloration of a metal surface, due to the formation of a thin layer of corrosion products.
Thermal phase. The period during which most of the fission product activity decays before the natural waters permeating the disposal medium contact the waste and activity leaching commences, i.e. the period necessary to allow the waste to cool before contacting the water, thus minimising the rate of leaching and, hence, the rate of release of the longer half-life actinide species.
20 Transmutation (nuclear). The process in which the longer lived radionuclides are converted to shorter lived or stable radionuclides by irradiating the long-lived wastes. Tuff. A rock composed of compacted volcanic ash. It is usually porous and often relatively soft.
Uniform corrosion. Corrosion proceeding at almost the same rate over the entire surface of the metal exposed to an aggressive environment.
Vitrification. Process in which liquid HLW is calcined to dryness, mixed with inert chemicals, melted, and cooled to form a stable, solid glassy waste product.
Waste, radioactive. Any material that contains, or is contaminated with, radionuclides at concentrations greater than clearance levels (i.e. levels greater than those deemed safe by national authorities), and for which no use is foreseen (this definition is purely for regulatory purposes, because from a physical viewpoint any material with concentrations equal to or less than clearance levels is still considered radioactive, although the associated radiological hazards are considered negligible).
Wasteform. The waste in its physical and chemical form after treatment and/or conditioning (resulting in a solid product) prior to packaging. The waste form is a component of the waste package.
Waste package. The product of conditioning that includes the wasteform and any container(s) and internal barriers (e.g. absorbing materials and liner), as prepared in accordance with requirements for handling, transportation, storage, and/or disposal.
Welding. Joining pieces of metal together by heating their edges and pressing them together. The corrosion resistance of weldments is often known to be inferior to that of their base metal. The following table summarises the welding procedures investigated as possible container closure techniques.
Welding procedure Abbreviation Description Id. Nr.(1)
EBW Electron Beam Welding 511 GTAW Gas Tungsten Arc Welding 141 MAGW Metal Active Gas Welding 135 FCAW Flux Cored Arc Welding 136 SAW Submerged Arc Welding 121 PAW Plasma Arc Welding 15
(1) Identification number of the welding procedure according to EN ISO 4063 (2000) “Welding and allied processes. Nomenclature of processes and Reference Numbers”.
21 LIST OF ABBREVIATIONS AND ACRONYMS
AECL Atomic Energy of Canada Limited AISI American Iron and Steel Institute ANDRA Agence Nationale pour la gestion des Déchets RadioActifs (French national agency for the management of radioactive waste) ANRC Amarillo National Research Center as as received ASTM American Society for Testing and Materials BAMBUS Backfilling and Sealing of Underground Repositories for Radioactive Waste in Salt BWR Boiling Water Reactor CEA Commissariat à l’Energie Atomique cc crevice corrosion CEBELCOR CEntre BELge d'étude de la CORrosion CERT Constant Extension Rate Test CNNC China National Nuclear Corporation COGEMA Compagnie Générale des Matières Nucléaires CPP Cyclic Potentiodynamic Polarisation measurements CSH Calcium Silicate Hydrate Cu-OF Oxygen-Free copper (designation used in the Spanish contai- ner corrosion studies) Cu-OFP Oxygen-Free copper containing Phosphorus (designation used in the Swedish and Finnish waste management programmes for oxygen-free copper micro-alloyed with 30-80 ppm phospho-rus) DBE Deutsche gesellschaft zum bau und Betrieb von Endlagen für Abfallstoffe GmbH (German Company for the construction and operation of ultimate storage facilities for nuclear waste) DEBORA DEvelopment of BOrehole seals for high-level RAdioactive waste DEFRA Department for Environment, Food & Rural Affairs (UK) dIGA,Max maximum depth of intergranular attack dpit,Max maximum depth of pitting attack EBW Electron Beam Welding technique ECORR corrosion potential ENP critical potential for pit nucleation EPP protection potential EG/BS Extension Gallery Bottom Shaft ENRESA Empresa Nacional de Residuos Radioactivos (Spanish national radioactive waste management company) EUROBITUM bituminised waste product conditioned at the former EUROCHIMIC facility (which is currently owned and operated by BELGOPROCESS) FCAW Flux Cored Arc Welding technique FoCa Fourges Cahaignes FoCa-clay a sedimentary clay from the Paris Basin FZK Forschungszentrum Karlsruhe
22 GNB Gesellschaft für Nuklear-Behälter GNF Gesellschaft zur Förderung der Naturwissenschaftlich- technischen Forschung in Berlin-Adlershof e.V. GTAW Gas Tungsten Arc Welding technique (also TIG, Tungsten Inert Gas) HADES High Activity Experimental Disposal Site HAZ Heat-Affected Zone HLW High-Level radioactive Waste HT/HP High Temperature / High Pressure IAEA International Atomic Energy Agency ILW Intermediate-Level radioactive Waste INSC International Nuclear Societies Council IUT Institut für Umwelttechnologien GmbH JAERI Japan Atomic Energy Research Institute JNC Japan Nuclear Cycle Development Institute LLNL Lawrence Livermore National Laboratory MAGW Metal Active Gas Welding technique MIC Microbially Influenced Corrosion MLW Medium-Level radioactive Waste MOX Mixed Oxide fuel NAGRA Nationale Genossenschaft für die Lagerung radioaktiver Abfälle (Swiss National Cooperative for the Disposal of Radioactive Waste) NEA Nuclear Energy Agency NEI Nuclear Energy Institute NERAC Nuclear Energy Research Advisory Committee NIRAS/ONDRAF Nationale Instelling voor Radioactief Afval en Verrijkte Splijtstoffen / Organisme National des Déchets RAdioactifs et des matières Fissiles enrichies (Belgian Agency for Radioactive Waste and Enriched Fissile Materials) NIREX Nuclear Industry Radioactive Waste Executive (the UK Radioactive Waste Management company) NRVB Nirex Reference Vault Backfill NUMO Nuclear Waste Management Organisation of Japan OCP Open Circuit Potential OCRWM Office of Civilian Radioactive Waste Management (programme of U.S. Department of Energy, DoE) OFP Cu Oxygen-Free copper containing P OM Optical Microscopy OPC Ordinary Portland Cement OPG Ontario Power Generation (one of Canada’s electricity produ- cers) OPHELIE On Surface Preliminary Heating simulation Experimenting Later Instruments and Equipments p polished PAMELA Pilotanlage Mol zur Erzeugung Lagerfähiger Abfälle PAW Plasma Arc Welding technique PNC Power Reactor and Nuclear fuel development Corporation (Japan) POLLUX a self-shielding gas-tight metal cask system developed in Germany for the final disposal of spent fuel
23 Posiva Finnish national Waste Management Agency PRACLAY PReliminAry demonstration test for CLAY disposal PTFE PolyTetraFluoroEthylene (Teflon®) PURAM Public Agency for Radioactive Waste Management (Hungary) PVDF polyvinylidenedifluoride PWR Pressurised Water Reactor RCW Real interstitial Clay Water SAW Submerged Arc Welding technique SCC Stress Corrosion Cracking SBW Synthetic Bentonite pore-Water SCK•CEN Studiecentrum voor Kernergie / Centre d'étude de l'Energie Nucléaire (Belgian nuclear research center) SCW Synthetic interstitial Clay Water SEM/EDS Scanning Electron Microscopy / Energy Dispersive Spectros- copy SF Spent Fuel SHE Standard Hydrogen Electrode SKB Svensk Kärnbränslehantering AB (Swedish Nuclear Fuel & Waste Management Company) SOC Synthetic Oxidised Clay Water SRB Sulphate Reducing Bacteria SSRT Slow Strain Rate Test SYNATOM the Belgian Society for Nuclear Fuels (a subsidiary from TRACTEBEL and ELECTRABEL, in charge of the supply of nuclear fuels and the management of irradiated nuclear fuels from the Belgian nuclear power plants) TIG Tungsten Inert Gas welding technique (also GTAW, Gas Tungsten Arc Welding) TVO Teollisuuden Voima Oy (Finnish nuclear power company) UOX Uranium Oxide fuel US DOE United States Department of Energy vCORR average uniform corrosion rate V/S brine volume to specimen surface area ratio VUB Vrije Universiteit Brussel XRD X-Ray Diffraction
24
PART A
REVIEW OF CONTAINER CORROSION INVESTIGATIONS DURING THE LAST 15-20 YEARS
A.1 INTRODUCTION
Currently, nuclear power plants provide approximately one fifth (16%) of the world's total consumption of electric power (in some countries, the contribution of nuclear power amounts to well over 50%, e.g. France, Belgium) [1,2]. A more detailed overview of the contribution (%) of nuclear power plants to the world's electricity production is summarized in Appendix A [3].
The intensive industrialisation worldwide requires an ever-increasing production of energy [4]. Nuclear energy presents a number of unique features in comparison with other energy sources (e.g. coal, oil, gas, wind, solar, hydroelectric, biomass, etc.), which affect its attractiveness as one of the many methods of supplying an energy- demanding world [5,6]: • nuclear energy avoids the wide variety of detrimental environmental problems arising from burning fossil fuels – coal, oil, and gas: nuclear power plants do not emit harmful gasses or particulate combustion products into the atmosphere causing global warming, acid rain, and air pollution [7-9]. • nuclear energy guarantees an abundant supply of fuel to satisfy the world's energy needs for may thousands of years to come without affecting the earth's natural resources sorely needed for other applications (fossil fuels are also the feedstock for producing plastics and organic chemicals, while uranium has little value for other purposes). The world's supply of coal, oil, and gas, on the other hand, is limited, to probably less than 100 years of energy consumption [7]. • a relatively small amount of the earth's natural resources is required to generate a large amount of electricity compared to other energy sources: one uranium nuclear fuel pellet – which has the size of the tip of your little finger – is equivalent to the energy provided by ~810 kg of coal, ~560 l of oil, or ~480 m³ of natural gas [10]. • the exploitation of uranium ore reduces strip mining, used for the extraction of coal, which adversely affects the surrounding vegetation [7,11]. • the space needed to accommodate a nuclear power plant is considerably smaller than that needed for other electrical plants with the same output of energy [12]. • nuclear energy reduces the dependency on external resources (e.g. foreign oil). This is an important factor for countries with limited natural resources [4].
Nuclear power, which was developed for the benefit of society, is also unfortunately leaving behind radioactive residues, which are an inevitable by-product [4,13]. Several radioactive species are produced during the nuclear fuel cycle. The highest content (>99%) of the radiotoxic and heat-generating species is concentrated within the spent fuel discharged from reactors (i.e. uranium fuel that has been used in a nuclear reactor and is no longer efficient in generating power [14]) and hence the majority of the effort has been directed towards the development of technologies for the safe disposal of heat-generating and long-lived nuclear waste. There are two principal options for the management of this type of nuclear waste [15]: • the spent fuel can either be reprocessed to recover plutonium and uranium for recycling. The resulting liquid high-level radioactive waste (HLW) can be conditioned (e.g. vitrified) and eventually disposed of, or • the spent fuel (SF) can be disposed of in a similar manner to vitrified HLW disposal without further treatment (direct disposal).
26 Radioactive waste differs from other industrial waste in two ways. First of all, the risk associated with radioactive waste decreases with time [16]: the radioactive waste decays (or disintegrates) exponentially with time and becomes non-radioactive [14]. The rate at which this decay occurs depends on the half-life of the radioactive isotopes [17]. Industrial wastes, for example from coal burning, include chemical carcinogens like beryllium, cadmium, arsenic, nickel, and chromium, which, unlike the nuclear wastes, last forever and, in this respect, can be considered far more dangerous [7]. Second, the overall volume of radioactive waste is much less than that of other chemical and industrial wastes. A 1000 MW coal power plant discharges approximately 300 tonnes of SO2 and 5 tonnes of fly ash every 24 hours directly to the atmosphere. A 1000 MW nuclear power plant produces ~500 m³ of waste per year, with an average density of 160 to 240 kg/m³ (giving a total weight of 80-120 tonnes), none of which is released into the atmosphere [16].
The suitable management of HLW/spent fuel from nuclear reactors is a major technological challenge within the various European Union Member States operating nuclear reactors, because it has an impact on the safety and public acceptance of nuclear energy and the further development of this technology [4]. The primary objective of safe disposal of HLW/spent fuel is to adequately protect human health and the environment from potential hazards arising from the ionising radiation emitted by the waste now and in the future, without imposing undue burdens on future generations. This involves the emplacement of the radioactive waste into a disposal facility that is intended to isolate the waste from man's environment until radioactive decay has reduced its toxicity to negligible non-hazardous levels and to prevent or limit releases of potentially harmful substances (toxic metals, radionuclides, organics) [15].
Deep geological disposal is the reference solution for high-level and long-lived waste management
The safe disposal of HLW/spent fuel represents a major technological challenge for all nations that operate nuclear reactors, and is a burning issue that has been studied for many years already [18].
Over the years, several possibilities for the disposal of heat-generating long-lived waste have been considered. At present, deep underground disposal in stable geological formations (rock salt, clay, granite, tuff, etc.) is one of the means that is being investigated extensively and, internationally, to dispose of nuclear waste in a feasible and safe manner.
In some countries, the deep geological disposal concept is also being considered as a possible option for disposing of ILW.
The multibarrier concept forms the backbone of deep geological disposal [19]
The strategy of deep geological disposal relies on a system of natural and engineered (man-made) barriers, which act in concert to contain the wastes and ensure that any radionuclides released from the waste return to the biosphere in
27 concentrations which do not pose an environmental hazard [20]. This joint action of several natural and engineered barriers is known as the multibarrier concept, which is schematically represented in FIGURE 1-1 [21]. The multibarrier concept is based on the principle that safety is not endangered even if some of the barriers fail, because in this case, the remaining undamaged barriers will be sufficient to guarantee adequate protection of the environment and people [22,23].
FIGURE 1-1. Schematic view of the multibarrier concept. Source: adapted from [21].
The various components of the multibarrier concept generally include [4,19,24-26]: • the engineered repository system (manmade barriers) - the waste package (waste, solidification material, canister, and possibly overpack), - the backfill/buffer material, and - the repository structures (e.g. concrete lining). • natural barriers - the geosphere (i.e. the geological host rock formation and the underlying- and overlying sediments).
The near-field environment of a geological repository for radioactive waste disposal comprises the excavated area near or in contact with the waste package (including backfill, filling and sealing materials) and those parts of the host rock formation whose characteristics have been or could be modified by the construction of the repository (e.g. excavation works) and the exploitation of the repository (e.g. heat emission of the radioactive waste) [27]. The far-field environment is that part of the host rock formation, at a sufficient distance from the repository, that has not been altered.
The waste package consists of the wasteform, the canister, and/or an additional container. The heat-generating long-lived radioactive waste is either encapsulated as it is produced (e.g. as spent fuel) or, after reprocessing, incorporated into a leach- resistant matrix – for HLW, glass is the favoured solidification material – before emplacement in metallic drums (the canister). The metallic canister may also be surrounded by an outer metallic vessel (the overpack). The waste container comprises the canister and possibly an overpack. The metallic waste containers are emplaced in the repository, which will be constructed several hundreds of metres
28 deep in stable geological formations (e.g. rock salt, clay, granite, tuff), where they will be embedded in a backfill material (e.g. bentonite clay, crushed salt, cement, etc. depending on the chosen option), which will fill the space between the containers and the surrounding rock [20].
The inherent insolubility of the waste form, the metallic container, and the backfill material will limit the ingress of water to the waste thereby restricting the rate at which radionuclides can pass into solution in any groundwater. Furthermore, by choosing a geological host formation in which water movement is extremely slow and where the constituent minerals readily absorb or bind with the radionuclides, it will be possible to delay the emergence of the more soluble radionuclides for a very long time [28].
A more detailed description of the different disposal concepts envisaged in various European countries operating nuclear power plants (e.g. Belgium, Finland, France, Germany, Spain, Sweden, UK) can be found in Appendix B.
The role of the metallic container as one of the principal engineered barriers
The primary role of the metallic container in the multibarrier system is to separate the wasteform from the disposal environment (geosphere), as long as possible, by acting as a barrier against the mobilisation of radionuclides. This is achieved in two ways, by the container providing a physical and a chemical barrier [29].
As a physical barrier, the disposal container is designed to hermetically enclose the radioactive waste at least until the temperature in the near-field has dropped to near the background level (i.e. the thermal phase, ~500–1,000 years). This is to prevent the repository environment from coming into contact with the waste during the period when the activity is highest and consequently its possible impact on the biosphere is at its most severe [29]. For disposal in rock salt, there are some additional reasons why the container must remain intact during the thermal phase: (i) the temperature of the glass will drop to below 100°C, thereby reducing leaching of radionuclides from the glass, and (ii) due to the convergence of the host rock to the container, intrusion of brines towards the container can be avoided over this timescale.
The metallic container must provide a particularly high corrosion resistance to the aggressive underground environment. The metallic container must also be sufficiently mechanically stable to withstand the external pressures (rock pressure, hydrostatic pressure) [29].
Additionally, the physical function of the disposal container is more amenable to public understanding than the concept of e.g. sorption as a mechanism for containing the radionuclides over a sufficiently long period of time [30].
In some cases, e.g. for carbon steel, under certain conditions where it actively corrodes, the metallic container can function as a chemical barrier by establishing a reducing geochemical environment, with low radionuclide solubilities and the possibility that radionuclides, initially mobilised, may be adsorbed on corrosion products of iron. Such effects begin to play a role when the container has been completely corroded and the HLW/spent fuel starts to react with the repository environment [29].
29 In the framework of this study, however, only the physical aspects of the engineered barrier and not the chemical aspects are being considered.
The metallic container acts as the primary barrier by providing an 'absolute and complete containment' of the radioactive substances [22,31]. The container will isolate the wasteform (fuel, HLW) from contact with groundwater/salt brine. If failure of the container occurs, following corrosion of the container, the waste matrix, if present, will provide the second containment measure: when it comes into contact with groundwater/salt brine, the release of radionuclides to the environment will be restricted [32]. If the waste is encapsulated, the radionuclides will be immobilised in a highly insoluble, corrosion resistant form such as glass, thereby retarding the leaching of radionuclides from the waste for tens or hundreds of thousands of years [31].
The gaps between the waste containers and the walls of the disposal facilities (e.g. boreholes or galleries), into which the waste containers will be placed, will be filled with a backfill material (in many countries, precompacted blocks of bentonite clay are the preferred option) [31]. The backfill should fulfil the following functions [4,22,31,33]: • restrict groundwater flow around the container, • provide mechanical and structural support to the container, thereby protecting it against minor movements in the rock, • retard the release of radioactive materials in the case of a failing container, and • prevent radionuclide migration through sorption of radionuclides. An adequate choice of the backfill material could also contribute to an enhanced transfer of heat from the waste to the surrounding rock [24].
The geosphere acts as a barrier in order to protect [4] • the artificial barriers (the waste form, the container, the backfill material) from natural disruptions and human intrusions, and • the biosphere by - maintaining a physically and chemically stable environment for the artificial barriers favourable for long-term waste isolation, and - retarding and restricting radionuclide transport from the repository to the biosphere (e.g. retention, dilution before reaching the biosphere).
The main potential pathways by which radionuclides could return from a closed waste repository to the bioshere are [26]: • transport in groundwater, • migration in gases evolved from materials in the repository, and • human intrusion into the repository.
Metallic materials are the favoured choice for the fabrication of the container
In general, the candidate container materials should possess the following properties [34]: • corrosion resistance with respect to the disposal environment to last for 500–1000 years before penetration, • long-term physical stability including resistance to radiation damage,
30 • good fabricability, including ease of sealing either by welding or mechanical means, • mechanical strength and toughness for handling, transportation, emplacement, and service (ability to withstand the pressures imposed by the underground host rock formation), • ready availability, and • cost effectiveness.
Over the years, several types of materials – such as metals, ceramics, polymers, and concrete – have been envisaged as candidate container materials. However, considering the current technological knowledge in materials science, only metallic materials can potentially fulfil all of the above-mentioned requirements [34]. Also, because the container must act both as a transportation vessel and as a physical barrier against the corrosive action of the repository environment, almost all of the materials being considered for its construction are metallic [35].
The main disadvantages of ceramic materials is their poor impact strength, which means that there is a significant risk that containers fabricated from such materials could be damaged during handling and emplacement operations, the sealing technique, thermal shock, creep, etc. A significant limitation for the use of polymeric materials is their relatively poor resistance to the combined effects of temperature, an aqueous environment, and radiation damage. Additionally, they also possess poor long-term aging characteristics. A major factor affecting the performance of concrete – and possibly reinforced concrete – is its long-term durability, which mainly depends on the interactions between the concrete and the natural waters permeating the disposal site, and on the corrosion of the metal embedded in it. Both cases lead to the formation of reaction (corrosion) products which are likely to have a greater volume than the reactants, thereby causing the concrete to crack. Also, the reaction products, resulting from the interactions between the concrete and the groundwater, may be leached out, thus weakening the concrete and increasing its porosity [34].
Metallic disposal containers: corrosion-allowance vs. corrosion-resistant materials
Evaluating the suitability of candidate container materials for the disposal of HLW/SF in deep geological formations (rock salt, clay, granite, tuff) and predicting the corrosion performance of waste containers over long periods of time have presented a worldwide challenge from a materials, engineering, and corrosion point of view for several decades.
The ideal container would be one fabricated from a metal that is inert in all foreseen environments (i.e. a metal which is thermodynamically stable in the disposal environment). In this respect platinum is an ideal candidate. However, from an economical point of view, the use of this metal is not feasible [34].
Over the years, two basic approaches have been considered to potentially satisfy the lifetime requirement imposed on the waste container, viz. the corrosion-resistant and the corrosion-allowance concept [34–38]:
31 • in the corrosion-resistant concept relatively thin and highly corrosion-resistant materials are relied upon to provide the required container lifetime. Corrosion- resistant materials are alloys (e.g. stainless steels, nickel-base alloys, titanium alloys) that exhibit a high resistance to corrosion attack in the expected disposal environments (i.e. the amount of corrosion is sufficiently low) [34]. These alloys usually owe their durability to the formation of protective adherent oxide films (passive surface layers) [35]. Such materials degrade at a very low corrosion rate and can therefore be used in comparatively small thicknesses [34,36]. With the use of corrosion-resistant materials, the risk of breakdown of the passive film, leading to localised corrosion phenomena (e.g. pitting corrosion, crevice corrosion) has to be taken into account, because this can cause sudden large and unpredictable increases in the rate of corrosion [37].
• in the corrosion-allowance concept, materials (e.g. carbon steel, copper) which corrode at a significant but relatively predictable corrosion rate are used [38]. These materials can be used if sufficiently thickness is foreseen to allow for the depth of corrosion expected during the desired lifetime [36].
In both concepts, the containers (overpacks) surrounding the HLW canisters must possess a wall thickness sufficient to withstand the hydraulic and geological pressure [38].
TABLE 1-1 represents a summary of the different repository concepts (host rock formation, maximum temperature criterion, candidate container material, wall thickness container, etc.) envisaged in the various countries operating nuclear power plants [39].
32 TABLE 1-1. Summary of the different repository concepts (host rock formation, maximum temperature criterion, candidate container material, wall thickness container, etc.) envisaged in various countries operating nuclear power plants, updated from Ref. [39].
(1) (2) Country Authority Host rock Tmax Candidate container material D (°C) (mm)
Belgium NIRAS/ONDRAF clay <100 stainless steel (AISI 316L hMo) n.d. Canada OPG granite 100 titanium / copper 6.35 / 25 China CNNC granite n.d. - - Finland posiva granite <100 cast iron (inner container) 60 copper (outer container) 50 France ANDRA clay, granite <100 carbon steel 55 (alternative: passive alloys) - Germany DBE salt 200 HLW: Cr-Ni steel 4-5 spent fuel: carbon steel 425 carbon steel plated with Ni- 3-4 alloy (Hast. C-4) / Ti 99.8-Pd Japan NUMO granite 100 carbon steel 300 (alternative: titanium) - Spain ENRESA clay, granite 100 carbon steel 100 Sweden SKB granite <100 cast iron (inner container) 50 copper (outer container) 50 Switzerland NAGRA granite n.d. carbon steel 250 UK(3) NIREX n.d. 100 stainless steel - USA US DOE tuff 250 stainless steel (AISI 316 or 316L) (inner 50 container) Ni- alloy (alloy C-22) (outer container) 10 Ti-grade 7 (drip shield) 20
(1) organisation responsible for radioactive waste management. (2) D : wall thickness of the container. (3) this refers to disposal of ILW. There is currently no UK concept for disposal of HLW. n.d. : not yet determined.
33 A.2 HOST ROCK FORMATIONS FOR HLW/SPENT FUEL DISPOSAL
A.2.1 Disposal in Salt
A.2.1.1 Characterisation of the host rock formation considered in Germany
Rock salt is an elasto-viscoplastic material that creeps under the influence of high stresses without significant fracturing. Because of this material behaviour it is capable of encapsulating any waste type. At the higher temperatures produced by heat generating high-level waste, the creep capability of rock salt increases and enables it to deform quickly around the waste containers and hence to isolation of the waste from the biosphere [40].
The rock salt formation considered in Germany is very dry (H2O-content < 0.01 wt.%). However, it is possible that under some conditions salt brines can be present in the disposal area. Such brines may originate from the thermal migration of brine inclusions in the rock salt or by accidents in the repository, e.g. brine inflow through an anhydrite layer [41-43].
To examine the influence of the brine composition on the corrosion behaviour of the candidate materials for the waste containers, three salt brines – all of which are representative for the German rock salt disposal environment – are used. The compositions and measured pH values of the three test brines are given in TABLE 2-1. Two of them (brines 1 and 2) are highly concentrated with MgCl2; the third one (brine 3) has a high concentration of NaCl [41].
TABLE 2-1. Composition and pH value of the salt brines, representative for the Gorleben rock salt repository environment, used in the corrosion experiments to support the German disposal concept [41].
Composition (wt. %) Brine pH (25°C) NaCl KCl MgCl2 MgSO4 CaCl2 CaSO4 K2SO4 H2O
1(a) 1.4 4.7 26.8 1.4 - - - 65.7 4.6
2 0.31 0.11 33.03 - 2.25 0.005 - 64.3 4.1
3 25.9 - - 0.16 - 0.21 0.23 73.5 6.5
(a) Q-brine.
34 A.2.2 Disposal in Clay
A.2.2.1 Characterisation of the argillaceous host rock formation and bentonitic backfill material considered in Belgium
In Belgium, the Boom clay layer is being considered as the host rock formation for final disposal of HLW/spent fuel. However, a final choice concerning the material to be used to backfill the voids between the metallic container and the disposal gallery walls has not been made. Two candidate backfill materials are envisaged: (i) Boom clay, recuperated during the excavation of the disposal galleries (after several thermal and mechanical treatments), or (ii) a mixture of 60% FoCa bentonite clay + 35% sand + 5% graphite (from hereon referred to as ‘PRACLAY’3 bentonite). Because of this unclarity, corrosion experiments have been performed in media simulating both environments.
The mineralogical composition of Boom clay (host rock formation) and FoCa clay (candidate backfill material) is summarised in TABLE 2-2 and TABLE 2-3, respectively.
TABLE 2-2. Mineralogical composition of Boom clay [44].
Weight% dry matter Remarks
Clay minerals 60 Illite 20 – 30 (% of clay minerals) Smectite 10 – 20 " Chlorite 5 – 20 " Kaolinite 20 – 30 " Illite/smectite mixed layers 5 – 10 " Chlorite/smectite mixed layers 5 – 10 " Quartz 20 Feldspars-K 5 – 10 Carbonates 1 – 5 Calcite present Siderite present Dolomite present Ankerite present Pyrite 1 – 5 Organic matter 1 – 5
Others: glauconite, calcium phosphate, rutile, anatase, ilmenite, zircon, monazite, xenotime.
3 the bentonite mixture of 60% FoCa clay + 35% sand + 5% graphite was also used in the OPHELIE mock-up experiment in the framework of the PRACLAY experiment. Bentonite is a generic term used to denote a specific group of several minerals. FoCa is a sedimentary clay from the Paris Basin, extracted in a site near to the Vexin region (Fourges-Cahaignes). The chemistry of the PRACLAY bentonite porewater was determined by mechanical squeezing after rehydrating the crushed bentonite mixture with deionised water. The mixture was allowed to equilibrate with the added water for 2 weeks. All manipulations were carried out in a low oxygen-containing (<100 mg/l) chamber.
35 The Boom Clay is a marine sediment of tertiary, Rupelian age, which is located in the northeast of Belgium and extending under the Mol-Dessel nuclear site at a depth between 180 and 280 meter. Its marine origin is responsible for a relative high organic matter content. It is a plastic ‘carbonate type clay’ with a CO2 partial pressure in the interstitial porewater of 10-2.42 atm. Under in situ conditions, Boom Clay is slightly alkaline, with a pH in the range 8.2-9.7, due to the presence of carbonate (calcite and siderite). It also exhibits moderately reducing properties, with Eh values ranging from about -250 to -350 mV (vs. Standard Hydrogen Electrode), due to the presence of pyrite, Fe2+, and humic substances [45-47].
TABLE 2-3. Mineralogical composition of FoCa clay [44].
Weight% dry matter Remarks
Clay minerals 80 – 88 (interstratified kaolinite/smectite) Kaolinite(1) 50 – 60 (% of clay minerals) Smectite (Dioctaedrical) 40 – 50 " Montmorillonite FeMg 30 (% of smectite) Beidellite FeMg 40 " Beidellite Fe 30 " Quartz 5 – 7 Carbonates – calcite 1.2 – 1.6 Goethite 5 – 7 Hematite 0.2 Gypsum 0.4 Organic matter 0.1
(1) In some batches, up to 4 wt.% of free kaolinite is present.
The main physico-chemical and thermo-hydro-mechanical properties of Boom clay and FoCa clay are summarized in TABLE 2-4. The mean chemical composition of EG/BS4 in situ interstitial Boom clay porewater and PRACLAY bentonite porewater are presented in TABLE 2-5.
4 the chemistry of the Boom clay porewater was determined from water collected by the Experimental Gallery Bottom Shaft (EG/BS) piezometer. This piezometer is installed at a depth of 260 m and is in contact with the Silty Layer of the Boom clay formation. This porewater represents the chemical composition of the underground Boom clay formation water under anaerobic and reducing conditions.
36
TABLE 2-4. Main physico-chemical and thermo-hydro-mechanical properties of Boom clay and FoCa clay [44,48-52].
Parameter Unit Boom clay FoCa clay
Particle size distribution % d < 2µm ('clay') 50 – 60 90 2µm < d < 60µm ('silt') 40 – 45 60µm < d < 100µm ('sand') 3 – 8
Physico-chemical characteristics
Water content % dry wt. 19 – 24 10 – 12 Average grain density g.cm-3 2.65 2.67 (2) Porosity % 36 – 40 29.6 Liquid limit w [%] 55 – 80 112 L Plastic limit w [%] 23 – 29 50 P (3) (4) CEC(1) 227 710 µeq.g-1 Exchangeable cations -1 Ca2+ µeq.g - 731 Mg2+ 108 65 Na+ 83 36 K+ 36 8
Thermo-hydro-mechanical characteristics
Thermal conductivity W.m-1.K-1 1.68 0.95 Specific heat J.kg-1.K-1 1400 1212(6) -1 -14 Hydraulic conductivity m.s 2.1 – 4.5×10-12 (5) 5.10 Swelling pressure MPa 0.9 13
(1) CEC = Cation Exchange Capacity. (2) calculated value for the bentonite mixture (60% FoCa clay + 35% sand + 5% graphite) at saturation (the backfill material exerted a swelling pressure of approximately 2MPa upon complete saturation) [44]. (3) the global CEC determined as the sum of the major cations CEC. -1 -1 AgTU-CEC = 300 µeq.g ; Ca-CEC = 240µeq.g present in CaCO3. (4) the CEC is measured by the sorption of cobaltihexammine and Cu-ethylenediamine. The global CEC is lower than the sum of the major cations CEC due to the presence of secondary mineral phases like calcite and gypsum. These phases are (partially) dissolved and measured with the exchangeable cations, increasing the measured CEC value [51]. (5) data from in situ field testing: vert. 2.1×10-12, horiz. 4.5×10-12 m.s-1. (6) m -4 specific heat at 25°C calculated from the polynomial Cp = 1.18794 + 8.58416.10 T + 4.79461.10-6T² [52]
37
TABLE 2-5. Mean chemical composition of EG/BS in situ interstitial Boom clay porewater and bentonite porewater [46,53].
Constituent EG/BS Boom clay Bentonite [mg.L-1] porewater porewater(1)
Anions Cl- 27.0 68.0 F- 3.6 1.56 Br- 0.49 n.d. 2- SO4 0.2 1633.0 - NO3 - <2.50 - HCO3 828.0 84.0 2- HPO4 3.8 <0.25 2- S2O3 - <2.50
Cations Ca 4.0 583.0 Mg 2.9 37.5 Na 408.0 183.0 K 11.0 8.40 Fe 0.9 <0.10 Al 0.08 0.45 Mn n.d. 0.145 B 7.5 n.d. Si 5.0 n.d.
TOC(2) 41.3 – 144.0 106.0 TIC(3) n.d. 6.82
(1) data of the deionised water/bentonite mixture after the first squeezing. (2) Total Organic Carbon content [mg/L C]. (3) Total Inorganic Carbon content [mg/L C]. n.d. : not determined.
A.2.2.2 Characterisation of the host rock formation considered in France
In France, the Callovo-Oxfordian clay formation is being studied as a possible host rock formation for the final disposal of HLW/spent fuel. The main characteristics of the Callovo-Oxfordian clay formation are summarised in TABLE 2-6.
The geochemical characterisation (pH, EH, chemical composition) of the various groundwaters considered as repository sites in France is summarised in TABLE 2-7.
38
TABLE 2-6. Summary of the main characteristics of the Callovo-Oxfordian Formation (current clay host rock formation considered in the French disposal concept) [54].
Host medium Description (clay formation)
General • Stratigraphic unit: middle Callovian to lower Oxfordian. (geological characteristics) • Thickness: 135 metres. • Burial: from 417 to 552 metres. • Mineralogy: clay minerals (illite/smectite ML: "R1" and "R0", illite, chlorite, kaolinite) = 40 to 45%; quartz ≈ 25%; carbonates ≈ 30%; other minerals (pyrite, feldspar) and organic matter ≤ 1%.
Hydraulic characteristics • Permeability: from 10-11 to 10-13 m/s. • Water content: 3 to 9%. • Porosity: 12% (± 1.8%). • Pore diameter: 2 to 4·10-8 m. • Hydraulic pressure: +340 to 342 m.
Geo-chemical characteristics • Water salinity: 2.7 to 7.2 g/L. • CEC: 17.5 meq/100g (10 to 25 meq/100 g). (other parameters such as pH and Eh have not yet been validated)
Solution transfer characteristics • Diffusion coefficient – cations: 1.2·10-12 m²/s (11·10-13 to 2·10-11). • Diffusion coefficient – anions: 8·10-11 m²/s (1·10-11 to 4·10-10).
Thermo-mechanical characteristics • Simple mechanical compressive strength: 22 to 28 MPa. • Uniaxial tensile strength: 0.9 to 5.4 MPa. • Young modulus: 2.3 to 11 GPa. • Rock density: 23.2 to 26.1 kN/m³. • Thermal conductivity: 1.35-1.65 W/m°C (horizontal). 1.8-2.0 W/m°C (vertical). • Heat capacity: 1.9·106 to 1.2·106 J/m³K
TABLE 2-7. Groundwater characterisation of the disposal site considered in France [55].
Constituent Site pH EH (mg/L) (mVSHE) - 2- - Cl SO4 HCO3
Est 7.1–8.2 -266 700–2,000 1,000–3,000 37–100
A.2.2.3 Characterisation of the argillaceous host rock formation and bentonitic backfill material considered in Spain
At present, clay and granite are being considered as potential host rock formations for the disposal of HLW/spent fuel in Spain. According to the Spanish reference concept for disposal of radioactive waste, the waste canisters will be surrounded by a clay barrier made up of highly compacted bentonite blocks [56]. The bentonite, currently considered as reference material to backfill the disposal galleries, originates from the province of Almería [57].
39
A.2.2.3.1 Characterisation of the argillaceous host rock formation
The main physico-chemical and hydro-mechanical properties of the argillaceous host rock formation are summarised in TABLE 2-8.
TABLE 2-8. Main physico-chemical and thermo-hydro-mechanical properties of the argillaceous host rock formation [57].
Parameter Unit argillaceous host rock
Physical properties In situ density g/cm³ 2.21 Dry density g/cm³ 1.92 Porosity % 31 Water content wt.% 15.14
Hydraulic properties Permeability m/s 4.2×10-12
Thermo-mechanical properties Thermal conductivity W/m·K 1.5 Specific heat J/kg°C 1,117 Geothermal gradient °C/m 0.03 Young Modulus MPa 572 Poisson Ratio (in situ) 0.3 Compression strength MPa 4.2 Effective friction angle ° 31 Effective cohesion MPa 0.39 Compressive modulus MPA 476.6 Tangential modulus MPA 220 Swelling pressure MPA 0.17
Considering a general geothermal gradient of 3°C/100m, the temperature of the argillaceous host rock formation at 260 m depth will be approximately 17.8°C, if an average temperature on the surface of 10°C is considered [58].
A.2.2.3.2 Characterisation of the bentonitic backfill material
The mineralogical and chemical composition of the bentonitic backfill material, considered in the Spanish reference concept, is given in TABLE 2-9 and TABLE 2-10, respectively. The main physico-chemical and hydro-mechanical properties of the bentonitic backfill material are summarised in TABLE 2-11.
40
TABLE 2-9. Mineralogical composition of the bentonitic backfill material [59].
Species Content (%)
Major minerals Smectite 92 ± 3 Plagioclase 2 ± 1 Quartz 2 ± 1 Cristobalite 2 ± 1 Tridimite traces Calcite traces Feldspars-K traces
Minor minerals Carbonates (calcite, dolomite) 0.60 ± 0.13 Organic matter (expressed as CO2) 0.35 ± 0.05 Sulphates (soluble, gypsum) 0.14 ± 0.01 Sulphates (low soluble, barite, celestite) 0.12 ± 0.05 Sulphides (pyrite) 0.02 ± 0.01 Chlorides (halite) 0.13 ± 0.02
TABLE 2-10. Chemical composition (expressed in oxides) of the bentonitic backfill material [59].
Major elements Content (%) Residuals Content (mg/L)
SiO2 58.71 ± 1.89 Ba 164 ± 25 Al2O3 17.99 ± 0.71 Sr 220 ± 23 FeO 0.25 ± 0.09 Ce 74 ± 5 Fe2O3 2.85 ± 0.12 Co 9 ± 3 MgO 4.21 ± 0.21 Cr 8 ± 2 MnO 0.03 ± 0.00 Cu 25 ± 9 CaO 1.83 ± 0.10 La 40 ± 2 Na2O 1.31 ± 0.09 Ni 21 ± 3 K2O 1.04 ± 0.05 V 16 ± 2 TiO2 0.23 ± 0.01 Y 25 ± 2 P2O5 0.02 ± 0.01 Zn 65 ± 4 + H2O 6.38 ± 0.33 Zr 43 ± 0 - H2O 6.58 ± 2.52 U 2.0 ± 0.5 CO2 (organic) 0.35 ± 0.05 Th 19 ± 1 CO2 (inorganic) 0.26 ± 0.06 Rb 41 ± 2 - SO2 (total) 0.21 ± 0.10 Cl 774 ± 140 - 2- F 0.16 ± 0.04 SO4 984 ± 65
41
TABLE 2-11. Main physico-chemical and thermo-hydro-mechanical properties of the bentonitic backfill material [57].
Parameter Unit host rock ‘granite’
Physical properties Dry density(1) g/cm³ 1.65
Hydraulic properties Permeability m/s 10-13
Thermo-mechanical properties Thermal conductivity W/m°C 0.75 Specific heat J/kg°C 1,150 Swelling pressure(2) MPa 6.0
(1) The compacted bentonite blocks fabricated to surround the waste canisters will have a density of 1.81 g/cm³ (porosity: 10%). (2) Swelling pressure for a dry density of 1.65 g/cm³.
The characterisation of the granitic host rock formation considered in Spain is discussed in section A.2.3.2.
42 A.2.3 Disposal in Granite
A.2.3.1 Characterisation of the host rock formation considered in Sweden/Finland
A summary of the ranges of groundwater constituents estimated for different times is presented in TABLE 2-12. The estimated values for the constituents of interest for bentonite porewater in the case of the Olkiluoto site is summarized in TABLE 2-13. These estimates are also considered valid for a Swedish site (but such high levels of chloride and methane have not been found yet in Sweden).
TABLE 2-12. Ranges of groundwater constituents estimated at different times [60].
Constituent Unit At closure, After closure and After closure up to infiltration into saturation 10,000 years unsaturated bentonite (up to 100 years)
pH 6−8 7−8 7−9
Redox mV oxic to −400 −150 to −308 −200 to −300
DIC(1) mol/L (0.1−16.4) ·10-3 (0.5−10.0)·10-3 (0.1−7.0)·10-3
Cl- mol/L (0.1−6.2) ·10-1 (0.2−1.6)·10-1 (0.06−4.2)·10-1
Na+ mol/L (0.1−2.8) ·10-1 (0.02−9.1)·10-2 (0.04−2.2)·10-1
Ca2+ mol/L (0.03−1.5) ·10-1 (0.03−0.2)·10-1 (0.005−1.0)·10-1
Mg2+ mol/L (0.4−1.0) ·10-2 (0.4−1.0)·10-2 (0.004−1.0)·10-2
K+ mol/L (1.3−7.7) ·10-4 (1.3−7.7)·10-4 (0.5−5.1))·10-4
2- -3 -3 -3 SO4 mol/L 0−6.3·10 0−5.8·10 0−5.2·10
HS- mol/L 0−3.0·10-4 0−3.0·10-4 0−0.9·10-4
+ -6 -4 -4 NH4 mol/L <5.5·10 , (0.03−1.7)·10 <0.6·10 if marine <1.7·10-4
-6 -2 -3 CH4 (g) mol/L <4.5·10 , 0.4·10 (0.004−17.9)·10 if saline <2.7·10-2
-5, -6 -5 H2 (g) mol/L <2.2·10 <4.4·10 <2.2·10 if saline <8.9·10-4
DOC(2) mol/L <1.7·10-4 <8.3·10-4 <1.7·10-4
(1) Dissolved Inorganic Carbon. (2) Dissolved Organic Carbon.
43 TABLE 2-13. Estimated constituent ranges at different times in bentonite porewater in the case of Olkiluoto (the type of bentonite considered in the Swedish/Finnish disposal concept is Wyoming MX-80 bentonite) [60].
Constituent Infiltrating Porewater in Porewater after groundwater at saturated bentonite closure up to closure (up to 100 years) 10,000 years
pH 6−8 7−9 7−9
Redox mV oxic to −250 −150 to −250 −200 to −280
DIC(1) mol/L (0.02−1.6)·10-4 no estimate no estimate
Cl- mol/L (0.3−6.2)·10-1 (0.3−6.2)·10-1 (0.06−4.2)·10-1
Na+ mol/L (0.2−2.8)·10-1 (3.0−5.0)·10-1 (3.0−4.0)·10-1
Ca2+ mol/L (0.3−1.5)·10-1 (0.4−4.0)·10-2 (0.4−4.0)·10-2
2- -3 -1 -1 SO4 mol/L 0−5.2·10 1.4·10 1.4·10
HS- mol/L 0−0.9·10-4 0−3.0·10-4 0−0.9·10-4
+ -6 -4 -5 NH4 mol/L <5.5·10 (0.03−1.7)·10 <5.5·10 <1.7·10-4 (2)
-6 -3 -3 CH4 (g) mol/L <4.5·10 <4.5·10 (0.004−17.9)·10 <2.7·10-2 (3)
(1) Dissolved Inorganic Carbon. (2) constituent value in the case of marine water. (3) constituent value in the case of saline water.
The assessment of long-term porewater evolution presented in TABLE 2-13 is based on modelling exercises [61-63], assuming that bentonite becomes fully saturated within the first 100 years, by which time the infiltrating groundwater will have become anoxic again. The porewater chemistry at saturation is based on the results from (i) short-term batch experiments studying the interaction of aqueous solutions with bentonite [64] and (ii) experiments with compacted bentonite [65].
A.2.3.2 Characterisation of the host rock formation considered in Spain
The main physico-chemical and thermo-hydro-mechanical properties of the granitic host rock formation are summarised in TABLE 2-14.
If granite were to be chosen as the host rock formation, then the chemistry of the near-field repository environment, which will be in direct contact with the metallic container, will consist of a bentonite buffered granitic groundwater. The chemical composition and pH of the synthetic bentonite buffered granitic porewater considered as reference media for the near-field repository environment in Spain is given in TABLE 2-15.
44 TABLE 2-14. Main physico-chemical and thermo-hydro-mechanical properties of the granitic host rock formation [57].
Parameter Unit granitic host rock
Physical properties Dry density g/cm³ 2.63
Hydraulic properties Hydraulic conductivity(1) m/s 10-8 – 10-11
Thermo-mechanical properties Young Modulus MPa 72,800 Poisson Ratio 0.31 Compression strength MPa 140 Tensile strength MPa 8.5 Thermal conductivity W/m°C 3.56 Thermal expansion coefficient(2) 1/°C 20×10-6 Specific heat J/kg°C 750 Geothermical gradient °C/10m 0.35
(1) measured in the bulk: 10-8 m/s (0-200 m) ; 10-11 m/s (200-700 m). (2) (25-250°C).
Considering a general geothermal gradient of 0.3°C/10m and an average temperature of 13°C on the surface, the temperature of the granitic host rock formation at 500 m depth will be approximately 30°C [57,66].
TABLE 2-15. Chemical composition (limits) and pH of the synthetic bentonite buffered granitic porewater (reference media in Spain) [67,68].
Constituent Composition (mg/L)
Cl- 6,550 ± 250 - NO3 110 ± 10 2- SO4 1,500 ± 30 - HCO3 27 ± 5
SiO2 8.3 ± 0.5 Br- 15 ± 1 Ca2+ 135 ± 10 K+ 20 ± 1 Mg2+ 600 ± 30 Na+ 3,750 ± 100 pH (25°C) 7.3
45 A.2.4 Disposal in Cementitious Backfill
After emplacement of the waste packages in the repository vaults, the space between them will be backfilled with a cementitious grout, known as Nirex Reference Vault Backfill (NRVB). NRVB is a mixture of water and cement (Ordinary Portland Cement, OPC), which contains several filler materials such as hydrated lime (calcium hydroxide) and fine limestone flour (calcium carbonate). During curing, the cement reacts with the water to form a cement gel which binds the material together. The material is porous and can contain a range of water concentrations: ‘free’ water within the cement gel pores is referred to as porewater and, for most OPC-based cements, has a pH in the range 12.5 to 13.3. The principal mineralogical components of fresh NRVB, hydrated under conditions representative of those expected in the repository, are [69,70]:
• sodium and potassium hydroxides, • calcium hydroxide (Ca(OH)2), • calcium silicate hydrate (CSH) gels, • calcium carbonate (CaCO3), and • hydrated calcium aluminates.
Calcium hydroxide is only slightly soluble and will be present in relatively large quantities in the solid state. The sodium and potassium hydroxides, on the other hand, are highly soluble and are expected to dissolve completely as the NRVB is saturated with water thereby conditioning the porewater to a pH value higher than that provided by Ca(OH)2. Once the initial porewater has been displaced by ingressing groundwater, the chemistry will be dominated by the Ca(OH)2 and CSH gels. The pH of NRVB porewater is expected to remain above 12.5 for over one hundred thousand years and above 10.5 for over a million years [69,70].
The chemical composition of the backfill porewater (NRVB), the near-field groundwater originating from the Sellafield site, and the near-field groundwater equilibrated with NRVB are given in TABLE 2-16 [69,71,72]. Sellafield was used as an example of a possible site for a repository, but the actual site of any repository has still to be decided.
After closure of the repository, groundwater will be able to resaturate all the excavated and backfilled spaces, and any remaining oxygen will be consumed by processes such as corrosion and microbial activity. The rate of oxygen consumption will be highly dependent on local factors such as the dryness of the backfill, which determines the rate of oxygen diffusion through the backfill, and the corrosion rates of the containers [69].
The incoming groundwater may contain significant concentrations of chloride ions (e.g. up to 18,500 mg/L in the case of the Sellafield site). The behaviour of chloride in cement is important because of the ion’s role in causing corrosion of metals. Cement is able to bind chloride ions and hence reduce the risk of corrosion. If sulphate is present, it can greatly reduce the chloride binding capacity of the cement. The addition of species that reduce the pH (e.g. silica fume, which is an industrial by-product consisting of fine silica particles), can cause an increase of the free chloride in the porewater and hence the risk of corrosion [69].
The ability of cementitious materials to form a cohesive solid with a high degree of chemical stability and a reserve of alkalinity is used to advantage in storing,
46 immobilising, and disposing of low and intermediate radioactive waste: the retention of many of the radiologically important isotopes as solids within cement, as a result of the very low solubilities of these isotopes in alkaline solutions such as cement porewater, will minimise the movement of radioactivity from a waste repository [69].
TABLE 2-16. Chemical composition of the backfill porewater (NRVB), the near-field groundwater (originating from the Sellafield Borehole 2) and the near-field groundwater equilibrated with NRVB [69,71,72].
Composition of NRVB Composition of near-field Composition of near-field porewater(1) groundwater(2) groundwater equilibrated Species with NRVB(3)
(mmol.L-1) (mg.L-1)(4) (mmol.L-1) (mg.L-1)(4) (mmol.L-1) (mg.L-1)(4)
Cations Na+ 89.1 2048.4 440.0 10115.5 390.0 8966.0 K+ 141.0 5513.4 4.9 191.6 3.9 152.5 Ca2+ 4.1 164.3 38.0 1523.0 54.0 2164.3 Mg2+ <0.1 <2.5 5.8 140.9 1.1×10-4 2.7×10-3 Fe <0.1 <5.6 1.4×10-1 7.8 1.6×10-1 8.9 Al 0.3 8.1 7.6×10-4 2.0×10-2 Cr <0.1 <5.2 Mn 3.5×10-2 1.9 2.5×10-2 1.4 Sr 0.3 26.3 3.8×10-1 33.3 0.3 26.3 Si 0.1 2.8 1.8×10-1 5.1 2.4×10-2 0.7 + NH4 0.1 1.8
Anions Cl- 0.9 31.9 510.0 18081.0 430.0 15244.8 Br- 0.4 31.2 3.2×10-1 25.6 2- SO4 0.8 76.9 12.0 1152.7 I- 1.2×10-3 1.5×10-1 F- 1.4×10-1 2.7 OH- 230.0 3911.7 2- S2O3 <0.1 <11.2 S2- <0.1 <3.2 11.0 352.7 - NO3 0.1 6.2 - NO2 <0.1 <4.6 2- CO3 0.8 48.0 2- SO3 <0.1 <8.0
pH 13.1 7.25 (<7.1)(5) 12.39
(1) ambient temperature, 600 days curing [69]. (2) near-field groundwater originating from the Sellafield site (Borehole 2; 1587 - 1602 m DET 7) [71]. (3) obtained by contacting the Sellafield Borehole 2 near-field groundwater with NRVB at 25°C [72]. (4) the values expressed in mg/L are derived from the values in mmol/L ([mg/L] = [mmol/L] × atomic weight ; atomic weight in mg/mmol). (5) the pH in brackets refers to the values measured experimentally at an on-site laboratory.
The presence of acid gases, such as carbon dioxide and sulphur dioxide, in the repository after closure can cause an increase of the corrosion rate of the container because these gases can react with the porewater in the cement thereby reducing the alkalinity [69]. The layers of material that commonly form at the steel-cement interface – the so-called ‘laitance’ layers – often act as a controlling factor for the rate of corrosion of steel in
47 cement, because they have been shown to limit the diffusion of species such as chloride and oxygen. In some cases (e.g. for Portlandite, calcium hydroxide), these layers can reduce the ability of chloride ion to initiate pitting corrosion because they can provide a reserve of hydroxide to counter local acidification caused by corrosion product hydrolysis [69].
In conclusion, the containers will initially be exposed to aerated, alkaline (pH 12.5-13) backfill porewater. Once the repository becomes anoxic and fully resaturated the containers will be exposed to an anoxic, high pH, aqueous phase (dominated by the chemistry of the ingressing groundwater) which may contain chloride ions and a mixture of other inorganic salts [69].
48 A.3 SCIENTIFIC BACKGROUND OF CORROSION DURING THE AEROBIC AND ANAEROBIC PERIOD
According to TABLE 1-1, it is evident that Fe-based materials, Cu, Ni-alloys, and Ti- alloys are considered today as the most promising candidate container materials. Therefore, the corrosion of these materials will be briefly described in this section.
A.3.1 Corrosion of Fe-based materials
The major component of steels (carbon steel, stainless steel) is iron. The thermodynamics of iron corrosion will therefore be explained in this section.
During the initial period after closure of the repository, the environment will be oxygenated as large volumes of air will be trapped in pores and voids due to the excavation works and the construction of the repository. During this period, no significant amounts of gas will be produced as a reaction product of the corrosion of the metallic overpack [73]. The oxygen will be consumed by a number of mechanisms, one being metal corrosion. The environmental conditions in the repository will therefore change progressively and become anoxic. After the oxygen has been exhausted, the nature of the corrosion processes will change and the formation of hydrogen gas will begin to dominate.
In the presence of oxygen, steel will corrode according to one of the following reactions (aerobic corrosion) [74–77]:
4 Fe + 3 O2 → 2 Fe2O3 (1)
2 Fe + 2 H 2O + O2 → 2 Fe(OH )2 (2)
The amount of available oxygen, the availability of water, the hydration conditions, and the reaction kinetics will determine which solid phase dominates [76]. The corrosion process described by reaction (2), is relatively fast in the case of carbon steel, whereas stainless steels are passivated by an oxide film. Reaction (2) produces ferrous hydroxide, which can, in the presence of oxygen, corrode further via the reaction:
4 Fe()OH 2 + 2 H 2O + O2 → 4 Fe(OH )3 (3) giving an overall reaction [75,78,79]:
4 Fe + 6 H 2O + 3 O2 → 4 Fe(OH )3 (4)
Due to the consumption of oxygen through aerobic corrosion (and probably other processes such as microbiological activity), the underground repository environment will eventually become anoxic. The time to achieve this anoxic state will depend on the rate and manner in which the repository saturates with water [77].
49 When anoxic conditions are attained, steel will continue to corrode by direct reaction with water by the following reaction (anaerobic corrosion) [75–80]:
Fe + 2 H 2O → Fe()OH 2 + H 2 ↑ (5)
Since Fe(OH)2 is thermodynamically unstable, it may be transformed to magnetite in a second reaction step known as the 'Schikorr' reaction [75,78,80]:
3 Fe()OH 2 → Fe3O4 + 2 H 2O + H 2 ↑ (6) giving an overall reaction:
3 Fe + 4 H 2O → Fe3O4 + 4 H 2 ↑ (7)
Although Fe3O4 is the thermodynamically favoured end product, Fe(OH)2 may actually dominate at temperatures below 60°C, since the conversion reaction (6) may be inhibited or at least delayed [75,76,81]. Smart et al. [82] have, however, observed passivation due to magnetite formation at temperatures as low as 30°C. Smailos et al. [83] observed the formation of a magnesium-containing ferrous hydroxide, (Fe,Mg)(OH)2, as the primary corrosion product in MgCl2-rich brines, without 2+ evidence of Fe3O4 indicating the inhibiting effect of Mg on the reaction kinetics of the formation of Fe3O4.
The Fe(OH)2-layer formed in alkali solutions (KOH) is rather porous and will most probably not lead to passivity but rather reduce the corrosion rate [84,85]. If the corrosion potential exceeds the equilibrium potential of magnetite formation, an impervious oxide film is formed on the metal surface as a protective coating against further corrosion. This effect is called passivation. In the passive state, metal dissolution is determined by the balance between slow passive film dissolution on the outer surface and film formation at the metal-film interface. These conditions occur for iron under anaerobic conditions in an alkaline environment [76].
From reactions (5), (6), and (7), it is evident that the anaerobic corrosion of the metallic overpack could result in the formation of substantial hydrogen concentrations in the underground repository during the anoxic period. If the production rate of hydrogen is higher than its migration rate through the backfill, the hydrogen activity could exceed the hydrostatic pressure in the repository, resulting in gas bubble formation. Such an occurrence could have a detrimental effect on nuclide transport in the near-field, and may even cause cracking of a concrete engineered barrier [77]. A build up of hydrogen pressure within the repository can lead to a reduced driving force for the corrosion reactions, thereby inhibiting the formation of Fe(OH)2 or Fe3O4. From thermodynamic calculations (free energy changes), it is known that the overpressure to suppress the production of ferrous hydroxide and magnetite is 140 and 80 MPa respectively. Therefore, in view of the limited hydrostatic pressure available in the repository, the hydrogen overpressure cannot suppress the corrosion reactions completely [79]. It has been shown experimentally [78,82] that pressures up to ~9 MPa have no effect on the corrosion potential of anaerobically corroding carbon steel wires in artificial Swedish groundwater.
50 A.3.1.1 Effects of environmental factors
A.3.1.1.1 Role of the chemistry of the aqueous environment in contact with the metallic materials
With regard to corrosion in neutral or alkaline media imposed by the surrounding materials (backfill/buffer), a key parameter is probably the redox potential. This parameter strongly influences the general corrosion kinetics of unalloyed and low- alloyed steels and determines the initiation or propagation of localised corrosion of passive alloys [86].
The presence of dissolved oxygen (and its possible supply) in the aqueous environment in contact with the metallic container is a major factor in determining the different kinetics or risks of corrosion. Consumption of this oxygen (by corrosion or through reacting with the minerals) will lead to a change in the conditions from ‘oxidising’, which is generally detrimental for metallic materials, to ‘reducing’, which is favourable in limiting the general or localised corrosion processes [86]. Generally, the addition of oxygen to the environment significantly increases the corrosion rate [87,88]. This can be understood because oxygen forms part of the overall electrochemical reactions occurring at the interface between the aqueous phase and the metal surface [89]: the corrosion reaction can only go on as quickly as the electrons, which are produced at the anodic sites through oxidation of the metal surface, can be consumed in a cathodic reaction. The presence of oxygen provides an additional reduction reaction thereby causing more electrons to be consumed and thus increasing the corrosion rate [87]. Temperature plays a dual role with respect to oxygen corrosion. It is generally known that increasing the temperature reduces the oxygen solubility in water and increases the corrosion rate (see section A.3.1.1.2). The corrosion rate reaches a maximum at around 80°C. However, beyond this temperature, the corrosion rate of carbon steel decreases because the reduced oxygen content limits the oxygen reduction reaction, thereby preventing occurrence of the iron dissolution process. This only applies to an open system, in which oxygen can be released from the system. For closed systems, in which oxygen cannot escape, corrosion will continue to increase linearly with temperature[89].
The presence of individual species in the aqueous near-field environment, even at trace levels, can have a major impact on the corrosion behaviour of the metallic container, either detrimental or beneficial [87,90]. Chloride is particularly detrimental to materials that depend on a passive film for their corrosion resistance (e.g. stainless steels). At high concentrations, the passive film can break down locally and result in pitting [87,90], while at low concentrations the anodic current in the passivity range is increased. For most metals, a critical concentration has been established below which pitting does not occur (depending on the type of metal, pH, T, etc.) [91]. For corrosion allowance materials, the presence of chloride is more likely to lead to general or uniform corrosion. Fluoride ions can also induce pitting corrosion in iron and steel. Because the ion radius of F- is smaller than the ion radius of O2-, F- can replace O2- in the oxide film, destroying passivity [91].
51 Sulphur-containing species can enhance the dissolution, block or retard the growth of the passive film, and cause passivity breakdown and pitting by enrichment at the metal-film interface [92,93]. In near-field disposal environments, based on argillaceous backfill materials 2- (bentonite), the presence of thiosulphate (S2O3 ) is likely as it is an intermediate compound in the oxidation of pyrite to sulphate [94,95]. Thiosulphate will decompose by disproportionation or electroreduction to form elemental sulphur and sulphide [96-99], which will adsorb on the stainless steel surface, thereby activating anodic dissolution [93,100]. Thiosulphate pitting is known to occur in a limited - 2- 2- potential range [97,101-103] and molar ratio ([Cl ] + [SO4 ] / [S2O3 ]) [101-104], as - 2- illustrated schematically in FIGURE 3-1. When the molar ratio ([Cl ] + [SO4 ] / 2- [S2O3 ]) is in the range 10 to 30 [101-103], the risk for thiosulphate pitting is most imminent.
FIGURE 3-1. Explanation of critical conditions of potential and solution composition for thiosulphate pitting of stainless steels [102].
When chloride and thiosulphate are present simultaneously, the initial breakdown of the passive film is caused by Cl- (thiosulphate does not play a major role in this step), whereas the presence of thiosulphate has a drastic effect on the subsequent step, viz. stabilising metastable pits, and thus effectively lowering the pitting potential [86,92].
The near-field disposal environment, surrounding the metallic container, can also contain beneficial species (corrosion inhibitors) that are added intentionally to the environment or in the form of naturally occurring species [87,90]. The most common occurring anions, present in an argillaceous disposal environment, that reduce the tendency to pitting (in chloride-containing environments) are sulphates, hydroxides, nitrates, carbonates, etc [88]. The inhibition mechanisms are thought to involve either competitive adsorption or the formation of more protective films. The inhibitive effect
52 generally becomes less pronounced in environments with a higher chloride content [87,88,90].
A.3.1.1.2 Role of temperature
The emplacement of waste packages in the repository will lead to a rise in temperature of the metallic materials and their environment. Depending on the nature of the packages and the repository concepts, this rise in temperature, particularly in contact with the container, can last several tens, hundreds, even thousands of years before reverting to the temperature imposed by the geological medium [86].
These temperature changes can influence corrosion in various ways [86,87,90]:
• it can cause desaturation or maintenance of a dry atmosphere in contact with the container, thus delaying or limiting the corrosion process.
• it can exacerbate corrosion rates in aqueous media. Corrosion is an activation- controlled chemical reaction, the rate of which is greatly affected by temperature. Generally, the corrosion rate increases significantly as temperature increases. The increased corrosion rate results from increased activation energy for chemical and electrochemical reactions, increased diffusion rates in the electrolyte, and increased transport through the electrolyte or environment and across films that may be formed on the metal surface [87,90].
• it can modify the corrosion products, which depending on their nature and/or structure can play a greater or smaller protective role in corrosion.
A.3.1.2 Role of pH
In alkaline environments (pH = 10-13.5), the corrosion rate of iron is expected to be negligibly low due to the ease with which a protective passive film is formed [105]. The influence of pH on the corrosion behaviour of a passive alloy in a chloride- containing environment is presented schematically in FIGURE 3-2. In general, passivity of metallic materials decreases with increasing acidity of the environment: the polarisation curves move to higher current densities in the passive range (Ipass) with decreasing pH, indicating an increase in corrosion rate if an environment establishes potentials in the passive range; as the pH drops, the passive potential range becomes smaller making it much more difficult to maintain passivity and finally pasivity is completely destroyed [86,105].
The effect of even small pH changes, however, cannot be interpreted independent of the possible influence of accompanying ions. These species may alter the anodic dissolution processes through effects on the kinetics of the interface reactions and by altering the physical and chemical structure of solid corrosion products [105].
The effect of pH of the bulk environment on the localised corrosion behaviour of passive Fe-based metallic materials has already been the subject of many investigations. It has been found that the pH can play a major role on the initiation process of pitting corrosion: a decrease in pH is usually accompanied by a
53 displacement of the pitting potential, ENP, in the active direction. In highly alkaline environments, the pitting potential is significantly displaced in the noble direction, which is in line with the known inhibiting effect of higher concentrations of OH- ions [88,106,107]. The pit propagation process seems to be rather independent of pH. This may be understood when considering that the hydrolysis reaction in the pit generates its own characteristic acidity, which is little influenced by the acidity of the bulk environment [88].
FIGURE 3-2. Influence of pH on the polarisation curve of a passive alloy in a chloride-containing environment. Source: adapted from [86].
A.3.1.3 Role of radiation
The presence of HLW/spent fuel will impose a radiation field on the immediate vicinity of the waste containers. This radiation field can exert an influence on the corrosion behaviour of the metallic materials in disposal relevant environments. Of the various types of radiation, only gamma radiation is present at the surface of container. The dose rate depends on the container design and the nature of the waste enclosed inside the container.
Gamma radiation can affect the environmental chemistry by producing oxidising and reducing reactive particles (radicals) and molecular species, which can change the rate and the mechanism of corrosion [108-114]. In pure water, the oxidising species • - will be radicals like OH and molecular species like O2, H2O2, and O2 . The reducing - species include H, eaq , and H2. In groundwater, other species may also be produced depending on its composition. In highly saline solutions, chloride species are - - produced (Cl2 , Cl2, ClO , etc.). In highly carbonate-containing solutions, the • carbonate radical, CO3 , may form. During the radiolysis of moist air, nitrogen oxides (and as a consequence, nitric acid) could be produced [60].
54 Furthermore, the absorption of gamma radiation in the semi-conducting protective oxide layers of metals can induce photoradiation effects [115], which can lead to a change in the corrosion rate by facilitating cathodic/anodic reactions at the oxide/electrolyte interface. Glass [115] concluded that the radiolytic products will exert the greatest effect on the corrosion process of metals.
55 A.3.2 Corrosion of Ni-based materials
The Pourbaix (potential-pH) diagram for the nickel-water system at 25°C is shown in FIGURE 3-3. Nickel can be considered to be a slightly noble metal, as its domain of thermodynamic stability has a small zone in common with that of water [116].
Passivation
Corrosion
Corrosion immunity
FIGURE 3-3. Pourbaix (potential-pH) diagram of the nickel- water system at 25°C (adapted from Ref. [117]).
According to the thermodynamic diagram (FIGURE 3-3), the corrosion resistance of nickel should depend on the pH and the presence of oxidising agents, for non- complexing solutions, as follows: Ni should be uncorrodible in neutral or alkaline solutions free from oxidising agents, slightly corrodible in acid solutions free from oxidising agents, and very corrodible in acid or very alkaline solutions containing oxidising agents. Neutral or slightly alkaline oxidising solutions should cover it with a layer of oxide [116]. These predictions are only part in agreement with the experimental facts. When there is no oxidising action, i.e. notably in the case of solutions not containing oxidising agents and in the absence of any anodic polarisation, nickel is hardly corroded at all, not only in neutral or alkaline solutions, but also in many acid solutions. This favourable behaviour in non-oxidising acid media is probably due, on the one hand, to the great irreversibility of the corrosion reaction Ni → Ni2+ + 2e- and, on the other hand, to the large hydrogen overpotential on nickel [116].
The position of the domain of stability of Ni(OH)2 in FIGURE 3-3 shows that it is a thermodynamically stable substance in the presence of water or neutral or slightly alkaline solutions free from oxidising or reducing agents. It readily dissolves in acid
56 solutions with the formation of nickelous ions Ni2+. In very alkaline solutions, it dissolves - as nickelite ions HNiO2 . According to FIGURE 3-3, the oxidation of nickelous hydroxide in alkaline media can give rise to the formation of Ni3O4, Ni2O3, and NiO2 [116]. Ni3O4 is thermodynamically stable in the presence of aerated water due to the fact that its stability domain has a large zone in common with that of water [116]. Ni2O3 is an oxidising agent, since its stability domain lies completely above line (b) (oxidation of H2O). It is therefore unstable in the presence of water which it tends to decompose with the evolution of oxygen. It dissolves in acid solutions with the formation of Ni2+ ions and the evolution of oxygen. It is insoluble in alkaline solutions [116]. NiO2 is an unstable substance which decomposes rapidly into Ni2O3, Ni3O4, and oxygen [116].
Compared to stainless steels, nickel can accommodate larger amounts of alloying elements in solid solution, mainly chromium, molybdenum, and tungsten. Therefore, nickel-base alloys generally can be used in more severe environments than the stainless steels [87]. Nickel generally shows good resistance to corrosion in the normal atmosphere, in natural freshwaters and in deaerated nonoxidising acids, and it has excellent resistance to corrosion in alkaline solutions (e.g. many corrosion problems involving caustic solutions are handled with nickel) [118,119]. Another important attribute is the large and rapid increase in stress-corrosion resistance as the nickel content of stainless alloys increases above 10%. Nickel is not resistant to strongly oxidising solutions, e.g. nitric acid and ammonia solutions. Nickel is often used for applications requiring strength at high temperatures. However, nickel and its alloys are attacked and embrittled by sulphur-bearing gases at elevated temperatures [119].
The localised corrosion susceptibility (pitting and crevice corrosion) of stainless steels and Ni-Cr-Mo alloys, in chloride-containing environments, depends on the same parameters: the pitting and crevice corrosion resistance improve with increasing chromium and molybdenum content. Nitrogen additions, in the range of 0.1-0.4%, is also known to exert a benificial influence on the corrosion resistance, at least for the higher alloyed grades [55].
A useful guide to the relative corrosion resistance (with respect to pitting and crevice corrosion) of stainless steels is the pitting resistance equivalent number (PREN), commonly given by [55,88]:
PREN = %Cr + 3.3 %Mo + 16 %N
The nominal chemical compositions of the main Ni-Cr-Mo alloys that are being considered as candidate container material, in various countries, for the final disposal of HLW/spent fuel in deep geological formations are listed in TABLE 3-1, classified in ascending order of their PREN-value. Some stainless steels have been added for comparison.
57 TABLE 3-1. Nominal chemical compositions of Ni-Cr-Mo alloys and stainless steels studied as candidate container material, in various countries, for the final disposal of HLW/spent fuel in deep geological formations [55].
Alloy Chemical composition (wt.%) PREN(4) designation
AISI(1) UNS(2) Fe Cr Ni Mn Mo Si C(3) S P Others
Stainless steels 304L S30403 bal 18.0-20.0 8.0-12.0 2.0 - 1.0 0.03 0.03 0.045 - 18 316L S31603 bal 16.0-18.0 10.0-14.0 2.0 2.0-3.0 1.0 0.03 0.03 0.045 - 25 904L N08904 bal 19.0-23.0 23.0-28.0 2.0 4.0-5.0 1.0 0.02 0.035 0.045 Cu: 1.5 35
Ni-Cr alloys 825 N08825 29.0 21.5 42.0 - 3.0 0.5 0.05 - - Cu: 2.0; Ti: 1.0 32 625 N06625 5.0 21.5 62.0 - 9.0 0.5 0.1 - - Nb: 4.0 52 C-22 N06455 3.0 22.0 56.0 - 13.0 0.08 0.015 - - W: 3.0 65 C-4 N06022 3.0 16.0 65.0 - 15.5 0.08 0.01 - - - 69
(1) American Iron and Steel Institute numbering system. (2) Unified Numbering System. (3) Maximum. (4) Pitting Resistance Equivalent Number (PREN = %Cr + 3.3 %Mo + 16 %N) [88]. bal: balance.
58 A.3.3 Corrosion of Cu-based materials
Copper is a relatively 'noble' metal. Copper is thermodynamically stable in oxygen-free water and does not corrode. However, in water containing complexing species such as - - 2- + Cl , HS , S2O3 , and NH4 , copper may no longer be immune to corrosion. The dissolved species that are particularly conducive to copper corrosion under geological repository conditions are sulphides and chlorides.
A.3.3.1 Copper corrosion in the presence of oxygen
In aqueous environments at ambient temperature, the corrosion product predominantly responsible for protection is cuprous oxide (Cu2O). This Cu2O film is adherent and is formed by the electrochemical processes, presented by the following half reactions [120,121]:
+ − 4 Cu + 2 H 2O → 2 Cu2O + 4 H + 4 e (anode : Cu dissolution) (10a) and
− − O2 + 2 H 2O + 4 e → 4 OH (cathode : O2 reduction) (10b) giving an overall reaction:
4 Cu + O2 → 2 Cu2O (Cu dissolves as Cu(I)) (10c)
Under repository conditions, upon saturation, the canister surface will be covered by a duplex corrosion product layer comprising an inner layer of Cu2O and an outer layer of electrically insulating Cu(II) salts, which would probably be either malachite (Cu2CO3(OH)2) or atacamite (CuCl2 ·3Cu(OH)2), depending upon the relative 2- - concentrations of CO3 and Cl in the porewater [121]. In general, this outer layer is porous and non-continuous, enabling further O2 reduction to proceed on the inner Cu2O layer. The rate of the interfacial reaction will depend on the electronic properties of the defected, semi-conducting Cu2O layer. Although the rate of the interfacial reaction will be slower than on a 'bare' Cu surface, the overall rate of O2 reduction on the canister surface is likely to be controlled by the rate of supply of O2, rather than by the rate of the surface reaction. A consequence of the presence of electrically insulating Cu(II) salts is the possible spatial separation of anodic and cathodic surface reactions, possibly leading to localised corrosion [121].
Cupric species are formed from the homogeneous irreversible oxidation of Cu(I) by O2 according to reaction (11) [121,122]:
− 4 Cu(I) + O2 + 2 H 2O → 4 Cu(II) + 4 OH (11)
The precise speciation of the Cu(I) and Cu(II) species will depend on the composition of the porewater. A significant fraction of the Cu(II) formed by reaction (11) will adsorb on the bentonite clay buffer. Other Cu(II) species will precipitate on the canister surface, resulting in a thickening of the outer layer of the duplex Cu2O/Cu(II) salt film formed on the canister surface during the unsaturated phase [121].
59
Cupric species can also be reduced on the canister surface, according to reaction (12) [121]:
− − − Cu(II) + 2 Cl + e → CuCl2 (12)
A.3.3.1.1 Influence of chloride in the presence of oxygen
The most appropriate description of the various chemical and electrochemical steps in the anodic dissolution of Cu in a chloride-containing environment is [121]:
− − Cu + Cl = CuCl ADS + e (13a)
− − CuCl ADS + Cl = CuCl2 (surface) (13b)
− diffusion − CuCl2 (surface) ⎯⎯→⎯⎯ CuCl2 (bulk) (13c)
Chloride ions play an important role in the formation and properties of surface films on Cu. Cuprous oxide (Cu2O) may form via a number of processes in chloride-containing environments. The initial stages of film formation involve a competition between Cl- and - OH for surface sites, followed by the loss of H2O [121]
− − CuCl ADS + OH = Cu()OH ADS + Cl (14a)
2 Cu()OH ADS → Cu2O + H 2O (14b)
The extent of Cu2O formation depends on the relative chloride and hydroxide - concentrations (i.e. pH) and the rate of mass transport: CuCl2 formation is favoured by higher chloride concentrations and higher rates of mass transport; higher pH favours Cu2O formation. Although these processes relate only to the formation of the first few monolayers of Cu2O, they are nevertheless important in determining the properties of the passivating interfacial Cu2O layer. The incorporation of CuCl "islands" in the surface Cu2O film creates defects that are believed to be initiation points for pitting [121].
Copper continues to dissolve through the thin surface layer, especially via defects and - cracks in the film. Hydrolysis of dissolved CuCl2 , or of CuCl produced by the precipitation of dissolved Cu(I), results in further Cu2O growth [121].
− + − 2 CuCl2 + H 2O → Cu2O + 2 H + 4 Cl (15a) or
+ − 2 CuCl + H 2O → Cu2O + 2 H + 2 Cl (15b)
In the presence of O2 in chloride-containing environments, an outer layer of precipitated CuCl2 ·3Cu(OH)2 forms on the Cu2O film. Two mechanisms have been proposed to account for the formation of the CuCl2 ·3Cu(OH)2 layer [121]:
60 • the CuCl2 ·3Cu(OH)2 layer is formed by the precipitation of dissolved Cu(II), formed by the homogeneous oxidation of Cu(I) by O2 (reaction (11)), once local super- saturation of the environment by Cu(II) is achieved. • other authors [123,124] suggest that the CuCl2 ·3Cu(OH)2 layer is formed through the oxidation of the underlying Cu2O film via the overall reaction (16):
− − 2 Cu2O + O2 + 2 Cl + 4 H 2O → CuCl2 ⋅3Cu(OH )2 + 2 OH (16)
Chloride ions will also affect the properties and stability of the precipitated films. The - 2- substitution of monovalent Cl ions for divalent O ions in the Cu2O lattice creates defects and enhances the semi-conducting properties of the surface film. Depending - - upon the [Cl ], however, Cu2O films formed in Cl solutions may be more susceptible to localised breakdown and pitting attack. At sufficiently high [Cl-], the surface layer may become so defected that it no longer protects the surface and Cu dissolves actively [121].
2- A.3.3.1.2 Influence of other anions (SO4 , ammonia)
Although Cl- may be the predominant groundwater species, the porewater will contain other anions, especially at short times prior to the influx of Cl-. A major porewater 2- - 2- constituent is SO4 . The dissolution behaviour of Cu in Cl /SO4 mixtures will follow the same mechanism as in Cl- solutions (i.e. reaction (13)) [121].
Ammonia may be present in the repository in small amounts, introduced either during vault construction, by microbial activity in the groundwater, or by gas-phase radiolysis of atmospheric N2. Ammonia is important for the stress corrosion cracking of Cu. Ammonia also forms strong complexes with Cu(I) and Cu(II) [121].
A.3.3.1.3 Mechanism for the corrosion of copper in compacted buffer material saturated with saline, O2-containing groundwater
FIGURE 3-4 shows the proposed mechanism for the corrosion of Cu in compacted bentonite saturated with saline, O2-containing groundwaters. Copper dissolves - - reversibly in the form of CuCl2 species. If this species is stable, i.e. at high [Cl ] and/or - low [O2], CuCl2 slowly diffuses away from the Cu surface. In low-salinity groundwaters and/or at high [O2], Cu(I) is irreversibly oxidised to Cu(II). Cupric species are removed from solution by adsorption and (to a greater extent than Cu(I)) by precipitation. These processes drive further dissolution by slowing down the rate of reduction of Cu(II) to Cu(I) [121].
61
FIGURE 3-4. Proposed mechanism for the corrosion of copper in compacted buffer material saturated with saline, O2-containing groundwater [121].
A.3.3.2 Copper corrosion in the absence of oxygen
A.3.3.2.1 Influence of chloride
In the absence of oxygen, according to thermodynamic considerations, a combination of high chloride content, low pH (pH < 5-6), and a high temperature (T = 80-100°C) can be unfavourable for the general corrosion of Cu [125–127]. However, the presence of other electron acceptors than protons is needed for any substantial corrosion to occur at pH 6 [126,127].
The corrosion of copper in chloride solutions, with water or protons as the only electron acceptor can be written as [121]
− + 1−n Cu (s) + n Cl + H = CuCln + ½ H 2 (aq) (17)
A.3.3.2.2 Influence of sulphide
Because sulphide minerals are present in many types of bentonite and in deep Fenno- Scandian Shield groundwaters (see section A.2.3.1, TABLE 2-1), and because cuprous sulphide (Cu2S) is thermodynamically stable at potentials below the H2/H2O equilibrium line [128], Cu canisters may be subject to corrosion in the presence of sulphide under the long-term reducing conditions expected to develop in the repository.
Cu forms a duplex bilayer corrosion product film in sulphide solutions, comprising an inner layer of Cu2S (variously reported to be between 0.4 nm [129] and 25-50 nm thick [130]) and a thicker outer layer of CuS. Cu2S is formed according to reaction (18) [131]:
− − 2 Cu + HS + H 2O → Cu2 S + H 2 + OH (18)
62 It is expected that a duplex Cu2O/Cu2S layer will be formed in a repository once the initially trapped atmospheric O2 has been consumed and the environmental conditions become more reducing. This assumption is confirmed by the results of a corrosion study on Cu-Ni alloys in deaerated seawater in the presence of sulphides [132].
63 A.3.4 Corrosion of Ti-based materials
The Pourbaix (potential-pH) diagram for the titanium-water system at 25°C is shown in FIGURE 3-5. Titanium is not a noble metal: in actual fact, its domain of thermodynamic stability does not have any portion in common with the domain of thermodynamic stability of water, and it lies well below the latter. If this metal generally exerts a high electrode potential, it is because a passivating film of oxide is formed on its surface [133].
FIGURE 3-5. Pourbaix (potential-pH) diagram of the titanium-water system at 25°C [134].
The excellent corrosion resistance of titanium alloys results from the formation of these very stable, continuous, highly adherent, and protective oxide films on the metal surface. Because titanium metal itself is highly reactive and has an extremely high affinity for oxygen, these benificial surface oxide films form spontaneously and instantly when fresh metal surfaces are exposed to air and/or moisture [134]. The study of the corrosion resistance of titanium can therefore essentially be reduced to a study of the corrosion resistance of the oxide film [135]. In fact, a damaged oxide film can generally reheal itself instantaneously if at least traces of oxygen or water (moisture) are present in the environment. However, anhydrous conditions in the absence of a source of oxygen may result in titanium corrosion, because the protective film may not be regenerated if damaged [134].
64 The nature, composition, and thickness of the protective surface oxides that form on titanium alloys depend on environmental conditions. In most aqueous environments, the oxide is typically TiO2, but may consist of mixtures of other titanium oxides, including TiO2, Ti2O3, and TiO. High-temperature oxidation tends to promote the formation of the chemically resistant, highly crystalline form of TiO2, known as rutile, whereas lower temperatures often generate the more amorphous form of TiO2 (anatase), or a mixture of rutile and anatase. Although these naturally formed films are typically less than 10 nm thick and are invisible to the eye, the TiO2 oxide is highly chemically resistant and is attacked by very few substances, including hot, concentrated HCl, H2SO4, NaOH, and (most notably) HF [134]. If the solution is moderately oxidising, the metal can be considered to be safe from corrosion. In fact, in seawater, which is a particularly dangerous medium for most metals, titanium resists corrosion [133].
When titanium is in the fully passive condition, corrosion rates are typically much lower than 0.04 mm/year, which is well below the maximum corrosion rate commonly accepted by designers. This very small, acceptable corrosion is attributed to the finite oxidation (typically TiO2 film growth) of titanium alloy surfaces. In many environments in which titanium is fully resistant, slight surface oxide growth may occur. This oxide growth manifests itself as coloured surfaces and very slight weight gain by test coupons [134]. Titanium and its alloys are fully resistant to water, all natural waters, and steam to temperatures in excess of 315°C. Slight weight gain is usually experienced in these benign environments, along with some surface discolouration at higher temperatures from finite passive film thickening. The typical contaminants encountered in natural water streams, such as iron and manganese oxides, sulfides, sulfates, carbonates, and chlorides, do not compromise the passivity of titanium. In media containing chloride levels greater than 1,000 mg/L at temperatures about 75°C, consideration should be given to possible crevice corrosion when tight crevices exist in service. Titanium alloys exhibit negligible corrosion rates in seawater to temperatures as high 260°C. Pitting and crevice corrosion are totally absent in ambient seawater, even if marine deposits form and biofouling occurs [134].
Titanium exhibits anodic pitting potentials, ENP, that are very high (>> 1 V): pitting potentials, versus Ag/AgCl reference electrode, are typically in the +5 to +10 V range for chlorides. Thus, pitting corrosion is generally not of concern for titanium alloys under normally operating repository conditions [134]. The protection potentials, EPP, of titanium alloys are also very high relative to the alloy corrosion potentials, e.g. EPP values of +5.6 V (vs. Ag/AgCl) have been reported for grade 7 alloy in saturated NaCl. This explains why titanium alloys will not experience spontaneous pitting attack in aqueous chloride media, even if oxidising species are present [134].
The titanium grade 7 alloy5 has extensively been investigated for potential applications in high-level nuclear waste disposal [134]. The effect of temperature and pH on crevice corrosion of titanium grade 7 in saturated NaCl brines is presented in FIGURE 3-6. This alloy provides crevice corrosion resistance to brine pH values as low as 0.7 and temperatures as high as 300°C. It should be cautioned that the presence of ions, which differ from Na+, such as Ca2+ and
5 Commercially pure titanium alloyed with 0.15% Pd.
65 Mg2+, can give rise to localised corrosion phenomena at temperatures deviating from normal crevice corrosion guidelines (T lower than the ones predicted in FIGURE 3-6) [134,135].
FIGURE 3-6. Temperature-pH limits for crevice corrosion of grade 7 titanium in saturated NaCl brines. Adapted from Ref. [135].
Titanium grade 7 is also immune to stress corrosion cracking (SCC) except in a few specific environments, which are not encountered in deep geological disposal media. These specific environments include anhydrous methanol/halide solutions, nitrogen tetroxide (N2O4), red fuming HNO3, and liquid or solid cadmium [134].
The coupling of titanium with dissimilar metals usually does not accelerate the corrosion of titanium. The exception is in strongly reducing environments in which titanium is severely corroding and not readily passivated. In this uncommon situation, accelerated corrosion may occur when titanium is coupled to more noble metals [134].
66 A.4 EXPERIMENTAL RESULTS
A.4.1 Corrosion studies relevant to the German disposal concept in rock salt
The experimental results presented in this section originate from investigations performed at FZK, GNF/IUT, and ENRESA/INASMET. The long-term immersion experiments and the in situ corrosion experiments were developed and performed at FZK. The electrochemical and radiochemical studies were performed at GNF/IUT. The SCC-experiments were performed at ENRESA/INASMET.
More detailed information can be found in References [29,30,36,38,41-43,68,83,136- 150,152-160].
To qualify packaging materials, comprehensive corrosion studies have been performed on a wide range of metallic materials in brines relevant to the repository conditions in deep rock-salt formations. These corrosion studies consist of laboratory- scale and in situ experiments in the Asse salt mine [136,137. The laboratory-scale experiments include • screening experiments to determine candidate container materials, • experiments to evaluate the long-term corrosion behaviour of the candidate metallic materials , and • experiments on real specimens originating from the POLLUX-disposal container for spent fuels.
Screening corrosion tests were peformed on a wide range of materials such as unalloyed and low-alloyed steels (0.15-0.17 wt.% C), nodular cast iron (3.7 wt.% C), Ni- resists (22-30 wt.% Ni), Cr-Ni steels, Ni-base alloys (Inconel 625, Incoloy 825, Hastelloy C-4), and the Ti-alloy Ti 99.8-Pd in order to make a selection of the most promising candidate metallic materials for further detailed studies [136,137]. The chemical composition of the materials tested in salt brines is summarised in TABLE 4-1.
On the basis of the results of the screening experiments, carbon steels, especially the fine-grained steel TStE 355 (for the corrosion-allowance packaging concept), and the passive corroding alloys Ti 99.8-Pd and Hastelloy C-4 (for the corrosion-resistant packaging concept) were identified as the most promising materials for the manufactering of long-lived containers for heat-generating nuclear wastes (vitrified high-level waste and spent fuel).
Subsequently, these materials – that were identified as the most promising candidate overpack materials – were subjected to more detailed investigations to determine the influence of several important parameters on its corrosion behaviour in disposal relevant salt brines, such as brine composition (MgCl2-rich and NaCl-rich brines), temperature (90, 150, 170, and 200°C), pH, gamma dose rate, welding (TIG, EB, SAW, surface welding), etc.
67 TABLE 4-1. Chemical composition of the materials tested in salt brines [29,41,43,83,138-143].
Composition (wt.%) Material Cr Ni Mo Ti Pd C Si Mn V O2 H2 Fe
TStE 3551 - - - - - 0.17 0.44 1.49 - - - bal.
TStE 4602 0.03 0.51 - - - 0.18 0.34 1.50 0.15 - - bal.
15 MnNi 6.33 0.04 0.79 - - - 0.17 0.22 1.59 - - - bal.
GS 16Mn54 - - - - - 0.16 0.66 1.51 - - - bal.
GGG 40.35 - - - - - 3.70 1.83 0.21 - - - bal.
Si-cast iron - - - - - 0.72 15.0 0.62 - - - bal.
Cr-Ni steel 1.43066 18.1 10.2 ------bal.
Cr-Ni steel 1.48337 22.4 13.8 ------bal.
Ni-Resist D28 2.39 22.0 - - - 2.65 2.40 1.14 - - - bal.
Ni-Resist D49 5.50 30.9 - - - 2.60 4.25 0.50 - - - bal.
Hastelloy C-410 15.4-16.8 bal. 15.2-15.9 0.33 - 0.006 0.05 0.09 - - - 0.05
Inconel 625 21.8 61.9 8.5 - - 0.03 0.3 0.1 - - - 0.05
Incoloy 825 21.3 43.3 2.9 - - 0.01 0.4 0.6 - - - 0.05
Ti 99.8-Pd11 - - - bal. 0.18 0.01 - - - 0.04 0.001 0.05
bal. : balance. 1 unalloyed carbon steel (material no. : 1.0566). 2 low-alloyed carbon steel (material no. 1.8915). 3 low-alloyed carbon steel (material no. 1.6210). 4 cast steel (material no. : 1.1131). 5 nodular cast iron. 6 austenitic stainless steel (AISI 304L). 7 austenitic stainless steel (AISI 309S). 8 Material no. : 0.7660. 9 Material no. : 0.7680. 10 Ni-alloy (material no. : 2.4610). 11 Ti-alloy (material no. : 3.7025.10).
According to the German concept, cavities remaining in the emplacement location will soon be closed automatically by convergence. The limited amount of brines that will be present in these cavities will be consumed very rapidly through corrosion with the metallic overpack and also the contact time brine/overpack will be very short [29]. The average composition (wt.%) of the rock salt of the Asse salt mine is presented in TABLE 4-2. Also, experiments in very agressive brines were performed for several reasons: • to demonstrate the contribution to safety made by waste packages, a number of unrealistically conservative assumptions were forced to be made, a.o. the existence of a permanent contact between the metallic overpack and the most highly corrosive types of brine, even though studies show that safety-related contacts between brines and HLW will exist only for very short periods of time [29]. • salt brines in the disposal area may originate from the thermal migration of brine inclusions in the rock salt and must be considered in accident scenarios, e.g. brine inflow through an anhydrite layer [41-43,144]. The compositions, measured pH-values, and O2-content of the three test brines are given in TABLE 4-3.
68 TABLE 4-2. Average composition (in wt.%) of the rock salt of the Asse salt mine [42,43].
+ + 2+ 2+ - 2- Na K Ca Mg Cl SO4 H2O
38.3 0.33 0.17 0.16 58.02 2.47 0.1
TABLE 4-3. Composition and pH-value, and O2-content of the salt brines, representative for the Gorleben rock salt repository environment, used in the corrosion experiments [41,83,136,137,144-146].
Composition (wt. %) Brine pH [O2] (25°C) (mg/L) NaCl KCl MgCl2 MgSO4 CaCl2 CaSO4 K2SO4 H2O
11 1.4 4.7 26.8 1.4 - - - 65.7 4.37-4.62-6 0.84-6-1.07-2.83
2 0.31 0.11 33.03 - 2.25 0.005 - 64.3 4.12-6 0.64-6-1.53
3 25.9 - - 0.16 - 0.21 0.23 73.5 6.52-6-6.93 1.24-6-4.93
1 Q-brine. 2 Reference [41]. 3 Reference [144]. 4 Reference [136,137]. 5 Reference [83]. 6 Reference [146]. 7 Reference [145].
A.4.1.1 Laboratory-scale corrosion experiments
A.4.1.1.1 Screening experiments
The screening experiments were performed in MgCl2-rich brines (Q-brine with 26.8 wt.% MgCl2 and a brine containing 45 wt.% MgCl2) at 170°C. The tested materials (see TABLE 4-1) interacted with the brines for a total period of up to 400 days. The criterion for the selection of the materials as candidate packaging materials was their resistance to local corrosion, selective corrosion, and stress corrosion cracking (SCC). For this, single plane specimens, crevice corrosion specimens (contact between two plane specimens), and U-bent specimens (to investigate the susceptibility to SCC) were investigated for corrosion attacks by surface profilometry and metallography [136,137].
The corrosion results can be summarised as follows [136,137]:
• for the Ni-Resists, D2 and D4, and the nodular cast iron GGG 40.3, a heavy intergranular corrosion was observed.
• the Cr-Ni steels and the Ni-base alloys Inconel 625 and Incoloy 825 suffered from severe pitting/crevice corrosion. In case of the Cr-Ni steels, also stress corrosion cracking occured in addition.
69 • the passively corroded alloys Hastelloy C-4 and Ti 99.8-Pd as well as the actively corroded unalloyed steels, fine-grained steel TStE 355 and cast steel GS 16Mn5, were resistant to pitting/crevice corrosion and stress corrosion cracking.
On the basis of the results of these screening experiments, the passively corroding alloys Ti 99.8-Pd and Hastelloy C-4, and the actively corroding steels TStE 355, TStE 460, and 15 MnNi 6.3 were selected as the most promising candidate packaging materials and investigated into more detail (see section A.4.1.1.2) [136,137].
A.4.1.1.2 Experiments to evaluate the candidate metallic materials
A.4.1.1.2.1 Background
Detailed investigations were performed on the most promising candidate container materials to determine the influence of several important parameters on their long- term corrosion behaviour in disposal relevant brines. These parameters were: • brine composition. To examine the influence of the brine composition on the corrosion behaviour, MgCl2-rich brines (brine 1, Q-brine, and brine 2) and NaCl- rich brines (brine 3), were used. These brines were selected because of their relevance to rock-salt disposal environments [41]. The chemical composition of the tested brines is given in TABLE 4.3. • pH. The pH of the MgCl2-rich brine (Q-brine) was varied between 3 and 7, and that of the NaCl-rich brine between 1 and 10. The various pH values were adjusted by addition of HCl or NaOH to the brines. It is well-known that the pH of the corrosion medium can significantly influence the corrosion behaviour of metals by forming or dissolving the protective oxide layers on the metal surface. The pH of the disposal relevant brines (rock-salt) is affected by a number of factors such as salt impurities, temperature, hydrolysis, etc [147]. - • salt impurities. The influence of the salt impurity B(OH)4 , the main gamma - 3+ radiolytic products H2O2 and ClO , and the container corrosion product Fe on the steel corrosion behaviour was investigated. The concentrations of the various species varied between 10-1 mol/L and 10-3 mol/L. In addition, the influence of the thermally released salt impurity H2S on the corrosion behaviour of the fine-grained steel TStE 355 was studied. For this, H2S was added to the Q-brine as Na2S.9H2O at concentrations of 25 mg/L, 100 mg/L, and 200 mg/L. These values correspond to the amounts released from 10 cm, 20 cm, and 40 cm thick ring-shaped salt elements around an HLW borehole [41,142]. • temperature. Long-term immersion experiments were performed at 90°C, 150°C, 170°C, and 200°C [148]. The maximum temperature on the surface of the container depends on the waste loading, the interim storage period of the waste packages, and the disposal concept (e.g. distance between the HLW boreholes). For the disposal in rock salt, temperatures in the range of 150-200°C are considered in Germany. For this reason, experiments have been performed at 150°C, 170°C, and 200°C. In addition, experiments were performed at the temperature of 90°C since this is the temperature that would be established after about 500-1000 years, depending on the disposal concept. The temperatures of
70 90°C and 170°C were also the reference temperatures considered in the early joint European programmes [140]. • gamma dose rate. The influence exerted by gamma radiation from HLW on the corrosion behaviour of the candidate metallic materials was studied by imposing a gamma radiation field in the range of 1-103 Gy/h. The dose rate of 103 Gy/h corresponds to the gamma dose rate on the surface of a thin-walled (thickness of about 5 mm) HLW-canister. A gamma dose rate of 10 Gy/h is relevant to thick- walled steel containers, whose corrosion and mechanical allowances give a wall thickness of about 15 cm. Such investigations are essential because the radiolytic - - products formed by the effect of radiation on salt brines (e.g. H2O2, ClO , and ClO3 ) might influence the corrosion process. The experiments under gamma irradiation were performed in the spent fuel storage pool of KFA Jülich [148,149]. • welding. In addition to specimens made of the parent material, also Tungsten Inert Gas welded (TIG-welded), Electron Beam welded (EB-welded), Submerged- Arc Welded (SAW-welded), and Metal Active Gas welded (MAG-welded) specimens were also tested. These welding procedures are considered as potential closure techniques for the container lid [83,139,141,142,155]. The submerged arc welding technique (SAW) is under discussion as a potential sealing technique for the POLLUX spent fuel disposal container, where a 3 to 4 mm thick corrosion protection layer of Hastelloy C-4 is applied by surface welding to a mechanically stable disposal container made of carbon steel [29,41].
After removal from the brines, the specimens were examined for general and localised corrosion. The general corrosion rate was calculated from the weight losses and the material density. Localised corrosion was characterised using surface profilometry and metallography, and by measurements of pit depths using an inductive gauge [41,83,144,146]. In addition, the resistance of fine-grained steel to stress corrosion cracking (SCC) was investigated in disposal relevant salt brines. For this purpose, statically loaded U-bend specimens and the slow strain rate technique (SSRT) at various strain rates (10-4-10-7 s-1) and temperatures were applied [83,148,150,155,157,160]. For the fine-grained steel TStE 355 and the low-alloyed carbon steels TStE 460 and 15 MnNi 6.3, regression analyses of the corrosion data were performed to determine the time dependence of the general corrosion [83]. In the case of the passive alloys Ti 99.8-Pd and Hastelloy C-4, no statistical analysis of the general corrosion data was carried out because of the negligible corrosion rates measured in the brines [41].
71 A.4.1.1.2.2 Corrosion of unalloyed and low-alloyed steels in brines
A.4.1.1.2.2.1 General and localised corrosion
A.4.1.1.2.2.1.1 Influence of brine composition and temperature on steel corrosion in brines
Both the temperature and the brine composition were found to have a significant influence on the corrosion rate of the steels investigated, as illustrated in TABLES 4-4 and 4-5.
The localised corrosion bahaviour (pitting corrosion and intergranular attack) of the fine-grained carbon steel TStE 355, the cast steel GS 16Mn5, the nodular cast iron GGG 40.3, the Ni-Resist D2 and D4, and the Si-cast iron in the MgCl2-rich Q-brine for a maximum period of 775 days at various temperatures without radiation (V/S6 = 5 mL/cm²) is summarised in TABLE 4-4.
TABLE 4-4. Pitting and intergranular corrosion of various materials in the MgCl2-rich Q-brine for a maximum period of 775 days without gamma irradiation (V/S = 5 mL/cm²) [140].
(1) (2) Material T Maximum dpit,Max dIGA,Max (°C) exposure (µm) (µm) time (days)
90 285 n.p. n.IGA TStE 355 170 722 n.p. n.IGA 200 550 n.p. n.IGA n.p. 90 586 n.IGA n.p. 170 589 n.IGA GS 16Mn5 n.p. 200 556 n.IGA 1,300 90 465 n.IGA GGG 40.3 >1,300 170 162 n.IGA n.p. 90 775 450 Ni-Resist D2 100 170 589 200 n.p. 90 775 1,600 Ni-Resist D4 100 170 589 >1,600
90 446 300 500 Si-cast iron 170 448 450 800
(1) dpit,Max : maximum depth of pitting corrosion. (2) dIGA,MAX : maximum depth of intergranular attack (IGA). n.p. : no pitting corrosion. n.IGA : no intergranular attack.
6 V/S : brine volume to specimen surface area ratio.
72 The corrosion rates of the fine-grained steel TStE 355 and the low-alloyed carbon steels TStE 460 and 15 MnNi 6.3 in various test brines at different temperatures (V/S = 5 mL/cm²) are presented in TABLE 4-5.
TABLE 4-5. Corrosion rates of fine-grained steel TStE 355 and the low-alloyed carbon steels TStE 460 and 15 MnNi 6.3 in various test brines at different temperatures (V/S = 5 mL/cm²) [41,83,136,137].
(1) Material Corrosion T vCORR medium (°C) (µm/year)
90 69.7 ± 1.8
170 199.4 ± 15.0 brine 1 200 462.8 ± 33.6
90 37.6 ± 15.0 (2) TStE 355 brine 2 170 307.9 ± 45.3 200 651.2 ± 155.2
90 5.1 ± 2.2 brine 3 170 46.0 ± 7.3 200 18.3 ± 3.6
brine 1 150 203.5 TStE 460(3) brine 2 150 65.4 brine 3 150 56.3
brine 1 150 117.3 15 MnNi 6.3(3) brine 2 150 94.0 brine 3 150 71.3
Brines 1 and 2 : MgCl2-rich ; brine 3 : NaCl-rich. (1) VCORR : average uniform corrosion rate (calculated from the weight losses and extrapolated linearly). (2) The data for the fine-grained steel TStE 355 originates from long-term immersion experiments lasting up to 4 years. (3) The data for the low-alloyed carbon steels TStE 460 and 15 MnNi 6.3 originates from long-term immersion experiments lasting up to 18 months.
From TABLE 4-4, it can be seen that except for both the unalloyed steels TStE 355 and GS 16Mn5, all other Fe-base materials are susceptible to pitting or intergranular corrosion.
From TABLE 4-5, it can clearly be seen that an increase in temperature greatly accelerated the corrosion rate of all the materials investigated. This behaviour is because the materials were corroding by a thermally activated process and the pH of the test solutions decreased with rising temperature. Acid corrosion is considered to be the dominant process because at the temperatures investigated the measured pH-values of the Q-brine were all less than 4 (pH (90°C) = 3.6) [136,137,140,148].
Metallographic examinations and surface profiles of corroded specimens revealed that the fine-grained steel TStE 355, the cast steel GS 16Mn5, and the low-alloyed carbon steels TStE 460 and 15 MnNi 6.3 were resistant to pitting and crevice corrosion in all three brines. Under all tested conditions, these steels were subject to
73 non-uniform general corrosion, which is attributed to inhomogeneities in the steel. However, the maximum penetration depth of this uneven corrosion attack corresponded to the values of the average thickness reduction (the initial surface roughness was about 20 µm) [29]. Characteristic optical micrographs of various steels after exposure to the MgCl2-rich Q-brine (brine 1) at 170°C and 150°C are shown in FIGURES 4-1 to 4-4.
200 µm 500 µm
FIGURE 4-1. Optical micrograph of the fine- FIGURE 4-2. Optical micrograph of the cast grained steel TStE 355 showing steel GS 16Mn5 showing slight non- uniform corrosion after 720 days uniform corrosion after 590 days exposure to the Q-brine at 170°C exposure to the Q-brine at 170°C (magnification: 100×) [140,148]. (magnification: 50×) [140,148].
100 µm 500 µm FIGURE 4-3. Optical micrograph of the low- FIGURE 4-4. Optical micrograph of the low- alloyed steel TStE 460 showing alloyed carbon steel 15 MnNi 6.3 uniform corrosion after 18 months showing slight non-uniform corrosion exposure to the Q-brine at 150°C, after 18 months exposure to the Q- (magnifcation: 100×) [83,150]. brine at 150°C, (magnification: 50×) [83,150].
The nodular cast iron GGG 40.3, the two Ni-Resists D2 and D4, and the Si-cast iron suffered pitting and intergranular corrosion after extended exposure periods (see
74 TABLE 4-4). In the case of the Ni-Resists, an intermetallic phase consisting of Cr, Fe, and Ni, which had been precipitated at the grain boundaries, was destroyed by corrosion, whereas in the case of the nodular cast iron, the iron phase rather than the graphite phase was attacked because it is much less noble than carbon. These effects caused deep holes in the Ni-Resist and cast iron specimens.
From TABLE 4-5, it can be seen that the lowest corrosion rates of the steels occurred in the NaCl-rich brine 3. In the MgCl2-rich brines 1 (Q-brine) and 2, significantly higher corrosion rates were obtained. The higher corrosivity of the MgCl2-rich brines in comparison to the NaCl-rich brine is attributed to its higher chloride concentration and to the presence of Mg2+. It appears that Mg2+, which is incorporated into the ferrous hydroxide layer replacing Fe2+, interferes with the normally expected conversion of Fe(OH)2 to Fe3O4 [83]. The corrosion product formed (Fe, Mg)(OH)2 appears to have little or no ability to protect the carbon steel from corrosion in contrast to Fe3O4, which is formed in the NaCl-rich brine. The acceleration of the steel corrosion in brines containing high amounts of MgCl2 is in good agreement with the results reported by Westerman et al. [151].
The corrosion products formed on the surface of the steel specimens were analysed by X-Ray Diffraction (XRD). The brine composition and the temperature influenced the nature of the corrosion products, as follows (all at 90°C): brine 1, α-Fe2O3 (hematite) and Fe3O4 (magnetite); brine 2, β-FeOOH (akaganeite); brine 3, Fe3O4 and β-FeOOH. At 170°C, a magnesium-containing ferrous hydroxide, (Fe,Mg)(OH)2 (amakinite), was identified as the primary corrosion product (without evidence of Fe3O4) in the MgCl2-rich brines 1 and 2, while the corrosion product in the NaCl-rich brine 3 was Fe3O4 [41,83,136,137,144]. The formation of (Fe,Mg)(OH)2 and Fe3O4 at high temperatures in MgCl2-rich and NaCl-rich brines, respectively, is also reported by Westerman et al. [151].
Linear regression results of the thickness reduction data
To determine the time dependence of the uniform corrosion, regression analyses of the thickness reduction data were performed. From the regression analyses, it was found that the following equations could be used to fit the experimental data [41,83,144]:
∆S = A + B ⋅t (19a)
∆S = A + B ⋅t + C ⋅t 2 (19b)
∆S = A + B ⋅t N (19c) where A, B, C, and N are the regression coefficients, t the time and ∆S the computed thickness reduction.
Parameter A can be interpreted as an indicator of an initial corrosion process that is influenced by the preparation of the specimens. This parameter is of minor importance for the long-term thickness reduction of the specimens. Parameter B gives a measure of the linear corrosion rate, whereas the other two coefficients (C,N)
75 can be taken as an indication of the formation of protective corrosion layers on the specimen surfaces [41,83,144]. The regression analyses of the time-dependent thickness reduction of the investigated steels in the brines at various temperatures show that a linear equation fits the experimental data well, under all conditions. The results of the linear regression analyses of the thickness reduction of the fine-grained steel TStE 355 and the low-alloyed carbon steels TStE 460 and 15 MnNi6.3 in the three brines at 90°C, 150°C, and 170°C (without irradiation) are compiled in TABLE 4-6. As an example, the experimental data and the linear regression lines of the three above-mentioned steels in the three brines at 150°C are plotted in FIGURE 4-5. From FIGURE 4-5, it can clearly be seen that the thickness reduction in the brines increases linearly with exposure time. The slope of the regression line represents the corrosion rate.
TABLE 4-6. Results of linear regression analyses for the thickness reduction (∆S) of the TStE 355, TStE 460, and 15MnNi 6.3 steels in various brines (without irradiation) at 90°C, 150°C, and 170°C with different brine volume to surface ratios (V/S = 2 and 5 mL/cm²) [36,41,83,144,146,149].
Steel Corrosion T V/S(2) N° of A Standard B Standard medium(1) (°C) (mL/cm²) spe- (µm) error of A (µm/year) error of B cimens (µm) (µm/year)
90 5 56 -2.2 2.9 69.7 1.8 1 150 5 - -3.58 11.75 226.8 16.62 170 5 16 24.5 8.4 199.4 15.0 170 2 16 8.0 2.5 47.1 2.5
90 5 13 5.0 10.8 37.6 15.0 TStE 355 2 150 5 - 45.5 11.72 183.8 18.63 170 5 7 -3.4 30.2 307.9 45.3 170 2 20 20.9 9.7 119.6 10.6
90 5 18 0.21 1.38 5.1 2.2 3 150 5 - 8.1 1.74 35.7 4.46 170 5 12 0.18 4.06 46.0 7.3 170 2 16 1.3 1.0 15.3 1.0
1 150 5 12 -0.48 5.69 203.5 5.77 TStE 460 2 150 5 12 85.2 8.56 65.4 8.68 3 150 5 12 -5.94 4.34 56.3 4.4
1 150 5 12 47.5 9.25 117.3 9.4 15MnNi6.3 2 150 5 12 57.8 6.86 94.0 6.96 3 150 5 12 -3.9 1.64 71.3 1.66
Regression equation : ∆S = A+B·t. The slope of the regression line (B) represents the corrosion rate. (1) Brines 1 and 2: MgCl2-rich ; brine 3: NaCl-rich. (2) V/S : brine volume to specimen surface area ratio.
76
FIGURE 4-5. Thickness reduction of the TStE 355, TStE 460, and 15 MnNi 6.3 steel grades vs. exposure time in the salt brines 1, 2, and 3 at 150°C [29,41].
A.4.1.1.2.2.1.2 Influence of pH on steel corrosion in brines
The influence of the pH of the brines on the corrosion behaviour of the fine-grained carbon steel TStE 355 was evaluated by means of long-term immersion experiments, lasting up to 1 year at 170°C. For the investigations, the pH of the MgCl2-rich brine 1 (‘Q-brine’) was varied between 3 and 7 and that of the NaCl-rich brine between 1 and 10. The various pH-values were adjusted by addition of HCl or NaOH to the brines. The brine volume to specimen surface ratio (V/S) was 2 ml/cm² [68,147,152].
77
In both brines and at all pH-values, the steel was resistant to pitting corrosion in the sense of an active-passive corrosion element (the carbon steel corroded only in the active state and no passivation occured; therefore, pitting, as is possible only in passively corroding materials, did not occur). Under the test conditions, non-uniform general corrosion was observed, and the thickness reduction of the specimens increased linearly with time. FIGURES 4-6 and 4-7 exemplify the linear time- dependence of the thickness reduction of the steel specimens in the two brines at 170°C and pH=5 [68,147,152].
FIGURE 4-6. Time-dependence of the thick- FIGURE 4-7. Time-dependence of the thick- ness reduction of the TStE 355 steel in ness reduction of the TStE 355 steel in the MgCl2-rich Q-brine at 170°C and the NaCl-rich brine at 170°C and pH=5 pH=5 [68,147,152]. [68,147,152].
The integral corrosion rates, calculated from the weight losses and the density of the steel, as a function of the initial pH-values of the brines, are shown in FIGURES 4-8 and 4-9.
FIGURE 4-8. Corrosion rate of the TStE 355 FIGURE 4-9. Corrosion rate of the TStE 355 steel as a function of pH in Q-brine at steel as a function of pH in the NaCl- 170°C [68,147,152]. rich brine at 170°C [68,147,152].
The data show that the initial pH values of the MgCl2-rich and NaCl-rich brines did not significantly affect the corrosion rate of the TStE 355 carbon steel at 170°C in the
78 pH-ranges of 3-7 and 1-5, respectively. In these pH ranges, the values of the corrosion rate (177-209 µm/year in the Q-brine and 42-50 µm/year in the NaCl-rich brine) differed between each other by at most about 20%, which is within the statistical variations of the measured values. At higher pHs in the NaCl-rich brine (pH=6-10), the corrosion rate of the steel clearly decreased to 26-28 µm/year. This observation may be explained by the formation of a very dense corrosion protection layer on the metal surface, which was observed during the metallographic examinations [68,147,152].
After termination of the experiments (under the conditions considered in this study), all the brines reached nearly the same pH-value irrespective of their initial pH (4.5-4.8 in the MgCl2-rich Q-brine and 5.4-5.7 in the NaCl-rich brine). This observation may be explained by the existence of a buffering capacity provided by the reaction of the corrosion products with the brine constituents. The experimental conditions (brine volume to specimen surface area ratio) were chosen as such that the H+ ions were not consumed due to the corrosion reactions [68,147,152].
A.4.1.1.2.2.1.3 Influence of chemical species on steel corrosion in brines
The influence of selected chemical species on the steel corrosion behaviour was examined by means of long-term immersion experiments, lasting up to 520 days, in the MgCl2-rich brine 1 (‘Q-brine’) and the NaCl-rich brine 3 at 90°C and 170°C. The following species were tested [68,152]: - • the salt impurity B(OH)4 , - • the main gamma radiolytic products of the brines, H2O2 and ClO , and • Fe3+, which can be generated by oxidation of the container corrosion product Fe2+, as a result of radiolysis or ingress of oxygen into the disposal area. The concentrations of the species added to the brines are given in TABLE 4-7. Both the individual and the synergistic effects of the chemical species on corrosion were investigated. The concentrations of the various species in the brines varied between 10-1 mol/L and 10-3 mol/L [68,152].
TABLE 4-7. Chemical species examined in the corrosion studies on TStE 355 steel in salt brines [68,152].
Species Added to the brine as Concentration (mol/L)
- -1 B(OH)4 H3BO3 1.4·10 3+ -2 Fe FeCl3·6H2O 3.5·10 -3 -2 H2O2 15% H2O2-solution 10 ; 10 ClO- 15% NaClO-solution 10-3
79
The integral corrosion rates of the fine-grained carbon steel TStE 355 over a test period of up to 1 year in the MgCl2-rich ‘Q-brine’ at 90°C and 170°C are shown in FIGURE 4-10. An increase in temperature from 90°C to 170°C caused a significant increase of the corrosion rates of the steel in all brines, especially in the pure brine (224 µm/year at 170°C compared to 70 µm/year at 90°C). At 90°C, the corrosion attack was non- uniform in all brines. Pitting corrosion in the sense of an active-passive transition was not observed. However, at 170°C, after long exposure periods to the brines containing added chemical species, localised corrosion was observed on the steel specimens. The maximum depth of the pits after 1 year’s exposure to the brines was up to 500 µm which, is clearly a higher depth of attack than the non-uniform corrosion of the specimens in the pure brine at 170°C (about 200 µm after 1 year) [68,152]. The addition of the chemical species to the MgCl2-rich Q-brine did not noticeably increase the corrosion rate of the steel, neither at 90°C nor at 170°C [68,152]: • at 90°C, the corrosion rates in the brines containing extra chemical species were 90-123 µm/year and were therefore only a factor of about 1.3-1.8 higher than in the pure brine (70 µm/year). • at 170°C, the corrosion rate was slightly increased to 254 µm/year in the presence of all species, compared to 224 µm/year in the pure brine. In fact, in the brines - 3+ containing B(OH)4 or Fe , the corrosion rates (15 µm/year and 155 µm/year, respectively) were clearly lower than in the pure brine. This behaviour is attributed to the formation of very dense protective layers on the specimens.
The integral corrosion rates of the fine-grained carbon steel TStE 355 over a test period of up to 520 days in the NaCl-rich brine 3 at 90°C and 170°C are shown in FIGURE 4-11. The addition of the chemical species to the NaCl-rich brine 3 increased the corrosion rate of the steel only at 90°C, whereas at 170°C, the corrosion rate was not significantly affected by the presence of the chemical species [68,152]: • at 90°C, the corrosion rate in the pure NaCl-rich brine was very small (5 µm/year). The very high corrosion rates in the brines containing all the chemical species simultaneously clearly indicates a synergistic effect. • at 170°C, the maximum corrosion rate occurred in the brine containing all species simultaneously (73 µm/year), which is only a factor of about 1.6 higher than those in the pure brine. After long exposure periods, the corrosion layer formed in the brine containing all species broke down locally, and non-uniform corrosion with a maximum penetration depth of 130 µm was observed after 1 year exposure time. In all other brines and at both test temperatures, the corrosion of the steel was nearly uniform.
80 FIGURE 4-11. Corrosion rates of TStE 355 steel in the NaCl-rich brine 3 with and FIGURE 4-10. Corrosion rates of TStE 355 without additions of chemical species steel in the MgCl2-rich Q-brine with at 90°C and 170°C [68,152]. and without additions of chemical species at 90°C and 170°C [68,152].
- - 3+ In addition to B(OH)4 , H2O2, ClO , and Fe , the possible influence of the thermally released salt impurity H2S on the corrosion behaviour of the fine-grained steel TStE 355 in the MgCl2-rich Q-brine (brine 1) and the NaCl-rich brine 3 was studied. The sulfide (0-200 mg/L H2S) was added to the brine as Na2S·9H2O.
It was found that sulfides do not influence noticeably the general corrosion behaviour of the steel TStE 355 (in the range 0-200 mg/L H2S) and that the material remained resistant to pitting corrosion, in the MgCl2-rich Q-brine as well as in the NaCl-rich brine 3. Both in the H2S-free and in the H2S-containing salt brines, the steel underwent non-uniform corrosion with approximately the same depth of attack (the maximum rates of penetration of deeper corrosion zones were only slightly higher than the general corrosion rates calculated from the mass losses) [41,142]. FIGURE 4-12 shows the non-uniform corrosion attack of the steel TStE 355 after 325 days of exposure to the pure Q-brine and a 200 mg/L H2S-containing Q-brine at 170°C.
81
100 µm 100 µm (a) (b)
FIGURE 4-12. Optical micrographs of the fine-grained carbon steel TStE 355 after 325 days
immersion in the MgCl2-rich Q-brine at 170°C [142]. (a) without H2S-addition (magnification: 200×). (b) with 200 mg/L H2S (magnification: 200×).
The results of the linear regression analyses of the thickness reduction of the fine- grained steel TStE 355 in the H2S-containing MgCl2-rich Q-brine at 170°C are compiled in TABLE 4-8. The thickness reduction of the fine-grained steel TStE 355 in -4 the NaCl-rich brine 3 containing 6×10 M Na2S (i.e. 200 mg/L H2S) at 150°C is presented in FIGURE 4-13.
TABLE 4-8. Results of linear regression analyses for the thickness reduction (∆S) of the fine- grained steel TStE 355 in the MgCl2-rich Q-brine with various contents of H2S (0- 200 mg/L) at 170°C (V/S = 5 mL/cm²) [83,144].
[H2S] A Stand. B Stand. (mg/L) (µm) error of A (µm/y) error of B (µm) (µm/year) FIGURE 4-13. Thickness reduction of the fine-grained steel TStE 355 in 0 24.5 8.4 199.4 15.0 the NaCl-rich brine 3 with and 25 42.8 4.9 192.5 8.0 without Na2S at 150°C [41]. 100 22.4 24.9 317.4 40.9 200 44.9 14.0 203.0 23.0
Regression equation : ∆S = A+B·t. The slope of the regression line (B) represents the corrosion rate. Test duration : 325 days.
82 As in the Na2S-free brine, a linear decrease in the thickness of the steel with exposure time was observed (see FIGURE 4-13) [41,144]: • in the MgCl2-rich Q-brine (170°C) (TABLE 4-8), the linear corrosion rates in the brines containing 25 and 200 mg/L H2S (193-203 µm/year) were in the same range as for the H2S-free brine (199 µm/year). The same was observed for the specimens tested in the brine containing 100 mg/L H2S up to 240 days. For longer exposure periods (up to 325 days), however, an unusally high thickness reduction was determined (317 µm/year) [144]. • in the NaCl-rich brine 3 (150°C) (FIGURE 4-13), the linear corrosion rate of 44 µm/year in the Na2S-containing (200 mg/L H2S) brine was only slightly higher than in the Na2S-free brine (35.7 µm/year) [41].
X-Ray Diffraction (XRD) analyses revealed that the corrosion products formed in the Na2S-containing brines were Fe3O4 and γ-Fe2O3, as in the Na2S-free brines [41].
A.4.1.1.2.2.1.4 Influence of gamma irradiation on steel corrosion in brines
The influence of gamma irradiation on the corrosion behaviour of TStE 355 steel was investigated at 90°C (in the MgCl2-rich Q-brine) and at 150°C (in the three test brines).
A.4.1.1.2.2.1.4.1 Investigations at 90°C (in the MgCl2-rich Q-brine)
The results of the investigations under irradiation (1-103 Gy/h) show that gamma dose rates of 1-102 Gy/h, corresponding to the value on the surface of a thick-walled packaging, did not affect the corrosion rate of the steels. On the contrary, at the very high dose rates of 103 Gy/h, which corresponds to a 5-mm thin-walled packaging, exerted a strong influence on the corrosion behaviour of the investigated steels [136,137,146]. The corrosion results (general and pitting corrosion) of the fine- grained carbon steel TStE 355, the cast steel GS 16Mn5, the nodular cast iron GGG 40.3, the Ni-Resists D2 and D4, and the Si-cast iron in the MgCl2-rich Q-brine (V/S = 5 mL/cm²) for a maximum period of 244 days at 90°C are summarised in TABLE 4-9. The corrosion rates for the fine-grained carbon steel TStE 355 obtained 3 in the MgCl2-rich Q-brine at 90°C and gamma dose rates of 1-10 Gy/h are plotted in FIGURE 4-14.
At a gamma dose rate of 103 Gy/h, all investigated steels suffered from weight losses which were considerably higher than in the absence of radiation (see TABLE 4-5 for comparison). The corrosion rates, calculated from these weight losses, increased under irradiation by a factor of about 4 for the nodular cast iron GGG 40.3, by a factor 10-20 for the fine-grained carbon steel TStE 355, the cast steel GS 16Mn5, and the Si-cast iron, and by a factor 40-60 for the two Ni-Resists D2 and D4. The calculated general corrosion rates have to be considered as average values because the steels were subjected to non-uniform corrosion accompanied by shallow pit formation. However, the maximum depth of attack of shallow pit formation was only slightly higher than the average general corrosion rate.
It was concluded that the effect of radiation on the corrosion of the investigated steels was strongly dependent on the concentration of radiolytic products in the brines.
83 - Strong oxidizing agents, such as H2O2 and ClO3 , are considered to be mainly responsible for the considerable increase in corrosion rate since they are known to act as cathodic depolarizers [146,153].
TABLE 4-9. General and pitting corrosion of the investigated steels in the MgCl2-rich Q- brine at 90°C and a gamma dose rate of 103 Gy/h [140].
(1) (2) (3) Materials tMax vCORR dpit,Max (days) (µm/year) (µm)
TStE 355 244 464.2 (4) GS 16Mn5 231 665.8 (4) GGG 40.3 173 165.2 220 Ni-Resist D2 160 157.4 160 Ni-Resist D4 160 77.1 200 FIGURE 4-14. Corrosion rates of TStE 355 in Si-cast iron 153 55.8 250 the MgCl2-rich Q-brine at 90°C with and without gamma radiation [38,136,137].
(1) Maximum exposure period. (2) Average uniform corrosion rate. (3) Maximum pit depth. (4) non-uniform corrosion with shallow pit formation.
A.4.1.1.2.2.1.4.2 Investigations at 150°C (in the three test brines 1, 2, and 3)
At 150°C, as for 90°C, the general corrosion of the fine-grained carbon steel TStE 355 increased linearly with the exposure time both in the unirradiated and irradiated brines. The results of the linear regression analyses for the general corrosion of TStE 355 in the three brines, expressed as thickness reduction, are compiled in TABLE 4-10.
84 TABLE 4-10. Results of linear regression analyses for the thickness reduction (∆S) of the fine-grained steel TStE 355 in irradiated and unirradiated brines at 150°C (V/S = 2 mL/cm²) [36,83,146,149].
Corrosion D(1) Number of A Standard B Standard medium (Gy/h) specimens (µm) error of A(2) (µm/year) error of B(2) (µm) (µm/year)
brine 1 0 16 8.0 2.5 47.1 2.5 10 19 34.1 10.6 72.6 11.0
brine 2 0 20 20.9 9.7 119.6 10.6 10 15 23.6 17.5 162.4 20.4
brine 3 0 16 1.3 1.0 15.3 1.0 10 18 1.6 1.1 13.5 1.8
Regression equation : ∆S = A+B·t. The slope of the regression line (B) represents the corrosion rate. Brines 1 and 2 : MgCl2-rich ; brine 3 : NaCl-rich. (1) D : gamma dose rate (Gy/h). (2) the standard error is based on the 95% confidence interval.
The imposition of a 10 Gy/h radiation field on the 150°C brine environment increased the linear corrosion rates of the steel specimens in the MgCl2-rich brines 1 and 2 from about 47 µm/year to 72 µm/year, and from 120 µm/year to 163 µm/year, respectively. In the NaCl-rich brine 3, the corrosion rate in the 10 Gy/h irradiated environment (13.5 µm/year) was very close to the value obtained in the uniradiated system (15.3 µm/year) [36,83,146,149].
FIGURE 4-15 shows the time-dependence of the thickness reduction of the fine- grained carbon steel TStE 355 in unirradiated and irradiated (10 Gy/h) brines (V/S = 2 mL/cm²) at 150°C. These results show a linear increase of the thickness reduction with exposure time.
85
(a) (b)
(c)
FIGURE 4-15. Thickness reduction of the fine-grained carbon steel TStE 355 at 150°C (V/S = 2 mL/cm²) with and without gamma radiation [146]. (a) in the MgCl2-rich brine 1 (‘Q-brine’). (b) in the MgCl2-rich brine 2. (c) In the NaCl-rich brine 3.
The fine-grained carbon steel TStE 355 is resistant to pitting corrosion in both the unirradiated and irradiated brines. Furthermore, a non-uniform general corrosion was observed in all brines. FIGURE 4-16 shows optical micrographs of TStE 355 after 166 days of exposure to the test brines at 150°C and 10 Gy/h [146].
86 (a) 100 µm
(b) 100 µm
(c) 100 µm
FIGURE 4-16. Optical micrographs of the fine-grained carbon steel TStE 355 after 166 days exposure to various tested brines at 150°C and 10 Gy/h [83,143,146]. (a) MgCl2-rich brine no. 1 (magnification: 200×). (b) MgCl2-rich brine no. 2 (magnification: 200×). (c) NaCl-rich brine no. 3 (magnification: 200×).
87 A.4.1.1.2.2.1.5 Influence of welding on steel corrosion in brines
The influence of Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding, and Submerged Arc Welding (SAW) technique, as potential container closure techniques, on the general and local corrosion behaviour of various steels was studied by means of long-term immersion experiments in the MgCl2-rich brines 1 and 2 and in the NaCl- rich brine 3 [83,154]:
• the general corrosion rates and the local corrosion results obtained for unwelded and welded (TIG- and EB-welding) TStE 355 steel specimens in the MgCl2-rich Q- brine at 150°C are compiled in TABLE 4-11. Also, the influence of a stress relief thermal treatment (2 hours at 600°C) and gamma irradiation of 10 Gy/h was investigated under these conditions [154]. • the influence of welding (TIG- and EB-welding) and heat treatment on the general and local corrosion of the fine-grained carbon steel TStE 355 in the NaCl-rich brine 3 at 150°C is presented in TABLE 4-12. The experiments in the NaCl-rich brine 3 were conducted without gamma radiation because previous investigations have shown that in this brine a gamma dose rate of 10 Gy/h does not affect the corrosion behaviour of the steel [83,154]. • the influence of submerged arc welding (SAW) on the general and local corrosion of TStE 460 and 15 MnNi 6.3 in the three test brines 1, 2, and 3 at 150°C is shown in TABLE 4-13.
TABLE 4-11. Influence of welding and heat treatment on the corrosion of the fine-grained carbon steel TStE 355 in the MgCl2-rich Q-brine at 150°C (without gamma irradiation and with a gamma dose rate of 10 Gy/h) [68,152,154].
(1) (2) (3) (4) Material condition ∆t ∆γ vCORR dMax (days) (Gy/h) (µm/year) (µm) base material weld HAZ(5)
unwelded 585 0 55.1 ± 6.0 60 - - TIG-welded(6) 306 0 173.0 ± 16.0 80 300 500 EB-welded(7) 306 0 171.8 ± 5.0 30 740 650
unwelded 550 10 72.6 ± 11.0 60 - - TIG-welded(6) 300 10 65.8 ± 40.0 180 1,350 1,250 EB-welded(7) 300 10 67.7 ± 41.0 90 1,200 1,500
TIG-welded + HT(8) 480 10 112.0 ± 49.0 100 100 120 EB-welded + HT(8) 480 10 94.7 ± 31.0 80 100 90
(1) ∆t : test time. (2) ∆γ : dose rate. (3) VCORR : average uniform corrosion rate. This is an intergral corrosion rate, calculated from the slope of the thickness reduction/exposure time curves. (4) dMax : maximum depth of local attack. (5) HAZ : Heat Affected Zone. (6) TIG : Tungsten Inert Gas weld. (7) EB : Electron Beam weld. (8) HT : Heat Treatment (the welded specimens were stress-relief thermal treated for 2 hours at 600°C).
88 TABLE 4-12. Influence of welding and heat treatment on the corrosion of the fine-grained carbon steel TStE 355 in the NaCl-rich brine 3 at 150°C (without gamma irradiation) [154].
(1) (2) (3) Material condition ∆t vCORR dMax (days) (µm/year) (µm) base material weld HAZ(4)
unwelded 585 15.3 ± 1.0 n.p. n.p. n.p. TIG-welded(5) 506 19.1 ± 1.1 n.p. n.p. n.p. EB-welded(6) 506 18.6 ± 0.7 n.p. n.p. n.p.
TIG-welded + HT(7) 463 22.6 ± 6.8 n.p. n.p. n.p. EB-welded + HT(7) 463 25.9 ± 7.0 n.p. n.p. n.p.
(1) ∆t : test time. (2) vCORR : average uniform corrosion rate. This is an intergal corrosion rate, calculated from the slope of the thickness reduction/exposure time curves. (3) dMax : maximum depth of local attack. (4) HAZ : Heat Affected Zone. (5) TIG : Tungsten Inert Gas weld. (6) EB : Electron Beam weld. (7) HT : Heat Treatment (the welded specimens were stress-relief thermal treated for 2 hours at 600°C).
TABLE 4-13. Average corrosion rates and maximum penetration depth (pitting attack) in the HAZ of submerged arc (SAW) welded TStE 460 and 15 MnNi 6.3 specimens after a maximum period of exposure of 18 months to the three brines 1, 2, and 3 [83,150,155].
(1) (2) (3) Material Corrosion ∆t vCORR dMax(HAZ) medium (months) (µm/year) (mm)
12 175.2 1.5 brine 1 18 203.5 2.0
12 90.0 2.2 brine 2 TStE 460 18 65.4 2.5
12 29.0 0.05 brine 3 18 56.3 0.04
12 127.5 1.8 brine 1 18 117.3 1.9
12 112.0 3.9 brine 2 15 MnNi 6.3 18 94.0 4.0
12 72.0 0.08 brine 3 18 71.3 0.05
brines 1 and 2 : MgCl2-rich brines ; brine 3 : NaCl-rich brine. (1) ∆t : Maximum exposure time. (2) VCORR : average uniform corrosion rate( of unwelded specimens). (3) dMAX(HAZ) : Maximum penetration depth (pitting attack) in the Heat Affected Zone (HAZ) of SAW-welded specimens.
89 In general, it can be stated that the TIG- and EB-welded specimens show a very similar corrosion behaviour in the MgCl2-rich Q-brine (TABLE 4-11). However, they exhibit a significantly lower corrosion resistance, especially to local corrosion, than the base material [68,152,154]:
• in the absence of gamma irradiation, the general corrosion rate of the TIG- and EB-welded specimens (171-173 µm/year) increased by a factor of about 3 compared to the unwelded specimens (55 µm/year). In the presence of 10 Gy/h gamma radiation, however, the general corrosion rates of the TIG- and EB-welded specimens (66-68 µm/year) are very close to the value of the unwelded specimens (73 µm/year). • for the unwelded specimens, a non-uniform general corrosion was observed. The welded specimens, on the other hand, suffered from severe local corrosion attack in the welds and the Heat Affected Zones (HAZ), which was enhanced by the presence of gamma irradiation. The maximum depth of such local attack reached values up to 500-700 µm (after 306 days) and 1,350-1,500 µm (after 300 days) in the absence and presence of gamma irradiation, respectively.
The application of a thermal heat treatment (2 hours at 600°C) eliminated the severe local corrosion attack in the welds and HAZ (TABLE 4-11): the thermally treated TIG- and EB-welded specimens showed non-uniform corrosion as was the case for the unwelded specimens (the maximum corrosion depth of 90-120 µm was of the same order of magnitude as the average corrosion rate). The corrosion behaviour of the thermally treated welded specimens was similar to that of the parent material [154].
The corrosion behaviour of the TIG- and EB-welded specimens exposed to the NaCl- rich brine is among one another very similar (as in the MgCl2-rich Q-brine) (TABLE 4-12): the corrosion rates of the welded specimens (18.6-19.1 µm/year) are very close to the value of the unwelded specimens (15.3 µm/year). The corrosion rates of the heat treated welded specimens (22.6-25.9 µm/year) are only slightly higher than the values of the unwelded (15.3 µm/year) and only-welded (19.1 µm/year) specimens. In a NaCl-rich brine, a heat treatment of the welds of steel containers is therefore not efficient because it does not improve its corrosion resistance [154]. The metallographic examinations of corroded specimens show that corrosion of the welded and heat treated specimens was uniform, as in the case of the unwelded specimens [154].
In the NaCl-rich brine 3, no influence of submerged arc welding (SAW) on the corrosion behaviour of the steels TStE 460 and 15 MnNi 6.3 was observed, even after 18 months of exposure (TABLE 4-13). The SAW specimens underwent a non- uniform corrosion attack as did the unwelded specimens, and the general corrosion rates corresponded to those obtained for the parent materials [29,83]. In the MgCl2-rich brines, on the other hand, a pronounced local corrosion attack was observed in the heat affected zones (HAZ) of the welded specimens. The depth of pitting attack was found to increase with the length of immersion, attaining values between 2 and 4 mm after 18 months. Therefore, to improve the corrosion resistance of the welded material, a heat treatment of the welds is recommended [29,83,160].
90 A.4.1.1.2.2.2 Stress corrosion cracking (SCC)
Long-term immersion tests (up to 4 years) and the slow strain rate technique (SSRT) have been used to investigate the susceptibility of several unalloyed and low-alloyed steels (TStE 355, TStE 460 and 15 MnNi 6.3) to stress corrosion cracking (SCC) in the MgCl2-rich Q-brine and the NaCl-rich brine (brine 3) at various temperatures (25°C, 90°C and 170°C). The immersion tests were performed using statically loaded U-bend specimens. The SSRT experiments were carried out at strain rates ranging from 10-4 s-1 to 10-7 s-1 using round tensile specimens. To facilitate the interpretation of the results from the SSRT experiments, additional investigations were carried out in argon (Ar), as an inert reference medium, for comparison.
Besides the parent material, welded specimens simulating a possible container closure technique were also tested. The considered welding techniques were Metal Active Gas Welding (MAGW), Electron Beam Welding (EBW) and Flux Cored Arc Welding (FCAW).
A.4.1.1.2.2.2.1 MgCl2-rich ‘Q’-brine
The results of the SSRT experiments for the unwelded TStE 355, TStE 460 and 15 MnNi 6.3 steel specimens in argon and Q-brine revealed that all steels investigated suffered a loss of ductility, which was found to be greater at 170°C than that observed at 90°C and 25°C. At 25°C, a very slight drop of the mechanical properties was manifested only at the slowest strain rate (10-7 s-1). At higher temperatures (90°C and 170°C), a clear decrease in the elongation, reduction of area, energy and true stress at fracture for all steels was observed in Q-brine, compared to the values in argon [36,83,136,137,161,162]. FIGURE 4-17 shows the results of the slow strain rate tests for the unwelded TStE 355, TStE 460 and 15 MnNi 6.3 steel specimens in argon (A) and Q-brine (Q) at 170°C and various strain rates (10-4-10-7 s-1).
Metallographic examinations showed that the hot-rolled and annealed unalloyed TStE 355 steel was resistant to SCC in the MgCl2-rich Q-brine up to 170°C. For the hot-rolled and annealed low-alloyed TStE 460 steel, no sensitivity to SCC was found at 25°C and 90°C, but at 170°C and a strain rate of 10-5 s-1, a slight susceptibility to SCC was observed (a minor degree of secondary cracking was encountered). For the forged and annealed low-alloyed 15 MnNi 6.3 steel, extensive secondary cracks were observed at 90°C (strain rate of 10-7 s-1) and at 170°C (strain rate of 10-5 s-1), indicating a high susceptibility to SCC in the Q-brine under these conditions [36,83,161,162]. Extensive lateral secondary cracking of 15 MnNi 6.3 steel in Q-brine is shown in FIGURE 4-18.
SEM examinations of the fracture surface showed a change from a fully ductile fracture (dimples), when specimens were tested in argon, to a more brittle one when tests were performed in Q-brine. The brittle nature of the fracture surface became greater as the strain rate was slower and with increasing temperature [36,83,161,162]. FIGURE 4-19 shows brittle features on the fracture surface of a TStE 355 specimen tested in Q-brine at 90°C and a strain rate of 10-7 s-1.
91 The welding procedures were not found to have any effect on the SCC behaviour of the tested materials. The fracture of the welded specimens was always located in the parent material [161,162].
FIGURE 4-17. Elongation, reduction of area, energy, yield strength, maximum load and true stress at fracture versus strain rate for the TStE 355, TStE 460 and 15 MnNi 6.3 steel grades tested at 170°C in argon (A) and the MgCl2-rich Q- brine (Q) [36].
92
(a)
(b)
FIGURE 4-18. Evidences of extensive lateral secondary cracking of the 15 MnNi 6.3 steel tested in Q-brine. (a) at 170°C and a strain rate of 10-5 s-1 [83]. (b) at 90°C and a strain rate of 10-7 s-1 (optical micrograph, magnification: 12×) [162].
FIGURE 4-19. SEM micrograph of the fracture surface of a MAGW welded TStE 355 steel specimen tested in Q-brine at at 90°C and a strain rate of 10-7 s-1 [162].
A.4.1.1.2.2.2.2 NaCl-rich brine ‘3’
The results of the SSRT experiments for the parent and welded (EBW and FCAW) TStE 355 steel specimens tested in argon (A) and in the NaCl-rich brine at 170°C and at various strain rates (10-4-10-7 s-1) are shown in FIGURE 4-20. Compared to the values in argon, there was a clear decrease of the elongation and reduction of area parameters when tested in the NaCl brine. The drop of these parameters was not significant at the highest strain rate (10-4 s-1), but it was important at the other
93 lower strain rates used in the tests. The values of the yield strength and maximum load parameters did not show important changes between the two media, for both parent and welded specimens [68,152,163,164].
Metallographic examinations showed a slight sensitivity to SCC of TStE 355 under all test conditions: secondary cracks, with a maximum crack depth of 130 µm, were observed when tested in NaCl brine at the slowest strain rate [68,152].
SEM examinations of the fracture surface showed a change from a fully ductile fracture surface with dimples formation (see FIGURE 4-20(a)) for specimens tested in argon to a more brittle fracture mode (see FIGURE 4-20(b)) when the tests were performed in NaCl brine. The brittle nature of the fracture surface became greater as the strain rate was slower [68,152,163,164].
FIGURE 4-20. Elongation, reduction of area, yield strength and maximum load values versus strain rate for the parent and welded TStE 355 steel tested at 170°C in argon (A) and NaCl-rich brine [68,152].
94
(a) (b)
FIGURE 4-21. SEM micrographs of the fracture surface of a TStE 355 steel specimen tested at 90°C and a strain rate of 10-7 s-1 [163]. (a) in argon. (b) in NaCl.
For both test environments (MgCl2 ‘Q’-brine and NaCl brine), the loss of ductility, mainly noticed in the elongation and reduction of area parameters and in a change from ductile fracture surface to a more brittle one, was thought to be attributed to the embrittling effect of the hydrogen produced on the specimens surface during the tests by the following reactions (due to the general corrosion phenomena) [36,83,162]:
Fe → Fe 2+ + 2e − (anodic reaction) 2H + + 2e − → 2H (cathodic reaction)
The hydrogen produced enters the material, mainly through the highly stressed zones of the specimens, and interacts with the microstructure resulting in a deterioration of the mechanical properties. This effect was specially manifested in the elongation and reduction of area parameters. This loss of ductility of the sudied steels was interpreted as a Hydrogen Assisted Stress Cracking (HASC) phenomenon. Only in the case of the 15 MnNi 6.3 and TStE 460 steels, in which secondary cracks were found, the loss of ductility was also attributed to stress corrosion cracking (SCC) [36,68,83,152,160-162].
A.4.1.1.2.3 Corrosion of Hastelloy C-4 in brines
The corrosion behaviour of Hastelloy C-4 specimens exposed to the ‘Q-brine’ (MgCl2- rich brine 1) at temperatures of 90°C, 170°C and 200°C for periods of up to 3 years has been investigated. The corrosion results are summarised in TABLE 4-14.
95 TABLE 4-14. Corrosion results for HASTELLOY C-4 after 3 years exposure to Q-brine at various temperatures [136,137].
(1) Temperature vCORR Pitting corrosion Crevice corrosion Stress corrosion (°C) (µm/year) (µm) (µm) cracking
90 0.20±0.10 n.p. 50 n.SCC 170 0.35±0.05 n.p. 150 n.SCC 200 0.90±0.10 200 200 n.SCC
(1) vCORR : average uniform corrosion rate. n.p. : no pitting corrosion. n.SCC : no stress corrosion cracking.
The uniform corrosion rate of Hastelloy C-4 increased with increasing temperature, but still remained fairly low, even at the maximum test temperature of 200°C: the corrosion rates varied between 0.2 µm/year and 0.9 µm/year depending on the temperature [136,137]. The negligible influence of temperature on the uniform corrosion of Hastelloy C-4 is illustrated in FIGURE 4-22.
FIGURE 4-22. Average corrosion rates of Hastelloy C-4 in Q-brine at different tem- peratures [148].
At 90°C and 170°C, the material was resistant to pitting and stress corrosion cracking (U-bend specimens) but suffered from crevice corrosion (at the contact surface between two plane specimens). At 200°C, besides crevice corrosion, pitting corrosion occurred to a depth of 200 µm in tests exceeding 3 years duration. The sensitivity of Hastelloy C-4 to local corrosion at elevated temperatures in Q-brine was also confirmed by electrochemical studies. These studies showed that an increase in temperature significantly reduces the stability of the passive layer, which develops spontaneously in this brine and causes local corrosion [136,137,148]. The influence of temperature on the pitting susceptibility of Hastelloy C-4 is also illustrated in FIGURES 4-23 (a) and (b).
96
(a) 100 µm (b) 200 µm
FIGURE 4-23. Optical micrographs of etched Hastelloy C-4 specimens in the as- delivered condition after immersion in the Q-brine [140,148,156]. (a) 3 years; 170°C; no γ-irradiation magnification: 200× (no pitting corrosion)
(b) 700 days; 200°C; no γ-irradiation magnification: 100× (maximum pit depth: 200 µm) (c) 606 days; 90°C; 103 Gy/h magnification: 50× (maximum pit depth: 1,000 µm) (c) 500 µm
The influence of gamma irradiation on the corrosion behaviour of Hastelloy C-4 in Q- brine at 90°C is shown in TABLE 4-15. The dose rate in the range 1 Gy/h - 100 Gy/h did not noticeably influence the uniform corrosion rate: at dose rates between 1 Gy/h and 100 Gy/h, the maximum uniform corrosion rates of 0.4 µm/year are low and roughly correspond to the value obtained in the absence of gamma irradiation. Moreover, Hastelloy C-4 remained resistant to stress corrosion cracking and local corrosion after a 12-month exposure period to the Q-brine at 90°C and a gamma dose rate of 1 Gy/h [38,136,137,142,157]. At higher dose rates of 10 Gy/h and 100 Gy/h, Hastelloy C-4 was still resistant to stress corrosion cracking, but was susceptible to pitting and crevice corrosion (a maximum pit penetration rate of 20 µm/year was detected). However, the penetration rate remained unchanged during the whole testing period of 12 months [38,142,157]. Increasing the dose rate accelerated the pitting corrosion from 20 µm/year at 10 Gy/h and 100 Gy/h to 1 mm/year at 10³ Gy/h [136,137,142]. The effect of a gamma irradiation dose rate of 10³ Gy/h on the pitting susceptibiulity of Hastelloy C-4 is illustrated in FIGURE 4-23 (c).
97 TABLE 4-15. Uniform corrosion rates and maximum pitting rates of Hastelloy C-4 after 1 year immersion in Q-brine at 90°C with and without gamma irradiation [38,141,142].
(1) (2) D vCORR Maximum pitting (Gy/h) (µm/year) corrosion rate (µm/year)
0 0.1 n.p. 1 0.05 n.p. 10 0.4 20 100 0.4 20 1,000 3.5 1000
(1) D : gamma dose rate. (2) vCORR : average uniform corrosion rate. n.p. : no pitting corrosion.
The effect of the presence of the salt impurity H2S on the corrosion behaviour of Hastelloy C-4 is shown in TABLE 4-16. Long-term immersion experiments exposing Hastelloy C-4 specimens to the Q-brine, without H2S and with 25 mg/L H2S, at 150°C for periods up to 12 months were performed. These experiments revealed that the corrosion resistance of Hastelloy C-4 is markedly reduced by the addition of H2S to the Q-brine [38,41,142]:
• in the H2S-free solution, Hastelloy C-4 exhibited a very low uniform corrosion rate of less than 1 µm/year and was resistant to pitting corrosion [142]. • in the H2S-containing brine, the uniform corrosion rate of Hastelloy C-4, despite an increase, still remained fairly low with a maximum of 4 µm/year. However, when 25 mg/L H2S was added to the test medium, pitting corrosion of 2.3 mm maximum occurred after an extended immersion period of 12 months [38,142]. • also, crevice corrosion was detected in both the H2S-free and the H2S-containing salt brines on the surface of the Hastelloy C-4 specimens below the PTFE threads by which the specimens had been suspended. The crevice attack was, however, more pronounced in the H2S-containing brine [41,142].
98 TABLE 4-16. Corrosion results for Hastelloy C-4 in Q-brine at 150°C in the absence and presence of H2S (25 mg/L) [142].
(1) Corrosion medium Testing time vCORR Maximum pitting Maximum rate of (months) (µm/year) corrosion rate crevice corrosion(2) (µm/year) (µm/year)
2 0.9±0.1 n.p. n.cc Q-brine 6 0.5±0.05 n.p. n.cc 12 0.7±0.3 n.p. 150
2 3.0±0.4 n.p. n.cc Q-brine + 6 4.1±1.1 40 1,000 25 mg/L H S 2 12 3.9±0.6 2,300 2,200
(1) vCORR : average uniform corrosion rate. (2) crevice attack occurred on the surfaces of the specimens below the PTFE threads by which the specimens had been suspended. n.p.: no pitting. n.cc : no crevice corrosion.
A.4.1.1.2.4 Corrosion of Ti 99.8-Pd in brines
The uniform corrosion behaviour of Ti 99.8-Pd exposed to the ‘Q-brine’ (MgCl2-rich brine 1) at temperatures of 90°C, 170°C, and 200°C for periods of up to 3.5 years has been investigated. The time-temperature behaviour of the uniform corrosion rates of Ti 99.8-Pd investigated in the Q-brine is shown in TABLE 4-17 and FIGURE 4-24. Under these test conditions, Ti 99.8-Pd corroded uniformly: the corrosion rates, which were already low initially, declined even further with exposure time to extremely low levels, which were of the order of 0.1-0.2 µm/year. The rate of attack was not influenced noticeably by the test temperature up to 200°C [30,136,137,140,148].
Ti 99.8-Pd was also resistant to pitting corrosion, crevice corrosion, and stress corrosion cracking (U-bend specimens) [30,136,137,140,148]. The optical micrographs in FIGURES 4-25 (a) and 4-25 (b) illustrate that the surface of Ti 99.8- Pd was attacked uniformly in the Q-brine (without gamma irradiation).
Ti 99.8-Pd also proved resistant to localised corrosion in the Q-brine under gamma irradiation [30,68,136,137,140, 145,148,152]:
• at 90°C and under a gamma irradiation field of 103 Gy/h, a value corresponding to the dose rate on the outer surface of a 5-mm thin-walled packaging, Ti 99.8-Pd corroded uniformly (localised corrosion was not detected), as illustrated in FIGURE 4-25 (c): the average uniform corrosion rate increased only slightly to 0.7 µm/year, compared to the corrosion rate calculated in the absence of irradiation (0.17 µm/year), and remained fairly constant over an immersion period of 609 days [30,136,137,140,148]. • the average corrosion rates calculated for Ti 99.8-Pd in irradiated and unirradiated MgCl2-rich Q-brine at 150°C are presented in TABLE 4-18. Without irradiation, the uniform corrosion rate remained between 0.02 µm/year and 0.7 µm/year, which is
99 negligibly low. The imposition of a 10 Gy/h gamma irradiation field did not increase the corrosion of Ti 99.8-Pd [68,145,152].
TABLE 4-17. Average uniform corrosion rates of Ti 99.8-Pd exposed to the Q-brine at various temperatures (90°C, 170°C, and 200°C) and testing periods [136,137,140,148].
(1) T Exposure vCORR (°C) period (µm/year)
559 days 0.17±0.03 90 1280 days 0.06±0.02
180 days 0.16±0.02 170 529 days 0.07±0.03 1280 days 0.04±0.02 FIGURE 4-24. Average uniform corrosion rates of Ti 99.8-Pd in Q-brine at 540 days 0.14±0.03 different temperatures [148]. 200 1280 days 0.15±0.02
(1) vCORR : average uniform corrosion rate, calculated from the weight losses (average value of three specimens) and extrapolated linearly.
TABLE 4-18. Uniform corrosion rates of Ti 99.8-Pd in Q-brine at 150°C with and without gamma irradiation field [68,152].
(1) (2) D Exposure time vCORR (Gy/h) (days) (µm/year)
63 0.09±0.04 125 0.02±0.02 0 191 0.73±0.01 268 0.02±0.01
63 0.06±0.06 174 0.14±0.02 10 244 0.62±0.01 356 0.21±0.02
(1) D : gamma dose rate (Gy/h). (2) vCORR : average uniform corrosion rate, calculated from the weight losses (average value of three specimens) and extrapolated linearly.
100
(a) 100 µm (b) 100 µm
FIGURE 4-25. Optical micrographs of etched Ti 99.8-Pd specimens in the as-delivered condition after immersion in the Q-brine [140,148,156]. (a) 368 days; 170°C; no γ-irradiation magnification: 200× (uniform corrosion) (b) 420 days; 200°C; no γ-irradiation magnification: 200× (uniform corrosion) (c) 609 days; 90°C; 103 Gy/h magnification: 200× (c) (uniform corrosion)
100 µm
The high corrosion resistance to uniform and localised corrosion exhibited by Ti 99.8- Pd in brines is attributable to the formation of stable protective surface layers of titanium oxides, which have been identified by means of surface analyses to consist mainly of TiO2. In general, a non-linear increase of the TiO2 layer thickness with corrosion time was observed. Electrochemical studies of the potential-time behaviour (E/t) and the current density-potential behaviour (I/E) of Ti 99.8-Pd in Q-brine at 25°C and 90°C revealed a wide stable passive zone for Ti 99.8-Pd. The I/E-plots showed Ti 99.8-Pd to be resistant to localised corrosion up to E=+3.0 V vs. Ag/AgCl (3M KCl) [136,137,145,148]. A similar film was also found on the specimens exposed to gamma irradiation. In this case, however, a second over-layer consisting of Mg and O was also detected, and the overall film thickness of 600 nm (formed on Ti 99.8-Pd specimens in Q-brine at 90°C and a gamma dose rate of 103 Gy/h after 609 days) was a factor of about 10 greater than that formed on the unirradiated specimens [30,136,137,140,148]. Evaluation of the crevice area of metal-to-metal specimens did not show any signs of pitting corrosion. However, inside the crevice area, a thin, multicoloured corrosion product having yellow, violet, and blue colours was observed. XPS-examination indicated that the colour differences inside the crevice were due to optical interference colours caused by varying oxide film thickness. Outside the crevice, a very thin corrosion film with a yellow colour was formed [145].
101 The TIG- and EB-welding did not reduce the high corrosion resistance of Ti 99.8-Pd in the Q-brine. As for the unwelded specimens, the TIG- and EB-welded specimens showed a negligible uniform corrosion rate (< 1 µm/year), and a high resistance to localised corrosion [68,152].
A.4.1.1.3 Experiments on real specimens originating from the POLLUX7-disposal container
A.4.1.1.3.1 Background
Gesellschaft für Nuklear-Behälter (GNB) has developed two container concepts for the disposal of spent fuel: the corrosion-resistant concept and the corrosion- allowance concept. In the corrosion-resistant concept, an inner steel container, ensuring mechanical stability to the rock pressure, is cladded with a 3- to 4-mm welded overlay of Hastelloy C-4 for corrosion protection. In the corrosion-allowance concept, GNB is considering a thick-walled container made of unalloyed or low- alloyed steel grades, respectively, whose long-term stability to corrosion and rock pressure is safe-guarded by a sufficient wall thickness. In both concepts, the disposal container is surrounded by an overpack of nodular cast iron (Globularer Grauguβ, GGG 40.3, which contains 3.63 wt.% carbon, 1.26 wt.% silicon, 0.18 wt.% manganese, 0.04 wt.% chromium, and nickel). In the corrosion-allowance concept, also the use of very thick-walled self-shielding steel containers without an overpack is possible. The overpack not only protects personnel during the operating phase of the repository but also shields the saline rock and any brine existing in the surroundings from gamma-radiation induced effects (radiolysis) during the postoperating phase [29,41].
The results of the corrosion studies on carbon steels related to the corrosion- allowance concept are presented in section A.4.1.1.2.2. In the following section the corrosion results obtained on surface-welded Hastelloy C-4 (corrosion-resistant concept) are discussed.
The specimens examined during these experiments were cut from large Model- POLLUX containers, which were fabricated in Germany under conditions corresponding exactly with those of the original POLLUX-containers (i.e. using the original welding technique, the original fabrication process, etc.). The Hastelloy C-4 specimens originated from a typical POLLUX surface weld on the low-alloyed carbon steel 15 MnNi 6.3 [41].
The surface-welded Hastelloy C-4 specimens were examined under conditions simulating an accidental intrusion of large amounts of brine into the disposal area, in order to create a worst case scenario. For the double POLLUX-container system, it was assumed that the welds will not be able to withstand the corrosion, thereby allowing brines to enter the annular gap between the shielding (cast iron) and the inner container (Hastelloy C-4) by intrusion through the corroded welds of the outer cast iron container lid [41]. The possibility of ‘denting’ (i.e. material straining due to a
7 The POLLUX-container system, which is used for the final disposal of spent fuel in rock-salt, is described in detail in APPENDIX B.
102 volume increase induced by the growth of the corrosion products) was not considered because the very thick cast iron outer container (> 200 mm) is expected to withstand the pressures exerted by the corrosion products.
Immersion tests lasting up to 18 months were performed at 150°C and at a brine volume to specimen surface ratio (V/S) of 5 mL/cm², which provides an excess of corrodent. All experiments were performed under static conditions without refreshing the brine. The pressure in the vessels during the tests was 0.4 MPa, which corresponding to the vapour pressure of the brine at 150°C [41].
A.4.1.1.3.2 Results and discussion
The evolution of the general corrosion rate for surface-welded Hastelloy C-4 in the MgCl2-rich brines 1 and 2, and the NaCl-rich brine 3, at 150°C, is plotted in FIGURE 4-26. The values are the average values from five specimens. The ranges of the corrosion rates represented as bars show the standard deviations of the measured values. In the NaCl-rich brine (brine 3), the corrosion rate of the material remained fairly constant during the whole test duration and was extremely low (< 0.1 µm/year). In the MgCl2-rich brines (brines 1 and 2), higher general corrosion rates were observed compared with the values of the NaCl-rich brine. However, the values of 1.5 and 2.0 µm/year determined for up to 12 months in brine 1 and for up to 8 months in brine 2, respectively, were relatively low. After longer test durations, the corrosion rates in these brines clearly increased. The final corrosion rates of the 18-month specimens amounted to 5 to 6 µm/year [41].
FIGURE 4-26. General corrosion rates as a function of time for surface-welded HASTELLOY C-4 in rock-salt brines at 150°C [41]: ○ MgCl2 -rich brine 1 (26.8 wt.% MgCl2). ● MgCl2-rich brine 2 (33 wt.% MgCl2). □ NaCl-rich brine 3 (26.9 wt.% NaCl).
103 The results of the metallographic examinations of specimens after 12 and 18 months of immersion in the three brines at 150°C are summarized in TABLE 4-19. FIGURES 4-27 (a), 4-27 (b), and 4-27 (c) show characteristic photographs after an 18-month exposure to the three brines. In the NaCl-rich brine, the material corroded uniformly. In the MgCl2-rich brines, however, the surface-welded Hastelloy C-4 suffered from severe pitting corrosion of up to 900 µm after 18 months of exposure to the brines. This explains the increase in the general corrosion rates and the fluctuation of the values for the specimens tested longer than eight months [41].
TABLE 4-19. Average uniform corrosion rates and maximum pit depths for surface-welded Hastelloy C-4 (long-term immersion tests in rock-salt test brines at 150°C) [41].
(1) (2) (3) Brine Test duration vCORR dpit, Max (months) (µm/year) (µm)
Brine 1 12 1.5 n.p. 18 3.0 300
Brine 2 8 2.0 500 18 5.0 - 6.0 900
Brine 3 12 0.07 ± 0.03 n.p. 18 0.06 ± 0.02 n.p.
-4 12 0.5 ±0.04 150 Brine 3 + 6×10 M Na2S 18 0.12 ± 0.01 200
(1) brines 1 and 2: MgCl2-rich brines ; brine 3: NaCl-rich brine. (2) vCORR : average uniform corrosion rate. (3) dpit, Max : maximum pit depth. n.p. : no pitting.
The maximum rates of penetration (pitting corrosion) observed in the long-term corrosion tests for two types of cladding are listed in TABLE 4-20.
TABLE 4-20. Maximum rates of penetration observed during the long-term corrosion tests on Hastelloy C-4 samples [29].
Maximum rate of penetration (µm/year) Medium Welding A(a) Welding B(b)
Brine 1 (26.8 wt.% MgCl2) 200 300
Brine 2 (33.0 wt.% MgCl2) 600 300
(a) welding on cold steel plate. (b) welding on preheated steel plate.
104
FIGURE 4-27. Optical micrographs of surface-welded Hastelloy C-4 specimens after 18 months of exposure to various rock-salt brines at 150°C [41]. (a) MgCl2-rich brine 1 (magnification: 50×), (b) MgCl2-rich brine 2 (magnification: 50×), and (c) NaCl-rich brine 3 (magnification: 200×).
In addition to the corrosion tests in pure brines, experiments were conducted in the NaCl-rich brine 3 that contained sulfide to examine a possible influence of this salt impurity on the corrosion behaviour of the surface-welded Hastelloy C-4. The sulfide -4 was added to the brine as Na2S.9H2O at a concentration of 6×10 M. The corrosion results obtained for surface-welded Hastelloy C-4 at 150°C in the NaCl-rich brine
105 -4 containing 6×10 M Na2S are also given in TABLE 4-19. The resistance of Hastelloy C-4 to corrosion is markedly reduced in the presence of Na2S in the brine:
• after an 18-month exposure period to the Na2S-containing brine, only a slight increase in the general corrosion rate (0.12 µm/year) was observed compared with the value in the Na2S-free brine (0.06 µm/year) [41]. • however, the addition of Na2S caused severe pitting corrosion, with a maximum depth of 250 µm after an 18-month exposure period, whereas the surface-welded Hastelloy C-4 did not suffer from pitting in the Na2S-free brine. This local attack started at the cut edges and propagated to the specimen surface. Moreover, crevice corrosion occurred below the PTFE threads by which the specimens had been suspended. The optical micrograph in FIGURE 4-28 illustrates the pitting corrosion susceptibility of Hastelloy C-4 in the Na2S-containing NaCl-rich brine [41].
FIGURE 4-28. Optical micrograph of surface-welded Hastelloy C-4 showing pitting corrosion in NaCl- -4 rich brine + 6×10 M Na2S after 18 months of testing at 150°C (magnification: 200×) [41].
A.4.1.1.3.3 Conclusions
In case of an intrusion of NaCl brines into the disposal area, carbon steel containers coated with surface-welded Hastelloy C-4 can provide a long-term corrosion protection of the waste products. Under the experimental conditions applied, Hastelloy C-4 is resistant to local corrosion, and its general corrosion rate is very low. However, in the presence of sulphides or large amounts of highly concentrated MgCl2-rich brines, long-term corrosion protection cannot be expected from such containers because of the susceptibility of welded Hastelloy C-4 to pitting corrosion. In order to improve the corrosion resistance of the welded Hastelloy C-4, a thermal stress relief treatment of the material is recommended [41].
106 A.4.1.2 In-situ corrosion experiments
Several long-term in-situ corrosion experiments have been performed in the Asse salt mine, simulating the HLW disposal conditions prevailing during normal operation of a repository and certain accidental scenarios. Both metal sheets and model containers, provided with selected container manufacturing characteristics, were examined. The investigations have included [43,142]:
• testing metal sheets in rock salt at rock temperature (32°C). This was the reference experiment. • investigation of metal sheets under normal operating conditions (HLW design temperature, gamma radiation field, limited amounts of brine thermally released into the boreholes) of a repository. In these experiments, the applied gamma dose rate (3×10² Gy/h) was about one order of magnitude higher than the value to be expected on the surface of thick-walled HLW packaging. These experiments were performed within the framework of the German/US Brine Migration Test. • testing carbon steel tubes (model containers) with and without a corrosion protection layer (Ti 99.8-Pd or Hastelloy C-4), at the HLW design temperature under simulated accident conditions, with intrusion of large amounts of brine into the boreholes. • long-term exposure experiments, investigating the corrosion behaviour of pre- selected candidate container materials under normal operating conditions. These experiments were performed within the framework of the DEBORA-experiment and the BAMBUS-experiment.
A.4.1.2.1 Investigations in rock salt at rock temperature
In this reference experiment, five Fe-base materials (fine-grained carbon steel TStE 355, nodular cast iron GGG 40.3, Ni-resists D2 and D4, and Si-cast iron), Hastelloy C-4, and Ti 99.8-Pd were investigated in loose rock salt for 3 years at 32°C (i.e. the rock temperature). The composition of the tested materials is given in TABLE 4-1. For nodular cast iron, Ni-resists, and Si-cast iron, only the parent materials were studied. For the most promising HLW packaging materials, fine- grained carbon steel, Hastelloy C-4, and Ti 99.9-Pd, specimens with a TIG (Tungsten Inert Gas) weld were added, so that the influence of welding on the corrosion behaviour could be investigated. Plane specimens were stored in small boreholes (50 mm diameter, 200 mm length) at a depth of 775 m in the Asse salt mine and covered with crushed salt [43,142]. The average composition of the rock salt is given in TABLE 4-21.
TABLE 4-21. Average composition (in wt. %) of the rock salt of the Asse salt mine [42,43,83,142,143,155].
+ + 2+ 2+ - 2- Constituent Na K Ca Mg Cl SO4 H2O
Chemical composition 38.3 0.33 0.17 0.16 58.02 2.47 0.1 (wt. %)
107 For all the tested materials the general corrosion rate was either below the detection limit or very small (< 1 µm/year). Localised corrosion was not detected on any of the materials tested, either in the as-received or welded condition [43,142].
A.4.1.2.2 Brine Migration Test (investigations in rock salt / limited amounts of brine at HLW design temperature)
In these experiments, five Fe-base materials (fine-grained carbon steel TStE 355, nodular cast iron GGG 40.3, Ni-resist D4, Si-cast iron, and cast steel), Ti 99.8-Pd, and Hastelloy C-4 were tested. Both unwelded and TIG welded specimens were examined for general and localised corrosion [142].
The experiments consisted of placing plane specimens in heated, cased boreholes of 2m depth, two of which accomodated Co-60 sources. The specimens were exposed to a steam atmosphere with salt constituents (a maximum of 140 ml brine migrated into the boreholes) and gases emanating from the rock salt (H2, O2, CO2, CH4, C2H6, C3H8). The specimens were exposed to temperatures between 120°C and 210°C, a maximum gamma dose rate of 3×10² Gy/h, and a maximum rock pressure of 28 MPa. The maximum testing period was 900 days [43,157].
The measured weight losses and the calculated corrosion rates for the tested materials are summarised in TABLE 4-22.
TABLE 4-22. Weight losses and corrosion rates of the materials tested in-situ in Asse rock salt with limited amounts of brine (German/US Brine Migration Test) [43,142,157].
Without γ-radiation With γ-radiation (exposure time: 900 days) (dose rate: 3×10² Gy/h) Material T (exposure time: 700 days) (°C)
(1) (1) Weight loss vCORR Weight loss vCORR (g/m²) (µm/year) (g/m²) (µm/year)
Ti 99.8-Pd 210 n.d. n.d. 1.4 0.16 Hastelloy C-4 210 n.d. n.d. 19.19 1.18 TStE 355 150 17.79 0.95 200.06 13.68 Cast steel 150 22.72 1.18 9.49 0.63 Ni-resist D4 150 5.43 0.29 3.23 0.22 GGG 40.3 120 n.d. n.d. 13.26 1.01 Si-cast iron 120 29.31 1.72 n.d. n.d.
(1) vCORR : average uniform corrosion rate. n.d. : not determined.
The corrosion rates of all materials not exposed to irradiation were <2 µm/year. These low corrosion rates were attributed to the very low amount of available brine (a
108 maximum of 140 mL brine had migrated into the borehole which, besides contacting the samples, was also spread over the large surfaces of the inserts). The γ-radiation field of 3×10² Gy/h did not result in increased corrosion rates, except for the fine- grained carbon steel TStE 355 (~14 µm/year). This increased corrosion rate was not believed to be caused by the effect of radiation, but it is probably attributable to the fact that the specimens were exposed to non-condensed brine (the brines trapped in the gas phase recondensed again) for an additional period of approximately 12 months (because the samples could not be retrieved simultaneously with the removal of the radiation source), at a mean temperature of about 70°C [43,142,157].
All the materials, except for the fine-grained steel TStE 355 exposed to irradiation, underwent uniform corrosion with and without irradiation. In the case of the fine- grained carbon steel TStE 355, non-uniform corrosion attack with a maximum penetration rate of ~25 µm/year was observed, which was attributed to corrosion through non-condensed brine [43,142,157].
No noticeable influence of TIG welding was observed on the corrosion behaviour of the tested materials [157].
The available results from these in-situ experiments showed much lower corrosion rates of the tested iron-base materials than those obtained in the laboratory-scale experiments with brine in excess. This difference is probably due to the difference in the available amount of brine: in the in-situ experiments, the specimens were exposed to a limited amount of brine, whereas in the laboratory-scale experiments, the specimens were exposed to an excess of brine [43,142].
A.4.1.2.3 Testing of welded tubes (model containers) in rock salt / brine at HLW design temperature
In this experiment, several carbon steel tubes, with and without a corrosion protection layer (Ti 99.8-Pd or Hastelloy C-4), were examined in heated boreholes in the Asse salt mine, to investigate the influence of selected container manufacturing characteristics (e.g. sealing/welding technique, application of a corrosion protecting layer). The experiments were performed under the conditions of a hypothetical inflow of NaCl-rich brine (brine no. 3; see TABLE 2-1) or MgCl2-rich brine (brine no. 1 or Q-brine; see TABLE 2-1) into the HLW boreholes during the initial disposal phase, while the annular gap between the container and the borehole wall was still open [43,83,143,155,158].
The tubes consisted of several cylindrical sections (50 mm long, 40 mm outside diameter, 10 mm thick) which were joined by electron beam welding. Some of the tubes were equipped with a corrosion protection layer of Ti 99.8-Pd or Hastelloy C-4 by means of explosion plating and EB-welding. The non-corrosion protected steel tube was subjected, after welding, to thermal treatment for 2 hours at 700°C, in order to simulate the cooling conditions applicable to the seam of a thick-walled container [43,142,155,158].
The tubes were placed into 2 m deep heated vertical boreholes in the rock salt Asse mine (at a depth of 775 m) and the 1 mm wide annular gap between the tubes and
109 the borehole wall was filled with either NaCl-rich brine no. 3 or MgCl2-rich brine no. 1 (Q-brine) [42,43,83,142,143,155].
A temperature gradient ranging from 90°C to 200°C was measured along the contact surface between the tubes and the borehole wall [42,43,142,143,158].
The unprotected steel tube showed the most severe non-uniform corrosion attack in the low temperature zones (90°C): a maximum rate of penetration of corrosion of 120 µm/year was measured. In these lower temperature zones, the steel tube was in contact with condensed evaporated salt brine. For the specimens taken from the higher temperature zones (140°C and 200°C), the maximum penetration rates were much lower, viz. 40 to 50 µm/year. In these higher temperature zones, the specimens were in contact with vapour. No pitting, crevice corrosion, or SCC were detected. EB- welding did not noticeably influence the corrosion behaviour of steel. The corrosion products consisted of a mixture of iron oxides (Fe3O4, Fe2O3, α-FeO(OH), and γ- FeO(OH)), with NaCl incorporated into the corrosion layer [43,142,158].
The steel tubes plated with a Ti 99.8-Pd layer were resistant to corrosion in both brines. The steel tubes with a corrosion protection layer of Hastelloy C-4 were resistant to localised corrosion in the NaCl-rich brine; however, in the MgCl2-rich brine, the Hastelloy C-4 surface layer was covered with small pits (1-2 µm in size), which converged into pits approximately 10 µm in diameter and about 10-15 µm deep [42,43,155].
A.4.1.2.4 Long-term testing under normal operating conditions at high temperature
Long-term in-situ corrosion experiments were performed in the Asse salt mine on materials which were identified either as (i) promising materials for the manufacture of long-lived containers for HLW/spent fuel (corrosion-allowance and corrosion- resistant concept), (ii) potential canister materials for vitrified HLW, or (iii) γ-shielding material for the POLLUX-container.
These experiments were performed in the Asse salt mine under conditions similar to those expected in the normal operating phase of the rock salt repository. During the experiments, the specimens were exposed to the salt brine and gases such as O2, CO2, CH4, and H2 (H2 was generated by corrosion of the heater casks; O2, CO2, and CH4 were thermally released from the rock salt into the boreholes) [159].
The corrosion results (general corrosion rates, pit depths) of the various materials investigated in the Asse rock salt environment at 190°C for 5.3 years (no gamma irradiation) are summarised in TABLE 4-23.
The materials Ti 99.8-Pd and Hastelloy C-4 corroded, as expected, at extremely small general corrosion rates (<0.1 µm/year). Furthermore, Ti 99.9-Pd was completely resistant to localised corrosion phenomena such as pitting or crevice corrosion. In the case of Hastelloy C-4, however, some small pits of about 10 µm depth were observed in the crevice between the specimens and the borehole salt wall after 5.3 years testing. The general corrosion of the Cr Ni steel 1.4833 was very low (about 1 µm/year), but the material exhibited severe pitting corrosion (up to 200 µm after 5.3 years) and stress corrosion cracking. Finally, the fine-grained
110 carbon steel TStE 355 showed a non-uniform general corrosion of 40 ± 20 µm/year (the maximum depth of this uneven attack corresponds to the values of the gravimetrically determined average thickness reduction of the specimens). The carbon steel TStE 355 was resistant to pitting and crevice corrosion [144,159].
TABLE 4-23. Corrosion results of the materials tested in-situ under the conditions of the HLW test
disposal (limited H2O/brine intrusion, T = 190°C, time = 5.3 years) [159].
(1) Material vCORR Pitting corrosion (µm/year) (µm)
Ti 99.8-Pd 0.01 ± 0.004 n.p. Hastelloy C-4 0.02 ± 0.01 10 Cr-Ni steel (1.4833) 0.7 ± 0.4 200 + SCC Carbon steel TStE 355 40 ± 20 n.p.
(1) vCORR : average uniform corrosion rate. n.p. : no pitting was observed. SCC : Stress Corrosion Cracking was observed.
A.4.1.2.4.1 Results from the DEBORA-Experiment
The corrosion results (general corrosion rates, pit depths) of the various materials investigated in the Asse rock salt environment at 180°C for 504 days (no gamma irradiation) are summarized in TABLE 4-24. These in-situ corrosion experiments were performed in the framework of the DEBORA (Development of Borehole Seals for High-Level Radioactive Waste) experiment. Besides unwelded (parent) specimens, TIG- and EB-welded specimens, simulating potential container closure techniques, were also tested.
For Ti 99.8-Pd and Hastelloy C-4, the general corrosion rate was below the detection limit. The corrosion rates of the Cr-Ni steels (304L and 309S) are about 1 µm/year, which are very low, and the values for the unwelded fine-grained carbon steel TStE 355 are in the range 3-17 µm/year, which are also relatively low for a corrosion- allowance material. The TIG- and EB-welding did not cause a significant decrease of the material’s corrosion resistance [138].
The Ti-alloy Ti 99.8-Pd was also found to be completely free of localised corrosion. The degree of corrosion of Hastelloy C-4 and the steels depended on the specimens’ location. For the specimens located in the lower ring, no localised corrosion (Hastelloy C-4) or very small amounts of localised corrosion (steels) occurred. A much higher corrosion rate was observed for the specimens located in the upper ring. For Hastelloy C-4 some slight pitting was observed, for the Cr-Ni steels strong pitting (30-120 µm) and IGA (> 2 mm) occurred, and for the carbon steel severe shallow pit formation (350 µm) was detected. The higher corrosion rate of the specimens located in the upper ring compared to those in the lower ring is attributed to a stronger release of brine from the rock salt [138].
111
TABLE 4-24. Average uniform corrosion rates and results from localised corrosion investigations of the in-situ specimens exposed to the Asse salt mine environment at 180°C for 504 days (DEBORA-experiment) [138].
(1) (2) Material Condition vCORR dMax (µm/year) upper ring lower ring upper ring lower ring
unwelded n.m. n.m. n.p. n.p. Ti 99.8-Pd TIG-welded(3) n.m. n.m. n.p. n.p. EB-welded(4) n.m. n.m. n.p. n.p.
unwelded n.m. n.m. 10 µm (pitting) n.p. Hastelloy C-4 TIG-welded n.m. n.m. EB-welded n.m. n.m.
304L(6) unwelded 1.2 ± 0.7 n.m. 120 (pitting) 7 µm (pitting) TIG-welded 1.1 ± 0.6 n.m. > 2 mm (IGA(5))
309S(7) unwelded 1.0 ± 0.6 n.m. 30 µm (pitting) 4 µm (pitting) TIG-welded 1.7 ± 0.1 0.06 ± 0.05 > 2 mm (IGA)
unwelded 17.1 ± 5.6 2.7 ± 1.7 350 µm (non-uniform 20 µm (non-uniform TStE 355 TIG-welded 24.2 ± 0.7 1.1 ± 0.6 general corosion) general corrosion) EB-welded 20.9 ± 0.6 2.0 ± 0.1
n.m. : not measurable. n.p. : no pitting attack. (1) vCORR : average uniform corrosion rate (determined from the measured weight losses and the material density of 2-4 specimens). (2) dMax : the maximum depth of corrosion attack (pitting / IGA) of the unwelded material. No preferential corrosion of the welds was observed. (3) TIG : Tungsten Inert Gas welding. (4) EB : Electron Beam welding. (5) IGA : intergranular attack. (6) DIN Werkstoff Nr. 1.4306. (7) DIN Werkstoff Nr. 1.4833.
A.4.1.2.4.2 Results from the BAMBUS II-Experiment
The corrosion results (general corrosion rates, pit depths) of the various investigated materials in the Asse rock salt environment at 90°C (liquid corrosion) and 180°C (vapour corrosion) for 3,740 days (no gamma irradiation) are summarized in TABLE 4-25 and TABLE 4-26, respectively. These in-situ corrosion experiments were performed in the framework of the BAMBUS II (Backfilling and Sealing of Underground Repositories for Radioactive Waste in Salt, Phase II) experiment. Besides unwelded (parent) specimens, TIG-, EB-, and SAW-welded specimens, simulating potential container closure techniques, were also tested [139].
The results of this experiment showed that the average general corrosion rates of all materials, at both temperatures, were negligible (0.02-1.7 µm/year). The welding did not noticeably increase the corrosion rate. Furthermore, Ti 99.8-Pd, Hastelloy C-4, Hastelloy C-22, 99.9% Ni, and cast iron were completely resistant to pitting corrosion at both temperatures. The unalloyed and low-alloyed steels at 180°C and the Cr-Ni
112 steels at 90°C and 180°C showed some pitting corrosion, but the maximum pit depths after more than 10 years were only 10-100 µm [139].
TABLE 4-25. Average uniform corrosion rates and maximum penetration depths (pitting attack) of the in-situ specimens exposed to the Asse salt mine environment at 90°C (liquid corrosion) for 3,740 days (BAMBUS II-experiment) [139].
(1) (2) Material Condition vCORR dMax (µm/year) (µm)
unwelded 0.01 ± 0.01 n.p. Ti 99.8-Pd TIG-welded(3) 0.02 ± 0.01 n.p. EB-welded(4) 0.02 ± 0.01 n.p.
unwelded 0.02 ± 0.01 n.p.
TIG-welded 0.02 ± 0.01 n.p. Hastelloy C-4 EB-welded 0.03 ± 0.01 n.p. SAW-welded(5) 0.02 ± 0.01 n.p.
unwelded 0.02 ± 0.01 n.p. Hastelloy C-22 TIG-welded 0.02 ± 0.01 n.p. EB-welded 0.02 ± 0.01 n.p.
99.99% Ni unwelded 0.01 ± 0.01 n.p. n.p. unwelded 0.07 ± 0.01 n.p. TIG-welded 0.11 ± 0.02 TStE 355 n.p. EB-welded 0.02 ± 0.01
unwelded 0.05 ± 0.01 n.p. 15 MnNi 6.3 SAW-welded 0.06 ± 0.03 n.p.
unwelded 0.04 ± 0.01 n.p. TStE 460 SAW-welded 0.07 ± 0.01 n.p. EB-welded 0.16 ± 0.05 n.p.
cast iron unwelded 0.09 ± 0.04 n.p.
unwelded 0.98 ± 0.90 100 304L(6) TIG-welded 0.90 ± 0.80 100
unwelded 0.30 ± 0.20 50 309S(7) TIG-welded 1.50 ± 1.40 60
(1) vCORR : average uniform corrosion rate (determined from the measured weight losses and the material density). (2) dMax : maximum depth of pitting attack. (3) TIG : Tungsten Inert Gas welding. (4) EB : Electron Beam welding. (5) SAW : Submerged Arc Welding (6) material no. 1.4306. (7) material no. 1.4833.
113 TABLE 4-26. Average uniform corrosion rates and maximum penetration depths (pitting attack) of the in-situ specimens exposed to the Asse salt mine environment at 180°C (vapour corrosion) for 3,740 days (BAMBUS II-experiment) [139].
(1) (2) Material Condition vCORR dMax (µm/year) (µm)
unwelded 0.02 ± 0.002 n.p. Ti 99.8-Pd TIG-welded(3) 0.02 ± 0.002 n.p. EB-welded(4) 0.04 ± 0.002 n.p.
unwelded 0.04 ± 0.001 n.p.
TIG-welded 0.05 ± 0.010 n.p. Hastelloy C-4 EB-welded 0.20 ± 0.010 n.p. SAW-welded(5) 0.20 ± 0.001 n.p.
unwelded 0.05 ± 0.002 n.p. Hastelloy C-22 TIG-welded 0.05 ± 0.003 n.p. EB-welded 0.02 ± 0.001 n.p.
99.99% Ni unwelded 1.7 ± 0.30 n.p. n.p. unwelded 0.07 ± 0.03 n.p. TIG-welded 0.09 ± 0.04 TStE 355 10 EB-welded 0.05 ± 0.02
unwelded 0.05 ± 0.03 - 15 MnNi 6.3 SAW-welded 0.05 ± 0.04 80
unwelded 0.04 ± 0.01 60 TStE 460 SAW-welded 0.06 ± 0.01 80 EB-welded 0.22 ± 0.12 60
cast iron unwelded 0.08 ± 0.06 n.p.
unwelded 0.03 ± 0.001 n.p. 304L(6) TIG-welded 0.02 ± 0.001 30
unwelded 0.02 ± 0.001 20 309S(7) TIG-welded 0.04 ± 0.001 60
(1) vCORR : average uniform corrosion rate (determined from the measured weight losses and the material density). (2) dMax : maximum depth of pitting attack. (3) TIG : Tungsten Inert Gas welding. (4) EB : Electron Beam welding. (5) SAW : Submerged Arc Welding. (6) material no. 1.4306. (7) material no. 1.4833.
A.4.1.2.5 Conclusions
The in-situ corrosion studies confirm that the alloy Ti 99.8-Pd and the fine-grained carbon steel TStE 355 exhibited a very high resistance to general and pitting corrosion in rock salt environments, and therefore are considered very promising
114 materials for the manufacture of long-lived containers for the disposal of HLW/spent fuel in rock salt [139]. Although the materials Hastelloy C-4 and C-22, the cast iron, and the Cr-Ni steels 304L (1.4306) and 309S (1.4833) showed a good corrosion resistance to general and local corrosion under the conditions of the BAMBUS II-experiment, it was found in previous studies that these materials could suffer from pitting and/or intergranular corrosion under more severe conditions. Therefore, containers made of Hastelloy C-4 and C-22, cast iron, and Cr-Ni steels (304L, 309S) could fail due to localised corrosion.
In general, the corrosion rates of the unalloyed/low-alloyed steels and cast iron under in situ conditions are significantly lower (< 0.1 µm/year) than the values observed in the laboratory immersion experiments (5-50 µm/year), as illustrated in TABLE 4-27. For Ti 99.8-Pd, Hastelloy C-4, and Hastelloy C-22, on the other hand, the in situ corrosion rates (0.002-0.02 µm/year) are very close to the values determined in the laboratory immersion experiments (0.06-0.2 µm/year). This is attributed to the fact that in the in situ experiments only very small amounts of water/brine were present as a liquid corrosion medium, while in the laboratory immersion experiments, large amounts of brine were used in order to simulate severe accidental conditions in the repository. The iron-base alloys corrode actively in the brines and, therefore, their corrosion rates strongly increase with increasing amount of brine (because in this case, the corrosion rate is controlled by the consumption of brine). Ti 99.8-Pd and the Ni-alloys, on the other hand, corrode passively in brines, by forming a very stable corrosion protective oxide surface layer. The corrosion of these materials is therefore not dependent on the amount of corrodent, but it depends only on the dissolution of the corrosion protective layer, which is very slow [139].
TABLE 4-27. Comparison of uniform corrosion rates obtained at 90°C under in situ conditions, similar to those expected in the normal operating phase of the rock salt repository, and in the laboratory in the NaCl-rich brine 3 [139].
Material vCORR (in situ) vCORR (laboratory) (µm/year) (µm/year)
Ti 99.8-Pd 0.01 ± 0.01 0.06 ± 0.02 Hastelloy C-4 0.02 ± 0.01 0.2 ± 0.1 Hastelloy C-22 0.02 ± 0.01 0.1 ± 0.05 TStE 355 0.07 ± 0.01 5.1 ± 2.2 Cast iron 0.08 ± 0.01 47.0 ± 5.0
vCORR : average uniform corrosion rate. Maximum exposure period (in situ experiments) : 3,740 days. The data originate from the BAMBUS II-experiment. Maximum exposure period (laboratory experiments) : 480 days.
115 A.4.2 Corrosion studies relevant to the Belgian disposal concept in clay
The experimental results presented in this section originate from investigations performed at SCK•CEN.
More detailed information can be found in References [36,38,47,68,152,160,165- 189].
A.4.2.1 Background information
To evaluate the suitability of a wide variety of metallic materials as candidate container material for the disposal of HLW/spent fuel, an extensive research programme was developed. This research programme consisted of in situ corrosion experiments [36,38,47,165-176], which were performed in the Boom Clay host rock formation, and laboratory-scale corrosion experiments [68,152,160,174-189], which were performed in artificially-made media that simulated the different conditions that were expected to develop in the course of the geological disposal in the underground near-field argillaceous environment.
A.4.2.1.1 In situ corrosion experiments
The in situ corrosion experiments were performed between 1986 and 1996 in the underground research facility, HADES, constructed at a depth of 225 metres in the Boom Clay layer.
The main goals of these experiments were (i) to evaluate the corrosion behaviour of a number of candidate container materials and (ii) to predict the long-term corrosion behaviour (for periods up to 500-1000 years) through medium-term interactions (up to 7.5 years). The reasoning at that time was that by implementing the results of the medium-term tests into an existing model, the lifetime and the wall-thickness of the container could be calculated: by extrapolating the uniform thickness reduction, calculated from the medium-term tests, to periods up to 1000 years, it was believed that the lifetime of the container could be predicted (at that time, the corrosion- allowance material carbon steel was considered as a primary candidate container material) [175].
The in situ corrosion experiments aimed at simulating a wide variety of conditions, which could occur in an underground repository, based on the then prevailing disposal concept [175]: (i) normal long-term repository conditions, (ii) accidental conditions, where gas might eventually penetrate through capilarities within the other engineered barriers before contacting the waste package, (iii) temporary storage of containers in an underground gallery in which ventilation and a non-uniform heat distribution would give rise to local condensation and vaporisation of clay water, and (iv) to investigate the effect of the large amounts of concrete that are used either as gallery wall or as candidate backfill on the corrosion behaviour of the container materials.
116 In order to simulate the different environments that were likely to develop in the course of the geological disposal period, three types of experiments were developed placing the metallic samples either in ‘direct contact with clay’ (type I, horizontal), in contact with a ‘humid clay atmosphere’ (type II, vertical), or in contact with a ‘concrete saturated humid clay atmosphere’ (type III, oblique) [175]. A detailed description of the in situ corrosion experiments (installation, recuperation technique, surface analysis, etc.) can be found in References [169,172,173]. The in situ corrosion experiments basically consisted of placing the metallic samples on the outside or inside of a stainless steel support tube, heated to the desired temperature by a retractable furnace (i.e. heating wires), which was situated in the interior of the tube, and exposing the metallic samples to Boom Clay for a predetermined duration. The support tube was sufficiently long (5 to 7 metres) to ensure that the experiments were performed in the undisturbed clay. At the end of the exploitation period, the corrosion test tubes were retrieved by the overcoring technique, a technique specially developed at SCK•CEN, and the samples were investigated [165-169,172,173,175]. FIGURE 4-29 presents a general view of an in situ corrosion experiment exposing the metallic samples directly to Boom Clay.
FIGURE 4-29. View of a corrosion test tube, developed for direct contact between the samples and the Boom Clay host formation, during installation in the underground research facility (right side: section with metallic samples ; left side: section with waste form samples) [172,173,175].
During the in situ corrosion experiments, only interactions with Boom Clay (under various states of aggregation: solid clay, humid clay atmosphere, or concrete saturated clay atmosphere) were studied. The experiments were performed at 16°C, 90°C, and 170°C: 16°C is the temperature of the Boom Clay host rock formation at 225 meters depth; 90°C was chosen as an intermediate temperature; 170°C was, at that time, considered the maximum temperature to which the outside of a container would be subjected to [36,165-167,169,172,173].
A total of twelve in-situ corrosion experiments were installed in the underground laboratory during the period 1985-94: four type I, six type II, and two type III corrosion
117 experiments [36,169,172,173]. The identification of the twelve in-situ corrosion experiments are listed in TABLE 4-28.
TABLE 4-28. Identification data of the twelve in-situ corrosion experiments installed in the SCK underground laboratory [166,167,169,170,172-174].
No. Type Target T Atmosphere Installation Date of Date of Retrieval Time of experiment experiment (°C) date reaching interrupting date exposure(*) target T heating (days)
1 I 170 CLAY 16.12.1985 16.10.1986 30.07.1991 12.09.1991 2096/1748 2 I 90 CLAY 21.10.1985 29.08.1986 02.06.1988 05.08.1988 1019/643 3 I 90 CLAY 06.02.1986 25.08.1986 13.081993 16.11.1993 2840/2545 4 I 16 CLAY 02.10.1985 02.10.1985 - 28.11.1990 1883/1883
5 II 170 He 11.1988 23.01.1989 23.07.1989 10.1989 ±334/174 6 II 90 AIR 17.11.1988 19.11.1988 17.05.1990 21.05.1990 550/544 7 II 90 He 18.11.1988 19.11.1988 21.04.1990 26.04.1990 524/518 8 II 16 AIR 11.1988 11.1988 - 14.09.1992 ±1399/1399 5b II 90 He 01.09.1993 09.07.1994 29.07.1996 08.08.1996 1072/751 8b II 16 He 06.05.1994 06.05.1994 - 10.10.1996 888/888
9 III 90 AIR 03.03.1989 30.08.1989 10.01.1994 06.05.1994(1) 1890/±765(2) 10 III 16 He 17.02.1993(3) 17.02.1993 - 19.02.1996(4) 1125/920(5)
(*) Total time of exposure / Effective time at target temperature. (1) 06.05.1994 : Retrieval date of the container and waste form samples situated on the inside of the tube (concrete-saturated clay atmosphere). 21.11.1994 : Retrieval date of the waste form samples situated on the outside of the tube (direct contact with clay). (2) Due to several technical difficulties, occurring during the exploitation of tube no. 9, the total time of exposure/effective time at target temperature (90°C) was only about 800 days; the temperature was in excess of 60°C for a total period of about 1020 days. (3) 17.02.1993 : Installation date of the container and waste form samples on the outside of the tube (direct contact with clay). 13.08.1993 : Installation date of the container samples on the inside of the tube (concrete-saturated clay atmosphere). (4) 19.02.1996 : Retrieval date of the container samples situated on the inside of the tube. 18.03.1996 : Retrieval date of the container and waste form samples situated on the outside of the tube. (5) 1125 : Total time of exposure/effective time at target temperature (16°C) of the container and waste form samples situated on the outside of the tube (direct contact with clay). 920 : Total time of exposure/effective time at target temperature (16°C) of the container samples situated on the inside of the tube (concrete- saturated clay atmosphere).
During the in situ corrosion programme, a large scatter of materials was screened as possible candidate container material, such as carbon steel, stainless steels (AISI 430, 1803 MoT, AISI 309, AISI 316, AISI 316Ti, UHB 904L, ...), Ni-alloys (Hastelloy C4, Inconel 625, etc.), Ti-alloys (IMI 115, Ti/0.2Pd, etc.), etc. The chemical composition of the candidate container materials tested during the in situ corrosion experiments is presented in TABLE 4-29.
In parallel with this programme other in situ programmes were (CERBERUS) and are still being elaborated (CORALUS) in order to try assessing the influence of radiation on the degradation of, for example, the metallic container [175].
118 TABLE 4-29. Chemical compositions of the candidate container materials tested during the in situ and laboratory corrosion programme [68,152,166-168,173,183,185,186,188].
Materials Chemical composition (wt.%) designation 1 (AISI ) Fe Cr Ni Mn Mo Ti Si Cu Pd C S P N Others
In Situ Corrosion Programme
EEG USINOR (C-steel) bal 0.02 0.04 0.04 <0.01 <0.01 0.27 0.015 - 0.11 <0.01 0.02 - Al: 0.04; Sn< 0.01
AISI 430 (f S.S.2) bal 16.50 - - - - 0.37 - - 0.045 - - 0.026 - 1803 MoT (f S.S.2) bal 18.07 0.28 0.28 2.07 0.36 0.27 - - - - - 0.009 - AISI 309 (a S.S.3) bal 23.0 13.0 2.0 - - 1.0 - - 0.20 - - - - AISI 316 (a S.S.3) bal 17.80 10.8 1.63 2.13 - 0.56 - - 0.03 - - 0.032 - AISI 316Ti (a S.S.3) bal 16-18 10-14 <2.0 2-3 >5×%C <1.0 - - <0.08 - - - - 0.014 UHB 904L (h.-a. S.S.4) bal 18.65 26.34 1.12 4.49 - 0.44 0.12 - - - 0.0497 - - Hastelloy C-4 (Ni-alloy) 0.50 15.6 68.0 0.22 15.4 0.11 0.03 - - 0.0207 - - - Co< 0.1 Inconel 625 (Ni-alloy) 3.57 12.2 bal 0.10 8.85 <0.02 0.2 0.02 - - - 0.023 Nb: 2.7; Ca: 0.005; Mg: 0.02; Al: 0.01; Co< 0.05 IMI 115 ( commercially pure Ti) - - - - - bal ------Ti 99.8-Pd (Ti-alloy) 0.06 - - - - bal - - 0.16 0.01 - - - V< 0.03; Al: 0.008; Sn: 0.07
Laboratory Corrosion Programme
TStE 355 (C-steel) bal 0.03 0.03 1.12 - 0.003 0.344 - - 0.18 0.002 0.01 0.005 Nb: 0.017
AISI 309S (a. S.S.3) bal 22.58 13.51 1.7 - - 0.33 - - 0.063 0.002 0.021 - - AISI 316L (a. S.S.3) bal 16.9 11.0 1.54 2.08 - 0.54 - - 0.017 0.001 0.032 - - AISI 316L hMo (a. S.S.3) bal 17.67 12.53 1.16 2.84 - 0.61 - - 0.015 0.001 0.030 - - AISI 316Ti (a. S.S.3) bal 16.8 10.7 1.08 2.05 0.3 0.40 - - 0.044 0.009 0.028 - -
UHB 904L (h.-a. S.S.4) bal 19.7 25.0 1.48 4.47 - 0.19 1.51 - 0.019 0.001 0.019 0.08 - Cronifer 1925hMo (h.-a. S.S.4) 45.45 20.6 24.85 0.92 6.4 - 0.30 0.86 - 0.005 0.002 0018 0.198 -
Hastelloy C-4 (Ni-alloy) 0.98 15.75 67.0 0.04 15.85 <0.01 0.02 - - 0.003 0.003 0.004 - Co: 0.01
Ti 99.8-Pd (Ti-alloy) 0.04 - - - - bal - - 0.16 0.01 - - <0.01 O2: 0.13; H2: 0.001
1 AISI : American Iron and Steel Institute numbering System. 2 f. S.S. : ferritic stainless steel. 3 a. S.S. : austenitic stainless steel. 4 h.-a. S.S. : high-alloyed stainless steel.
A.4.2.1.2 Laboratory-scale corrosion experiments
The main goals of these experiments were (i) to demonstrate that a suitable material can be selected as the container material and (ii) to investigate the influence of various parameters, such as the composition of the disposal environment (chloride, sulphate, thiosulphate), temperature, oxygen content and the presence of radiolytic products (hydrogen peroxide), on the susceptibility to localised corrosion (with an emphasis on pitting corrosion) of various candidate container materials.
For the Belgian disposal concept, the nature of the underground argillaceous - 2- 2- environment (Cl , SO4 , S2O3 , etc.) and the nature of the studied materials (passive
119 metals), have led to considering pitting corrosion as being of particular importance for container integrity because it is the form of attack that could lead to the greatest container area penetration and hence the greatest release of radionuclides [175,184].
During the laboratory corrosion programme, a wide variety of materials were investigated as candidate container materials, some of which were also tested during the in situ corrosion programme, such as carbon steel, stainless steels (AISI 309S, AISI 316Ti and UHB 904L), Ni-alloy (Hastelloy C4), and Ti-alloy (Ti/0.2Pd). Also some new materials were chosen, based on their favourable pitting resistance characteristics, e.g. stainless steels (AISI 316L, AISI 316L hMo and Cronifer 1925 hMo) [175,184]. The chemical composition of the various materials tested during the laboratory corrosion programme are listed in TABLE 4-29 (section A.4.2.1.1).
The laboratory-scale corrosion experiments consisted of electrochemical experiments [68,152,175,184] and long-term immersion tests [175,179,184].
A.4.2.1.2.1 Electrochemical experiments
The electrochemical experiments [68,152,175,184] involved (i) performing cyclic potentiodynamic polarisation (CPP) measurements and (ii) monitoring the free corrosion potential, ECORR as a function of time, in media representative of the underground repository conditions. Experiments monitoring the evolution of ECORR were performed to gain a more accurate value of the free corrosion potential under conditions representative for the underground disposal because the open circuit potential, OCP, which is derived from the CPP-curves, can be influenced by too many parameters, such as oxygen content, surface conditions, mass transport, etc., to be considered as the actual value of ECORR under repository conditions.
The protection potential, EPP, which is derived from the CPP-curves, was considered as a conservative lower-bound value of the critical potential for the occurrence of pitting corrosion (i.e. that in the potential domain below this potential, pitting will never take place). The approach adopted in these studies involved the comparison of this critical potential to the free corrosion potential, ECORR: pitting can only be initiated if the free corrosion potential of the metal, ECORR, exceeds the protection potential, EPP.
Cyclic potentiodynamic polarisation measurements were performed to investigate the influence of several important parameters on the long-term pitting behaviour of the most promising candidate container materials. These parameters were:
• chemical composition of the near-field underground disposal environ- ment. At the time of the editing of this report, two materials to backfill the disposal galleries were under investigation, viz. a clay-based material, coming from the excavation of the underground repository in the Boom clay host rock formation, which would be re-used after several treatments, and a bentonite-based material, consisting of 60% FoCa clay, 35% sand and 5% graphite. A detailed summary of the chemical composition of both the in situ interstitial Boom Clay porewater and the bentonite porewater have already been given elsewhere (TABLE 2-5, section A.2.2.1).
120 Electrochemical experiments were performed in media (solutions and slurries) containing much higher concentrations of chloride, sulphate and thiosulphate than the ones expected under normal repository conditions. The chloride concentration was varied between 27 and 50,000 mg/L. The sulphate concentration was varied between 0.2 and 5,400 mg/L: 0.2 mg/L represents the sulphate content in the undisturbed Boom Clay formation (anoxic conditions); 216 and 5,400 mg/L are representative for the sulphate content of slightly oxidised and fully oxidised Boom Clay under oxic conditions, respectively; 1700 mg/L represents the sulphate content in the bentonite backfill material. The influence of thiosulphate was investigated in the range 2–200 mg/L. - - 2- Because F , HCO3 and HPO4 were either present in very low concentrations or were not expected to have a major influence on the pitting susceptibility of the investigated materials, they were not considered in the test matrix. • temperature. Three different temperatures were considered: 140°C, being the temperature of the near field surrounding the containers in the first phase after disposal, 16°C, being the temperature of the Boom Clay formation at 225 metres below ground level, and 90°C, being an intermediate temperature. • oxygen content. The oxygen content reflects the aerobic and anaerobic phases of the disposal. The experiments under oxic conditions were performed in media that were in equilibrium with air (1 atm.). The experiments under anoxic conditions were performed in a glove box, under a controlled and inert (Ar) atmosphere. The content of the dissolved oxygen in the test media was always kept below 10 ppb. • radiolytic products. The maximum level of hydrogen peroxide (H2O2) that can be expected at the surface of a 30 mm thick stainless steel container was estimated -2 8 to be 8.10 mol/L . This maximum H2O2-concentration was calculated by Prof. Dr. Marx according to the kinetic method reported in [152] and based on the Belgian disposal concept existing at the time of the editing of this report. The hydrogen peroxide concentration was varied in the range 8.10-3–8.10-1 mol/L.
A detailed description of the parameters investigated during the electrochemical experiments is summarised in TABLE 4-30.
The synthetic media were prepared from analytical-grade chemicals and deionised water. Chloride was added to the test media as sodium chloride (NaCl) and potassium chloride (KCl). Sulphate additions were made using sodium sulphate (Na2SO4), potassium sulphate (K2SO4) and magnesium sulphate (MgSO4.7H2O). Thiosulphate was added to the test media as sodium thiosulphate (Na2S2O3). The pH of all test media was situated between 6.0 and 6.5 (no pH adjustment was made). Slurries with an L/S ratio9 of 90/10 and 30/70 were used.
8 MARX G., 2002. Personal communication (e-mail: July 3, 2002). 9 ratio of the liquid phase (in ml) to the solid phase (in grams).
121 TABLE 4-30. Overview of the parameters studied during the electrochemical corrosion experiments (cyclic potentiodynamic polarisation experiments).
- 2- 2- Test environments Test T [Cl ] [SO4 ] [S2O3 ] [H2O2] conditions (°C) (mg/L) (mg/L) (mol/L)
Solutions SOCW(1) oxic 16, 90, 140 50-50000 216, 5400 2-200 0 oxic 90 1000 216 0 8.10-3-8.10-1 SICW(2) anoxic 16, 90 27-50000 0.2 20-200 0 SBW(3) oxic 16, 90,140 100-50000 1700 20-200 0 anoxic 16, 90 100-50000 1700 20-200 0 RCW(4) anoxic 16, 90 ~27 <0.1 n.d. 0
Slurries based on Boom clay oxic, anoxic 16, 90 100-10000 0.2 0 0 based on bentonite(5) oxic, anoxic 16, 90 100-10000 1700 0 0
(1) SOCW : Synthetic Oxidised Boom Clay Water solutions. (2) SICW : Synthetic Interstitial Boom Clay Water solutions. (3) SBW : Synthetic Bentonite Water solutions. (4) RCW : Real Interstitial Boom Clay Water solutions. (5) bentonite denotes a mixture of 60% smectite clay ‘FoCa’, 35 % sand and 5% graphite. n.d. : not determined.
A.4.2.1.2.2 Long-term immersion tests
The long-term immersion tests [175,179,184] are to be regarded as complementary to the in situ corrosion experiments. The immersion tests, however, had the advantage of being much more flexible for studying the medium-term interactions between the candidate container materials and backfill material under a wide variety of conditions. A detailed description of the parameters investigated during the long- term immersion tests is summarised in TABLE 4-31.
TABLE 4-31. Overview of the parameters studied during the long-term immersion tests.
No. of T Condition Backfill material Saturation media experiment (°C)
A 140 oxic bentonite(1) deionised water B 140 anoxic bentonite(1) SICW1(3) C 90 anoxic bentonite SICW1(3) D 90 anoxic Boom clay deionised water E 16 anoxic Boom clay deionised water F 140 oxic Foca clay(2) deionised water
(1) bentonite denotes a mixture of 60% smectite clay ‘FoCa’ , 35% sand and 4% graphite. (2) FoCa clay denotes a mixture of 65% smectite clay ‘FoCa’ and 35% sand, i.e. bentonite without graphite. (3) - 2- SICW1 : Synthetic Interstitial Boom Clay Water solution containing a.o. 27 mg/L Cl and 0.2 mg/L SO4 .
Only AISI 316L hMo grade stainless steel was subjected to long-term immersion testing (periods lasting as long as 4 years). The samples were tested in the ‘as-
122 received’ and ‘polished’ condition. The polished samples were wet-polished to a 600-grit finish. Also samples with an artificial crevice were added.
A.4.2.2 Experimental results
A.4.2.2.1 In situ corrosion experiments
A.4.2.2.1.1 Carbon steel
The average general corrosion rates and the maximum pit depths of carbon steel, loaded on the type I (direct contact with clay), type II (humid clay atmosphere) and type III (concrete saturated clay atmosphere) corrosion experiments, are summarised in TABLE 4-32. The average general corrosion rate was calculated after chemical cleaning according to the guidelines of the procedure ASTM G1-90. The maximum pit depth was determined by optical microscopy.
The following conclusions can be drawn from the results given in TABLE 4-32 [166- 168,173,174]:
• the uniform corrosion rate increased with increasing temperature, e.g. for the experiments exposing the samples directly to Boom clay, the uniform corrosion rate increased from 1.8 µm/year at 16°C to 8.6 µm/year at 170°C. • the uniform corrosion rate decreased with increasing period of exposure, e.g. for the experiment exposing the samples directly to Boom clay at 90°C, the uniform corrosion rate decreased from 7.7 µm/year to 4.6 µm/year after 1.7 and 7.0 years interaction, respectively. • carbon steel was susceptible to pitting corrosion, with maximum pit depths observed in the parent material up to 240 µm (after a 2 year exposure period at 90°C). • the samples tested in the ‘as received’ condition were attacked more severely than those tested in the ‘polished’ condition (within the duration of the experiments, which is much shorter than the envisaged disposal period. The uniform corrosion rates of the ‘as received’ samples were 1.5 to 13 times higher than for the ‘polished’ samples. Pits formed on the ‘as received’ samples were, generally, 2.5 to 6 times deeper than those observed on the ‘polished’ samples. • the weld region and the heat affected zone were slightly more susceptible to pitting corrosion than the parent material: for the experiment exposing the samples directly to Boom clay at 90°C for 7 years, pits up to 90 and 130 µm deep were measured in the parent material and the weld region, respectively. • the susceptibility to pitting of the weld became even more pronounced in the presence of a γ-radiation field (400 Gy/h): at some places, the weld material was reduced to about 280 µm (initial thickness = 2,450 µm) after a 5 year exposure period at 80°C, as illustrated in FIGURE 4-30. The influence of γ-radiation was investigated in the CERBERUS-experiment. However, a drawback of this experiment was that only one welded carbon steel sample was included. Drawing definite conclusions based on the results of one single test could therefore turn out to be unsound. Therefore further welded carbon steel samples are being tested in
123 the presence of a radiation field in another in situ corrosion experiment (CORALUS). This experiment is still ongoing and results are currently not available. Nevertheless, the observations made during the CERBERUS- experiment stress the fact that welds could turn out to be the weakest link from a corrosion point of view and therefore deserve a more in-depth investigation (corrosion resistance, filler metals, stress relief treatments, etc.).
CORROSION PRODUCTS
280 µm
WELD METAL
FIGURE 4-30. Optical micrograph (magnification: 50×) of the corrosion attack of a carbon steel weld exposed to clay for 5 years at 80°C and a γ dose rate of 400 Gy/h [47,171].
• it was generally assumed that the humid clay and the concrete saturated humid clay atmosphere were less aggressive environments than a direct exposure to Boom clay. However, no appreciable difference in the degree of corrosion attack was observed. The values for vcorr and dmax of the carbon steel samples in all three tested environments were of the same order of magnitude. For the tests at low temperature (16°C), the samples tested in a humid clay and concrete saturated clay atmosphere were attacked slightly more severely than those tested in direct contact with clay. This was believed to be attributed to the combination of the following effects:
- aqueous solutions (clay moisture) have a higher electrical conductivity (i.e. lower electrical resistivity) than solid media (humid clay): the ions must migrate through the electrolyte to supply the metal surface with the electron donors or acceptors necessary for the corrosion reaction to proceed. The concentration of these salts in the electrolyte is important. Resistivity is a measure of the concentration of the ions and how easily they move through the environment. As a consequence, a high resistivity suggests a low corrosion rate because of the low rate of ion diffusion. - the corrosive environment in the experiments exposing the samples to either a humid clay (type II experiments) or a concrete saturated clay atmosphere (type III experiments) was created by penetration of Boom clay porewater through a stainless steel (type II) or a concrete (type III) filter. A thin film of electrolyte (clay water) was then formed on the surface of the samples through condensation. The environmental conditions so created comprised alternating
124 periods of condensation (immersion) and drying (exposure to air). These alternating conditions were, however, not controlled. The occurrence of these cycles of wetting and drying of the samples is well known to accelerate corrosion. - the corrosion behaviour of steel greatly depends on the availability of oxygen. The amount of dissolved oxygen in clay moisture is much higher than in solid clay. Moreover, in solid clay the oxygen will be consumed more rapidly due to the oxidation of pyrite.
TABLE 4-32. Calculated average uniform corrosion rates and maximum pit depths of the carbon steel samples obtained from the in situ corrosion experiments [47,166-168,170,173,174,176].
(4) (5) Corrosion T Exposure γ dose vCORR dMAX medium (°C) period rate (µm/year) (µm) (year) (Gy/h)
170 4.7 0 8.59 n.d. 90 1.7 0 7.68 240 (parent mat.) 90 7.0 0 4.65 90 (parent mat.) 130 (weld region) (1) 1 16 4.7 0 1.81 100-120 (parent mat.) 16(6) 3.0 0 8.57 (as) 260 (as) 5.47 (p) 45 (p) 80(7) 5.0 400 3.8 150 (parent mat.) 20.3 2,170 (weld region)
90 2.0 0 9.03 (as) 180 (as) 3.24 (p) 30 (p) 8.09 (cc) 245/105 (cc)(8) 2(2) 16 2.5 0 5.13 (as) 195 (as) 0.40 (p) 75 (p) 5.18 (cc) 245/65 (cc)(8)
90 2.0 0 0.73 (as) 60 (as) (3) 3 16 2.5 0 9.48 (as) 235 (as) 1.31 (p) 50 (p)
(1) corrosion medium 1: direct contact with Boom clay (type I corrosion experiments). (2) corrosion medium 2: contact with a humid clay atmosphere (type II corrosion experiments). (3) corrosion medium 3: contact with a concrete saturated clay atmosphere (type III corrosion experiments). (4) vCORR: average uniform corrosion rate. (5) dmax: maximum pit depth. (6) samples mounted on the outside of a type III corrosion experiment. (7) results originating from the CERBERUS-experiment. (8) maximum pit depth measured at the edge of the crevice / maximum pit depth measured inside the crevice. as : samples tested in the ‘as received’ condition (no suface treatment was applied). p : samples tested in the ‘polished’ condition. cc : samples with an artificial crevice.
125 Surface analysis of the carbon steel samples (optical microscopy and scanning electron microscopy) revealed that the morphology of the corrosion attack was that of hemispherical pitting. In many occasions, pits had coalesced resulting in the appearance of general corrosion. Also the presence of ‘secondary’ pitting (i.e. the initiation of smaller pits inside already existing pits) was observed, together with intergranular attack at the bottom of the pits [165,166,170,173]. It was believed that the succession of wide pits with abraded edges, giving the surface of the carbon steel samples an aspect of general pitting, is formed by a combination of two effects [167,170]:
• as oxygen will be consumed relatively quickly by the corrosion reactions of the container material after closure of the borehole in the underground repository and as the steel does not build up an effective protective layer, existing pits may coalesce, resulting in a rough, coalesced pit appearance of the steel surface. Theoretically, pit coalescence (shallow pits) could occur if the pits were not surrounded by passive regions on the metal surface [190]. • formed pits are eventually covered with corrosion products that restrict the flow of ionic species into or out of the pits. Deposition of corrosion products at the mouth of the pits (due to the lower solubility of corrosion products in the higher pH environment outside the pits) can inhibit mass transfer. Thus, a pit eventually becomes a small crevice, thereby rounding off the pit edges. This results in a succession of wide and shallow pits [191].
The surface morphology of the corrosion attack was found to be different according to the test temperature, as illustrated in FIGURE 4-31 [36,166,167,170,173]:
• samples corroded for five years at 16°C (initial clay temperature) and samples corroded for two years at 90°C showed evidences of a rather similar type of general pitting (see FIGURE 4-31(a) and FIGURE 4-31(b), respectively). The observations at 16°C revealed maximum pit depths of up to 100-120 µm. The samples tested at 90°C for 2 years showed zones of attack extending typically over 2 to 3 cm². The maximum depth of this form of localised attack was 240 µm. • at 170°C however, a different surface morphology was observed (see FIGURE 4-31(c)). Although the surface still showed signs of general pitting, it was covered with a succession of wider pits with abraded edges. At 170°C, in many cases, the observations by OM of the unetched polished cross section of ultrasonically cleaned samples revealed the presence of cracks in the oxide layers. The origin of these cracks is still somewhat unclear – dehydration due to the high temperature may be a possible cause – but the unprotective nature of the corrosion layers is evident.
126
(a) (b) (c)
FIGURE 4-31. SEM-micrographs of chemically cleaned carbon steel samples after direct expo- sure to Boom clay [36,166,167,170,173]. (a) for 5 years at 16°C, (b) for 2 years at 90°C, and (c) for 5 years at 170°C.
The differences in surface morphology of the carbon steel samples exposed to Boom Clay at various temperatures and for various lengths of time was explained as follows: during the initial aerobic phase of the experiment, corrosion proceeded predominantly, if not only, by pitting, the rate of which was influenced by the test temperature. When oxygen was consumed (i.e. during the subsequent anaerobic phase), corrosion proceeded uniformly, thereby displacing the metal surface parallel to itself or even progressively eroding its irregularities. The similarity in the appearances of the samples corroded for 5 years at 16°C and for 2 years at 90°C could be explained by the fact that although the initial pitting damage would have been higher at 90°C than at 16°C, it was afterwards counterbalanced by a slower but longer homogeneous corrosion for the experiment run at clay temperature (16°C). For the test at 170°C, the influence of the temperature on the kinetics of corrosion and the duration of the experiment will both have been responsible for the eroded aspect of the surface [166,170].
The results of the XRD-analysis [165-167,170] of the corrosion products formed on the carbon steel samples after direct exposure to Boom Clay at 16°C, 90°C, and 170°C are summarised in TABLE 4-33. The presence of SiO2 in the corrosion layer originated from the clay. XRD-analysis also suggested that the nature of the corrosion products is only marginally affected by temperature. These results suggested that the modifications in the surface appearance, as also seen by stereomicroscopy and SEM, were caused by variations in the environment chemistry, either with time or with temperature, and by kinetic factors.
127 TABLE 4-33. Results of XRD analyses on 'blank' carbon steel samples and carbon steel samples exposed to Boom Clay at 16°C, 90°C, and 170°C ("x" = certain; "?" = likely; "??" = possible) [165-167,170].
Phases
Formula Description blank 16°C 90°C 170°C
Fe3O4 magnetite x x x x
α-Fe2O3 hematite x ? ? ?
γ-Fe2O3.H2O lepidocrocite ?? ? x -
γ-Fe2O3 maghemite - - x x α-FeOOH goethite - ?? x ?
SiO2 quartz - x x x
CaCO3 calcite - - - ?
Fe(1-x)S pyrrhotite - - - ?
x : certain. ? : likely. ?? : possible.
Based on the nature of the corrosion products detected in the corrosion layers, the following reaction mechanism was proposed [165]:
• the Fe-oxide products detected in the corrosion layers reflected Fe3+ as the primary ion in the anodic corrosion reaction. • under oxic conditions (immediately after emplacement of the samples in the borehole), these anodic reaction products will react, in neutral and alkaline - solutions, with OH to yield Fe(OH)3, which undergoes a progressive dehydration as follows:
Fe()OH → FeOOH + H O 3 2
2 FeOOH → Fe2O3 ⋅ H 2O → Fe O + H O 2 3 2 • under anoxic conditions (after oxygen consumption), Fe corrodes according to:
Fe + 2 H 2O → Fe()OH 2 + H 2
3 Fe()OH 2 → Fe3O4 + 2 H 2O + H 2
To be able to interpret the influence of the solid clay phase on the formation of the corrosion layers, a comparison was made with corrosion layers formed on 'blank' carbon steel samples, i.e. carbon steel samples stored in the laboratory under ambient conditions and for approximately the same length of time [36,166,167]. The EPMA micrographs of the cross section of carbon steel samples exposed for 5 years
128 to ambient conditions ('blank' sample) and to clay at respectively 16°C and 170°C is shown in FIGURE 4-32.
(a) (b) (c)
FIGURE 4-32. EPMA micrographs of the cross section of carbon steel samples [166,167]. (a) 'blank' sample (exposed to ambient conditions in the laboratory). (b) exposed to clay at 16°C for a period of 5 years. (c) exposed to clay at 170°C for a period of 5 years.
The corrosion layer on the ‘blank’ carbon steel samples appeared to be on average about 10 to 20 µm thick. Moreover, the layer was subdivided into 3 sub-layers, of which the middle one had a higher Fe/O ratio than the 2 others. It could correspond with a layer that was initially protective – possibly magnetite-containing (Fe3O4) – and had cracked, thereby allowing corrosion to proceed further, forming a hematite-type corrosion product (Fe2O3) [36,166,167]. This multi-layered structure of the corrosion layer was less obvious after interaction with clay for 5 years at 16°C and had practically disappeared after 5 years at 170°C, as illustrated in FIGURE 4-32. In both cases, the corrosion layer was approximately 30 to 50 µm thick in total. FIGURE 4-32(b) suggests the presence of two layers of iron oxide. Here again, one oxide layer clearly had a higher Fe/O ratio than the other. It can also be seen that this layer was cracked. This could again suggest that this layer was probably protective initially, but that its rupture had allowed localised corrosion underneath it. Indeed, a rather large pit can be detected on the metallic surface just at the interface of this thin layer [36,166,167].
A.4.2.2.1.2 Stainless steels
The average uniform corrosion rates of the 1803 MoT, AISI 430, AISI 309, AISI 316, AISI 316Ti and UHB 904L grade stainless steels, loaded on the type I (direct contact with clay), type II (humid clay atmosphere) and type III (concrete saturated clay atmosphere) corrosion experiments, are summarised in TABLE 4-34.
The uniform corrosion rates were situated between 0.003 µm/year and 0.15 µm/year. The weld material (GTAW-weld) was not found to be more susceptible to corrosion than the parent material. The uniform corrosion rates were not affected by the presence of a γ-radiation field (400 Gy/h) [47,165,166,170,171,175].
129 Surface analyses (stereomicroscopy, optical microscopy and scanning electron microscopy) indicated that all stainless steels remained unaffected by interaction with clay between 16°C and 170°C (no pitting attack was observed) [47,165,166,170,171,173].
A.4.2.2.1.3 Ni- and Ti-based alloys
The average uniform corrosion rates of the Ni-based alloys (Hastelloy C-4 and Inconel 625) and the Ti-based alloys (IMI 115 and Ti 99.8-Pd), loaded on the type I (direct contact with clay), type II (humid clay atmosphere) and type III (concrete saturated clay atmosphere) corrosion experiments, are summarised in TABLE 4-35.
The uniform corrosion rates of the Ni-based alloys were situated between 0.003 µm/year and 0.018 µm/year. The uniform corrosion rates of the Ti-based alloys were situated between 0.005 µm/year and 0.06 µm/year. The weld material (GTAW- weld) was not found to be more susceptible to corrosion than the parent material. The uniform corrosion rates were not affected by the presence of a γ-radiation field (400 Gy/h) [47,165,166,170,171,175].
Surface analyses (stereomicroscopy, optical microscopy and scanning electron microscopy) indicated that all Ni- and Ti-based alloys remained unaffected by interaction with clay between 16°C and 170°C (no pitting attack was observed) [47,165,166,170,171,173].
A.4.2.2.1.4 Conclusions
The main shortcoming of the in situ corrosion experiments was that the uniform corrosion rates, calculated from the weight losses, reflected an average value over the entire exposure period, i.e. the aerobic and anaerobic phases, because it was impossible to experimentally separate both phases [173,175].
The weight loss technique did not turn out to be particularly well suited for the measurement of very low uniform corrosion rates, because this technique was found to be susceptible to experimental error: the weight losses from coupons tended to be very small, even after long exposure periods, which resulted, on occasions, in weight gains (as indicated in TABLE 4-34 and TABLE 4-35 by negative vCORR-values) rather than weight losses. Another possible explanation for the weight gains is the growth (thickening) of the protective passive layer in the course of the experiment.
130 TABLE 4-34. Calculated average uniform corrosion rates of the stainless steels investigated during the in situ corrosion experiments [47,165,166,168,171].
(4) Corrosion Material T Exposure γ dose vCORR medium (°C) Period rate (µm/year) (year) (Gy/h) Identification Condition
as + w 16 4.7 0 -0.012 1803 MoT as + w 90 1.7 0 0.060 as + 90 7.0 0 0.015 as + ww 170 4.7 0 0.152 AISI 309 as 80 5.0 400 -0.010 as + w 16 4.7 0 -0.013 as + w 90 1.7 0 0.007 (1) AISI 316 1 as + w 90 7.0 0 -0.005 as + w 170 4.7 0 0.013 as 80 5.0 400 0.003 AISI 316Ti as + w 80 5.0 400 0.000 as + w 16 4.7 0 -0.030 as + w 90 1.7 0 -0.017 UHB 904L as + w 90 7.0 0 -0.005 as + w 170 4.7 0 -0.030 as 90 2.0 0 -0.003 p 90 2.0 0 0.013 as 16 2.5 0 0.022 AISI 430 p 16 2.5 0 0.011 as 90 2.0 0 -0.138 as 16 2.5 0 -0.005 1803 MoT as 16 2.5 0 0.021 as 90 2.0 0 0.009 p 90 2.0 0 0.003 AISI 309 ht 90 2.0 0 0.015 2(2) as 16 2.5 0 0.015 p 16 2.5 0 0.021 as 90 2.0 0 0.026 p 90 2.0 0 0.009 AISI 316Ti ht 90 2.0 0 0.132 as 16 2.5 0 0.020 p 16 2.5 0 0.017 as 90 2.0 0 -0.009 p 90 2.0 0 -0.004 UHB 904L as 16 2.5 0 0.016 p 16 2.5 0 0.003 as 90 2.0 0 -0.138 AISI 430 as 16 2.5 0 -0.005 AISI 309 as 90 2.0 0 -0.011 (3) as 90 2.0 0 0.006 3 AISI 316Ti as 16 2.5 0 0.005 as 16 2.0 0 0.017 UHB 904L p 16 2.5 0 0.012
(1) corrosion medium 1: direct contact with Boom Clay (type I corrosion experiments). (2) corrosion medium 2: contact with a humid clay atmosphere (type II corrosion experiments). (3) corrosion medium 3: contact with a concrete saturated clay atmosphere (type III corrosion experiments). (4) vCORR: average uniform corrosion rate (negative values correspond to weight gain). as : samples tested in the ‘as received’ condition (no surface treatment was applied). as + w : samples, containing a GTAW-weld, tested in the ‘as received’ condition. p : samples tested in the polished condition. ht : samples tested after having received a heat treatment.
131 TABLE 4-35. Calculated average uniform corrosion rates of the Ni- and Ti-based alloys investigated during the in situ corrosion experiments [47,165,166,168,171].
(4) Corrosion Material T Exposure γ dose vCORR medium (°C) Period rate (µm/year) (year) (Gy/h) Identification Condition
as + w 16 4.7 0 -0.050 as + w 90 1.7 0 -0.016 HASTELLOY C-4 as + w 90 7.0 0 -0.021 as + w 170 4.7 0 -0.080 as 80 5.0 400 0.010 as + w 80 5.0 400 -0.002 as + w 16 4.7 0 -0.004 as + w 90 1.7 0 -0.040 INCONEL 625 as + w 90 7.0 0 -0.009 as + w 170 4.7 0 -0.063 1(1) as + w 16 4.7 0 0.050 as + w 90 1.7 0 -0.120 as + w 90 7.0 0 -0.016 IMI 115 as + w 170 4.7 0 -0.080 as 80 5.0 400 -0.020 as + w 80 5.0 400 0.010 as + w 16 4.7 0 -0.005 as + w 90 1.7 0 -0.080 as + w 90 7.0 0 -0.070 Ti 99.8-Pd as + w 170 4.7 0 -0.155 as 80 5.0 400 0.005 as + w 80 5.0 400 -0.003 as 90 2.0 0 0.003 p 90 2.0 0 0.018 HASTELLOY C-4 as 16 2.5 0 0.000 p 16 2.5 0 0.023 as 90 2.0 0 0.017 (2) p 90 2.0 0 -0.017 2 IMI 115 as 16 2.5 0 0.005 p 16 2.5 0 -0.010 as 90 2.0 0 0.000 p 90 2.0 0 -0.010 Ti 99.8-Pd as 16 2.5 0 0.008 p 16 2.5 0 -0.009 as 90 2.0 0 0.003 HASTELLOY C-4 as 16 2.5 0 0.014 p 16 2.5 0 0.006 as 90 2.0 0 -0.021 3(3) IMI 115 as 16 2.5 0 0.018 p 16 2.5 0 -0.032 as 90 2.0 0 0.058 Ti 99.8-Pd as 16 2.5 0 0.061 p 16 2.5 0 0.016
(1) corrosion medium 1: direct contact with Boom Clay (type I corrosion experiments). (2) corrosion medium 2: contact with a humid clay atmosphere (type II corrosion experiments). (3) corrosion medium 3: contact with a concrete saturated clay atmosphere (type III corrosion experiments). (4) vCORR: average uniform corrosion rate (negative values correspond to weight gain). as : samples tested in the ‘as received’ condition (no surface treatment was applied). as + w : samples, containing a GTAW-weld, tested in the ‘as received’ condition. p : samples tested in the polished condition.
132 A.4.2.2.2 Laboratory-scale corrosion experiments
A.4.2.2.2.1 Determination of ECORR (free corrosion potential) based on experimental results
The time dependence of ECORR for five experiments, three in bulk solution and two in compacted bentonite, is presented in FIGURE 4-33 and summarised in TABLE 4-36:
• in synthetic oxidised Boom Clay water (SOCW) containing 1,000 mg/L chloride and 216 mg/L sulphate at 90°C under oxic conditions. • in synthetic oxidised Boom Clay water (SOCW) containing 1,000 mg/L chloride and 216 mg/L sulphate at 30°C under anoxic conditions. This media is not relevant to the underground disposal environment, but was merely added for comparison with the previous one. • in synthetic interstitial Boom Clay water (SICW) containing 27 mg/L chloride and 0.2 mg/L sulphate at 30°C under anoxic conditions. • compacted bentonite (saturated with deionised water) at 22°C under oxic conditions. • compacted bentonite (saturated with deionised water) at 30°C under anoxic conditions.
The experiments under oxic conditions were performed in media with a dissolved oxygen concentration of ~8 ppm (the media were in equilibrium with air, under atmospheric pressure, for the entire duration of the experiment). The experiments under anoxic conditions were performed in media with a dissolved oxygen concentration of less than 10 ppb (the experiments were performed in a glove box, under a controlled inert Ar-atmosphere). The experiments in compacted bentonite were performed at room temperature (the temperature in the laboratory was about 22°C ± 1°C; the local temperature in the glove box was about 30°C ± 2°C) [184].
TABLE 4-36. Time dependence of ECORR for various environment/condition compbinations [184].
(1) (2) Corrosion media Environmental T OCP(t0) OCP(te) Time condition (°C) (mVSHE) (mVSHE) (days) Id. No. state
SOCW(3) solution oxic 90 +40 +320 3 SOCW(3) solution anoxic 30 -310 +120 7 SICW(4) solution anoxic 30 -80 -10 6 bentonite(5) slurry oxic 22 +210 +265 120 bentonite(5) slurry anoxic 30 -170 -120 115
(1) open circuit potential measured at the beginning of the experiment. (2) open circuit potential measured at the end of the test duration. (3) Synthetic Oxidised Boom Clay Water containing 1,000 mg/L chloride and 216 mg/L sulphate. (4) Synthetic Interstitial Boom Clay Water containing 27 mg/L chloride and 0.2 mg/L sulphate. (5) compacted bentonite saturated with deionised water (with a liquid to solid ratio of 30/70).
133 700 400 SOCW (oxic) 500 200 compacted bentonite (oxic) SOCW (anoxic) 300
vs. SHE vs. SICW (anoxic) 0 (mV) (mV) vs. SHE vs. (mV) 100 compacted bentonite (anoxic)
CORR CORR
-200 E E -100
-400 -300 01234567 0.0 20.0 40.0 60.0 80.0 100.0 120.0 Time (days) Time (days)
(a) (b)
FIGURE 4-33. Time dependence of ECORR for AISI 316L hMo grade stainless steel - 2- (a) in solutions: (i) SOCW containing 1,000 mg/L Cl and 216 mg/L SO4 (oxic condition, - 2- 90°C); (ii) SOCW containing 1,000 mg/L Cl and 216 mg/L SO4 (anoxic - 2- condition, 30°C); (iii) SICW containing 27 mg/L Cl and 0.2 mg/L SO4 (anoxic condition, 30°C). Adapted from [180,184]. (b) in compacted bentonite: (iv) oxic condition, 22°C; (v) anoxic condition, 30°C. Adapted from [184,189].
The same trend could be observed for all three experiments performed in solutions (FIGURE 4-33(a)), viz. ECORR increased with time. ECORR changed most rapidly during the first 10-12 hours after immersing the samples into the solutions. Under anoxic conditions, both in the SOCW- and SICW-solutions, ECORR initially decreased slightly, followed by a sharp increase. The largest increase was observed in SOCW under anoxic conditions: an increase of 430 mV was found over the course of 7 days. The smallest increase was observed in SICW under anoxic conditions: ECORR only shifted about 70 mV to a more noble value. In SOCW under oxic conditions, ECORR increased from an initial value of +36 mVSHE to a value of +320 mVSHE over the course of 3 days. In SOCW, under anoxic conditions, ECORR reached a stationary value of approximately +120 mVSHE after about 7 days. In SICW, ECORR reached a stationary value of approximately -10 mVSHE after about 6 days [184].
In compacted bentonite (FIGURE 4-33(b)), under oxic conditions, ECORR initially decreased from +652 mVSHE to +208 mVSHE. After this steep decrease, ECORR gradually shifted to a more positive value. After about 120 days, ECORR had reached a value of approximately +265 mVSHE. The potential peak after about 75 days was believed to be caused by the initiation and repassivation of one or more pits. However, examination of the samples after the experiment did not reveal any signs of pitting [184,189]. The curve recorded under anoxic conditions displayed several gaps. This was due to the fact that on several occasions, the computer data-acquisition system crashed and the recorded data could not be recovered. After about 115 days, a sudden increase of ECORR was observed (the potential shifted from -120 mVSHE to +80 mVSHE). However, at the end of the experiment, it was found that the reference electrode had dried out. These data were therefore not considered reliable. Under anoxic conditions, ECORR initially decreased from +117 mVSHE to -173 mVSHE. However,
134 despite the fluctuations, a trend of an increasing ECORR was readily apparent. After about 115 days, ECORR had reached a value of approximately -120 mVSHE [184,189].
A.4.2.2.2.2 Pitting corrosion studies
The susceptibility of the different container materials to pitting corrosion was determined by means of cyclic potentiodynamic polarisation (CPP) measurements. For each alloy/solution combination, at least three CPP scans were recorded, and from the obtained polarisation curves, average values of the pit nucleation potential, ENP, and the protection potential, EPP, were determined. All CPP measurements were recorded at a scan rate of 1 mV/s (forward and backward scan) and the scan was reversed when the current density reached a value of 1 mA/cm².
The specimens had an exposed surface area of ~2 cm² and were wet-ground to a 600-grit finish (SiC-paper). After grinding, the specimens were cleaned with ethanol, rinsed in deionised water and dried in hot air. The test solutions were de-aerated with purified nitrogen for 15 minutes prior to each experiment.
After termination of the CPP-experiments, the surface of each sample was investigated and the number of pits (pit distribution), pit diameter and pit depth were determined. Pit depths were measured microscopically. A digital micrometer was attached to the microscope in such a way that the movement of the stage relative to the microscope body could be measured. The pit depths were determined by recording the difference between the positions of focus on the surfaces of adjacent regions of corroded (bottom of pit) and uncorroded (specimen surface at the lip of pit).
A.4.2.2.2.2.1 Experimental results
In the framework of various national and EC corrosion studies, a vast amount of pitting and protection potential data were gathered under different conditions:
• oxic conditions, 16°C (see TABLE 4-37). • oxic conditions, 90°C (see TABLE 4-38). • oxic conditions, 140°C (see TABLE 4-39). • anoxic conditions, 16°C (see TABLE 4-40). • anoxic conditions, 90°C (see TABLE 4-41).
135 TABLE 4-37. Pitting potentials (ENP) and protection potentials (EPP) for various candidate container materials (versus SHE) in SOCW and SBW under oxic conditions (16°C).
Chemical composition C-steel Stainless steel Ni-alloy Ti-alloy
- 2- 2- Media [Cl ] [SO4 ] [S2O3 TStE 355 316L 316L hMo 316Ti 309S 904L 1925hMo HAST. C-4 Ti 99.8-Pd (mg/L) (mg/L) ] (mg/L)
100 216 0 u.c. n.p. n.p. n.p. n.p. n.p. +1487* (+883) n.p. n.p. 1,000 216 0 u.c. +1008 (+408) n.p. +1145 (+330) +1450* (+206) n.p. +1440* (+894) n.p. n.p. 10,000 216 0 u.c. +735 (+372) +1470* (+64) +855 (+202) +1045* (+114) n.p. +1383* (+957) n.p. n.p. 20,000 216 0 ------50,000 216 0 ------100 5,400 0 u.c. n.p. n.p. n.p. +1467* (+985) n.p. +1347* (+936) n.p. n.p. 1,000 5,400 0 u.c. n.p. n.p. n.p. +1467* (+938) n.p. +1377* (+942) n.p. n.p. * SOCW(1) 10,000 5,400 0 u.c. +815 (+322) +1280 (+6) +905 (+188) +1040* (+111) n.p. +1347* (+923) n.p. n.p. 20,000 5,400 0 ------50,000 5,400 0 ------1,000 216 2 u.c. +962 (+428) n.p. +1162 (+305) +1427* (+116) n.p. +1487* (+908) n.p. n.p. 1,000 216 20 u.c. +1018 (+368) n.p. +862 (+235) +1322* (-) n.p. +1447* (+909) n.p. n.p. 1,000 216 50 u.c. +982 (+355) n.p. +1075 (150) +1483* (-) n.p. +1370* (+831) n.p. n.p. 1,000 216 100 u.c. +885 (+195) n.p. +1017 (+205) +1297* (-) n.p. +1287* (+742) n.p. n.p. 1,000 216 200 - - n.p. - - n.p. - n.p. n.p.
100 1,700 0 - - n.p. - - n.p. - - - SBW(2) 1,000 1,700 0 - - n.p. - - n.p. - - - 10,000 1,700 0 - - +1305* (+48) - - n.p. - - -
The data in bold represent the pitting potential (ENP). The data between brackets represent the protection potential (EPP). All reported potentials are given versus SHE (Standard Hydrogen Electrode). u.c. : uniform corrosion. n.p. : no pitting. (1) Synthetic Oxidised Boom Clay Water. (2) Synthetic Bentonite Water. * the pitting potential, ENP, was determined as the potential where the current density, I, reached a value of 20 µA/cm².
Source: SMAILOS E., CUÑADO M.A., KURSTEN B., AZKARATE I. and MARX G., to be published. “Long-Term Performance of Candidate Materials for HLW/Spent Fuel Disposal Containers,” Final Report to the EC for the period November 2000 – January 2004.
136 TABLE 4-38. Pitting potentials (ENP) and protection potentials (EPP) for various candidate container materials (versus SHE) in SOCW and SBW under oxic conditions (90°C).
Media Chemical composition C-steel Stainless steel Ni-alloy Ti-alloy
- 2- 2- [Cl ] [SO4 ] [S2O3 TStE 355 316L 316L hMo 316Ti 309S 904L 1925hMo HAST. C-4 Ti 99.8-Pd (mg/L) (mg/L) ] (mg/L)
100 216 0 u.c. n.p. +1195 (+165) +1240* (+217) +1290* (+142) +1122 (+333) n.p. n.p. n.p. 1,000 216 0 u.c. +473 (+68) +556 (+109) +590 (+110) +431 (+85) +782 (+196) +1070* (+176) n.p. n.p. 10,000 216 0 u.c. +305 (+32) +394 (+7) +377 (+72) +282 (+15) +408 (+98) +627 (+91) n.p. n.p. 20,000 216 0 ------50,000 216 0 ------100 5,400 0 u.c. n.p. n.p. n.p. n.p. n.p. +1230* (+734) n.p. n.p. 1,000 5,400 0 u.c. +1230* (+189) +1220* (+189) +1217* (+176) +1096 (+139) +1310* (+240) +1220* (+263) n.p. n.p. SOCW(1) 10,000 5,400 0 u.c. +432 (+125) +433 (+18) +347 (+27) +265 (-110) +416 (+145) +944 (+121) n.p. n.p. 20,000 5,400 0 ------50,000 5,400 0 ------1,000 216 2 u.c. +580 (+132) +517 (+94) +514 (+122) +370 (+127) +706 (+134) +862* (+162) n.p. n.p. 1,000 216 20 u.c. +552 (+107) +554 (+97) +592 (+30) +88 (+58) +784 (+161) +837* (+157) n.p. n.p. 1,000 216 50 u.c. +483 (+40) +524 (+11) +465 (0) +128 (-126) +760 (+165) +790* (+141) n.p. n.p. 1,000 216 100 u.c. +513 (+37) +546 (+40) +556 (-118) +169 (-161) +721 (+40) +675* (-2) n.p. n.p. 1,000 216 200 - - +491 (-104) - - +892 (+110) - n.p. n.p.
100 1,700 0 - - n.p. - - n.p. - - - SBW(2) 1,000 1,700 0 - - +730 (+56) - - +835 (+294) - - - 10,000 1,700 0 - - +486 (-144) - - +390 (-36) - - -
The data in bold represent the pitting potential (ENP). The data between brackets represent the protection potential (EPP). All reported potentials are given versus SHE (Standard Hydrogen Electrode). u.c. : uniform corrosion. n.p. : no pitting. (1) Synthetic Oxidised Boom Clay Water. (2) Synthetic Bentonite Water. * the pitting potential, ENP, was determined as the potential where the current density, I, reached a value of 20 µA/cm².
Source: SMAILOS E., CUÑADO M.A., KURSTEN B., AZKARATE I. and MARX G., to be published. “Long-Term Performance of Candidate Materials for HLW/Spent Fuel Disposal Containers,” Final Report to the EC for the period November 2000 – January 2004.
137 TABLE 4-39. Pitting potentials (ENP) and protection potentials (EPP) for various candidate container materials (versus SHE) in SOCW and SBW under oxic conditions (140°C) [247].
Chemical composition C-steel Stainless steel Ni-alloy Ti-alloy
Media - 2- 2- [Cl ] [SO4 ] [S2O3 ] TStE 355 316L 316L hMo 316Ti 904L Hastelloy C-4 Hastelloy C-22 Ti 99.8-Pd (mg/L) (mg/L) (mg/L)
100 216 0 u.c. +491 (-124) +569 (-110) +530 (-159) +1023 (+10) n.p. n.p. n.p. 1,000 216 0 u.c. +214 (-156) +266 (-199) +265 (-177) +511 (-181) n.p. n.p. n.p. 10,000 216 0 u.c. +94 (-247) +172 (-153) +124 (-309) 326 (-151) +363 (-208)(3) n.p. n.p. (3) 20,000 216 0 u.c. +52 (-226) +87 (-284) +90 (-336) +256 (-191) +330 (-222) n.p. n.p. 50,000 216 0 u.c. -15 (-372) -2 (-325) +74 (-350) +150 (-246) +342 (-183)(3) +334 (-101) n.p. 100 5,400 0 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 5,400 0 u.c. +555 (-114) +680 (-144) +549 (-145) +581 (-189) n.p. n.p. n.p.
10,000 5,400 0 u.c. +121 (-187) +147 (-336) +218 (-412) +425 (-279) n.p. n.p. n.p. (3) SOCW(1) 20,000 5,400 0 u.c. +62 (-190) +136 (-333) +141 (-344) +347 (-397) +401 (-188) n.p. n.p. 50,000 5,400 0 u.c. (3) +348 (-118) n.p. -8 (-276) +69 (-279) +133 (-343) +197 (-204) +301 (-204) 1,000 216 20 u.c. n.p. n.p. n.p. +301 (-128) +301 (-154) +269 (-153) +391 (-202) 1,000 216 50 u.c. n.p. n.p. n.p. +239 (-145) +267 (-213) +276 (-180) +522 (-134) 1,000 216 100 u.c. n.p. n.p. n.p. +74 (-190) +282 (-202) +285 (-172) +475 (-165) 1,000 216 200 u.c. n.p. n.p. n.p. +69 (-184) +153 (-204) +277 (-207) +554 (-128) 1,000 5,400 20 u.c. n.p. n.p. n.p. +570 (-125) +401 (-129) +499 (-148) +950 (-36) 1,000 5,400 50 u.c. n.p. n.p. n.p. +614 (-127) +626 (-149) +418 (-213) +789 (-37) 1,000 5,400 100 u.c. +904* (-183) +686 (-174) +683 (-278) +714 (-66) n.p. n.p. n.p. 1,000 5,400 200 u.c. n.p. +589 (-238) +391 (-216) +374 (-278) +671 (-101) n.p. n.p.
100 1,700 0 u.c. n.p. +876 (-232) +854 (-156) +810 (-171) +847 (-154) tarnish n.p. 1,000 1,700 0 u.c. n.p. n.p. +629 (-278) +555 (-239) +461 (-251) +599 (-198) tarnish 10,000 1,700 0 u.c. n.p. n.p. +199 (-319) +231 (-289) +225 (-308) +142 (-278) tarnish 20,000 1,700 0 u.c. n.p. n.p. +59 (-272) +161 (-267) +114 (-346) +96 (-267) tarnish (2) 50,000 1,700 0 u.c. tarnish n.p. n.p. SBW -20 (-388) +57 (-302) +40 (-321) -5 (-359) 1,000 1,700 20 u.c. n.p. n.p. n.p. +580 (-260) +501 (-244) +366 (-226) +440 (-210) 1,000 1,700 50 u.c. n.p. n.p. n.p. +675 (-255) +646 (-244) +399 (-302) +467 (-229) 1,000 1,700 100 u.c. n.p. n.p. n.p. +456 (-240) +639 (-266) +424 (-285) +512 (-233) 1,000 1,700 200 u.c. n.p. n.p. n.p. +472 (-266) +555 (-234) +586 (-303) +326 (-217)
The data in bold represent the pitting potential (ENP). The data between brackets represent the protection potential (EPP). All reported potentials are given versus SHE (Standard Hydrogen Electrode). u.c. : uniform corrosion. n.p. : no pitting. (1) Synthetic Oxidised Boom Clay Water. (2) Synthetic Bentonite Water. (3) these specimens were found to be subjected to crevice corrosion (evidences of attack was only observed at the interface between the specimen and the mounting resin). * the pitting potential, ENP, was determined as the potential where the current density, I, reached a value of 20 µA/cm².
138 TABLE 4-40. Pitting potentials (ENP) and protection potentials (EPP) for various candidate container materials (versus SHE) in SICW, SBW and RCW under anoxic conditions (16°C).
Chemical composition C-steel Stainless steel Ni-alloy Ti-alloy
Media - 2- 2- [Cl ] [SO4 ] [S2O3 ] TStE 355 316L 316L hMo 316Ti 904L Hastelloy C-4 Hastelloy C-22 Ti 99.8-Pd (mg/L) (mg/L) (mg/L)
27 0.2 0 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 100 0.2 0 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 0.2 0 u.c. +1199 (+151) +1186 (+269) +1083 (+48) n.p. n.p. n.p. n.p. 10,000 0.2 0 u.c. +759 (+157) +1314 (+123) +857 (+77) n.p. n.p. n.p. n.p. (1) 20,000 0.2 0 u.c. +801 (-15) +1136 (+17) +769 (-64) n.p. n.p. n.p. n.p. SICW 50,000 0.2 0 u.c. +616 (+8) +1156 (-3) +727 (-31) n.p. n.p. n.p. n.p. 1,000 0.2 20 u.c. n.p. n.p. +1036 (-22) n.p. n.p. n.p. n.p. 1,000 0.2 50 u.c. n.p. n.p. +1043 (-16) n.p. n.p. n.p. n.p. 1,000 0.2 100 u.c. n.p. n.p. +978 (+11) n.p. n.p. n.p. n.p. 1,000 0.2 200 u.c. n.p. n.p. n.p. n.p. n.p. +1220 (+18) +1184 (-24)
100 1,700 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 10,000 1,700 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p. (2) 20,000 1,700 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p. SBW 50,000 1,700 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 20 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 50 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 100 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 200 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p.
(3) RCW 27 <0.1 0 n.p. n.p. n.p. n.p. n.p. n.p. n.p. n.p.
The data in bold represent the pitting potential (ENP). The data between brackets represent the protection potential (EPP). All reported potentials are given versus SHE (Standard Hydrogen Electrode). u.c. : uniform corrosion. n.p. : no pitting. (1) Synthetic Interstitial Boom Clay Water. (2) Synthetic Bentonite Water. (3) Real Interstitial Boom Clay Water.
139 TABLE 4-41. Pitting potentials (ENP) and protection potentials (EPP) for various candidate container materials (versus SHE) in SICW, SBW and RCW under anoxic conditions (90°C).
Chemical composition C-steel Stainless steel Ni-alloy Ti-alloy
Media - 2- 2- [Cl ] [SO4 ] [S2O3 ] TStE 355 316L 316L hMo 316Ti 904L Hastelloy C-4 Hastelloy C-22 Ti 99.8-Pd (mg/L) (mg/L) (mg/L)
27 0.2 0 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 100 0.2 0 u.c. +914 (+243) n.p. +715 (+120) n.p. tarnish n.p. n.p. 1,000 0.2 0 u.c. +844 (+83) +645 (+30) +684 (+122) +798 (+151) tarnish tarnish n.p. 10,000 0.2 0 u.c. +420 (-50) +504 (-7) +463 (-49) +560 (+116) tarnish tarnish n.p. SICW(1) 20,000 0.2 0 u.c. +295 (-92) +497 (-100) +428 (-80) +548 (+61) tarnish tarnish n.p. 50,000 0.2 0 u.c. +223 (-136) +401 (-121) +312 (-115) +420 (-102) tarnish tarnish n.p. 1,000 0.2 20 u.c. +578 (+31) +694 (+39) +737 (+88) +898 (+161) n.p. n.p. n.p. 1,000 0.2 50 u.c. +617 (+39) +724 (+14) +656 (+17) +842 (+185) n.p. n.p. n.p. 1,000 0.2 100 u.c. +738 (-21) +805 (-73) +691 (+8) +891 (+223) n.p. n.p. n.p. 1,000 0.2 200 u.c. +576 (-103) +774 (-73)) +636 (-146) +818 (+195) n.p. n.p. n.p.
100 1,700 0 u.c. n.p. n.p. n.p. n.p. n.p. n.p. n.p. 1,000 1,700 0 u.c. +722 (-7) +681 (-105) +431 (-3) +980 (+132) n.p. n.p. n.p. 10,000 1,700 0 u.c. +416 (-109) +482 (-130) +229 (-79) +578 (-66) n.p. n.p. n.p. 20,000 1,700 0 u.c. tarnish n.p. n.p. SBW(2) +187 (-146) +261 (-64) +304 (-88) +348 (-88) 50,000 1,700 0 u.c. +213 (-181) +343 (-102) +342 (-133) +407 (-80) tarnish n.p. n.p. 1,000 1,700 20 u.c. +819 (-64) +632 (-49) +853 (-52) +834 (+34) n.p. n.p. n.p. 1,000 1,700 50 u.c. +709 (-26) +606 (-45) +321 (-223) +426 (-13) n.p. n.p. n.p. 1,000 1,700 100 u.c. +776 (-70) +879 (+24) +609 (-175) +875 (+85) n.p. n.p. n.p. 1,000 1,700 200 u.c. n.p. n.p. n.p. +772 (-13) +693 (-154) +629 (-126) +871 (+122)
(3) 27 <0.1 0 u.c. n.p. n.p. n.p. RCW n.p. n.p. n.p. n.p.
The data in bold represent the pitting potential (ENP). The data between brackets represent the protection potential (EPP). All reported potentials are given versus SHE (Standard Hydrogen Electrode). u.c. : uniform corrosion. n.p. : no pitting. (1) Synthetic Interstitial Boom Clay Water. (2) Synthetic Bentonite Water. (3) Real Interstitial Boom Clay Water.
140 A.4.2.2.2.2.2 Discussion
A.4.2.2.2.2.2.1 Carbon steel
The corrosion behaviour of the corrosion-allowance material carbon steel was that of active corrosion in all tested media (even at chloride levels up to 50,000 mg/L) [68,152,160,183].
A.4.2.2.2.2.2.2 Stainless steels
To facilitate the interpretation of the ENP- and EPP-data with respect to the long-term pitting behaviour, the value of the free corrosion potential (ECORR), determined in their respective test environments and corresponding conditions, was added to the graphs as a dotted horizontal line.
A.4.2.2.2.2.2.2.1 Influence of alloy composition
The various stainless steels that were studied as candidate container material for the disposal of HLW/spent fuel in an argillaceous disposal environment were ranked according to their pitting resistance, as ilustrated in FIGURE 4-34.
AISI 309S < AISI 316L, AISI 316L hMo, AISI 316Ti < UHB 904L < CRONIFER 1925hMo
increasing pitting resistance
FIGURE 4-34. Ranking of various stainless steels studied as candidate container material in an argillaceous disposal environment according to their pitting resistance.
AISI 309S had the lowest resistance to pitting, while Cronifer 1925hMo exhibited the highest resistance to pitting, i.e. the ENP-values were always found to be several hundreds of millivolts more noble for Cronifer 1925hMo compared to the other stainless steels. The difference in pitting susceptibility was less pronounced for EPP. AISI 316L, AISI 316L hMo and AISI 316Ti had comparable pitting charcateristics, i.e. the ENP-values generally differed less than 150 mV.
A.4.2.2.2.2.2.2.2 Influence of temperature
Pitting was initiated more easily with increasing temperature, i.e. ENP and EPP shifted towards lower potentials (towards active region) as the temperature was increased. This is illustrated for AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grade stainless steel in SOCW (216 mg/L sulphate) containing 1,000 mg/L and 10,000 mg/L chloride in FIGURE 4-35.
141 1300 1600
1100
1200 900
700 800 ) ) SHE SHE 500
(mV ECORR = +320 mVSHE (mV
E E ECORR = +320 mVSHE 400 300
100 0
-100
-300 -400 0 20 40 60 80 100 120 140 160 0 20 40 60 80 100 120 140 160 Temperature (°C) Temperature (°C)
♦ ENP, 316L ■ ENP, 316L hMo ▲ ENP, 316Ti ● ENP, 904L
◊ EPP, 316L □ EPP, 316L hMo EPP, 316Ti ○ EPP, 904L
(a) (b)
FIGURE 4-35. Influence of temperature on ENP and EPP for 316L, 316L hMo, 316Ti and 904L 2- grade stainless steel in SOCW containing 216 mg/L SO4 under oxic conditions. (a) SOCW containing 1,000 mg/L Cl-. (b) SOCW containing 10,000 mg/L Cl-.
The influence of temperature on the susceptibility of AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grade stainless steel to pitting corrosion in SICW 2- containing 1,700 mg/L SO4 under anoxic conditions is presented in FIGURE 4-36. The ENP-values (FIGURE 4-36(a)) remained well above the free corrosion potential (-10 mVSHE) for all tested alloys at 16°C and 90°C (up to 50,000 mg/L chloride). At 16°C, long-term pitting problems are expected to occur in solutions containing 20,000 mg/L chloride since the EPP-values (FIGURE 4-36(b)) dropped below the free corrosion potential. Increasing the temperature to 90°C, however, caused the EPP- values to drop below the free corrosion potential at much lower chloride levels (approximately 10,000 mg/L).
142 1400 300
1100 200 ) 800 ) 100 SHE SHE (mV (mV
ECORR = -10 mVSHE NP PP E 500 E 0
200 -100
ECORR = -10 mVSHE
-100 -200 0 10000 20000 30000 40000 50000 0 10000 20000 30000 40000 50000 [chloride] (mg/L) [chloride] (mg/L)
♦ 316L, 90°C ■ 316L hMo, 90°C ▲ 316Ti, 90°C ● ENP, 904L, 90°C ◊ 316L, 16°C □ 316L hMo, 16°C 316Ti, 16°C
(a) (b)
FIGURE 4-36. Influence of temperature on ENP and EPP for 316L, 316L hMo, 316Ti and 2- UHB 904L grade stainless steel in SICW containing 1,700 mg/L SO4 . (a) influence on ENP. (b) influence on EPP.
A.4.2.2.2.2.2.2.3 Influence of chloride
The pitting resistance decreased with increasing chloride content. The ENP-values were drastically shifted towards more active potentials with increasing chloride concentration.
The influence of chloride on ENP and EPP for 316L, 316L hMo, 316Ti and UHB 904L grade stainless steel under oxic (SOCW) and anoxic (SICW and SBW) conditions is shown in FIGURE 4-37 and FIGURE 4-38, respectively. In the absence of oxygen ([O2] < 10 ppb, anoxic conditions), the ENP-values remained well above the free corrosion potential (ECORR) for all test media (even up to 50,000 mg/L chloride) and for all tested stainless steels (FIGURE 4-38). In the presence of oxygen ([O2] Q 8 ppm, oxic conditions), the ENP-values dropped below the free corrosion potential for test media containing approximately 10,000 and 1,000 mg/L chloride at 90 and 140°C, respectively (FIGURE 4-37). The protection potential, EPP, was situated below the protection potential for all test media and for all tested stainless steels in the presence of oxygen, while under anoxic conditions the EPP-values remained more noble than the free corrosion potential for test media with a chloride content lower than 10,000-20,000 mg/L.
143 1300 1000
1100 800
900 600 ) 700 ) E = +320 mV SHE CORR SHE SHE 400 (mV 500 (mV E E 200 ECORR = +320 mVSHE
300 0
100 -200
-100 -400 0 2000 4000 6000 8000 10000 0 10000 20000 30000 40000 50000 [chloride] (mg/L) [chloride] (mg/L)
♦ ENP, 316L ■ ENP, 316L hMo ▲ ENP, 316Ti ● ENP, 904L
◊ EPP, 316L □ EPP, 316L hMo EPP, 316Ti ○ EPP, 904L
FIGURE 4-37. Influence of chloride on ENP and EPP for 316L, 316L hMo, 316Ti and UHB 904L grade stainless steel in SOCW containing 216 mg/L sulphate under oxic conditions. (a) 90°C. (b) 140°C.
1000 1100
800 900
700 ) 600 )
500 400 300 (mV vs.SHE (mV vs.SHE 200 E E 100 ECORR = -10 mVSHE ECORR = -120 mVSHE 0 -100
-200 -300 0 10000 20000 30000 40000 50000 0 10000 20000 30000 40000 50000 [chloride] (mg/L) [chloride] (mg/L)
♦ ENP, 316L ■ ENP, 316L hMo ▲ ENP, 316Ti ● ENP, 904L
◊ EPP, 316L □ EPP, 316L hMo EPP, 316Ti ○ EPP, 904L
(a) (b)
FIGURE 4-38. Influence of chloride on ENP and EPP for 316L, 316L hMo, 316Ti and UHB 904L grade stainless steel in SICW and SBW under anoxic conditions at 90°C. (a) SICW containing 0.2 mg/L sulphate. (b) SBW containing 1,700 mg/L sulphate.
144 An approach, based on estimating the maximum chloride level below which pitting is not expected to occur, was introduced as a tool to determine the suitability of stainless steels as candidate container material in argillaceous disposal environments containing chloride.
The maximum chloride level below which pitting is not expected to occur was calculated based on the linear relationship between the pitting potential, ENP, and the chloride concentration, [Cl-], of the type [88,91]
− ENP = A − B ⋅ log[Cl ] (20) where B is a constant depending on the alloy composition, supporting electrolyte composition, measurement technique, etc.
From equation (20), it is possible to estimate the maximum chloride level below which pitting is not expected to occur immediately in environments relevant for various underground repository conditions (conditions evolving in the course of the disposal period). The maximum chloride level is determined as the intersection between the regression line and the actual value of ECORR.
As an example, the regression lines and the fitting results for AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grade stainless steels in SOCW 2- containing 216 and 5,400 mg/L SO4 under oxic conditions (140°C) are presented in FIGURE 4-39.
♦ AISI 316L ▲ AISI 316L hMo ● AISI 316Ti ■ UHB 904L
(a) (b)
- FIGURE 4-39. Linear fitting of (ENP, log[Cl ]) plots for 316L, 316L hMo, 316Ti and UHB 904L grade stainless steel in SOCW (synthetic oxidised Boom Clay water) at 140°C [247]. (a) SOCW containing 216 mg/L sulphate. (b) SOCW containing 5,400 mg/L sulphate.
145 The estimated chloride levels above which pitting is expected to occur in environments relevant for various conditions in argillaceous formations are presented in TABLE 4-42.
TABLE 4-42. Estimated chloride levels above which pitting is expected to occur in environments relevant for various disposal conditions in argillaceous formations. Adapted from [247].
2- - Material Experimental T [SO4 ] [Cl ]max conditions (°C) (mg/L) (mg/L)
AISI 316L 140 216 0.56 × 10³ AISI 316Ti SOCW 140 216 0.98 × 10³ AISI 316L hMo AEROBIC 140 216 1.22 × 10³ UHB 904L 140 216 10.80 × 10³
AISI 316L 140 5,400 4.0 × 10³ AISI 316Ti SOCW 140 5,400 5.9 × 10³ AISI 316L hMo AEROBIC 140 5,400 6.8 × 10³ UHB 904L 140 5,400 20.6 × 10³
AISI 316L 140 1,700 6.9 × 10³ AISI 316Ti SBW 140 1,700 7.6 × 10³ AISI 316L hMo AEROBIC 140 1,700 9.0 × 10³ UHB 904L 140 1,700 6.5 × 10³
AISI 316L 90 216 8.1 × 10³ AISI 316Ti SOCW 90 216 9.2 × 10³ AISI 316L hMo AEROBIC 90 216 9.6 × 10³ UHB 904L 90 216 18.3 × 10³
AISI 316L 90 5,400 13.8 × 10³ AISI 316Ti SOCW 90 5,400 10.7 × 10³ AISI 316L hMo AEROBIC 90 5,400 13.9 × 10³ UHB 904L 90 5,400 12.8 × 10³
AISI 316L hMo SBW 90 1,700 80.5 × 10³ UHB 904L AEROBIC 90 1,700 22.1 × 10³
AISI 316L 90 0.2 0.27 × 106 AISI 316Ti SICW 90 0.2 1.8 × 106 AISI 316L hMo ANAEROBIC 90 0.2 67.8 × 106 UHB 904L 90 0.2 5.7 × 106
AISI 316L 90 1,700 0.35 × 106 AISI 316Ti SBW 90 1,700 108.3 × 106 AISI 316L hMo ANAEROBIC 90 1,700 1.8 × 106 UHB 904L 90 1,700 0.62 × 106
The following conclusions were drawn from TABLE 4-42 with respect to waste disposal in argillaceous formations:
• all four of the tested stainless steels seem to be adequately resistant to pitting corrosion in an argillaceous environment under oxic conditions: under the most severe conditions, i.e. a high temperature (140°C) and a low degree of oxidation 2- (216 mg/L SO4 ), the maximum level of chloride above which pitting is expected to occur for AISI 316L is approximately 560 mg/L, which is more than 6 times higher than the expected chloride content of the near-field environment for the Belgian
146 disposal concept (~90 mg/L Cl-). The other three stainless steels can even resist - environments containing higher amounts of chloride ([Cl ]MAX for AISI 316Ti, AISI 316L hMo and UHB 904L was estimated to be 980, 1220 and 10800 mg/L respectively). • as the near-field environment becomes more oxidised, i.e. the further in the disposal period, the tested stainless steels become more resistant to pitting: - [Cl ]MAX is increased with a factor of about 6-7 (140°C) when the sulphate content increases from 216 to 5,400 mg/L. This effect is most pronounced at 140°C and is almost negligible at 90°C. • the oxygen level has a significant influence on the pitting resistance. As oxygen will be consumed in the underground repository, the tested stainless steels - become much more resistant to pitting: [Cl ]MAX is increased with a factor of about 25, for AISI 316L hMo and UHB 904L, when the conditions change from oxic to anoxic. • when the temperature is increased from 90 to 140°C, the maximum chloride level - above which pitting is expected to occur decreases drastically: [Cl ]MAX decreases 2- approximately 15, 10, and 8 times (216 mg/L SO4 ) for AISI 316L, AISI 316Ti, and AISI 316L hMo respectively. In highly oxidised clay environments 2- - (5,400 mg/L SO4 ), the temperature has a much lower influence on [Cl ]MAX. UHB 904L is less susceptible to temperature.
A.4.2.2.2.2.2.2.4 Influence of sulphate
The influence of the sulphate content was studied by combining the data obtained in SOCW (216 mg/L and 5,400 mg/L sulphate) and SBW (1,700 mg/L sulphate). Under oxic conditions, the free corrosion potential, ECORR, in SOCW and SBW is +320 mVSHE and +265 mVSHE, respectively. A value of +320 mVSHE was used for the interpretation of the polarisation curves because it is the most noble value and therefore imposes the most severe condition. This value is indicated on the curves as a horizontal dotted line.
FIGURE 4-40 shows the influence of the sulphate concentration on ENP and EPP for AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grade stainless steel in clay water solutions containing 1,000 mg/L and 50,000 mg/L chloride at 140°C. Sulphate was found to have an inhibitive effect on the pitting corrosion of the tested stainless steels: the pitting potentials, ENP, increased with increasing sulphate concentration. The ENP-values became more noble than the free corrosion potential, ECORR (+320 mVSHE), when 1,700 mg/L sulphate (or higher) was added to the clay water solutions. The inhibitive effect of sulphate was less pronounced with increasing chloride concentration, i.e. the ENP-values were independent of the sulphate concentration when 20,000 mg/L chloride (or higher) was added to the test solutions (FIGURE 4-40(b)). It was also found that sulphate had no appreciable influence on the protection potential, EPP, regardless of the chloride content.
This inhibitive effect of sulphate on the initiation of pitting was also observed by other researchers [103,107]. Leckie et al. [107] believed that the inhibitive effect of sulphate is attributed to the competitive adsorption with chloride. Also, during the pitting process, the pit environment becomes strongly enriched in anions, such as chlorides and sulphates, by electromigration thereby creating a potential gradient when ionic current flows between the passive surface (cathode) and the pit (anode). Passivation
147 of the pit is prevented if chlorides predominate in the pit [104]. Garner et al. [102] also found that anions concentrate inside pits roughly in proportion to their concentration in the bulk environment.
800 800
600 600
400 ECORR = +320 mVSHE 400 ECORR = +320 mVSHE ) ) SHE SHE 200 200 (mV (mV E E
0 0
-200 -200
-400 -400 0 1000 2000 3000 4000 5000 6000 0 1000 2000 3000 4000 5000 6000 [sulphate] (mg/L) [sulphate] (mg/L)
♦ ENP, 316L ■ ENP, 316L hMo ▲ ENP, 316Ti ● ENP, 904L
◊ EPP, 316L □ EPP, 316L hMo EPP, 316Ti ○ EPP, 904L
(a) (b)
FIGURE 4-40. Influence of sulphate on ENP and EPP for 316L, 316L hMo, 316Ti and 904L grade stainless steel in clay water containing various levels of chloride at 140°C. (a) 1,000 mg/L chloride. (b) 50,000 mg/L chloride.
A.4.2.2.2.2.2.2.5 Influence of thiosulphate
Within the investigated test matrix, the presence of low amounts of thiosulphate (2 mg/L) had no appreciable influence on the pitting behaviour of AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grades stainless steel. The addition of higher amounts of thiosulphate (in the range 20-200 mg/L) did not result in a sudden drop of the pitting potential: all ENP-values remained more noble than the free corrosion potential, as illustrated in FIGURE 4-41. The protection potential, EPP, remained independent of the thiosulphate content.
148 ♦ ENP, 316L ■ ENP, 316L hMo ▲ ENP, 316Ti ● ENP, 904L
◊ EPP, 316L □ EPP, 316L hMo EPP, 316Ti ○ EPP, 904L
(a) (b)
FIGURE 4-41. Influence of thiosulphate on ENP and EPP of AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grade stainless steel in clay water solutions under various conditions. (a) SBW (1,000 mg/L chloride, 1700 mg/L sulphate); oxic conditons; 140°C. (b) SICW (1,000 mg/L chloride and 0.2 mg/L sulphate); anoxic conditions; 90°C.
The influence of thiosulphate on the pit growth rate of AISI 316L hMo and UHB 904L grades stainless steel at 90°C was studied by performing modified pit propagation rate tests. These tests consisted in first polarising the sample to a potential of 1,000 mV (i.e. a potential well above ENP) for five minutes and then holding the potential at 400 mV (i.e. a potential between ENP and EPP) for two hours, in order to allow the initiated pits to propagate. The main advantage of this test is that all specimens were pitted for the same length of time. After the tests, the number and depth of the formed pits were determined as a function of the thiosulphate concentration of the test solutions. The results for UHB 904L grade stainless steel are summarised in TABLE 4-43. Thiosulphate only seemed to have an influence on the initiation process of pitting but not on the propagation: the number of formed pits increased with increasing thiosulphate concentration, whereas the depth of the formed pits remained of the same order of magnitude [184,187].
149 TABLE 4-43. Influence of thiosulphate on the pit density of the formed pits for UHB 904L grade stainless steel in SOCW under oxic conditions at 90°C [184,187].
- 2- 2- [Cl ] [SO4 ] [S2O3 ] Pit density Pit depth (#pits/cm²) (µm)
1000 216 0 90 6-12 1000 216 2 225 5-22 1000 216 20 260 2-9 1000 216 50 200 4-22 1000 216 100 320 5-21
A.4.2.2.2.2.2.2.6 Influence of hydrogen peroxide (gamma radiolytic product)
The pitting potentials (ENP) and protection potentials (EPP) for AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grades stainless steel in SOCW containing varying hydrogen peroxide concentrations at 90°C under oxic conditions are presented in TABLE 4-44.
TABLE 4-44. Influence of hydrogen peroxide (H2O2) on ENP and EPP of AISI 316L, AISI 316L hMo, AISI 316Ti and UHB 904L grades stainless steel in SOCW at 90°C under oxic conditions.
Chemical composition OCP ENP EPP Material - 2- 2- [Cl ] [SO4 ] [S2O3 ] (mVSHE) (mVSHE) (mVSHE) (mg/L) (mg/L) (mol/L)
1,000 216 0 -531 473 68 1,000 216 0.8 +693 742 589 AISI 316L 1,000 216 0.08 +484 579 306 1,000 216 0.008 +425 528 278
1,000 216 0 -489 556 109 1,000 216 0.8 +673 n.h. n.h. AISI 316L hMo 1,000 216 0.08 +539 684 303 1,000 216 0.008 +393 624 310
1,000 216 0 -82 590 110 1,000 216 0.8 +659 819 549 AISI 316Ti 1,000 216 0.08 +487 640 309 1,000 216 0.008 +433 571 275
1,000 216 0 -74 782 196 1,000 216 0.8 +643 n.h. n.h. UHB 904L 1,000 216 0.08 +524 784 481 1,000 216 0.008 +359 741 378
All reported potentials are given versus SHE (Standard Hydrogen Electrode). OCP : Open Circuit Potential (derived from the CPP-curves). n.h. : no hysteresis (the CPP-curves did not show a positive hysteresis, which is typical for the presence of pitting corrosion; surface analyses did however reveal signs of pitting).
Source: SMAILOS E., CUÑADO M.A., KURSTEN B., AZKARATE I. and MARX G., to be published. “Long-Term Performance of Candidate Materials for HLW/Spent Fuel Disposal Containers,” Final Report to the EC for the period November 2000 – January 2004.
150 The addition of hydrogen peroxide seemed to have a dual effect:
• on the one hand, H2O2 seemed to have a beneficial effect on the pitting susceptibility because the ENP- and EPP-values were significantly more positive in the solutions containing H2O2. This could suggest that H2O2 promotes passivation. It was also observed that the ENP-values increased with increasing H2O2- concentration. This is indicative for an inhibiting effect of H2O2 on pitting (this effect was less pronounced than for sulphate). • on the other hand, H2O2 seemed to exert a detrimental effect on the pitting susceptibility because the addition of H2O2 caused a sharp increase of the open circuit potential (OCP). Consequently, the potential difference between ENP and OCP, which can be considered as a parameter that indicates the risk for pitting, was reduced significantly (in some cases the difference was less than 100 mV). Also, the protection potential, EPP, is situated below the OCP when H2O2 was added to the solutions. Under these circumstances, a local damage of the passive layer (either due to pitting or mechanically) will probably never be able to repassivate.
A.4.2.2.2.2.2.3 Ni-alloys
Under oxic conditions, Hastelloy C-4 and Hastelloy C-22 resisted pitting in all test media up to 90°C. At 140°C, Hastelloy C-4 and Hastelloy C-22 suffered from crevice corrosion when chloride was added to the test solutions in the range 20,000– 50,000 mg/L, i.e. signs of attack were observed only at the interface between the specimen and the mounting resin, as illustrated in FIGURE 4-42, while the specimen’s surface remained unaffected [247].
Under anoxic conditions, Hastelloy C-4 and Hastelloy C-22 resisted pitting in all test media (even up to 50,000 mg/L chloride). However, at 90°C, the CPP-curves occasionally showed a positive hysteresis, which is typical for the onset of pitting corrosion (in SICW-solutions containing chloride concentrations in the range 100- 50,000 mg/L and in SBW-solutions containing 20,000 mg/L chloride and higher). This hysteresis was believed to be the result of an alteration of the passive layer because analysis of the specimen’s surface after the experiments only revealed a tarnished surface (i.e. a yellow-orange discoloration), as illustrated in FIGURE 4-43, without any evidences of pitting attack [154].
151
mounting resin specimen
FIGURE 4-42. SEM-micrograph (magnification: 1,500×) of an uncleaned Hastelloy C-4 specimen after CPP-testing (140°C) in 2- - SOCW containing 216 mg/L SO4 and 50,000 mg/L Cl . Signs of attack were only observed at the interface between the specimen and the mounting resin. Adapted from [247].
FIGURE 4-43. General view (optical micrograph, magnification: 50×) of a tarnished surface of a Hastelloy C-22 specimen after CPP-testing in SICW containing 50,000 mg/L Cl- at 90°C under anoxic conditions [154].
152 A.4.2.2.2.2.2.4 Ti-alloys
Only Ti 99.8-Pd was resistant to pitting corrosion under all tested conditions (temperature up to 140°C and chloride levels up to 50,000 mg/L).
153 A.4.3. Granitic environment
A.4.3.1 Corrosion studies relevant to the Swedish and Finnish disposal systems in granite
The following discussions in this chapter are based on the findings and conclusions of a recent joint SKB and Posiva report dealing with copper corrosion under expected conditions in a geologic repository in granitic rock [60].
More detailed information can be found in References [60,129,205–246].
When assessing the extent of corrosion attack, King et al. [60] considered three distinct periods during the container service life:
• corrosion prior to water saturation. This comprises two periods: (i) corrosion during dry storage (i.e. during the period prior to disposal) and (ii) corrosion during the period after emplacement of the copper canister in the disposal hole and before saturation of the backfill. • corrosion during water saturation. It is believed that an uneven increase in the degree of water saturation around the canister will occur over the first several years of saturation. This inhomogeneous increase in water content in the bentonite backfill will probably result in an uneven swelling of the buffer. As a consequence of this, the gap may close in some areas while it remains open in others. Sites where the bentonite first contacts the copper canisters are latent sites for corrosion pits. • corrosion after water saturation. Following saturation of the repository, the environment surrounding the canister will evolve from initially warm and oxic, to eventually cool and anoxic. As a consequence, the corrosion behaviour of the canister will also change with time from an initial period of relatively fast general corrosion accompanied by possible localised corrosion, to a long-term steady-state condition of a low rate of general corrosion with little or no localised attack. Therefore, the impact of general corrosion and various forms of localised corrosion on the long-term performance of copper canisters has been assessed. The forms of localised corrosion the copper canisters will be most likely susceptible to under relevant repository conditions are pitting corrosion, crevice corrosion, ants-nest corrosion, and stress corrosion cracking (SCC). Also, microbially influenced corrosion (MIC) and corrosion induced by radiation have also been considered.
A.4.3.1.1 Corrosion prior to water saturation
The consequences of atmospheric corrosion for the performance of copper canisters were assessed by means of literature surveys for (i) dry storage (i.e. the period between encapsulation of the fuel inside the copper canister and disposal in the geologic repository; in extreme cases this period could last up to 2 years) and (ii) the period following the emplacement of the copper canister in the disposal hole and before saturation of the backfill. During both periods, the canister will be exposed to air (either at the encapsulation plant or at the repository site) [60]. The formation of a water film on the metal surface is of fundamental importance to atmospheric corrosion. The aqueous phase acts as an electrolyte for electrochemical
154 reactions, where the cathodic reaction is the reduction of an oxidant from the atmosphere [60].
The literature surveys all suggested that corrosion attack during the unsaturated phase will probably be uniform in nature and corrosion will have a negligible effect on the canister lifetime [60]: - the total corrosion attack after two years of dry storage is believed to be less than 1 µm. The most likely corrosion product will be a copper oxide. - from experiments exposing copper to an external atmosphere at ambient temperature with similar levels of contaminants (see TABLE 4-45), a corrosion rate of 100 to 300 µm per year has been estimated for repository conditions at 90°C (if the supply of oxygen is unlimited). The dominant surface species will probably be Cu2O. Pitting corrosion is considered not to be possible under these conditions. - the depth of corrosion, based on mass balance considerations, is estimated to be a maximum of 300 µm evenly distributed over the canister surface. - attempts to model the corrosion of copper, taking into account mass transport by diffusion and flow, equilibrium reactions and kinetic processes at the bentonite/canister interface, indicate a conservative corrosion rate of 7 µm/year for oxic conditions.
TABLE 4-45. Maximum levels of pollutants (during storage prior to water saturation either in a surface facility or in the repository prior to saturation) that are detrimental to copper. The levels of air pollution are comparable to those in an urban atmosphere [60].
Contaminant Maximum level (µg/m³)
SO2 100 NO2 75 NH3 <20 H2S <3
A.4.3.1.2 Corrosion during water saturation
A 0.35 m thick layer of highly compacted bentonite with a low hydraulic permeability will be used to surround the canisters in the deposition holes. This bentonite is not saturated during the installation. About 25% of the total amount of water in the buffer must be provided by the host rock. The form and the extent of the corrosion during this phase in the canister service life will depend on (i) the wetting process leading to the final water saturated conditions around the canisters, (ii) the form of its progress, and (iii) its duration [60].
Initially, the bentonite blocks will have a water ratio of 17%. The increased surface temperature of the canister (90°C) will create a temperature gradient in the bentonite resulting in a redistribution of the water and a lowering of the water ratio close to the canister to about 10% [60]. A model developed by Börgesson and Hernelind [204], describing the wetting process of the bentonite, predicts full water saturation after about 12 years. However, this model
155 was based on the assumption that there will be no gap between the bentonite blocks and the canister surface. As a consequence, no definite conclusions can be made concerning how and when the bentonite will contact the canister.
As mentioned earlier, the uneven swelling of the bentonite (due to an inhomogeneous increase in water content) could lead to locally increased corrosion rates (i.e. at sites where there will be a good supply of oxygen, where there is an electrolyte present, and where there is a short distance to a site where electrochemical corrosion of copper could take place). However, since the gaps close gradually as the bentonite swells, the location of these sites will not be constant. The resulting pit depth will be determined by the ratio between the rate of the general corrosion close to the cathodic sites and the rate of growth of the few corrosion pits. Averaged over the whole canister surface, the corrosion can still not exceed the 300 µm limit determined by the amount of available oxygen (see A.4.3.1.1) [60].
The main results, derived from experiments studying the corrosion of the canister during water saturation, are summarised below: • in-situ experiments in the Äspö Hard Rock Laboratory [205,206], exposing copper coupons to bentonite clay (placed in deposition holes drilled into the rock), revealed an average corrosion rate of 3 µm/year. The corrosion attack was somewhat uneven, but pitting was not observed. The corrosion products were found to be a.o. Cu2O and Cu2CO3(OH)2 (malachite). • laboratory immersion experiments [207] exposing copper to compacted bentonite/sand mixtures saturated with a saline synthetic groundwater (0.97 M Cl-) revealed corrosion rates in the range of 30-50 µm/year. Cupric chloride crystals (CuCl2 ·3Cu(OH)2) were formed on top of an oxide layer (possibly Cu2O). • long-term corrosion experiments [208] exposing copper to bentonite and - 2- - groundwater ([Cl ]: 60-7,600 mg/L; [SO4 ]: 60-6,200 mg/L; [HCO3 ]: 570-600 mg/L; T= 80°C), showed that dissolution of copper took place initially when there was oxygen present in the system. But, once the system had become anoxic, the corrosion virtually stopped. In these experiments, no signs of pitting were observed, because, it is believed, the corrosion potential had dropped to levels where pitting corrosion could not be sustained.
A.4.3.1.3 Corrosion after water saturation
During the period following saturation of the repository, the environmental conditions are expected to evolve from warm and oxic to cool and anoxic. As a consequence the corrosion behaviour of the copper canister will also change with time [60]:
• initially, general corrosion will be supported by the reduction of the atmospheric O2 trapped in the bentonite. Redox conditions will be relatively oxidising and the corrosion potential (ECORR) of the canister surface will be relatively positive. Localised corrosion is possible during this period. • as the initially trapped O2 is consumed, the rate of corrosion will become limited by diffusion of O2 to the canister surface. Eventually, conditions will become anoxic and corrosion will be supported by the reduction of H2O if sulphide is present in the clay and groundwater. The rate of corrosion is expected to be limited by the rate of supply of sulphide to the canister surface, and to fall to very low levels indefinitely.
156 Only general corrosion is expected under anoxic conditions, but only if sulphide is present.
In interpreting the experimental results from canister corrosion studies relating to the post-saturation phase, the following assumptions were made [60]:
• the compacted bentonite is fully saturated with groundwater, which implies that the mass transport of species to and from the canister surface is by diffusion through water-filled pores (the more rapid transport of species such as O2 through vapour- filled pores was not considered) [209]. • the bentonite adjacent to the canister surface is always saturated, so that the supply of H2O does not limit the extent of corrosion. • as groundwater saturates the buffer, the porewater Cl- concentration will gradually increase. • initially trapped atmospheric O2 will create oxidising conditions, but as this oxygen is consumed by (i) corrosion of the canister, (ii) reaction with oxidisable mineral impurities and sulphide in the clay, and (iii) microbial activity, the conditions will become anoxic and remain so indefinitely.
A.4.3.1.3.1 General corrosion
A.4.3.1.3.1.1 Determination of ECORR (corrosion potential) and iCORR (corrosion rate) based on modelling studies
FIGURE 4-44 shows the predicted range of ECORR values for copper in compacted clay for dissolved O2 concentrations between 8 ppm (aerated water) and 8 ppb and for chloride concentrations between 0.001 mol/L and 1 mol/L. This figure was constructed using a steady-state ECORR model for the corrosion of copper in oxygen containing chloride solutions [210], under mass-transport conditions similar to those for a copper canister surrounded by compacted clay (1 cm thick clay layer, O2 and - CuCl2 diffusion coefficients a factor of 100 lower than in solution). The predicted ECORR varied between -27 mVSCE and -376 mVSCE depending on the chloride concentration.
157
FIGURE 4-44. Evans' diagram showing the dependence of the corrosion potential of copper on dissolved oxygen and chloride ion concentrations in the presence of compacted clay. Based on the steady-state model of King et al. [210] with a clay - layer of 1 cm thick and O2 and CuCl2 diffusion coefficients a factor 100 smaller than in bulk solution.
FIGURE 4-45 shows the effect of O2 concentration on ECORR (FIGURE 4-45(a)) and iCORR (FIGURE 4-45(b)). Both ECORR and iCORR decrease with decreasing [O2]. The dependence of ECORR and iCORR on log [O2] in 1 mol/L chloride is consistent with transport control of both anodic (Cu dissolution) and cathodic (O2 reduction) reactions. Thus, ECORR decreases by 2.3RT/F V for each factor of ten decrease in [O2] (i.e. by 59 mV/decade) [211] and the rate of corrosion is first order with respect to [O2] (i.e. dlogiCORR/dlog[O2] = 1). The dependence of iCORR on [O2] does not necessarily indicate that the rate of corrosion is limited by oxygen transport, since the - rate of the anodic transport step (the diffusion of CuCl2 away from the Cu surface) - also decreases with [O2] (due to the decrease of the interfacial [CuCl2 ] with decreasing ECORR). The extent of transport control of both reactions diminishes with - - decreasing [Cl ]. In 0.001 mol/L Cl , for example, ECORR decreases by only 47 mV per decade decrease in [O2] (FIGURE 4-45(a)) and dlogiCORR/dlog[O2] = 0.80 (FIGURE 4- 45(b)), both indications that the O2 reduction reaction is partially controlled by the rate of the interfacial reaction (reaction 10b, section A.3.3).
The steady-state model of King et al. [210] was used to predict the dependence of - ECORR and iCORR on [Cl ] in compacted clay. FIGURE 4-46 shows the dependencies - on [Cl ] derived from the Evans' diagram in FIGURE 4-44. The dependence of ECORR - on [Cl ] varies from -105 mV/decade in environments containing 8 ppm O2 to -116 mV/decade in the presence of 8 ppb O2 (FIGURE 4-46(a)). For complete transport control of both anodic and cathodic reactions, the predicted dependence - would be -118 mV/decade at 25°C, i.e. dECORR/dlog[Cl ] = 2 (2.3RT/F), where the factor of 2 corresponds to the complexation of Cu(I) by two Cl- ions (Cl- concentrations were used instead of the more correct Cl- activities). The change in - [Cl ] dependence of ECORR with decreasing [O2] reflects the increasing transport- limitation of the O2 reduction reaction. The corrosion current density is only weakly - - 0.033 - 0.23 dependent on [Cl ], varying from [Cl ] for 8 ppb O2 to [Cl ] in 8 ppm O2 - (FIGURE 4-46(b)). The relative independence of iCORR on [Cl ] suggests that the
158 overall corrosion rate is largely mass-transport limited by the supply of O2 to the Cu - - surface. Mass-transport limitation by the diffusion of CuCl2 away from, or of Cl to, - the Cu surface would result in iCORR being proportional to [Cl ]². The much smaller - predicted dependence of iCORR on [Cl ] is consistent with rate control largely by the supply of O2, especially at lower [O2].
(a) (b)
FIGURE 4-45. Predicted dependencies of the corrosion potential (ECORR) and corrosion current density (iCORR) on oxygen concentration based on the data in FIGURE 4-44 [210]. (a) predicted dependence of ECORR on [O2]. (b) predicted dependence of iCORR on [O2].
(a) (b)
FIGURE 4-46. Predicted dependencies of the corrosion potential (ECORR) and corrosion current density (iCORR) on chloride concentration based on the data in FIGURE 4-44 [210]. - (a) predicted dependence of ECORR on [Cl ]. - (b) predicted dependence of iCORR on [Cl ].
159 A.4.3.1.3.1.2 Determination of ECORR (corrosion potential) and iCORR (corrosion rate) based on experimental results
Immersion experiments performed in compacted bentonite-sand buffer material (to - simulate conditions in a Canadian repository) as a function of [O2] and [Cl ] showed a dependence of the corrosion rate on the diffusion of Cu away from the coupon surface [207,209,212]. The dependence of the corrosion rate on [O2] was lower than that predicted using the steady-state model of King et al. [210] (see FIGURE 4-45(b)). 0.47 On average, the corrosion rate was proportional to [O2] , as opposed to the linear dependence predicted by the steady-state electrochemical model. There are two reasons for the difference between the predicted and observed reaction order: • it is likely that the rate of O2 diffusion was higher in the compacted bentonite than assumed in the steady-state prediction, both because diffusion would have occurred under transient as opposed to steady-state conditions and, possibly, because the bentonite was not totally saturated. • a fraction of the O2 was consumed by the homogeneous oxidation of Cu(I) (reaction 11, section A.3.3), a reaction not included in the steady-state model. Oxygen consumption by reaction 11 would not have been significant in tests at room temperature on which the model was based, since the rate of oxidation is twenty times slower at 25°C compared with the experimental temperature of 95°C used in the immersion tests.
Long-term corrosion tests on copper coupons in compacted bentonite performed at the Äspö Hard Rock Laboratory revealed an estimated mean corrosion rate after one year’s exposure of 3 µm/year [205]. As observed by Aaltonen and Varis [208], copper diffused into the surrounding bentonite. No indications of pitting attack were observed.
The effect of groundwater salinity on the corrosion rate of Cu in compacted bentonite has been reported by King and Kolář [209]. Copper coupons were sandwiched between plugs of Na-bentonite compacted to a dry density of 1.2 Mg/m³ and saturated with synthetic groundwaters of three different salinities ([Cl-] = 0.17 mol/L, 0.97 mol/L, and 2.5 mol/L) and two different [O2] (air or 0.2 vol.% O2/N2 mixture). The experiments were conducted at 95°C for periods between 10 days and 6 months. The experiment was designed such that well-defined 1-D mass-transport conditions were maintained throughout. After recovery of the coupons (after visual examination), the precipitated corrosion products were removed and the corrosion rate was determined from the weight loss of the coupon. FIGURE 4-47 shows the results of the tests that were performed in all three groundwaters. The results suggest that the corrosion rate decreases with increasing salinity. This conclusion was rationalised on the basis of a Cu-transport rate- determining step [207]. The experimental corrosion rates in FIGURE 4-47, especially at the lower [O2], are similar to those observed by Karnland et al. [205] in long-term pilot scale corrosion tests at Äspö (~3 µm/year).
160
FIGURE 4-47. Dependence of the corrosion rate of copper in compacted clay on the salinity of the groundwater. Tests conducted at 95°C for 90 days in groundwater-saturated bentonite clay (dry density = 1.2 Mg/m³). The ground- waters were either initially aerated or equilibrated with a 0.2 vol.% O2/N2 (low [O2]) mixture. In some tests, 0.5 wt.% Fe powder was added to the dry clay prior to compaction [60].
An experiment, originally developed to determine the effect of sulphate-reducing bacteria on the corrosion of a copper canister in a Canadian repository, also revealed the effect of sulphide on ECORR [213]. Copper electrodes were exposed to a 1 mol/L NaCl solution under well-controlled mass-transport conditions, either by rotating the electrode in bulk solution or by placing a 0.1 cm thick layer of compacted bentonite between the electrode and the bulk solution. The ECORR of the electrode was measured as the purge gas was changed or as sulphide ions were added to the bulk electrolyte The solution was initially aerated, but the purge gas was sequentially changed to 2 vol.% O2/N2, 0.2 vol.% O2/N2, and 100% Ar (nominally deaerated). In - two experiments, sufficient Na2S was added to the solution to give a bulk [HS ] of either 10 µg/g or 100 µg/g.
FIGURE 4-48 shows the time dependence of ECORR for three experiments, one in bulk solution and two with the compacted clay electrode. The ECORR of the rotating electrode in bulk solution responded quickly to changes in the purge gas, decreasing with decreasing [O2] (curve (a)). Under these conditions, the anodic reaction is limited - by the rate of transport of CuCl2 away from the electrode and the cathodic reaction is limited by the rate of the interfacial reduction of O2 at ECORR [210]. Upon the addition - of 10 µg/g HS , ECORR drops immediately by ~500 mV to a value of ca. -900 mVSCE. - Increasing [HS ] to 100 µg/g resulted in a further 60 mV drop in ECORR. The - precipitous drop in ECORR upon the first addition of HS was explained in terms of a - switch in the anodic reaction from dissolution as CuCl2 to the formation of a Cu2S or CuS surface film, and a switch in the cathodic reaction from the reduction of residual O2 to the reduction of H2O. In the presence of compacted clay, a similar decrease in ECORR is observed upon the addition of HS-, but more slowly as a consequence of the restricted mass-transport conditions. The ECORR of the compacted clay electrode also decreased as the [O2] decreased (curves (b) and (c), FIGURE 4-48). The magnitude of the decrease
161 suggests complete transport control of the anodic and cathodic reactions. Upon the - addition of 10 µg/g HS (point D, curve (c)), ECORR decreased by 10-20 mV but then stabilised for the next 75 hours. The addition of 100 µg/g HS- (point E, curve (c)) caused a further decrease in ECORR by 50-60 mV followed by a second plateau of ~40 hours. These small decreases in ECORR were thought to be due to the - consumption of residual O2 in the clay layer by reaction with HS . After the second plateau period, however, ECORR dropped steeply by -600 mVSCE to a final steady-state value virtually identical to that observed in bulk solution at that [HS-].
FIGURE 4-48. Variation of the corrosion potential of a copper/compacted clay electrode and of a copper rotating disc electrode in 1 mol/L NaCl solution as a function of oxygen and sulphide concentration [213]. (a) Rotating disc electrode in bulk solution (the time axis for the first 2 h of this experiment has been expanded by a factor of 50 for clarity); (b) copper/compacted clay electrode in O2-containing solution only; (c) copper/compacted clay electrode with various [O2] and sulphide additions. The arrows represent the times at which either the atmosphere was changed or sulphide additions were made according to: (A) 2 vol.% O2/N2; (B) 0.2 vol.% O2/N2; (C) Ar; (D) 10 µg/g sulphide; (E) 100 µg/g sulphide.
The variation of ECORR for the clay-covered electrode in FIGURE 4-48 may be similar to that expected for a canister in a Swedish/Finnish repository, albeit on a much shorter timescale. Initially, ECORR will be determined by the relative kinetics of the anodic dissolution of Cu and of the reduction of O2. As the trapped O2 is consumed, ECORR will decrease. At some stage during the evolution of the repository environment, the nature of the reactions on the copper surface will change. In the presence of sulphide, the anodic and cathodic reactions will change to the formation of copper sulphide films and the evolution of hydrogen. A relatively rapid decrease in ECORR by several hundred mV can be expected, with an ultimate ECORR value in the - range of -0.80 VSCE to -1.00 VSCE, depending upon the concentration of HS at the canister surface.
162 A.4.3.1.3.1.3 Influence of welding on the corrosion properties
There is no evidence that the weld region should suffer higher corrosion rates than the rest of the canister shell:
• Fennell et al. [214] found no evidence for intergranular corrosion of Cu-OFP10 in bentonite-equilibrated Äspö groundwater with [Cl-] of 20 mg/L and 20,000 mg/L. These experiments were carried out at room temperature in aerated solution with a small-grained material (50-100 µm) and with large-grained material (a few hundred micrometers) produced by heat treatment of the fine-grained material in order to simulate the heat affected zone in a weld. • Ryan et al. [215] found no preferential attack at the weld region of electron-beam welded copper samples exposed to compacted buffer material at 100°C for periods of up to 5 years.
A.4.3.1.3.1.4 Influence of increased pressure at repository depth
An increase in pressure from 0.1 MPa to 5 MPa will result in a <10% change in the - equilibrium [CuCl2 ] and a correspondingly small change in the corrosion rate. Mor and Beccaria [216] observed a small effect of pressure on the anodic and, especially, the cathodic reactions involved in the corrosion of copper in seawater. The effect of pressure on the cathodic reaction is partially due to an increase in the concentration of dissolved O2 with increasing O2 partial pressure. Although this might also occur in a repository, due to the increase in pressure due to the development of the hydrostatic head, this effect neither increases the overall amount of O2 nor the maximum extent of corrosion.
A.4.3.1.3.1.5 Influence of methane
Several studies [217–219] investigating the corrosion of copper in natural gas showed no evidence that methane itself would be deleterious towards copper alloys. The presence of high concentrations of methane in deep Finnish groundwaters (up to 600 mL/L (0.027 mol/L)) is therefore expected not to have any effect on the corrosion of copper canisters.
A.4.3.1.3.2 Localised corrosion
The near-field environment in the repository will initially be oxidising. During these early stages of the repository life, the copper canister will most likely be prone to pitting corrosion. This type of localised corrosion has therefore deserved most attention in the Swedish/Finnish corrosion programme.
Other forms of localised corrosion, such as crevice corrosion (see section A.4.3.1.3.2.2) and ants-nest corrosion (see section A.4.3.1.3.2.3) are believed to be unlikely to occur under repository conditions.
10 Cu-OFP is a designation used in the Swedish and Finnish waste management programmes for oxygen-free copper micro-alloyed with 30-80 ppm phosphorus.
163
A.4.3.1.3.2.1 Pitting corrosion
Pitting corrosion of Cu canisters was evaluated using (i) a literature survey (comparing data gained from other applications, (ii) experimental studies performed under conditions simulating those expected in a repository, and (iii) modelling studies (see section A.5.2.2).
A lot of useful information in relation to the pitting behaviour of copper can be found in the literature; the pitting of copper water pipes is the most studied field. Although pitting studies of copper in potable waters are of limited use for predicting the pitting behaviour of copper canisters due the difference in salinity between potable waters and saline porewaters, the literature can be useful because the proposed mechanisms can provide some insight into the possibility of localised corrosion in a repository.
Various forms of pitting corrosion of copper have been described:
• type I pitting. This form is associated with cold, hard and moderately hard waters - 2- - free of naturally occuring inhibitors, but containing HCO3 , SO4 , Cl , and O2 and on copper pipes with a residual surface carbon film remaining from the manufacturing process [220]. • type II pitting. This form proceeds in hot potable waters (>60°C) with a pH < 7.4, and - 2- a [HCO3 ]/[SO4 ] ratio < 1 [221–224]. Type II pitting tends to produce pits with a larger depth/width ratio than the approximately hemispherical pits characteristic of type I pitting.
Pitting studies are usually divided into three phases, viz. birth (initiation), propagation, and death (i.e. the three phases in the life of a pit). The initiation is often associated - with the formation of CuCl underneath a porous Cu2O layer. The higher [Cl ] encountered in deep groundwaters may either make pit initiation more likely, or may induce so many defects in the Cu2O layer that the surface dissolves generally (active dissolution) as opposed to locally as pits (passive behaviour). In order for pits to propagate, oxygen is a pre-requisite. Pitting of Cu water pipes is only sustained because of the high [O2] in fresh water and because it is continually replenished by the movement of water in the pipe. This would not be the case for pits on Cu canisters, both because of the limited amount of O2 available and because of the restricted mass- transport conditions which will limit the supply of O2 to the corrosion sites. Therefore, pits on Cu canisters will be far more likely to die than those on Cu water pipes.
- FIGURE 4-49 shows the variation of the pitting (breakdown) potential, ENP, with log [Cl ] - - - for Cu in Cl and Cl /HCO3 mixtures from a number of literature studies. The data originate from electrochemical experiments conducted potentiodynamically and have been corrected to zero scan rate to aid inter-comparison. No correction has been made for the variation of ENP with pH. The wide variation in the data is a consequence of - (i) the use of different HCO3 concentrations in the various studies and (ii) the fact that - ENP is a distributed, stochastic parameter. For all the studies, it was found that in Cl - - - /HCO3 mixtures, HCO3 promotes passivation, whereas Cl promotes pitting.
164
- - FIGURE 4-49. Dependence of the pitting potential ENP on [Cl ] in Cl solutions and in - - Cl /HCO3 mixtures at 25°C. (○) Ref. [225]; (▲) Ref. [226], (●) Ref. [227,228]; - - (■) Ref. [229]; (×) Cl /HCO3 mixtures [230]; (-○-) Ref. [231]. Lines fitted to - - (1) Ref. [129], (2) Ref. [229], (3) Ref. [230] Cl /HCO3 data.
EPP values of 0.37 VSCE and 0.07 VSCE have been reported [226] for Cu in solutions - - -4 containing 0.14 mol/L HCO3 , 0.028 mol/L Cl , and either 2.10 mol/L or 0.01 mol/L, respectively.
A number of corrosion studies have been performed under conditions that simulate the canister near-field environment soon after emplacement of the canisters and saturation of the buffer material. Copper coupons have been exposed to compacted buffer material wetted by (initially) aerated saline porewaters, and exposed for extended periods of time (up to 2 years), usually at elevated temperature [205,207,208,212,232]. Despite the relative aggressiveness of the conditions in such tests (oxidising, high - - [HCO3 ], and high [Cl ]), no evidence for pitting was observed.
A.4.3.1.3.2.2 Crevice corrosion
Crevice corrosion is not believed to present a threat to the integrity of Cu canisters because (i) the hydrolysis of Cu(I), especially when complexed by Cl-, and local acidification in crevices is unlikely to occur [233] (ii) the formation of Cu(II) requires the presence of O2 which is unlikely to occur in occluded regions, such as crevices, where O2 access is restricted.
In long-term irradiated corrosion tests under simulated conditions of a Canadian repository, no crevice corrosion was observed on either creviced U-bend of creviced planar samples [215].
A.4.3.1.3.2.3 Ants-nest corrosion
'Ants-nest' corrosion is a form of corrosion specific to Cu [234], which results in localised dissolution of the material to produce a honeycomb-like ants-nest
165 appearance. This type of corrosion, however, only occurs under very specific conditions (e.g. in air conditioning equipment), which are unlikely to prevail in a repository.
A.4.3.1.3.3 Stress corrosion cracking (SCC)
Stress corrosion cracking (SCC) can only occur under the combined action of the following factors: (i) a material susceptible to SCC, (ii) a tensile stress, and (iii) a suitable aggressive environment. Under repository conditions (Sweden/Finland), these three pre-requisites for SCC are present [60]:
• the candidate container material copper, and especially those containing P, has been found, in the past, to be highly susceptible to SCC. • tensile stresses on the container surface are likely during various stages in the evolution of the repository environment, either due to external loads or from residual manufacturing stresses. • from the literature, the only three species known to cause SCC of Cu are ammonia, acetate, and nitrite ions. Ammonia and acetate may be present in the near-field environment of a Swedish/Finnish repository as a natural constituent of the groundwater and, possibly, as a result of human activity during construction [235]. Nitrite ions are not expected to occur in the repository environment.
A.4.3.1.3.3.1 Influence of ammonia
Saario et al. [235] proposed a lower threshold ammonia concentration of ~0.5 mg/L. SCC was not observed on OFP Cu in a bentonite equilibrated synthetic groundwater containing 0.5 mg/L ammonia [236]. Slow strain rate tests (SSRT) [237] on base metal and EB-welded OFP Cu (45 wt. ppm P), conducted in deaerated synthetic saline and brackish Olkiluoto groundwaters (1-100 mg/L ammonia, 3,700 or 17,000 mg/L Cl-, 1 mg/L sulphide) at 80°C, revealed no SCC.
A.4.3.1.3.3.2 Influence of acetate
There is no information available in the literature regarding the minimum acetate concentration for cracking, although the threshold concentration, if there is one, must be below 0.05 mol/L [238].
A.4.3.1.3.3.3 Influence of chloride
Chloride ions may inhibit SCC of Cu in environments containing all three SCC agents (ammonia, acetate, and nitrite) because of the effect on the formation and stability of - - Cu2O films. Cl ions inhibit Cu2O film formation by competing with OH ions for surface sites (Cu(OH)surf is a precursor to Cu2O formation). Furthermore, the - incorporation of Cl ions into an existing Cu2O layer (as may form on a canister surface prior to emplacement in the repository) increases the defect density of the oxide and tends to promote general dissolution over localised corrosion. A defected
166 Cu2O layer will be less protective and, therefore, less likely to support SCC, than a - Cu2O film formed in Cl -free environments [60]. FIGURE 4-50 shows the results of constant extension rate tests (CERT) with OFP Cu - in 0.1 mol/L NaNO2 with additions of varying amounts of Cl to the solution [239]. The specimens were maintained at a constant current density of 1 µA/cm² (to simulate the flux of oxidant to the canister) and the corresponding potential recorded. The crack velocity decreased slightly with increasing [Cl-] for chloride concentrations of up to 0.01 mol/L. At higher [Cl-], the crack velocity fell to that in air, which corresponds to pure ductile tearing (inspection of the fracture surfaces also confirms this change from brittle to ductile behaviour) and the measured potential dropped ~250 mV to more active values. Thus, the decrease in SCC susceptibility was a result of the shift from passive to active dissolution.
FIGURE 4-50. Effect of [Cl-] on the crack velocity and potential of OFP Cu speci- mens in 0.1mol/L NaNO2 at room temperature. The OFP specimens were maintained at a constant current density of 1 µA/cm² and the corresponding potential is shown in the figure [239].
A.4.3.1.3.3.4 Influence of temperature
Copper becomes less susceptible to SCC with increasing temperature. The apparent decrease in SCC susceptibility with increasing temperature may be the result of several factors. First, in chloride solutions, increasing temperature promotes general dissolution of the surface, resulting in negative shifts in ECORR. Second, there is some evidence to suggest that the range of potentials for SCC shifts to more positive values with increasing temperature.
A.4.3.1.3.3.5 General conclusions concerning SCC
There are a lot of indications to suggest that copper is not susceptible to SCC in deep Fenno-Scandian groundwaters at elevated temperature:
167
• SCC was not observed on Cu in natural and synthetic groundwaters at 80°C [240– 242]. • SCC was not observed on Cu alloys in irradiated J-13 water (a dilute synthetic groundwater based on conditions at the Yucca Mountain site in the U.S.) and moist air at either 95°C or 105°C. • Beaver and Durr [243] failed to induce SCC in J-13 water or the vapour phase above it at 90°C.
A.4.3.1.3.4 Microbially influenced corrosion (MIC)
The worst case scenario for corrosion of the copper canister would be if sulphate reducing bacteria (SRB) formed biofilms on the canisters or grew intensively in the buffer close to the canister.
During the initial phase, the temperature in the repository will be elevated, with a maximum canister surface temperature of 90°C. This is, however, not an absolute constraint for micro-organisms. Sulphate-reducing bacteria may survive. The γ- radiation field at the canister surface will not be high and thus will not have any marked effect on the survival of microbes. The low availability of water in the buffer (i.e. the water content relative to groundwater) will, however, provide a constraint on the likelihood of long-term survival. Overall, the conditions close to the canister will not be conducive to the survival of microbes.
A.4.3.1.3.5 Corrosion induced by radiation effects
The radiation, emanating from the canister, will be dominated by the decay of 137Cs, which has a half-life of 30 years. With the canister design proposed by SKB and Posiva, the surface dose rate immediately after encapsulation will be about 0.5 Gy/h. With the half-life of 30 years, the dose rate will have dropped to negligible levels after a few hundred years [60].
A.4.3.1.3.5.1 Corrosion in air in the presence of γ-radiation
The influence of radiation will be negligible even at dose rates considerably higher than the maximum surface dose rate for the canister. The corrosion rates that are measured in the presence of γ-radiation are not higher than one would expect for corrosion of copper in unirradiated moist air (see section A.4.3.1.1).
In experiments in which OF Cu was exposed to an air/water vapour mixture at 150°C and to vapour saturated air at 95°C in a γ-radiation field of approximately 103 Gy/h (i.e. more than a factor of 1,000 times higher than the expected maximum surface dose rate on the canister), corrosion rates of about 1 µm/year and 4 µm/year at 150°C and 95°C, respectively, were measured after an exposure period of 6 months [244].
168 During experiments performed in dry air and in air with 40% relative humidity at 94°C (dose rate of 210 Gy/h) and 152°C (dose rate between 110–150 Gy/h), corrosion rates ranging from about 1 mm/year to 3 µm/year have been observed.
Marsh calculated a production rate of 0.002 mol/year for nitric acid in the repository [245]. This small quantity is believed to have a negligible effect on the life of the canister.
A.4.3.1.3.5.2 Corrosion in water in the presence of γ-radiation
The available data tend to indicate that γ-radiation will not cause enhanced corrosion rates. On the contrary, lower corrosion rates in the presence of radiation have been observed at least for dose rates in the range of 10–100 Gy/h.
Experiments exposing copper to fresh water of low ionic strength irradiated with 103 Gy/h for about 200 days revealed corrosion rates of about 2 µm/year [244].
Experiments exposing copper to a saline solution (1 mol/L chloride) at 150°C in the presence (dose rates in the range 14–27 Gy/h were used) and absence of radiation, showed corrosion rates a factor of about 4 lower in the presence of radiation. King and Litke [246] attributed this to the more protective nature of the surface film formed in irradiated solutions.
169 A.4.3.2 Corrosion studies relevant to the Spanish disposal concept in granite (investigations performed at ENRESA/INASMET)
The experimental results presented in this section originate from investigations performed at ENRESA/INASMET.
More detailed information can be found in References [67,68,152,154,160,162- 164,247-251].
A.4.3.2.1 Experimental procedure
The corrosion behaviour of several metallic materials such as carbon steels (AE-355D, TStE 355, TStE 460, 15 MnNi 6.3), Ni-Cr steels (AISI 316L), Ni-base alloys (Hastelloy C-22), and Cu-base alloys (Cu-OF, Cu10Ni, Cu30Ni) has been studied to evaluate their use as possible candidate materials for the fabrication of metallic containers for the final disposal of HLW/spent fuel in granite geological formations.
The chemical compositions of the materials tested during the experimental corrosion studies relevant to the Spanish disposal concept in granite are summarised in TABLE 4-46. Besides specimens of the parent material (unwelded specimens), welded specimens (simulating possible container closure techniques) were tested. The various welding techniques that were studied are described in TABLE 4-47.
Experiments were performed in two geological media that are representative of the near-field repository environment in granite, according to the Spanish disposal concept: • granitic-bentonite water. The chemical composition of the synthetic bentonite buffered granitic groundwater is presented in TABLE 2-15 (section A.2.3.2). During the experimental programmes, the chloride content was varied between 6,500 and 50,000 mg/L. • saturated bentonite. Crushed bentonite was saturated with deionised water (water content: 25.5 wt.%).
General and localised corrosion phenomena, such as stress corrosion cracking (SCC), pitting and crevice corrosion, were investigated. General corrosion studies on the carbon steels were carried out in both media at different temperatures by means of gravimetric methods (weight loss measurements). FIGURES 4-51 (a) and 4-51 (b) show a general view of the test assemblies used to evaluate general corrosion in granitic-bentonite water and saturated bentonite, respectively. The SCC resistance was evaluated by means of the Slow Strain Rate Testing (SSRT) technique at different temperatures and slow strain rates (10-4-10-7 s-1). The testing equipment for SSRT consisted of constant extension rate tensile testing machines of 50 KN capacity and selectable crosshead speed within the range of 0.1-10-6 mm/s (FIGURE 4-52). The test specimens were inserted in Hastelloy C-276 autoclaves. Additional comparative tests were performed in argon as an inert reference media [68,152,154]. Finally, the crevice corrosion behaviour of the Ni-Cr steels was studied using artificial crevice assemblies, according to ASTM specifications G-78. These assemblies were formed by bolting together flat test specimens with crevice formers made of Teflon
170 (FIGURE 4-53). The resulting assemblies were immersed in autoclaves containing the granitic-bentonite water [154,163].
(b)
(a)
FIGURE 4-51. Test assembly used to evaluate the general corrosion in (a) granitic-bentonite water [163], and (b) saturated bentonite [67].
FIGURE 4-52. General view of the SSRT FIGURE 4-53. Crevice corrosion assembly test set-up used to evaluate stress showing metal specimens and Teflon corrosion cracking (SCC) [163]. gaskets [163].
171 TABLE 4-46. Chemical composition of the materials tested during the experimental corrosion studies relevant to the Spanish disposal concept in granite [67,164,247].
Materials Composition (wt.%)
Fe Ni Cu Cr Mo Mn Si Al C P S N Others
Carbon steels AE-355D (DIN 1.0570) bal. - - - - 1.4 0.40 0.040 0.17 0.020 0.012 0.006 - TStE 355 DIN 1.5066) bal. - - - - 1.5 0.41 0.036 0.16 0.017 0.002 0.006 - TStE 460 (DIN 1.8915) bal. - - - - 1.4 0.35 0.027 0.14 0.010 0.001 0.560 V: 0.17 15 MnNi 6.3 (DIN 1.6210) bal. 0.82 - - - 1.5 0.23 0.027 0.18 0.012 0.003 - -
Ni-Cr steels AISI 316L (ASTM A-240) bal. 11.5 - 17.4 2.2 1.3 0.31 - 0.021 0.029 0.002 - -
Ni-base alloys Hastelloy C-22 (UNS N06022) 2.8 bal. 0.03 21.8 13.3 0.22 0.22 0.014 0.01 <0.005 - W: 3.0 ; Co: 0.38 ; V: 0.16
Cu-base alloys Cu-OF (UNS C10200)1 <0.005 - bal. ------<0.01 0.0015 - Pb: <0.01 ; Bi: <0.01 ; Zn: <0.005 Cu10Ni (UNS C70600) 1.0-1.8 9.0-11.0 bal. - - <1.0 ------Pb: <0.05 ; Zn: <1.0 Cu30Ni (UNS C71500) 0.70 31.2 bal. - - 1.0 <0.01 - 0.011 <0.01 0.005 - Pb: <0.01 ; Zn: <0.01
bal.: balance. 1 Cu-oxygen free.
TABLE 4-47. Welding procedures investigated as possible container closure techniques [68,152].
Welding procedure Target material for welding
Abbreviation Description Id. Nr.(1)
EBW Electron Beam Welding 511 carbon steels, Ti-, Ni-, and Cu-base alloys GTAW Gas Tungsten Arc Welding 141 Ni-Cr steels, Ni-, and Cu-base alloys MAGW Metal Active Gas Welding 135 carbon steels FCAW Flux Cored Arc Welding 136 carbon steels
(1) Identification number of the welding procedure according to EN ISO 4063 (2000) “Welding and allied processes. Nomenclature of processes and Reference Numbers”.
172 A.4.3.2.2 Carbon steels
A.4.3.2.2.1 General and localised corrosion
The general corrosion rates, as a function of exposure time, for the fine-grained carbon steel TStE 355 in granitic-bentonite water at 90°C and 120°C under anoxic conditions are presented in FIGURE 4-54. These results originate from long-term immersion experiments performed in autoclaves under anoxic conditions. The anoxic conditions were assured by purging an argon stream over the test solutions before closing the autoclave lid (the only oxygen present at the starting of the test was the oxygen dissolved in the test solution). The general corrosion rates, as a function of exposure time, for the fine-grained carbon steel TStE 355 in saturated bentonite (saturated with deionised water) at 50°C, 75°C, and 100°C are presented in FIGURE 4-55. The values of the general corrosion rates in bentonite-buffered granitic groundwater at 90°C are added for comparison.
In both environments, the general corrosion rate diminished asymptotically with exposure time to a stationary value [67,163]: • from FIGURE 4-54, it can be seen that the general corrosion rate reached a constant value of approximately 6 and 14 µm/year in granitic-bentonite water at 90°C and 120°C, respectively [163]. • from FIGURE 4-55, it can be observed that the temperature has no signifcant effect on the final general corrosion rate after 18 months of exposure to saturated bentonite. For all three test temperatures (50°C, 75°C, and 100°C), the general corrosion rate was about 10 µm/year after 18 months. The temperature, however, influences the rate with which this constant value is reached: the lower the temperature, the faster the relatively low constant corrosion rate is reached [67].
The general corrosion rates (calculated from weight loss measurements) and the results of the local corrosion investigations (optical microscopy and SEM) for the fine- grained carbon steel TStE 355 exposed to saturated bentonite at 50°C, 75°C, and 100°C are compiled in TABLE 4-48. Metallographic examinations (optical microscopy and SEM) revealed that the fine- grained carbon steel TStE 355 specimens were resistant to pitting corrosion in granitic-bentonite water. However, in the saturated bentonite media, the specimens were found to be susceptible to non-uniform corrosion with maximum pit depths up to 100 µm (18 months exposure at 75°C). This is attributed to the inhomogeneity of the test electrolyte (hydrated and compacted bentonite) surrounding the metallic specimens thereby creating potential cells, which results in the formation, on the specimens’ surface, of zones with a different corrosion potential. This will lead to a local attack as opposed to the uniform attack observed in a homogeneous medium such as a solution (bentonite buffered granitic groundwater) [67].
173
50 Corrosion rate (µm/year)
S355 Granitic water 90ºC
40 S355 Granitic water 120ºC
30
20
10
0 02468101214 Exposure time (months)
FIGURE 4-54. General corrosion rates for the fine-grained carbon steel TStE 355 exposed to granitic-bentonite water for 12 months at 90°C and 120°C [163].
Corrosion rate (µm/year)
Granite water 90 ºC Water sat-Bent 50 ºC Water sat-Bent 75 ºC Water sat-Bent 100 ºC
Testing time (months)
FIGURE 4-55. General corrosion rates for the fine-grained carbon steel TStE 355 exposed to saturated bentonite at 50°C, 75°C, and 100°C for 18 months. The data obtained in granitic-bentonite solutions at 90°C are added for comparison [67].
174 TABLE 4-48. Influence of exposure time on the average uniform corrosion rate, maximum pit depths, average pit depths, and pitting factor calculated for the fine-grained carbon steel TSTE 355 exposed to saturated bentonite at various temperatures (50°C, 75°C, 100°C) [248-250].
T ∆t(1) General Local corrosion (°C) (months) corrosion (2) (3) (4) vCORR dMax dAv Pitting (µm/year) (µm) (µm) factor(5)
1 32.0 - - -
3 11.0 15 2.8 5.4 50 6 11.0 20 5.5 3.6 9 10.0 43 7.5 5.7 18 6.7 56 10.5 5.9
1 45.0 - - -
3 17.0 30 4.3 7.0 75 6 20.0 44 10.0 4.4 9 13.0 53 9.8 5.4 18 10.0 98 15.0 6.5
1 57.0 30 4.75 6.3
3 34.0 40 8.5 4.7 100 6 24.0 58 12.0 4.8 9 14.0 60 10.5 5.7 18 6.3 50 6.3 7.9
(1) ∆t : exposure period. (2) vCORR : average uniform corrosion rate. (3) dMax : maximum depth of pitting attack. (4) dAv : average depth of pitting attack. (5) ratio of the maximum depth of pitting attack to the average depth of pitting attack.
The welds and the HAZ (Heat Affected Zone) of EBW- and FCAW-welded TStE 355 carbon steel specimens were NOT attacked by local corrosion in granitic-bentonite water at 90°C. At 120°C, on the other hand, pits up to 280 µm and 130 µm deep for FCAW- and EBW-welded specimens, respectively, were detected at the interface weld/HAZ after 12 months exposure [163]. This is illustrated in FIGURES 4-56 and 4-57.
175
PIT
PIT
200 µm 200 µm FIGURE 4-56. Optical micrograph of the FIGURE 4-57. Optical micrograph of the HAZ weld/HAZ-interface region of a FCAW- of a EBW-welded TStE 355 carbon welded TStE carbon steel specimen steel specimen after 12 months after 12 months exposure to granitic- exposure to granitic-bentonite water bentonite water (magnification: 100×) (magnification: 100×) [163]. [163].
XRD-analyses of the black oxide layer formed on the specimens exposed to the granitic-bentonite water (90°C and 120°C) identifid this layer as Fe2O3 (hematite). However, some discrepancies exist on the formation of this product: it could have been formed due to the interaction of the formed corrosion products with air during the XRD analysis instead of during the corrosion test. The corrosion products formed on the specimens exposed to saturated bentonite (50°C, 75°C, and 100°C), on the other hand, mainly consisted of FeCO3 (siderite), with minor traces of Fe3O4 (magnetite) and Fe2O3 (hematite). The formation of FeCO3 can be explained by the presence of CO2 in the saturated bentonite. FIGURE 4-58 presents a general view of the carbon steel TStE 355 samples exposed to saturated bentonite for 9 months at 50°C, 75°C, and 100°C [67,68,163,164,248].
FIGURE 4-58. General view of TStE 355 samples after a 9-months exposure period to saturated bentonite at 50, 75, and 100°C [248].
176 A.4.3.2.2.2 Stress Corrosion Cracking (SCC)
FIGURE 4-59 shows a comparison of the fracture between TStE 355 carbon steel specimens tested in Ar and granitic-bentonite water at 90°C under anoxic conditions and a strain rate of 10-7 s-1 (the anoxic conditions were assured by purging an argon stream over the test solutions before closing the autoclave lid; the only oxygen present at the starting of the test was the oxygen dissolved in the test solution). SEM- examinations of the fracture surface showed transgranular fracture modes for specimens tested in the granitic-bentonite water, whereas a fully ductile fracture surface was observed for specimens tested in Ar [67,68,163,164].
FIGURE 4-59. Macrographs of two TStE 355 tensile specimens tested in Ar and granitic-bentonite water at 90°C and a strain rate of 10-7 s-1 [163].
Values of elongation, reduction of area, and maximum load obtained in the slow strain rate tests (SSRT) on TStE 355 carbon steel in Ar and granitic-bentonite water at 90°C, and strain rates of 10-4-10-7 s-1 are given in FIGURE 4-60. Compared to the results in Ar, the steel suffers a loss of ductility in granitic-bentonite water, which is mainly noticed in the reduction of area parameter. The drop of these parameters is not significant at the highest strain rate (10-4 s-1) but it is important at any of the other slower strain rates used in the tests. This is attributed to the longer exposure periods to the corrosive medium for the slower strain rates. There is no important change in the maximum load parameter in granitic-bentonite water [68,152,163,164].
This loss of ductility was also observed in salt brines (MgCl2-rich ‘Q’-brine and NaCl- rich brine). It was thought to be attributed to the embrittling effect of the hydrogen produced on the specimens surface during the tests due to the general corrosion phenomena suffered by the steel and it has already been explained in detail in section A.4.1.1.2.2.2.2.
177
80 Elongation (%) 800 Maximum Load (MPa)
60 600
40 400
20 200
0 0 1E-04 1E-05 1E-06 1E-07 1E-04 1E-05 1E-06 1E-07 Strain Rate (1/s) Strain Rate (1/s)
100 Reduction of Area (%)
80
355 A 90ºC 60 355 G 90ºC 355(EBW) A 90ºC 355(EBW) G 90ºC 40 355(FCAW) A 90ºC 355(FCAW) G 90ºC
20
0 1E-04 1E-05 1E-06 1E-07 Strain Rate (1/s)
FIGURE 4-60. Elongation, maximum load, and reduction of area values vs. strain rate for the parent and welded TStE 355 steel tested at 90°C in Ar and granitic-bentonite water [68,152]. A: argon; G: granitic-bentonite water; EBW: Electron Beam Welding; FCAW: Flux Cored Arc Welding
Metallographic investigations carried out on the parent and FCAW-welded TStE 355 carbon steel specimens revealed no clear signs of sensitivity to SCC in granitic- bentonite water. Areas of localised corrosion were present on the lateral surface of TStE 355 specimens when tested at the slowest strain rate of 10-7 s-1. Optical micrographs of a parent and an FCAW-welded TStE 355 specimen tested in granitic- bentonite water at 90°C and a strain rate of 10-7 s-1, respectively are shown in FIGURES 4-61(a) and 4-61(b). It was however found that these areas of localised corrosion have a crack length to crack width ratio very close to 1. This indicates rather a localised corrosion phenomena due to the local and repetitive breaking of the oxide layer than secondary cracking. The crack length/crack width ratio parameter gives an idea of the sharpness of the crack and is useful to quantify and compare the resistance of the tested materials to SCC. The term ‘crack’ is used in a wide sense including both secondary cracks and localised corrosion due to the breaking of the oxide layer [68,152,163,164].
178
100 µm 100 µm (a) (b)
FIGURE 4-61. Optical micrographs of TStE 355 carbon steel specimens tested in granitic- bentonite water at 90°C and a strain rate of 10-7 s-1 [163]. (a) parent specimen (magnification: 100×). (b) FCAW-welded specimen (magnification: 100×).
A.4.3.2.3 Ni-Cr steels
A.4.3.2.3.1 General and crevice corrosion of 316L stainless steel
AISI 316L stainless steel (parent, EBW-, and GTAW-welded specimens) tested in granitic-bentonite water at 90°C was not attacked by any type of localised corrosion after an exposure period of 12 months.
No appreciable general corrosion rates (< 0.1 µm/year), nor sensitivity to localised corrosion, was observed for the AISI 316L stainless steel specimens exposed to saturated bentonite at 50°C, 75°C, and 100°C for periods up to 18 months, as illustrated in FIGURE 4-62.
179 FIGURE 4-62. General view of AISI 316L samples after an 18-months exposure period to saturated bentonite at 100°C [248].
A.4.3.2.3.2 Stress corrosion cracking (SCC) of 316L stainless steel
FIGURE 4-63 shows a comparison of the fracture between AISI 316L stainless steel specimens tested in Ar and granitic-bentonite water at 90°C and a strain rate of 10-7 s-1. Secondary cracks typical for stress corrosion cracking have not been observed under any of the test conditions. Several AISI 316L stainless steel specimens tested in the granitic-bentonite water at the slowest strain rate of 10-7 s-1 show isolated pits near the fracture zone [68,152,163,164].
Granite
Pit
Argon
FIGURE 4-63. Macrographs of two AISI 316L stainless steel tensile specimens tested in Ar and granitic-bentonite water at 90°C and a strain rate of 10-7 s-1 [68,152].
The results of the SSRT-tests obtained for the parent, EBW-, and GTAW-welded specimens in Ar and granitic-bentonite water at 90°C and strain rates in the range 10-5-10-7 s-1 are given in FIGURE 4-64. The values of elongation, reduction of area,
180 yield strength, and maximum load are very close to those obtained in Ar, thus indicating no loss of ductility. The fracture of the welded specimens was always located in the base material [68,152,164].
100 Elongation (%) 100 Reduction of Area (%)
80 80
60 60
40 40
20 20
0 0 1E-04 1E-05 1E-06 1E-07 1E-04 1E-05 1E-06 1E-07 Strain Rate(1/s) Strain Rate (1/s)
316 Argon 90ºC 316 Granite 90ºC
316(EBW) Argon 90ºC 316(EBW) Granite 90ºC
316(GTAW) Argon 90ºC 316(GTAW) Granite 90ºC
500 Yield Strength (MPa) 750 Maximum Load (MPa)
400 650
300 550 200
450 100
0 350 1E-04 1E-05 1E-06 1E-07 1E-04 1E-05 1E-06 1E-07 Strain Rate (1/s) Strain Rate (1/s)
FIGURE 4-64. Elongation, reduction of area, yield strength, and maximum load values vs. strain rate for the parent and welded AISI 316L stainless steel tested at 90°C in Ar and granitic- bentonite water [68,152].
SEM-examinations show a fully ductile fracture surface with the formation of dimples for the specimens tested in the granitic-bentonite water (see FIGURE 4-65(a)). It was also observed that gamma irradiation (10 Gy/h) had no influence, either on the mechanical properties or on the resistance to SCC. FIGURE 4-65(b) shows a SEM- micrograph of the fracture surface of an AISI 316L stainless steel sample after SSRT- testing in granitic-bentonite water at 90°C and with 10 Gy/h irradiation (strain rate of 10-7 s-1): the ductile fracture is similar to the ones observed for the samples without irradiation (FIGURE 4-65(a)) [68,152,163,164].
181
(a) (b)
FIGURE 4-65. SEM-micrographs of the fracture surface of AISI 316L stainless steel specimens tested in granitic-bentonite groundwater at 90°C and a strain rate of 10-7 s-1. (a) without irradiation (magnification: 750×) [68,152]. (b) with 10 Gy/h radiaiton (magnification: 750×) [163].
A.4.3.2.4 Cu-base alloys
During the last few years, Cu-base alloys were investigated as alternative container materials in the event that the candidate reference materials, i.e. the actively corroding carbon steel and the passively corroding Ti 99.8-Pd, prove to be inadequate as a corrosion barrier for the final disposal of HLW/spent fuel.
A.4.3.2.5.1 General and localised corrosion of Cu-OF and Cu30Ni
Both Cu-OF and Cu30Ni exhibited some degree of surface corrosion, with corrosion rates < 1 µm/year when exposed to saturated bentonite for periods up to 18 months (50°C, 75°C, and 100°C) [67,248,250]. FIGURE 4-66 presents a general view of Cu- OF and Cu30Ni specimens exposed to saturated bentonite for 18 months at 100°C. Cu-OF corroded uniformly, independent of the temperature. For Cu30Ni, on the other hand, the type of corrosion depended on the temperature [67,248]: • at 50°C, the specimens were corroded uniformly. • at 100°C, the corrosion proceeded non-uniformly with maximum penetration depths up to 25 µm after 18 months. • at 75°C, a mixed situation was observed.
182
(a) (b)
FIGURE 4-66. General view of Cu-base samples after 18 months of exposure to saturated bentonite at 100°C. (a) Cu-OF [67,248]. (b) Cu30Ni [248].
XRD-analysis revealed that the protective film formed on Cu-OF, tested in saturated bentonite at temperatures up to 100°C, consisted primarily of cuprous oxide (CuO2) [248,250]. For Cu30Ni, a mixture of corrosion products Cu2O, NiO and Ni(OH)2, as well as (Cu,Ni)-oxides were detected. Also, corrosion products rich in sulphur were detected on Cu30Ni. The effect of microbially influenced corrosion should therefore not be excluded [67,248,250].
A.4.3.2.5.2 Stress Corrosion Cracking (SCC) of Cu-OF and Cu30Ni
Neither of the Cu-based alloys experienced a deterioration of mechanical properties when tested in granitic-bentonite water compared to Ar, as illustrated in TABLE 4-49 [67,163,164].
TABLE 4-49. Results of the SSRT-experiments for Cu-OF and Cu30Ni tested in Ar and granitic- bentonite water at 90°C and a strain rate of 10-7 s-1 [163,164].
Parameter Cu-OF(1) Cu30Ni Argon / 90°C Granite / 90°C Argon / 90°C Granite / 90°C
Elongation (%) 4.8 4.3 30.0 31.8 Reduction of area (%) 10.8 9.8 82.5 79.6 Maximum load (MPa) 208.0 199.0 274.0 255.0 Time to fracture (hours) 133.0 120.0 859.0 907.0 Secondary cracks no no no yes Type of fracture brittle (intergr.) brittle (intergr.) ductile ductile (intergr.)
(1) Cu-OF : Oxygen-free Cu in half-hard condition (HB 97).
183 The appearance of the crack surface is, however, quite distinct for both Cu-base alloys, as can be seen in FIGURE 4-67 [67,163,164]: • the Cu-OF specimens cracked in a brittle manner (crack propagation was intergranular). No difference was detected between the specimens tested in Ar and those tested in granitic-bentonite water, indicating that Cu-OF is not susceptible to SCC under these conditions (granitic-bentonite water and 90°C). Cu-OF was tested in a hardened condition (HB 97) [163]. • the Cu30Ni specimens cracked in a ductile manner, both in Ar and in granitic- bentonite water. However, on the specimens exposed to granitic-bentonite water many secondary stress-corrosion cracks were detected up to 230 µm deep (see FIGURE 4-68), indicating a high susceptibility to SCC [67,163,164].
A new set of tests were performed on hardened and softened Cu-OF with other testing conditions: the strain rate was 2×10-7 s-1 and the aggressivity of the testing media was increased by adding 50,000 mg/L chloride. The material was softened by thermal treatment up to a hardness of HB 49. No evidence of sensitivity to stress corrosion cracking was found. The results are summarised in TABLE 4-50 [251].
TABLE 4-50. Results of the SSRT-experiments for Cu-OF tested in Ar and granitic-bentonite water (with 50,000 mg/L Cl-) at 90°C and a strain rate of 2×10-7 s-1 [251].
Parameter Cu-OF(1) Cu-OF(2)
Granite Argon Granite + + 50,000 mg/L Cl- 50,000 mg/L Cl-
90°C / 2×10-7 s-1 90°C / 2×10-7 s-1 90°C / 2×10-7 s-1
Elongation (%) 8.4 53.0 52.0 Reduction of area (%) 13.0 64.0 64.0 Maximum load (MPa) 220.0 177.0 163.0 Time to fracture (hours) 130.0 912.0 936.0 Secondary cracks no no no Type of fracture brittle ductile + brittle ductile + brittle
(1) Cu-OF, half-hard condition (HB 97). (2) Cu-OF, soft condition (HB 49).
184
(a) (b)
FIGURE 4-67. Macrographs of Cu-base tensile specimens tested in Ar and granitic-bentonite water at 90°C and a strain rate of 10-7 s-1. (a) Cu-OF (brittle fracture) [163,164]. (b) Cu30Ni (ductile fracture) [164].
200 µm
FIGURE 4-68. Optical micrograph of the lateral surface of a Cu30Ni specimen tested in granitic-bentonite water at 90°C and a strain rate of 10-7 s-1, showing secondary cracks up to 230 µm deep (magnification: 70×) [163].
Also, SEM-analysis of the fracture surface of Cu30Ni specimens tested in granitic- bentonite water at 90°C and a strain rate of 10-7 s-1 revealed a difference in fracture morphology between the periphery and the centre of the specimen, as illustrated in FIGURE 4-69: the presence of brittle zones, with intergranular crack growth, were detected along the periphery of the specimens. In the centre of the specimens, on the other hand, the fracture surface was found to be ductile, which is similar to the morphology observed on the specimens exposed to Ar [164].
185 FIGURE 4-69. SEM-micrographs of the fracture surface of a Cu30Ni specimen tested in granitic-bentonite water at 90°C and a strain rate of 10-7 s-1 [164]. (a) general view of the fracture surface (magnifica- tion: 35×). (b) Detailed view of the brittle zone (close to the lateral edge of the specimen) (magnification: 350×). (c) Detailed view of the ductile zone (centre of the specimen) (magnification: 350×).
186 A.4.4 Corrosion studies performed in cementitious backfill material
An overview of the current knowledge of the general and localised corrosion behaviour of candidate container materials in cementitious environments seems appropriate for the following two reasons:
• the disposal concepts/repository designs described in APPENDIX B are also considered in some countries (such as the UK), for the disposal of ILW and LLW. In this case, a cementitious backfill type of material would be used. • also, in the disposal concepts for HLW/spent fuel in geological argillaceous or granitic formations, large amounts of concrete will be present as a structural material used in the building of the repository [252]. These concrete structures (high pH, Na, K, Ca(OH)2) are likely to play a very important role in defining the chemistry of the near-field environment contacting the metallic container, and thereby influencing its corrosion behaviour, in the long-term.
The experimental results presented in this section are obtained from corrosion investigations, performed at Serco Assurance (a former division of AEA Technology, Harwell) and under contract to UK Nirex Ltd., in environments relevant for the near field of a deep repository, in which a cementitious backfill is used (for the disposal of ILW).
The Nirex concept for the disposal of intermediate-level and some low-level radioactive waste envisages immobilisation and packaging of wastes in stainless steel containers [253]. Grade 316L has been selected for thin-section components because of its superior resistance to localised corrosion. Grade 304L offers economic advantages when used for thicker sections (e.g. lifting features and flanges on packages with bolted lids) [254]. However, in a few cases, carbon steel containers are considered [253]. Also, large quantities of carbon steel will be present in the planned repository, as it will be used for several structural purposes. The chemical composition of the materials tested in cementitious environments is summarised in TABLE 4-51.
Under repository conditions, both uniform and localised corrosion are possible at various stages, depending on the precise physical and chemical environment surrounding the metallic container. The most common forms of localised corrosion are pitting corrosion, crevice corrosion, stress corrosion cracking, and microbially influenced corrosion [255]. Electrical coupling with other metals can lead to accelerated corrosion, by a process known as galvanic coupling, and the presence of radiation fields can affect the corrosion behaviour of metals by producing highly reactive species in-situ. These possible degradation modes will be addresed in this section.
187 TABLE 4-51. Ranges of the chemical compositions of the materials tested in cementitious environ- ments [82,252,253,255,258,261,265,272,274].
Chemical composition (wt%) Materials C Si Mn P S Cr Ni Mo Fe
Stainless steels AISI 316L 0.0019-0.055 0.10-0.55 1.09-1.69 0.024-0.033 0.001-0.028 16.30-17.04 10.23-11.90 2.08-2.40 bal AISI 304L 0.025 0.03-0.18 1.46 0.027 0.026 18.00 8.0-9.65 - bal
Carbon steel BS4360 grade 43A 0.15-0.30 0.08-0.28 0.60-0.90 0.010-0.018 0.017-0.06 <0.01-0.07 <0.02-0.1 <0.01-0.02 bal
bal. : balance.
A.4.4.1 Carbon steel
A.4.4.1.1 Uniform corrosion
A.4.4.1.1.1 Aerobic corrosion behaviour
The uniform corrosion rates under aerobic and alkaline conditions are approximately 0.08 µm/year at 30°C and 1.6 µm/year at 80°C [256].
A.4.4.1.1.2 Anaerobic corrosion behaviour
The anaerobic corrosion rate data was found to differ significantly according to the test method used (electrochemical measurement techniques, weight loss techniques, volumetric gas measurement method, or gas chromatography).
Four methods have been used to quantify anaerobic corrosion rates [257-260]:
• electrochemical measurement techniques (AC impedance and coulostatic pulse measurements). In the AC impedance technique, the corrosion rate is calculated from the charge transfer resistance (determined from the impedance plots) using the Stern-Geary equation. The coulostatic pulse technique involves the measurement of the response of the metal/electrolyte system to a fixed amount of electrical charge and its subsequent relaxation when the pulse is completed [257]. • weight loss techniques. • volumetric gas measurement method (gas cell technique). In this technique, the uniform corrosion rate is correlated with the amount of hydrogen gas evolving from the corrosion of the steel. The hydrogen production is measured through the displacement of a column of liquid [257-260]. • gas chromatography. An inert gas stream is passed over the top of a cell, which contains the test electrolyte and the corroding steel, thereby enabling hydrogen gas evolving from the corrosion processes to be carried to, and quantified in, a gas chromatograph [258].
188 Electrochemical techniques are not particularly well suited to the measurement of low anaerobic corrosion rates over long timescales (>1 year) since they can suffer from oxygen ingress through electrical connections to the test cell and tend to have detection limits >0.1 µm/year. Weight loss experiments are not suitable for the measurement of corrosion rates that are likely to vary with time. Also, the rate of metal dissolution associated with anaerobic corrosion is generally low and weight losses from coupons tend to be small, even after long exposure times. Consequently, under such circumstances, the technique is susceptible to experimental error, which has resulted, on occasions, in weight gains rather than weight losses [258]. Because of the limitations of these test methods, the best technique for measuring steel corrosion rates in simulated repository environments are likely to be those designed to measure the volume of hydrogen produced as a result of the corrosion process [258].
In preliminary work, Naish et al. [257,258] used both electrochemical measurement techniques (AC impedance and coulostatic pulse measurements) and weight loss measurements to determine anaerobic corrosion rates of carbon steel. Measurements were made on steel samples that had been embedded in various cement grouts (some containing chloride up to 20,000 mg/L) and stored in an anoxic glovebox. The cement blocks were kept either immersed in water or at 90% relative humidity, with the experimental temperature being in the range of 20°C to 30°C. Potentials below the water reduction potential were reached over a period of time varying from a few days to about 2.5 years from casting, indicating that there was a small amount of oxygen initially present in the cement blocks that was subsequently reduced.
It was found that the anaerobic corrosion rates derived from the electrochemical measurement techniques were up to one order of magnitude higher than those determined by the weight loss technique: • the electrochemical corrosion measurements yielded anaerobic corrosion rates ranging from 0.3 to 5.5 µm/year. The lower of these values relates to the chloride- free mixes and the higher value to the chloride containing mixes. • weight loss measurements, performed after two year immersion, showed anaerobic corrosion rates ranging from 0.23 to 1.44 um/year. The high anaerobic corrosion rates (5.5 µm/year) in the chloride-containing media (electrochemical measurements) were not confirmed by weight loss measurements, which showed no discernible difference between chloride-free and chloride-containing samples.
The anaerobic corrosion rates determined with the gas cell technique were found to be considerably lower than those observed using other experimental techniques such as electrochemical and weight loss measurements [256,258]. The long-term corrosion 11 rates in anoxic alkaline conditions (NaOH, Ca(OH)2, NaOH + Ca(OH)2, groundwater , NRVB12) at 30°C and 50°C in the absence of chloride were <0.1 µm/year. At 80°C, corrosion rates were ~1 µm/year after 2,500 hours and still falling (the anaerobic corrosion rate of carbon steel was found to decrease rapidly with time) [258].
11 the simulated groundwater in this study contained 0.358 g/L NaHCO3, 1.404 g/L Na2SO4, 0.0276 g/L NaBr, 2.02 g/L CaCl2.2H2O, 0.528 g/L MgCl2.6H2O, 0.0374 g/L SrCl2.6H2O, 0.188 g/L KCl, 18.2 g/L NaCl. 12 the NRVB (Nirex Reference Vault Backfill) mixture used in this study contained 26 wt.% OPC (Ordinary Portland Cement), 9.9 wt.% lime, 29 wt.% limestone flour, 35.1 wt.% A.R. (analytical reagent) water. The water to solid ratio was 1:2.85.
189 The decrease in the corrosion rate of carbon steel with time was accompanied by the development of a black film, believed to be magnetite. The addition of chloride to the alkaline test solutions did not have a major effect on the corrosion rate at 30°C and 80°C, but there was generally an increase in corrosion rate at 50°C [258]. The estimation, from the gas cell technique, of the anaerobic corrosion rate for carbon steel in NRVB at 50°C is presented in FIGURE 4-70. The highest corrosion rate in NRVB was ca. 4.82 µm/year at 72 hours, but it then decreased to ca. 0.02 µm/year after 11,100 hours. In NRVB containing chloride the highest corrosion rate was ca. 3.15 µm/year at 50 hours, but this value then decreased, with some perturba-tions, to ca. 0.04 µm/year after 11,000 hours[258].
FIGURE 4-70. Anaerobic corrosion rates of carbon steel specimens in NRVB with and without 20,000 mg/L chloride (T = 50°C) derived from the gas cell technique [258].
Hydrogen overpressure, in the range 100 KPa (1 atm) to 10 MPa (100 atm), was found to have no discernible effect on the anaerobic corrosion rate of carbon steel in anoxic
190 alkaline solutions [257,259]. This observation was confirmed by other investigations: Smart et al. [78,82] found that hydrogen pressures up to ~9 MPa had no effect on the corrosion potential of anaerobically corroding carbon steel wires in artificial Swedish groundwater.
Decreasing the pH of the test solutions was found to increase the corrosion rate of carbon steel in anoxic solutions [258]. A decrease from pH 13 to pH 4 increased the corrosion rate by one to two orders of magnitude, as illustrated in FIGURE 4-71. However, in the Nirex repository, the backfill would have an extremely large buffering capacity, so it is highly unlikely that corrosion processes would significantly shift the pH of the porewater.
(a)
(b)
FIGURE 4-71. Influence of pH on the anaerobic corrosion rates (derived from the gas cell technique) of carbon steel at 50°C [258]. (a) pH = 13 (0.1M NaOH solution with and without the addition of 0.35 wt.% NaCl; (b) pH = 4 and 7 (artificial groundwater, unbuffered).
In highly alkaline anoxic solutions (pH = 13; 0.1M NaOH), the corrosion rate of carbon steel, in chloride-free solutions, at the end of the tests (after ~15,000 hours) was ~0.04 µm/year. At pH 7, the anaerobic corrosion rate fell with time to ~0.2 µm/year (after ~8,500 hours). In pH 4 groundwater, the corrosion rate decreased to ~0.6 µm/year (after 8,600 hours). These long-term corrosion rates are about 5 (pH 7) or 15 (pH 4) times the rates recorded in pH 13 solutions. It was found that the pH of the test solutions increased during the tests to a value of 8-9, due to proton reduction [258].
191 Under anoxic conditions, it was found that the solid phase barrier did not significantly retard the corrosion rate [258]. This result is in contrast to that found in aerated concrete samples, where the corrosion rate can be one tenth of that measured for steel in solution. This difference is believed to be because under aerated conditions the concrete acts as a barrier to the mass transport of oxygen to the steel surface. However, water, which provides the cathodic current under anoxic conditions, is always present in excess at the steel surface.
A good agreement was found between the anaerobic corrosion rates obtained from the gas chromatography measurements and those calculated from the gas cell technique at 80°C [258].
In conclusion, the current best estimate for the long-term anaerobic corrosion rate of carbon steel in alkaline media is of the order of <0.1 µm/year at 30°C [256,258] and <1 µm/year [256] at 80°C.
A.4.4.1.2 Localised corrosion
The critical chloride concentration and critical temperature for the initiation of pitting corrosion of carbon steel embedded in a cementitious backfill (NRVB) have been determined. Long-term tests were conducted at temperatures ranging from ambient up to 80°C and chloride porewater concentrations up to 100,000 mg/L (measured at room temperature) [253]. Carbon steel pitted at a relatively low chloride concentration, equivalent to approximately 6,000 to 28,000 mg/L chloride in the porewater at room temperature. The pitting susceptibility increased with increasing temperature: at 50°C, the critical chloride concentration decreased to below 4,000 mg/L and less than ~300 mg/L chloride in the porewater was required to cause pitting at ~80°C [253].
The maximum depth of penetration of corrosion was determined on carbon steel samples originating either from long-term exposure tests with exposure periods of up to four years (carbon steel plates were embedded in various cement mixes) or from electrochemical experiments (carbon steel discs were held under galvanostatic control) [261]. The maximum pit depths were measured using the incremental grinding technique13 and the results were analysed using extreme value statistics to extrapolate the results from the relatively small samples used for the tests (long-term exposure tests: ~300 cm²; electrochemical tests: ~8 cm²) to the surface area of real containers (~4 m²). The results for carbon steel plates embedded in NRVB at ambient temperature, 50°C, and 80°C are summarised in TABLE 4-52. After embedding the plates in NRVB, containing 2 wt.% CaCl2.2H2O, the blocks were fully immersed in saturated aqueous Ca(OH)2 containing 10,000 mg/L chloride. There was little tendency for the depth of attack to increase with increasing exposure period. In all instances pits tended to spread laterally rather than deepen. No strong correlation was found between the maximum pit penetration depths and temperature. The most probable maximum pit depths predicted with extreme value statistics were generally significantly less than the
13 the incremental grinding technique involved the sequential removal of 25 µm thick layers of metal from the surface of the samples. After each traverse of the grinder, the presence (or absence) of corrosion was recorded. And this was repeated until all races of corrosion was removed.
192 depths predicted from an empirical pit growth law. Nevertheless the rate of corrosion appears to be sufficient to penetrate the walls of thin-walled carbon steel containers within a few years [261].
TABLE 4-52. Measured maximum penetration depths (incremental grinding technique) and estimated maximum penetration depths (using extreme value statistics and pit growth law) for carbon steel plates fully embedded in NRVB and immersed in saturated Ca(OH)2 solution containing 10,000 mg/L chloride at ambient temperature, 50°C, and 80°C [261].
(2) T Test Observed PMAX on 4m² PMAX from pit (°C) (1) (mm) (3) duration PMax growth law (years) (mm) (mm)
AT(4) 1.55 3.94 7.08 9.23 AT 2.55 2.08 5.08 12.0
50 1.55 2.36 3.71 9.21 50 2.53 2.21 4.90 11.96
80 1.53 3.99 n/a(5) 9.17 80 2.49 1.12 2.59 11.86
(1) measured maximum pit depths on carbon steel samples using the incremental grinding technique. (2) estimated maximum pit depths extrapolated to the surface area of a container (4 m²) using extreme value statistics : -ln[-ln(1 - s/S)] = B(x - A/B) (s = 4 cm²; S = 4 m²; x = extrapolated pit depth). (3) estimated maximum pit depths predicted from a pit growth law : Pmm = 7.3 · t 0.53 (Pmm is the pit depth that has only a probability of 0.1 of being exceeded and t is the time in years). (4) ambient temperature. (5) not applicable (the data from the extreme value statistical analysis were inadequate to predict reliable pit depths).
It was observed that the initiation and propagation of localised corrosion in cementitious environments depended upon the ability of cement to immobilise chloride ions [261]. Experiments were performed in which chloride ions were deliberately added to various cements in which carbon steel plates were embedded. The maximum penetration depths for carbon steel plates, together with information about the porewater composition ([Cl-] and [OH-]) are presented in TABLE 4-53. The results were compared to control plates, i.e. plates that were embedded in their respective cement mixes with no added chloride. Pits were found to propagate deeper in the presence of added chloride. Although some corrosion also occurred on the plates that were in contact with cements that were virtually free of chloride. The deepest pit penetrations were observed in Sulphate Resistant Portland Cement (SRPC). This was believed to be associated with the high chloride concentrations in the porewater. In the 8:2 Blast Furnace Slag (BFS)/OPC cement and in grout 2, there was little difference in pit depth between the control samples and the samples with added chloride. However, in these cements, the porewater chloride concentrations were lower than in the SRPC samples, due to their higher chloride binding capacity [261].
193 TABLE 4-53. Measured pit penetration depths on carbon steel samples exposed to various cement mixes, together with the porewater analyses [261].
(1) t Type of cement [Cl-] in mixing Measured PMax Porewater analysis (months) mix water(2) (mm) (mg/L) Side (a) Side (b) [Cl-] [OH-] (mg/L) (mg/L)
SRPC(3) 25000 0.94 0.66 ~12500 1360 6 (4) 8:2 BFS/OPC 5000 0.38 0.46 840 4352
SRPC 0 0.18 0.13 9 211 8:2 BFS/OPC 0 0.41 0.51 37 88 grout 2(5) 0 0.69 0.66 190 58 24 SRPC 25000 0.79 0.94 10600 75 8:2 BFS/OPC 5000 0.33 0.64 590 88 grout 2 4000 0.51 0.69 1550 35
The high hydroxyl concentrations, for the samples after six months, indicate that the pH of the porewater samples was greater than 12.5. A decrease of pH of the samples taken after 24 months may have been caused by carbonation. (1) test duration. (2) chloride content of the water used to prepare the cement mixes. Different mix water concentrations were used so that the chloride concentration in the final cement mix was constant at 1 wt.% OPC. (3) 100% Sulphate Resistant Portland Cement (SRPC). (4) Blast Furnace Slag (BFS) / Ordinary Portland Cement (OPC). (5) 8:2 BFS/OPC-FLA (FLA: Fine Limestone Aggregate).
In conclusion, pitting corrosion of carbon steel containers in a deep repository is likely, but it may be limited by the supply of oxygen once the repository re-saturates, or by the availability of water prior to re-saturation. Both the aeration period calculations [253] and experimental measurements of pit growth rates (experiments in which carbon steel samples were embedded in a range of cements, including NRVB, and held in either a moist air or saturated environment) [260] suggested that for a fully saturated backfill, the rate of oxygen diffusion would be too slow to support corrosion for more than a few years, during which time pits may propagate a few millimeters into the steel [256,261].
A.4.4.2 Stainless steel
The corrosion of stainless steel, in many environments, including cements and concrete, is normally controlled by the presence of a passive oxide film, which limits the corrosion rate to a very low value.
A.4.4.2.1 Uniform corrosion
A.4.4.2.1.1 Aerobic corrosion behaviour
The uniform corrosion rates under aerobic and alkaline conditions are approximately 0.03 µm/year at 30°C and 0.5 µm/year at 80°C [256].
194 The corrosion rates of 316L stainless steel covered with a passive film have been determined at various pH values ranging from 2 to 13 and for temperatures of 30°C, 50°C, and 80°C, using an electrochemical method [262]. No chloride was added to the test media. The data are shown in TABLE 4-54 and FIGURE 4-72. These data showed that below pH 7 the passive corrosion rates increased rapidly with increasing acidity, but were relatively insensitive to pH under alkaline conditions, with a 2-3 times increase in corrosion rate at the highest pH investigated (pH 13), compared to pH 7. The data also showed that uniform corrosion is accelerated with increasing temperature: the effect of temperature was highest at the most acidic pH (pH 2). The data collected at pH 2 at 50°C and 80°C have to be treated with caution as localised pitting corrosion was observed during these tests: visual examination of the stainless steel electrode tested at 80°C revealed 4 small pits. It was believed that pitting occurred due to chloride contamination (from the SCE reference electrode), although a quantification of the chloride concentration was not performed.
TABLE 4-54. Passive current measurements for 316L stainless steel under oxic conditions at various pH (1.9-13) and temperatures (30-80°C) [262].
pH Temperature Passive current Equivalent passive [°C] [nA/cm²] corrosion rate(1) [µm/year]
30 4.5(±0.8 0.05 13 50 13.3(±4.2) 0.15 80 70.0(±12) 0.80
30 2.7(±0.2) 0.03 11 50 8.5(±0.9) 0.10 80 19.8(±3.8) 0.23
30 2.8(±0.2) 0.03 9.2 50 5.8(±0.3) 0.07 80 18.9(±1.9) 0.22
30 2.3(±0.4) 0.03 6.4 50 3.8(±0.3) 0.04 80 18.4(±1.8) 0.21
30 105(±10) 1.2 4.0 50 160(±10) 1.9 80 220(±30) 2.6
30 3900(±250) 45 1.9 50 20,000(2)(±5,000) 232 80 >100,000(2) 1160
(1) the corrosion rates are calculated based on the measured passive currents. (2) localised pitting corrosion occurred.
195 10000
1000
100 [µm/year]
10
1 Corrosion rate
0.1
0.01 02468101214 pH
30°C 50°C 80°C
FIGURE 4-72. Passive corrosion rates, calculated from passive current measurements, for 316L stainless steel under oxic conditions at various values of pH (1.9-13) and temperature (30°C, 50°C, and 80°C).
A.4.4.2.1.2 Anaerobic corrosion behaviour
Anaerobic corrosion rates ranging from 0.1 µm/year to 0.001 µm/year have been estimated for stainless steel in alkaline media:
• the highest corrosion rates (0.1 µm/year) were determined from electrochemical experiments in which the stainless steel was polarised at a potential close to the oxygen reduction potential, so that its reduction was no longer thermodynamically possible (the reduction of dissolved oxygen results in a small cathodic current that would artificially reduce the value of the recorded passive current) [262]. • the lowest corrosion rates were obtained using the gas cell technique. Experiments were performed to measure the hydrogen gas generation rate from 316 stainless steel wires immersed in highly alkaline (pH~13) and anoxic solutions. Comparison with control experiments, which contained no metal, showed very similar results [258,259]. This indicated that if any corrosion of the stainless steel occurred, it was at such a low rate that gas evolution was too low to distinguish from the control cell background levels. This corresponds to a corrosion rate of less than about 0.001 µm/year [254]. The electrochemical open circuit potential of a stainless steel wire in a gas cell containing 0.1M NaOH and 0.35 wt% NaCl at 30°C was measured as being below the hydrogen evolution potential (~-960 mV vs SCE) [260] and measurements of the hydrogen dissolved in the steel were found to show slightly enhanced levels compared to uncorroded samples. This suggests that the hydrogen produced by very slow anaerobic corrosion was dissolved into the
196 stainless steel, rather than released as a gas. These experiments indicate that stainless steel is strongly passive under these conditions [254,260]. Periodic inspection of the wires during the course of the experiments showed that the wires remained bright throughout the tests [258]. • a realistic upper limit for the anaerobic corrosion rate of stainless steel in alkaline repository environments of 0.01 µm/year has been reported [256,258-260,263].
Hydrogen overpressure, in the range 100 KPa (1 atm) to 10 MPa (100 atm), was found to have no discernible effect on the anaerobic corrosion rate of stainless steel in anoxic alkaline solutions [259].
The corrosion rate of 316L stainless steel in simulated cement porewater under anoxic conditions was also found to be relatively insensitive to pH within the range 6.4 to 13. At pH 4, the corrosion rate was about one order of magnitude higher [262].
In conclusion, the current best estimate for the long-term anaerobic corrosion rate of stainless steel at all repository temperatures is <0.01 µm/year [256,260,263].
A.4.4.2.2 Localised corrosion
A.4.4.2.2.1 Pitting corrosion
The critical chloride concentration and critical temperature for the initiation of pitting corrosion of stainless steel embedded in a cementitious backfill (NRVB) have been determined. Long-term tests were conducted at temperatures ranging from ambient up to 80°C and chloride porewater concentrations up to 100,000 mg/L (measured at room temperature) [253]. In cementitious environments, the onset of pitting is determined by the ratio of the chloride and hydroxide concentrations in the porewater. Pitting of 304L stainless steel only occurred at 60°C and a porewater chloride concentration of >50,000 mg/L (in this case, the molar ratio of chloride to hydroxide at room temperature was estimated to be over 29 at pH 13). No pitting at all was observed on any of the 316L specimens tested in NRVB with up to ~100,000 mg/L chloride in the porewater at room temperature, or ~50,000 mg/L chloride in the porewater at 45°C or 70°C [253].
Using electrochemical test methods, pitting breakdown potentials and repassivation potentials were determined for 304L and 316L stainless steel as a function of pH (3-13) in simulated crevice solutions (20,000 mg/L chloride). The results are presented in TABLE 4-55 and FIGURE 4-73 [253].
197 TABLE 4-55. Corrosion potentials (ECORR), pitting potentials (EP), and repassivation potentials (ER) for 304L and 316L stainless steels (versus SCE) in deaerated 20,000 mg/L chloride solutions at 50°C and at various pH (3-13) [253].
304L 316L pH ECORR EP ER ECORR EP ER [mVSCE] [mVSCE] [mVSCE] [mVSCE] [mVSCE] [mVSCE]
13(1) -850 500 -330 -840 520 to 530 -200 11 -460 to -510 0 to 40 -320 to -310 -450 to -570 40 to 50 -340 to -210 9 -300 to -400 -130 to -80 -320 to -310 -400 to -500 -30 to -20 -270 to -250 7 -400 to -500 -60 to -50 -310 30 to 60 -270 to -230 5 -200 to -60 -370 to -290 -50 to -20 -310 3 -200 to -460 active to -140(2) -170 to -200 -330 to -440 -80 to -70 -140 to -130
(1) the pH 13 solutions contained only 12,000 mg/L chloride. (2) some 304L specimens at pH 3 were observed to display active corrosion behaviour rather than passivation followed by pitting.
(a)
(b)
FIGURE 4-73. Pitting breakdown and repassivation potentials for 304L and 316L stainless steel in deaerated 20,000 mg/L chloride solutions at various pH [253]. (a) pitting breakdown potentials. (b) repassivation potentials.
198 The pitting potentials for 304L and 316L stainless steel decreased with increasing temperature (only measured for pH 13), and increased with increasing pH (over the pH range 3-13). The pitting potentials were lower for 304L compared to 316L, reflecting the greater protection provided by molybdenum in the 316L stainless steel. The repassivation potentials were not significantly sensitive to changes in pH and temperature [253].
As the corrosion potentials were significantly below the pitting potentials, it was concluded that these results indicate that the passive films present on 304L and 316L stainless steel provide sufficient protection to prevent pit initiation in cementitious porewater except in very high chloride concentrations and at high temperatures [253].
A.4.4.2.2.2 Crevice corrosion
The resistance of grade 304L stainless steel to crevice corrosion in alkaline environments has been investigated in experiments in which sections of 304L stainless steel waste container lids were embedded in a cementitious grout (NRVB) with porewater chloride concentrations up to 50,000 mg/L, at 40°C and 80°C [254]. No crevice corrosion was observed at 40°C at porewater chloride concentrations up to 20,000 mg/L. At 80°C, the critical chloride concentration for the onset of crevice corrosion was close to 20,000 mg/L: crevice corrosion was produced within a few weeks when 20,000 mg/L chloride was present, but not with 15,000 mg/L.
The alkaline environment provided by NRVB (porewater pH ~13) is expected to result in an increase in the critical chloride concentration for crevice corrosion [254]. However, modelling studies could not conclusively dismiss the possibility of a critical crevice solution (CCS14) being attained at pH 13 [265,266]. Experimentally the addition of alkali, outside an active 316L crevice initiated at neutral pH, was shown to reduce but not eliminate the rate of crevice corrosion [254,265].
A.4.4.2.2.3 Stress Corrosion Cracking (SCC)
A literature-based assessment [267] of the likelihood of stress corrosion cracking of austenitic stainless steel waste containers in NRVB concluded that chloride-induced SCC was a distinct possibility. Stainless steels are also known to be susceptible to SCC in alkaline conditions, but it was concluded that in the absence of concentration effects and with temperatures below 80°C this form of SCC was unlikely. No other potential causes of SCC in NRVB were identified.
The possibility of chloride-induced SCC was supported by the examination of an unstressed bullet specimen of 304L which split after being embedded for two years in NRVB containing ~100,000 mg/L chloride in the porewater at 80°C. The examination revealed signs of failure by SCC, which occurred under the influence of residual stresses present in the specimen. The specimen had also suffered from severe pitting corrosion and it is possible that the pits acted as initiation sites for the subsequent SCC [255].
14 a critical solution composition that results in the onset of crevice corrosion.
199 Both chloride and thiosulphate are very likely to be present in a repository – chloride from the groundwater and thiosulphate from the cement grouts based on Blast Furnace Slag (BFS) that may be used to encapsulate some of the waste. The effect of mixed alkaline solutions of chloride and thiosulphate ions on the susceptibility of 304L and 316L stainless steels to SCC at 80°C was investigated by means of slow strain rate testing (SSRT) [252,255]. The SSRT specimens used were fitted with crevice formers to induced crevice corrosion and a galvanostatically controlled current was passed to ensure that crevice corrosion persisted. Summaries of the test results obtained for non- welded 304L and 316L stainless steels (i.e. stainless steel which had not been subjected to a simulated weld heat treatment) are presented in TABLE 4-56 and TABLE 4-57, respectively. The main conclusions are [69]:
• SCC was not observed in 316L or 304L stainless steel in solutions containing 17,750 mg/L Cl- at 80°C at pH 7 or 11. 2- - • the addition of 3,360 mg/L S2O3 to solutions containing 17,750 mg/L Cl led to SCC in 316L and 304L stainless steel at 80°C for both neutral and pH 12 conditions. The severity of SCC was found to increase with decreasing pH. 2- • in solutions containing 7,000 mg/L S2O3 at pH 12, SCC was observed for 304L stainless steel with only 1 mg/L chloride present. For 316L stainless steel, cracking was observed when chloride was added at levels of 100 mg/L and higher. • in solutions containing 15,000 mg/L Cl-, the minimum concentrations of thiosulphate required for cracking of 304L and 316L stainless steels to occur were >10 mg/L and > 100 mg/L, respectively. • failure occurred by both transgranular SCC (TGSCC) and intergranular SCC (IGSCC).
Suleiman and Newman [268] have found that SCC of austenitic stainless steels occurred in chloride-thiosulphate mixtures at pH 12.5, but not at pH 13.
200 TABLE 4-56. Summary of results from slow strain rate tests on 304L stainless steel in chloride / thiosulphate solutions at 80°C [255].
pH Chloride Thiosulphate Result(1,2) [mg/L] [mg/L]
17750 0 no SCC 7 17750 3360 SCC
17750 3360 SCC 17750 3360 SCC 0 3360 no SCC 20000 500 SCC 12 20000 100 SCC 20000 10 no SCC 1 7000 SCC 0 7000 no SCC
0 3360 no SCC 13 17750 3360 SCC
The test specimens were fitted with PTFE crevice formers and a galvanostatically controlled anodic polarisation at a current of 12.5 µA/cm² was applied. (1) specimens that failed by cleavage and/or intergranular failure were classified as suffering SCC. (2) conclusions drawn from SEM observations.
TABLE 4-57. Summary of results from slow strain rate tests on 316L stainless steel in chloride / thiosulphate solutions at 80°C [255].
pH Chloride Thiosulphate Result(1,2) [mg/L] [mg/L]
17750 0 no SCC 7 17750 3360 SCC
10 17750 3360 SCC
11 17750 0 no SCC
17750 3360 no SCC 17750 3360 SCC 17750 3360 SCC 15000 7000 SCC(3) (3) 12 1000 7000 SCC 100 7000 SCC(3) 1 7000 no SCC 20000 1000 SCC(3) 20000 500 SCC(3) 20000 100 no SCC
The test specimens were fitted with PTFE crevice formers and a galvanostatically controlled anodic polarisation at a current of 12.5 µA/cm² was applied. (1) specimens that failed by cleavage, intergranular and/or non-ductile failure were classified as suffering SCC. (2) conclusions drawn from SEM observations. (3) non-ductile failure was clear from the visual examination (no SEM analysis).
201 Stress-assisted intergranular corrosion sometimes occurs in high carbon grade stainless steels which have been welded, due to the production of chromium carbides and the formation of chromium-depleted grain boundary regions in the heat affected zone (HAZ). The problem can be greatly reduced by using low carbon grade stainless steels with a cross section of less than 5 mm. Thicker section material may be susceptible to SCC because of the longer time involved in cooling and the greater likelihood of prcipitating chromium carbides [267].
A.4.4.2.3 Galvanic corrosion
Most intermediate level wastes will be conditioned by encapsulation in a cementitious grout, and therefore information about the galvanic corrosion between the different metals, present in the waste, and the container is required.
Galvanic series have been measured for a range of different metals, which might arise in radioactive waste streams, in naturally aerated and chloride-free 3:1 PFA/OPC15 cement at 30°C and 80°C. The materials tested included graphite, Zircaloy, zirconium, 304L and 316L stainless steels, Cromweld CR12, aluminium, aluminium alloy 5083 (4.5% Mg), Magnox (magnesium alloy), lead, zinc, copper, iron, carbon steel, and uranium. The results at 30°C are presented in FIGURE 4-74 [269].
Stainless steel was not adversely affected by galvanic coupling with other waste materials, with the possible exception of graphite. Grade 304L and 316L stainless steels behaved almost identically [269]. When coupling grade 304L and 316L stainless steels, the galvanic currents fell to below 0.01 µA, which is equivalent to less than 0.04 µm/year, after two weeks [269]. Additional work was carried out to investigate the possible adverse effects of coupling stainless steel and graphite in a cementitious grout by electrically coupling specimens embedded in grout [270] and by setting up galvanic crevice specimens in artificial cementitious porewaters [271]. These tests showed that even graphite did not accelerate the corrosion of stainless steel to any significant degree except in the presence of high concentrations of chloride and oxygen at high temperatures (80°C).
Carbon steel directly exposed to cement showed only a small amount of general corrosion, suggesting that they had remained fairly passive. However, evidences of crevice corrosion were observed on the carbon steel samples [269].
The magnitude of the effect of galvanic coupling will depend on a number of factors, including the potential difference between the metals or alloys involved in the couple, the relative areas of the metals coupled, and the conductivity of the electrolyte between them. As many wasteform encapsulation cements and grouts have a high resistance to ionic electric current it is likely that the distances over which galvanic effects can operate will be limited to localised parts of the container. The galvanic couples which arise in waste containers may be temporary, because the galvanic couple may occur through a point contact which is lost once the material has corroded away.
15 PFA/OPC : mixture of pulverised fuel ash (PFA) and ordinary Portland cement (OPC).
202
FIGURE 4-74. Galvanic series for a range of metals in 3:1 PFA/OPC cement at 30°C [269].
A.4.4.3 Passive film formation on carbon steel and stainless steel
The corrosion of carbon steel and stainless steel in alkaline environments (pH~10-13) under oxic conditions, including cements and concretes, is normally controlled by the presence of a passive oxide film, which limits the corrosion rate to a very low value [86,259,272,273]. If the pH falls below 10 for any reason, the protective film on carbon steel becomes unstable, at least at normal environmental temperatures [273]. At higher temperatures, for example during the thermal phase of nuclear waste storage, this lower pH limit would probably shift to lower values although that remains to be quantified.
The formation of a passive film on carbon steel and stainless steel under anoxic conditions was investigated in highly alkaline media (0.1M NaOH, pH = 13) and in near-neutral solutions (artificial groundwater, pH = 6.9). Experiments were carried out in which a completely bare steel surface was produced in situ, either by cutting a piece of wire (guillotine electrode technique) or by fracturing a specimen (shear electrode technique). A combination of open-circuit potential and potentiostatic methods under reducing conditions indicated that the passive film rapidly reformed, as illustrated in FIGURE 4-75. FIGURE 4-75 shows the open circuit potential versus time profile obtained on exposing a freshly generated carbon steel and stainless steel surface to 0.1 mol/L NaOH. Initially, the open circuit potential, for both carbon steel and stainless
203 steel, fell rapidly indicating an increase in the corrosion rate, which is consistent with the removal of a passive oxide film. The potential then moved in a positive direction to the original value of the old surface (after about 16 hours). The fact that the open circuit potential of the stainless steel remained positive of those observed with carbon steel at all times reflects its higher corrosion resistance. These experiments, however, did not completely exclude the possibility that the passivation of the steel surface had been caused by reaction of the steel with the small amount of oxygen gas remaining in solution in the water [254,272].
(a) (b)
FIGURE 4-75. Evolution of the open circuit potential of a fresh carbon steel and stainless steel surface exposed to 0.1M NaOH (pH = 13) (guillotine electrode assembly) [272]. (a) carbon steel. (b) stainless steel. The more negative the potential the higher the corrosion rate. Region A-B is for the old surface, region B-C is for the fresh surface.
Complementary experiments using isotopically labelled water, i.e. water containing an oxygen-18 tracer, combined with Secondary Ion Mass Spectroscopy (SIMS), used to analyse the surface of the steel, have confirmed that under anoxic conditions bare 316L stainless steel surfaces reform the passive film by direct reaction with water (rather than with dissolved oxygen) [254,274]. These results suggest that the surfaces of steel waste containers in an oxygen-free deep cementitious repository will be covered by a protective film, leading to very low corrosion rates [260,274].
In combination, these experiments have shown that the surfaces of stainless steel containers will remain passive throughout the anaerobic period of repository evolution and that they will have a negligible impact on gas production [254,260].
From the literature, it is evident that the oxide films formed on carbon steel under aerobic and anaerobic conditions are distinctly different [275]. The aerobic film mainly consists of hematite, Fe2O3, whereas the anaerobic film is magnetite, Fe3O4. The presence of a thin black oxide film, identified as magnetite by Laser Raman Spectroscopy, was confirmed by analysis of corrosion products formed on carbon steel bullets exposed for 2 years to deaerated cement mixes at room temperature. These layers were also found to be stable in air [77,257]. Similarly, the black corrosion
204 products formed on carbon steel wires in anoxic gas cells containing alkaline artificial porewater was identified as predominantly magnetite, which was more crystalline at 50°C and 80°C compared to 30°C [260]. There was also an indication of amorphous FeO(OH) and/or Fe(OH)2 at 30°C [260].
No such films have, as yet, been identified on stainless steel, probably because the films formed on stainless steel are several times thinner than those on carbon steel, making their detection much more difficult [272]. Experiments conducted with stainless steel in artificial groundwater suggest that two types of protective films form at pH 7. The first film to form does not require the presence of dissolved oxygen in the solution. However, the second layer, which takes longer to form, either requires the presence of dissolved oxygen or involves a phase change in the first film [272]. A possible candidate for the first film, which provides the majority of the protection, is Cr2O3, the formation of which is thermodynamically possible in the complete absence of oxygen, and is known to provide protection against corrosion in neutral and alkaline environments [275]. An alternative explanation for the observed lack of corrosion of stainless steel is that a layer of nickel is generated on the surface, which, according to Pourbaix M. [275], is immune to corrosion in neutral and alkaline environments under anaerobic conditions.
The experimental results suggest that protective films, probably oxides, form on both carbon and stainless steels under electrochemical conditions similar to those expected within a deep repository for radioactive waste during the anoxic stage of its evolution. The thickness required for these films to provide a significant degree of protection against corrosion has been estimated at approximately 25 monolayers for carbon steel and 5 monolayers for stainless steel [272].
205 A.5 MODELLING
A.5.1 Modelling studies of carbon steel and low-alloyed steel containers
Regression analyses of the thickness reduction data of carbon and low-alloyed steels in salt brines at various temperatures showed a linear increase of the thickness reduction with exposure time (see TABLE 4-6 and FIGURE 4-5).
Mathematical modelling, which was based on empirical corrosion models for carbon steels in aqueous and vapour phases and developed in the framework of the Yucca Mountain Project [276], was used to predict the general corrosion of the fine-grained carbon steel TStE 355 exposed to saturated bentonite at different temperatures (50, 75, 100°C). The following equation was used [276]:
−1 2 ln Pmedia = b0 + b1 ⋅ ln t + b2 ⋅ T + b3 ⋅T (21) with Pmedia : average penetration (µm) T : temperature (K) t : time (years) b0, b1, b2, b3 : constants
The corrosion data (average penetration depth) of carbon steel exposed to saturated bentonite, obtained experimentally at different temperatures and test durations (see TABLE 4-48), were fitted to the mathematic equation (21), using the ‘least squares solution’, giving the following values of the constants [67,248-250]: b0 = 39.7056 ; b1 = 0.43412 ; b2 = -9084.67 ; b3 = -0.000092
FIGURE 5-1 shows the extrapolation results of the average penetration depth versus the disposal time. According to this extrapolation, estimated average penetration depths for carbon steel containers after 250 years, will be lower than 200 microns [250].
Marsh et al. [277-282] have carried out long-term immersion tests on carbon steel coupons immersed for up to 30,000 hours at 90°C in chloride and carbonated -1 - solutions (0.1M NaHCO3, 1000 mg.L Cl , pH = 8.4). The maximum pit depths were measured using the incremental grinding technique and the results were analysed using extreme value statistics (Gumble distribution function) to extrapolate the results from the relatively small samples used in the tests to the surface area of real disposal containers. They found that the maximum pit depth (P) varied with time (t) according to the following semi-empirical exponential law:
P = k . t n (22) with k a constant dependent on the material and environmental conditions, whereas n is a numerical constant (mostly situated between 0.4 and 0.6). Over the years, different values for k and n have been found (see TABLE 5, part B).
206 200
175 298 ºK 323 ºK 348 ºK 150 373 ºK
125
100
75
50
Average penetration (microns) penetration Average 25
0 0 50 100 150 200 250 300 Time (years)
FIGURE 5-1. Estimated average penetration depth (µm) of the fine-grained carbon steel TStE 355 in saturated bentonite at different temperatures [250].
A.5.2 Modelling studies of copper canisters
A.5.2.1 Modelling studies to predict the long-term general corrosion behaviour of Cu canisters in granitic environments
Two approaches have been used to predict the long-term general corrosion behaviour of Cu canisters under repository conditions [60]:
• mass-transport limited approach. Canister lifetime predictions on a thermo- dynamic description of the corrosion process(es) generally involve an assumption of rapid interfacial kinetics and rate control by the rate of (diffusive) mass transport. Thermodynamic models represent a "worst-case" assessment because of the assumption of infinite interfacial kinetics, and therefore produce a conservative estimate of the corrosion rate. • kinetic approach. Kinetically based lifetime prediction models combine the finite kinetics of interfacial reactions with possible limitation by mass transport to and from the corroding surface.
A lot of similarities exist between the Swedish/Finnish and the Canadian corrosion programmes. However, while thermodynamically based models are used in the Swedish/Finnish programmes, the assessment of the rate of general corrosion in the Canadian programme has been based on both thermodynamic (mass-transport limited) and kinetic models.
A.5.2.1.1 Mass-transport limited approach
In the early 1980’s, SKB developed a mass-transport limited model for predicting the extent of corrosion of copper canisters in a Swedish repository due to sulphide [60].
207 This model was based on the following conservative assumptions [283]:
• all the sulphide present in the bentonite, surrounding the canister, plus that formed by microbial activity in the deposition hole, is consumed in the first 1,000 years. • after this period, sulphide is assumed to diffuse to the canister from the tunnel and from the groundwater. • the sources of sulphide in the tunnel are the sulphide impurities in the bentonite and sulphide produced by microbial activity. • in the groundwater, the maximum sulphide concentration is assumed to be 1 mg/L, corresponding to that present naturally in the groundwater plus that produced by microbial activity.
The predicted depth of corrosion as a function of time due to sulphide from these various sources is given in TABLE 5-1.
TABLE 5-1. Predicted time dependence of the depth of corrosion (in mm) due to sulphide from various sources [283].
Source of sulphide Exposure time (years)
103 104 105 106
Deposition hole: 1. from bentonite 0.032 0.032 0.032 0.032 2. from microbial activity in bentonite 0.023 0.023 0.023 0.023
Tunnel: 1. from bentonite 3.6 × 10-5 3.2 × 10-4 3.2 × 10-3 0.032 2. from microbial activity in tunnel 9.0 × 10-6 5.5 × 10-5 5.5 × 10-4 5.5 × 10-3
Groundwater: 1. present naturally in groundwater 9.1 × 10-5 8.6 × 10-4 8.6 × 10-3 0.086 2. from microbial activity in groundwater 9.1 × 10-5 8.6 × 10-4 8.6 × 10-3 0.086
In the early 1990’s, Werme et al. [284] re-considered the conservative assumption that all the sulphide in the deposition hole was consumed within 1,000 years. Using a 1-dimensional sulphide consumption model, assuming instantaneous consumption of sulphide on the canister surface, they estimated that the sulphide in the deposition hole would be consumed in 850,000 years. The total amount of sulphide present in the deposition hole was higher than that assumed in the early 1980’s, and gave a maximum depth of corrosion of just over 1 mm (after 850,000 years).
Recently, a new model for the general corrosion of Cu in the presence of Cl under oxygen-free conditions has been developed. In this model, the following reaction is used to describe the corrosion process, in saline waters, of copper to aqueous copper complexes under the liberation of dissolved molecular hydrogen [60]:
+ − 1−n Cu (s) + H + n Cl ↔ ½ H 2 (aq) + CuCln (23)
Reaction (23) can go in both directions. In a semi-closed system, such as the repository, a maximum rate for reaction (23) is given by the rate with which the
208 1-n corrosion products (CuCln and H2) diffuse away from the copper surface. In the absence of other effects, the corrosion rate would be expected to decrease with time as the respective concentration gradients at the canister surface decrease as the two species diffuse through the bentonite. Any process, however, that maintains a high concentration gradient at the canister surface, will be capable in maintaining a high 1-n corrosion rate. Two such processes are the precipitation of CuCln by solid sulphides and the oxidation of H2 by reaction with Fe(III). In bentonite (MX-80), sulphide (~0.1 wt.%) is mostly present as pyrite and Fe (~4 wt.%) is predominantly present as substituted Fe(III) in the montmorillonite lattice and as Fe3O4 particles. 1-n The diffusive flux of CuCln and H2 away from the canister, and hence the corrosion rate, was calculated using a mass-transport reaction model with cylindrical geometry. FIGURE 5-2 shows the predicted corrosion rate and depth of corrosion as a function of time for 1.0 mg reactive sulphide and 3.0 mg reactive Fe(III) per kilogram bentonite. The corrosion rate increases with time initially as the Cl- diffuses to the canister surface from the rock (a process that takes of the order of 300-400 years). After this time, the corrosion rate is predicted to decrease with time as the - concentration gradients of the dissolved CuCl2 and H2 become less steep. The sulphide and Fe(III) content of the bentonite nearest the canister are reduced by - reaction with CuCl2 and H2, respectively, although only a relatively small fraction of the total inventory is consumed within the first 6,000 years shown in FIGURE 5-2. The total depth of corrosion is negligible, amounting to only ~0.06 µm after 6,000 years in bentonite at pH 7 with 1.0 M NaCl [60].
FIGURE 5-2. Corrosion depth and corrosion rate as a function of time (pH = 7.0, T = 50°C, 1.0 mg reactive sulphide and 3.0 mg reactive Fe(III) per kilogram bentonite) [60].
Higher sulphide and/or Fe(III) concentrations would lead to higher predicted corrosion - rates, since reaction between these species and Cucl2 and H2, respectively, would maintain the initially high gradients at the canister surface, for a longer time. However, even with a 3,000-fold increase in Fe(III) content and a 100-1,000-fold increase in sulphide content, the maximum depth of corrosion due to reaction (22) is predicted to be <6 µm after 50,000 years in bentonite at pH 7 with 1.0 M NaCl [60].
209 A.5.2.1.2 Kinetic approach
Over the years, three different models have been developed:
• the Cu(II) mass-transport limited model for oxidising conditions [285,286]. • the steady-state kinetic model of King et al. [210]. • the transient kinetic Copper Corrosion Model (CCM) [287–289].
The Cu(II) mass-transport limited model. This model was developed based on the results of corrosion experiments in compacted buffer material under oxidising conditions [207]. Under these conditions, the diffusion of Cu(II) away from the Cu surface limited the rate of corrosion, which is given by
½ 2A ρ m ⎛ ε τ D ⎞ Rate = Cu d 0 ⎜ e f 0 ⎟ (24) ½ ⎜ ⎟ ρCu ()π t ⎝ r ⎠ with ACu : atomic mass of Cu ρCu : density of Cu ρd : dry density of the buffer material m0 : total concentration of cu in the buffer immediately adjacent to the canister surface t : time εe : effective porosity for mass transport of the buffer τf : tortuosity factor of the buffer D0 : bulk-solution diffusion coefficient of Cu(II) r : buffer capacity factor for Cu(II)
The minimum predicted lifetime for a 25 mm thick Cu canister (only 16 mm of which was used as the corrosion allowance) in a Canadian repository was of the order of 30,000 years. This resulted in an extremely conservative assessment of the canister lifetime because in the model the assumptions are used that (i) the corrosion rate does not become limited by the supply of O2 to the canister surface and (ii) infinitely fast interfacial kinetics exist.
The steady-state kinetic model. A mixed-potential model, capable of predicting ECORR and iCORR for a wide range of mass transport and environmental conditions, was developed based on data from short-term electrochemical studies on the anodic - dissolution of Cu in Cl solutions [210] and on the reduction of O2 on Cu [290,291] in combination with steady-state mass-transport expressions. This model has never been used to predict canister lifetimes, but has instead been used to interpret the results of experimental studies. A distinct advantage of the steady-state model over the previous model, however, is its ability to predict ECORR, in addition to iCORR. This is a powerful tool for predicting the long-term localised corrosion behaviour of the canister (by comparing ECORR to critical potentials for localised corrosion).
The transient kinetic Cu corrosion model (CCM). This model is currently the most advanced general corrosion model developed in the Canadian programme [287– 289]. The CCM-model is based on the experimentally determined reaction scheme for the corrosion of Cu in compacted buffer with O2-containing saline groundwater as represented in FIGURE 5-3. The reactions shown in FIGURE 5-3 are assumed to
210 occur in a spatial grid, bounded on the left-hand boundary by the canister surface and on the right-hand boundary by, generally, a major water-bearing fracture. These reactions constitute an electrochemical mixed-potential model, enabling the time dependence of ECORR and iCORR of the canister to be calculated. In this model, the interfacial electrochemical reactions as well as the near-field processes (e.g. mass- transport, adsorption/desorption, precipitation/dissolution, chemical and redox processes, variations in groundwater salinity) are taken into account.
FIGURE 5-3. Mechanism for the general corrosion of copper in compacted buffer material with O2-containing saline groundwater [292].
FIGURE 5-4 shows the predicted time dependence of ECORR and iCORR of a Cu canister in a Canadian repository. The calculations are based on the CCM-model. For the simulation in FIGURE 5-4, the trapped O2 in the buffer and backfill was predicted to be consumed in approximately 2,600 years. Both ECORR and iCORR decrease during this initial period as the trapped O2 is consumed. The mean predicted corrosion rate is ~0.3 µm/year, which is an order of magnitude lower than, that observed by Karnland et al. [205] during the long-term corrosion tests in compacted bentonite at the Äspö Hard Rock Laboratory. In the simulation, however, the buffer is assumed to be completely saturated, wheras the buffer in the pilot scale test may have been partially saturated, which is found to drastically increase the rate of O2 diffusion [209,293].
211
FIGURE 5-4. Predicted time dependence of the corrosion potential (ECORR) and corrosion current density (iCORR) of a copper canister in a Canadian repository [292].
FIGURE 5-5 shows that the general corrosion of the canister ceases once all of the initially trapped O2 is consumed. The maximum depth of corrosion is ~11 µm. Of the total amount of atmospheric O2 trapped in the vault initially, only ~17% causes corrosion of the canister, 17% is consumed by the oxidation of Cu(I) to Cu(II), ~16% is consumed by reactions with Fe(II) minerals, and the remaining 50% is assumed to reside in inaccessible pores in the buffer and backfill materials. Although the Cu(II) produced by oxidation of Cu(I) can support corrosion (reaction (12), see section A.3.3.1), the rate is too slow due to the strong adsorption of Cu(II). Even if all of the Cu(II) were reduced on the canister surface, the maximum additional corrosion would only amount to a further 11 µm [60].
FIGURE 5-5. Predicted time dependence on the extent of corrosion of a copper canister in a Canadian repository. A charge density of 15 C/cm² is equivalent to 11 µm of corrosion [60].
212 A.5.2.2 Modelling studies to predict the long-term pitting behaviour of Cu canisters
The approach taken to model the long-term localised corrosion behaviour of a Cu canister depends on the nature of the available data [60]:
• mechanistic models, which can provide information about either the initiation or growth of pits, can be developed if the pitting mechanism is known. • the probability of pit initiation can be predicted if Eb and ECORR data are available. • predictions of pit growth on canisters can be made based on the measured pit depths of coupons from shorter exposure periods.
Taxén [294,295] has described a model for the growth of pits based on mass transport and chemical equilibrium principles. Continued growth of the pit is contingent on the transport of reactants to the base of the pit where the anodic reaction is located and the transport of dissolved Cu out of the pit (see FIGURE 5-6). Both solid and dissolved corrosion products are assumed to form within the pit. If sufficient Cu is not transported out of the pit, the corrosion products in the pit become so dense and non-porous that growth is stifled and the pit dies.
FIGURE 5-6. Schematic illustration of the site of a corrosion pit in copper with aqueous species diffusing and migrating [294,295].
Mass transport into and out of the pit is described by a combination of diffusion and migration and is simulated by a 1-D mass-balance equation. The pit geometry is simulated by a series of hemispherical shells (with a constant surface area to thickness ratio; the pits are assumed to grow at an equal rate in all directions) describing both the pit in the metal and the cap of porous corrosion products (see FIGURE 5-7). Chemical reactions between various species is assumed to be fast relative to the rate of mass transport, so that equilibrium chemical conditions can be used.
213
FIGURE 5-7. Description of the site of a corrosion pit as consisting of thin shells [294,295].
The fraction of corrosion products precipitating as a solid compared with the fraction that is transported out of the pit is decisive in determining whether the pit can continue to grow. If sufficient corrosion products are transported out of the pit, local acidification within the pit is necessary to avoid precipitation of the corrosion products. Acidification results from the formation of Cu2O according to
+ − 2 Cu + H 2O = Cu2O + 2 H + 2 e (25)
The equilibrium expression for reaction (25) defines a combination of potential and pH for the formation of Cu2O (the lower the pH the higher the potential, and vice versa).
FIGURE 5-8 shows the predicted conditions of potential and [Cl-] for pit propagation.
FIGURE 5-8. Potential-[Cl-] diagram identifying conditions for pit propagation. The thick line indicates that the fraction of the oxidised copper which is transported away from the site of the oxidation as aqueous species is equal to 0.4, the thin line indicates conditions corresponding to the transported fraction equal to 0.5 [294,295].
214 The major conclusions, drawn from these calculations, about the role of various components in the water in the pitting process are summarized below [294,295]:
• pitting of copper is less likely to occur at high pH values. The pH of the bulk water outside the corrosion pit has a small influence on the minimum potential for pit propagation. • pitting is less likely to occur at higher temperatures: for a given porewater composition, the difference between the minimum potential for propagation of a corrosion pit and the upper potential for stability of cuprous oxide increases with temperature. • pitting is more likely to occur in waters with high chloride concentrations: the difference between the minimum potential for propagation of a corrosion pit and the upper potential for stability of cuprous oxide decreases with increasing chloride concentration. • pitting is less likely to occur in water with high carbonate concentrations: a high carbonate concentration may increase the value of the minimum pitting potential and decrease the value of the upper stability potential for cuprous oxide. • pitting is more likely to occur in waters with high sulphate concentrations. • oxygen at low concentrations can also give potentials higher than the minimum pitting potential. • pits where the precipitation occurs mainly outside the cavity have higher growth rates. • pitting of copper has been observed in waters with a composition and temperature such that the minimum pitting potential is in a range where cuprous oxide is not stable at the pH of the bulk water. • pitting of copper is possible in a wide range of solution compositions. In some waters a corrosion pit will not propagate unless the cuprous oxide at the external surface is stabilised or if there is electronic contact with a conducting, more noble phase. • in waters with chloride contents approaching that of sea water, pitting is possible with high propagation rates at high pH values.
Taxén [296] has also developed a pit propagation model for Cu in reducing conditions in the presence of sulphide. In this model, the pit growth is assumed to be limited by the rate of supply of HS- to the base of the pit by diffusion. For a bulk [HS-] of 1 µg/g and a pre-existing pit depth of 0.5 cm, the pit was predicted to grow by a further 0.5 to 1.1 cm in the subsequent 105 years.
Another approach to model the localised corrosion behaviour of the Cu canister is to predict the probability of pit initiation based on observed ENP values. Based on the data reported in FIGURE 4-49, the minimum observed ENP value is –0.12 VSCE. The data reported in FIGURE 5-4 (Canadian situation) is used as a conservative estimate of ECORR of a Cu canister in a Swedish/Finnish repository since ECORR in a Swedish/Finnish repository is likely to be more negative because of more rapid consumption of the trapped atmospheric O2 by reaction with sulphide minerals that are expected to be present in the clay. Comparison of the observed ENP values with the predicted ECORR value suggests that pitting will not initiate on Cu canisters in a repository [60]:
• the predicted ECORR is a minimum of 60 mV more negative than the minimum ENP: the maximum predicted value of ECORR (FIGURE 5-4) is 0.18 VSCE after 3 days,
215 which then decreases to –0.22 VSCE (i.e. 100 mV below the minimum ENP) after ~5 months. During the subsequent ~ 2,600 year long aerated phase, ECORR remains even lower. • even if pits were to initiate on the canister surface, comparison of EPP data (0.07- 0.37 VSCE; see section A.4.3.1.3.2.1) with ECORR suggests that pit growth will not be sustained.
A final approach to predict the extent of localised corrosion on Cu canisters is based on observed pit depths. Because pitting has not been observed on Cu exposed to simulated repository conditions, pit depth data has been taken for this purpose from literature studies of the long-term burial of Cu alloys [297] and from an analysis of pit depths on archaeological artifacts [298]. Whilst the environmental conditions and Cu alloys are different from those in a repository, these studies offer the great advantage of having been conducted over long periods of time (14 and 3,000 years for the long-term burial test and analysis of Cu archaeological artifacts, respectively). These pit depth data can be used in two alternative ways to predict the lifetime of a Cu canister:
• use of the pitting factor, PF (adopted in Swedish/Finnish programme). PF is the ratio of the maximum pit depth (as measured from the original surface) to the depth of general corrosion and has a value > 1 (PF = 1 corresponds to general corrosion). The maximum pit depth on a Cu canister can be estimated by multiplying the depth of general corrosion by PF. Based on the data from the long-term burial tests [297], a conservative PF value of 25 was estimated [220,299]. Based on the analysis of the archaeological artifacts [298], a more realistic PF of 5 has been estimated [283,284,300]. • use of extreme-value statistics (adopted in the Canadian programme) [285,289]. FIGURE 5-9 shows the maximum pit depth as a function of time, based on a probability of 10-11 that the pit will be deeper than the given depth (i.e. a probability of <10-6 for any of the ~105 canisters in the Canadian repository considered in the analysis). For the assumed probability, the maximum pit depth of any of the canisters is 7.6 mm after a period of 106 years [301,302].
FIGURE 5-9. Predicted maximum pit depth on a copper canister as a function of time (assumed canister surface area 5.76 m²) [301,302].
A major disadvantage of the pitting factor and extreme-value analysis approach is that these models do not allow for pit death: propagation is assumed to continue indefinitely
216 regardless of the evolution in the repository environment, albeit at a diminishing rate in the extreme-value approach [60].
A.5.2.3 Modelling studies to predict the long-term SCC behaviour of Cu canisters
As with other forms of corrosion, the probability of SCC of a Cu canister will diminish with time as the repository environment evolves: since all SCC mechanisms involve some degree of oxidation of the metal, the probability of SCC will diminish with time as conditions become anoxic. Because of this, the difficult task of predicting the SCC behaviour of components with design lifetimes of the order of 100,000 years is reduced to a simpler task of predicting the SCC behaviour over the much shorter duration of aggressive conditions (perhaps of the order of tens to hundreds of years) [60].
Predicting the long-term SCC behaviour of Cu canisters is based on either the concept of threshold conditions or on the limited-propagation approach.
Concept of threshold conditions. Saario et al. [235] are developing a SCC model based on the concept of threshold parameters for SCC. Threshold values are defined for e.g. concentration of SCC agents (ammonia, acetate, nitrite), critical potential (ESCC), content of compositional elements of the Cu canister material (e.g. P), and stress. The basic idea behind this approach is that SCC will not occur provided the actual value is below the threshold value.
- The threshold NO2 concentration for the SCC of pure Cu alloys is approximately 46 mg/L [241,303]. In comparison, the maximum nitrite concentration in deep Swedish/Finnish groundwaters is ~0.01 mg/L. From the same studies, ~0.1 VSHE was determined as the critical potential (ESCC) in nitrite environments, which is much more positive than ECORR. Anttila et al. [304] found EH values situated between -0.2 and - 0.3 VSHE (the thermodynamic redox potential represents the maximum value of ECORR). The situation for acetate and ammonia is less definitive, and additional experimental studies are in progress to confirm that SCC of Cu canisters will not occur.
The proposed Cu-alloy for the fabrication of the canister contains 30 to 80 ppm P (P is added to improve the creep strength and creep ductility at higher temperatures). This amount exceeds the threshold P content for SCC in ammonia environments. Thompson and Tracy [305] report a ductility minimum at a P content of ~0.014 wt.% (140 ppm), with a measurable decrease in ductility for P contents as low as 40 ppm (see FIGURE 5-10). Sato and Nagata [306] found a threshold P content of between 50 and 80 ppm (see FIGURE 5-11).
217
FIGURE 5-10. Dependence of time-to- FIGURE 5-11. Effect of P content and ap- failure on P content [305]. plied stress on SCC of Cu [306].
There is relatively little information available regarding the threshold stress for SCC of Cu in relevant environments. Only two values were reported by Saario et al. [235]: 140 MPa in 1 mol/L NaNO2 at room temperature and 40 MPa in moist ammonia atmospheres. It is likely that the outer Cu shell will undergo some plastic deformation, making it difficult to claim that the sum of the applied and residual stresses will not exceed the threshold stress for SCC.
Limited-propagation approach. Pettersson and Oskarsson [307] calculated a maximum crack depth of less than 10 mm, after 100,000 years, based on crack growth rate measurements and using the limited-propagation approach (an ‘acceptable’ crack growth of 3×10-12 mm/s was used).
The approach being taken in Canada is to argue that the necessary environmental conditions for SCC will not persist for a sufficient period of time, if at all, to cause failure of the canister by SCC. Although this approach can be used in conjunction with either the concept of threshold conditions for SCC or for the limited-propagation argument, to date, the only prediction made in Canada has been based on the limited-propagation approach. Maximum crack depths between 11 µm and 11 mm, for anodic to cathodic surface area ratios of 1 (general corrosion) and 1,000, respectively, have been predicted [239,308,309].
A.5.2.4 Lifetime predictions of Cu canisters
The predicted Cu canister lifetimes from various international corrosion programmes, based on several approaches (combined mass-balance/mass-transport approach, steady-state mass-transport reaction model), are summarised in TABLE 5-2.
218 TABLE 5-2. Comparison of predictions of long-term corrosion behaviour and canister lifetime from various international programmes [60].
Country General Localised MIC(1) SCC(2) Predicted Ref. corrosion corrosion lifetime
Sweden/ 0.05 mm in 106 yrs 0.05 mm in 106 yrs - - > 106 yrs [62] Finland(a) (realistic) (realistic) 4 mm in 106 yrs 18 mm in 106 yrs (conservative) (conservative)
Sweden/ 0.35 mm in 106 yrs 0.35 mm in 106 yrs SRB(3) assumed to maximum possible nitrite > 106 yrs [283], (a) 2- - Finland (realistic) reduce SO4 to HS concentration below thres- [284], 1.4 mm in 106 yrs hold for SCC [310] (condervative)
Sweden/ 0.33 mm in 106 yrs 0.33 mm in 106 yrs SRB(3) assumed to SCC does not occur > 106 yrs [60] (a) 2- - Finland (realistic) reduce SO4 to HS based on threshold 1.3 mm in 106 yrs in tunnel and potential and concen- (condervative) groundwater only trations of SCC agents, and because creep is faster than SCC
Canada(b) 0.011 mm in 106 yrs 6 mm in 106 yrs limited impact: maxi- SCC not included > 106 yrs [301] mum additional wall because of limited period loss of 1 mm in 106 of stress, absence of SCC yrs agents, general lack of oxidant, and inhibitive effects of Cl- Japan 9-13 mm in 103 yrs, 18-26 mm in 103 yrs SRB(3) assumed to none given [311] 2- depending on repo- based on pitting reduce all SO4 to maximum concentrations sitory design factor of 3 HS- of ammonia, nitrite, and 2 mm in 103 yrs acetate less than thres- based on extreme- hold concentration value analysis
(a) Reference canister wall thickness of 50 mm. (b) Reference canister wall thickness of 25 mm. (1) MIC : Microboilogical Influenced Corrosion. (2) SCC : Stress Corrosion Cracking. (3) SRB : Sulphate Reducing Bacteria.
219 APPENDIX A - Contribution (%) of nuclear power plants to the world's electricity production [3]
(the figures are valid up to April 2001)
220 APPENDIX B - The different disposal concepts envisaged in various countries operating nuclear power plants (e.g. Belgium, Finland, France, Germany, Spain, Sweden, UK)
B.1 Current disposal concept in Belgium
The current Belgian disposal concept for HLW/spent fuel consists of a horizontal network of concrete galleries constructed in a thick plastic clay formation. The currently considered clay formation is located underneath the SCK•CEN site at Mol, viz. the Boom clay formation, which is situated at a depth between 180 m and 280 m. The repository would be constructed at mid-height of the host rock formation, i.e. about 230 m depth [44,312].
A schematic view of the current Belgian repository design is shown in FIGURE B-1. The 800 m long parallel disposal galleries, with an internal diameter of about 2 m, would be constructed perpendicular to larger transport galleries, with an internal diameter of about 3.5 m. The disposal galleries would be partitioned into four segments each 200 m in length. The parallel transport galleries would be accessible via two 6 m diameter vertical access shafts, which FIGURE B-1. Schematic view of the proposed Belgian repository design. Adapted from would be connected by means of [313]. a 400 m long connecting gallery. The spacing between the galleries would be sufficiently large (e.g. 40 m) to avoid (i) excessive temperatures in the zone surrounding the disposal galleries and (ii) significant warming of the overlying aquifer [44,312].
FIGURE B-2 represents a detailed view of the cross section of a disposal gallery (indicating the three different metallic barriers) for the final storage of HLW and spent fuel in the present Belgian repository concept. A long metallic tube (the disposal tube) would be placed in the center of the disposal galleries and they would be backfilled with either Boom Clay or a bentonite-type material. Each of the 0.43 m outer diameter metallic vessels (the canister), into which the radioactive waste is poured, would be placed in a 0.02 to 0.03 m thick sealed metallic container (the overpack). These overpacks would then be pushed into the 0.01 m thick disposal tube by means of remotely operated robots. The disposal tube allows for placement of the 1 m thick backfill layer prior to placement of the HLW glass canisters [44,312].
221 The overpack has a design lifetime of approximately 500 years (for vitrified HLW) to 2000 years (for spent fuel) during which it should remain intact. The material presently considered for construction of the overpack is the stainless steel AISI 316L hMo [44, 312]. The HLW canisters would be placed in the disposal galleries after saturation of the backfill, and at least 60 years after discharging the related spent fuel from the nuclear reactor [44].
FIGURE B-2. Schematic representation of the cross section of a disposal gallery [313].
222 B.2 Current disposal concept in France
The host rock formation currently envisaged in France is the Callovo-Oxfordian clay formation, which has a thickness of about 135 m (extending from 417 to 552 m depth). The repository would be constructed around the middle of the clay formation. A single level concept has been chosen to allow a significant thickness of clay to retain radionuclides on both sides of the disposal structure [54].
As repository design, for the disposal of both HLW and spent fuel, a concept with horizontal galleries is preferred to a concept with vertical boreholes. A network of horizontal galleries is favoured because it (i) provides a better technical and economical performance, (ii) minimises the volume of clay that has to be excavated, and (iii) takes up a smaller underground area for the construction of the disposal site [314].
B.2.1 Concept for the disposal of vitrified HLW
The main features of the disposal concept for HLW are summarised in TABLE B-1.
TABLE B-1. ANDRA preliminary design for vitrified HLW disposal [314].
Repository design for vitrified Description HLW
Horizontal disposal galleries • the horizontal disposal concept is schematically represented in FIGURE B-3. The waste packages are placed in the centre of the disposal galleries, which are connected to larger tunnels (for transportation and handling of the waste packages). • the technical aspects of the repository concept for the disposal of HLW (horizontal disposal galleries) are given in TABLE B-2. • two options are considered: - with a clay-based engineered barrier (Ø ≈ 2.5m) (see FIGURE B-4). - without a clay-based engineered barrier (Ø ≈ 0.7m) (see FIGURE B-5). • a clay-based plug ( ~3m thick) is foreseen to hydraulically and physico- chemically seal each disposal tunnel. • a steel liner is foreseen to facilitate the emplacement and retrieval of the waste packages.
Thermal loading • the target temperature on the outside of the waste packages should be kept below 100°C (~90°C) to allow for the uncertainties on the thermal data and its calculation.
Overpack • the primary waste package, fabricated from AISI 309S stainless steel, will be surrounded by a 55 mm thick carbon steel overpack • the metallic overpack is designed to withstand for several of hundreds of years (~1,000 year). This period corresponds to the thermal phase, during which the activity of the fission products and actinides decrease rapidly. The metallic container should also remain intact during this period to enable the retrievability of the waste packages during the exploitation period should this turn out to be necessary.
223
FIGURE B-3. Schematic representation of the horizontal gallery concept [315].
TABLE B-2. Technical aspects of the repository concept for the disposal of HLW [315].
Type of waste C0 C1 C3
Cooling period [year] (1) 60 60 Thermal power at disposal [W] +200 ~500 ~600 No. of waste packages per tunnel 15 4 2 Distance between tunnels [m] 12 12-15 12-15 Density of waste packages [pack./ha] 400 150 100
FIGURE B-4. Schematic representation of a horizontal tunnel (with a clay-based engineered barrier) for the disposal of vitrified HLW [314].
224 The disposal tunnels have a diameter of ~2.5 m. The gap between the clay formation and the metallic lining is filled with an expanding clay-based buffer material. This option can meet, to a certain extent, the uncertainties of the current knowledge of the disturbed clay behaviour (caused by the engineering operations necessary to construct the underground disposal tunnels) by increasing the hydraulic and physico-chemical robustness of the disposal site (e.g. the clay-based buffer creates an environment around the waste packages which can ensure confinement of the radionuclides, in case of a premature failure of the metallic container) [314].
FIGURE B-5. Schematic representation of a horizontal tunnel (without a clay-based engineered barrier) for the disposal of vitrified HLW [314].
In the option without a clay-based engineered barrier [314], the disposal tunnels are excavated to a diameter of ~0.7 m, which corresponds to the size of the waste packages increased with • the thickness of the metallic lining, • the gap between the metallic lining and the clay front, and • the gap between the metallic lining and the waste package. This option has the advantages of simplifying the concept and minimising the amounts of clay that have to be excavated (thereby improving the compactness of the underground disposal site) [314].
B.2.2 Concept for the disposal of spent fuel
The main features of the disposal concept for spent fuel are summarised in TABLE B-3.
225 TABLE B-3. ANDRA preliminary design for spent fuel disposal [314].
Repository design for spent fuel Description
Horizontal disposal galleries • the same concept is chosen as for the disposal of vitrified HLW, viz. the concept with horizontal disposal tunnels. Only the option with a clay-based engineerd barrier is considered. The concept is schematically represented in FIGURE B-6. • a clay-based plug ( ~3m thick) is foreseen to hydraulically and physico- chemically seal each disposal tunnel. • a steel liner is foreseen to facilitate the emplacement and retrieval of the waste packages.
Thermal loading • the target temperature on the outside of the waste packages should be kept below 100°C.
Overpack • 4 UOX assemblies or 1 MOX assembly per canister. The container design is shown in FIGURE B-7. • the UOX or MOX assemblies are placed in a primary stainless steel canister, which will be surrounded by a 130 mm thick carbon steel overpack. • the metallic overpack is designed to remain watertight for ~10,000 years (the thermal phase lasts much longer than for vitrified HLW because of the presence of actinides, 241Am in particular).
FIGURE B-6. Schematic representation of a horizontal tunnel (with a clay-based engineered barrier) for the disposal of spent fuel [314].
226
FIGURE B-7. Schematic representation of the metallic container developed for the disposal of spent fuel [314].
227 B.3 Current disposal concept in Sweden/Finland
The repository in Sweden and Finland is based on the Swedish KBS-3 design (see FIGURE B-8). After encapsulating the fuel in a canister, the canisters are emplaced individually in vertical boreholes drilled in the floors of deposition tunnels feeding off central tunnels at about 500 m depth in the crystalline, granitic bedrock. The space between the canisters and the wall of the borehole will be filled with precompacted blocks of bentonite. When the bentonite absorbs water from the surrounding bedrock it will exert an intense swelling pressure and completely fill all void space in the near vicinity of the canister with bentonite clay. The tunnels and shafts will eventually be backfilled with a mixture of crushed rock and bentonite, and sealing plugs will be emplaced to block specific FIGURE B-8. Schematic drawing of the Swedish/Finnish deep repository con- transport pathways for groundwater. cept in accordance with the KBS-3 design. Source: Ref. [35].
The conceptual canister design adopted is a copper canister with a steel insert (see FIGURE B-9): the spent nuclear fuel is planned to be encapsulated in spheriodal graphite cast iron canisters that have an outer thick shield made out of copper. The copper thickness will be about 50 mm and the steel thickness also about 50 mm. The copper is believed to provide a very good corrosion resistance to the geochemical environment foreseen in a deep repository in Sweden. The steel insert is designed to withstand the normal mechanical loads that will prevail on the canister in the repository such as hydrostatic pressure from groundwater (i.e. 7 MPa) and the pressure from swelling of the bentonite (i.e. 7 MPa). The combined thickness of steel and copper should be enough to prevent any significant radiolysis of water outside the canister after deposition in wet bentonite clay. The canister design aims at providing a FIGURE B-9. Exploded view of the copper corrosion lifetime of at least 100,000 years canister with steel insert. in the repository. The maximum allowed surface temperature has been set to 100°C
228 and the maximum allowed surface dose rate to 1 Gy/h. The fuel in the canister should also remain subcritical even if water enters the canister. The total weight of a canister with fuel (12 BWR fuel assemblies or 4 PWR assemblies) will be about 25 tonnes. In total some 4,000 canisters will be required for the spent fuel arising from the Swedish reactors up to 2010.
According to the current concept, the compacted bentonite blocks placed in the deposition hole will have an initial degree of saturation of 85%. Following closure, the repository host rock and excavation damaged zone (EDZ) are expected to return to saturated conditions over a few years. The full saturation and subsequent homogenisation of bentonite in the deposition hole are expected to occur between 6 and 35 years, but it could also extend over a longer period, even some hundreds of years depending, for example, on the access of water to the deposition hole. At full saturation, the target bentonite density is 2 Mg/m³ and the hydraulic conductivity is very low, less than 10-12 m/s.
Calculations, assuming dry bentonite and an initial temperature of the bedrock of 11°C, indicate a maximum canister surface temperature of 90°C. Saturation of the bentonite is likely to reduce this maximum canister temperature by up to 15°C. The temperature conditions in the repository system will approach the natural state after approximately 10,000 years when the decay heat has declined to less than one percent of its original value.
229 B.4 Current disposal concept in Germany
More detailed information can be found in References [40,319].
Amongst the various concepts studied over the past thirty years, two approaches that are quite distinct from each other have been developed in Germany taking the existing technical variants and the waste type ratios (HLW or spent fuel) into account:
• the HLW, remaining from the reprocessing of spent fuel, will be vitrified in Cr-Ni steel canisters, which, if necessary, would be surrounded by a corrosion-resistant container (overpack). The steel canisters will be disposed of in about 300 to 600 m deep and 0.6 m wide boreholes reaching vertically down from a disposal level at a depth of some 880 metres below ground (FIGURE B- 10). About 200 canisters will be lowered into each borehole. In order to transfer the weight load of the canister stack to the surrounding rock mass, the remaining annulus around the canisters is to be backfilled with crushed salt. At the top of the borehole, a seal consisting of crushed salt FIGURE B-10. The borehole concept for will be placed [40]. the emplacement of vitrified HLW [40]. • for the direct disposal of spent fuel, large self-shielding gas-tight POLLUX carbon steel casks have been developed, which will be emplaced in horizontal drifts (see FIGURE B-11) of about 200 m long, 4.5 m wide, and 3.5 m high. Each drift wil host ~20 casks one behind the other. The different emplacement drifts will be separated by dams impervious to water. Following the emplacement of a POLLUX cask, the remaining voids in the drift will be backfilled with crushed salt. FIGURE B-12 shows the basic design of the double POLLUX FIGURE B-11. The drift concept for the direct disposal cask. The POLLUX disposal disposal of spent fuel (POLLUX casks) [40]. cask consists of a 160 mm thick mechanically stable inner cask, made of fine-grained steel (15 MnNi 6.3), which, if necessary, can be entirely surrounded by a 3 to 4 mm corrosion protection layer (Hastelloy C-4). This protection layer is applied by surface welding. The inner cask is shielded against gamma and neutron radiation by a 265 mm surrounding ductile cast iron (GGG 40.3) container and a polyethylene filler on an AlMgSi 0.5 carrier material. Each of these casks weigh about 65 t and are to be used for interim storage as well as for final disposal [40,41,319].
230
FIGURE B-12. Schematic representation of the double ‘POLLUX’ disposal cask [319].
In both concepts, the creeping of the surrounding rock will lead to compaction of the initially loose material thereby sealing the waste from the biosphere. An additional sealing of the waste by buffer materials is not considered necessary [40].
The repository design that is based on the combined borehole and drift emplacement of both HLW-canisters and POLLUX casks is, nowadays, considered the most promising [40]. A typical layout of a HLW repository is shown in FIGURE B-13.
231 FIGURE B-13. Typical layout of a high-level waste repository in a salt formation in Germany [40].
The underground mine with its disposal boreholes and drifts keeps a safety distance of 200 metres from the edges of the salt dome. Shaft 1 will be used as air intake shaft and shaft 2 as exhaust shaft and for transportation of waste packages. The layout of the disposal fields and drifts is mainly determined by the actual geology at the disposal level and the maximum allowable temperatures of 200°C at the interface waste canisters/host rock [40].
Depending on the disposal and the local conditions in the geologic formations close to the disposal sites, the intrusion of brine to the radioactive waste must be expected to occur only in a few hundred or even thousand years, or perhaps never. Assuming an ingress of brine, the contact of waste with brine may continue to exist forever (no outlets) or may be interrupted after a very short period of time (brine being forced out by ‘convergence’ of the salt). In the most probable scenario, cavities remaining in the emplacement location will soon be closed automaticaly by convergence. Although the most probable case will be that no brine at all will penetrate to the waste, it is not possible to exclude with absolute certainty the development of a pathway that would allow the brine to reach the waste packages (e.g. as a result of an ingress from a brine inclusion not yet discovered), whereupon this brine would first corrode the waste packages and then become contaminated with radionuclides and leave the emplacement site.
To demonstrate the contribution to safety made by waste packages, we are forced to make a number of conservative assumptions. This is because the precise development of the safety-related properties of the repository cannot be predicted in detail, and hence, the worst case must always be assumed. For instance, it must be assumed that based on the corrosion rates measured in the laboratory, the container will have corroded through at a certain point in time, even if it is doubtful whether there is enough
232 brine near the emplacement site to make a thick-walled container corrode. Also, in determining the contribution to safety made by the waste form, we must assume permanent contact with the most highly corrosive types of brine, even though studies show that safety-related contacts between brines and HLW will exist only for short periods of time.
233 B.5 Current disposal concept in Spain
Both granite and clay are the two main options that are being studied in Spain as candidate geological formation to host an underground repository for the final disposal of HLW/spent fuel.
Cylindrical carbon steel canisters (4.54 m long, 0.10 m thick, and an external diameter of 0.90 m), filled with glass beads, are used to encapsulate the spent fuel elements. These canisters are emplaced horizontally in drifts and surrounded by a clay barrier made up of highly compacted bentonite blocks. The canisters are separated by a distance of 2 m and 2.5 m in the granite and clay repository concept, respectively [321].
The repository concept, developed by ENRESA, for disposal both in granite and clay are rather similar:
• repository concept in granite. The repository design for disposal in a granitic host rock formation is schematically presented in FIGURE B-14. The repository
FIGURE B-14. Repository design for granite [320].
consists of two disposal areas, located at a depth of 500 m and accessible by means of two access tunnels. Three access shafts and a ramp connect the surface with these galleries and the rest of the underground facilities. The disposal galleries (500 m long) are constructed perpendicular to these access tunnels, with an in-between distance of 35 m. The disposal galleries have a diameter of 2.40 m. A perforated steel liner with a thickness of 0.02 m, used to guide the emplacement of the canisters, will be placed in the centre of the disposal galleries. The gap of 0.75 m between the canisters and the gallery walls is filled with highly compacted bentonite blocks (1.65 g/cm³ dry density, buffer material). The access tunnels, access shafts, and ramp are filled with a mixture of bentonite and sand (80%-20%, backfill material). The disposal galleries and access tunnels are sealed with 6 metres of bentonite blocks with a higher compactation (1.83 g/cm³ dry density) and a plug of concrete with a length of 3 m. Additional seals are located in several
234 places along vertical tunnels and the ramp [66,321]. A scheme of a disposal gallery is shown in FIGURE B-15.
Liner 2 m
2.40 m
Access tunnel Concrete plug Bentonite plug Canisters
FIGURE B-15. Schematic view of the disposal gallery [321].
• repository concept in clay. The repository design for disposal in an argillaceous host rock formation is schematically presented in FIGURE B-16. Accesses to the disposal area (three shafts and a ramp), the type and geometry of the buffer
FIGURE B-16. Repository design for clay [320].
material, and seals are similar to the granite repository concept. The repository in clay is to be constructed at a depth of 250 m. The canisters, which are identical to the granite option, are disposed of in horizontal galleries with a length of 500 m and are seperated by a distance of 2.5 m. The distance between the galleries is 50 m. The diameter of the disposal galleries is larger than in the granite repository concept, viz. 3 m, in order to accomodate the sustaining structure made of concrete voussoirs with a thickness of 30 cm. FIGURE B-17 shows a schematic representation of a cross section of such a concrete sustaining structure (with bentonite blocks) in a clay gallery. Backfilling of the access tunnels is made with
235 the clay extracted during excavation of the facilities, compacted to a density of 1.6 g/cm³ [58,321].
Bentonite blocks Concrete voussoirs
Liner
Clay rock
Canister
FIGURE B-17. Schematic view of a cross section of a concrete sustaining structure (with bentonite blocks) in a clay gallery [321].
236 B.6 Current disposal concept in the United Kingdom
At present, a government policy to manage the High Level Waste (HLW) in the UK for the long-term does not exist. The vitrified HLW, produced from reprocessing, in the UK is currently being stored in stainless steel tanks, which are continuously cooled, next to the treatment plant at Sellafield. This HLW shall be stored above ground for at least 50 years before disposal. This will allow the considerable heat to dissipate and the radioactivity to reduce to a level at which it could be safely disposed in a similar manner to Intermediate Level Waste (ILW) [322,323].
NIREX has developed a repository concept, based on the phased deep geological disposal, with the aim to form a containment system for emplaced solid ILW and certain LLW. This concept, illustrated in FIGURE B-18, makes use of multiple barriers (both engineered and natural) [26]: • physical containment would be achieved by immobilisation and packaging of wastes in steel or concrete containers; • geological isolation by emplacement of the waste packages in vaults excavated deep underground within a suitable geological environment; • chemical conditioning by backfilling the vaults with a cement based material (the Nirex Reference Vault Backfill - NRVB) at a time determined by future generations; • geological containment achieved through final sealing of the repository at a time determined by future generations.
FIGURE B-18. Illustration of the NIREX Multi-Barrier Repository Con- cept [26].
The key features of the reference design concept, developed for a Sellafield repository, are shown in FIGURE B-19. In the deep repository concept, developed by NIREX, for the disposal of intermediate level and certain low level wastes, the wastes (packaged in steel or concrete containers, usually with a cement grout) are placed in large purpose- built vaults. These vaults are excavated at depth, currently expected to be between 300 and 1000 m below ground level, in a stable geologic environment. The space between
237 the waste packages in the repository vaults would be backfilled with a cement-based material (the NIREX Reference Vault Backfill –NRVB), which is designed to create uniform, alkaline chemical conditions. Prior to backfilling, which will only take place after a period that can last up to several hundred years, the emplacement operations are readily reversible and the waste packages could be retrieved if required [26]. The underground vaults would be accessed from the surface by means of either shafts or drifts (inclined tunnels). It is envisaged that three access ways would be needed to connect the surface facilities to the underground areas [26]: - two shafts would be used for ventilation and access for staff and materials; - the waste packages would be transported by a rack and pinion rail system through a drift (this system is similar to those used throughout the world for freight and passenger transport in hilly or mountainous terrain). In case of an emergency, the drift cold also provide a second means of egress for construction workers.
FIGURE B-19. Schematic view of the NIREX repository design [26].
238
PART B
EVALUATION, DISCUSSION AND NEEDS FOR FUTURE R&D
B.1 COMMON SCIENTIFIC APPROACH
A similar scientific approach was adopted to study the long-term corrosion behaviour of candidate container materials for the disposal of long-lived HLW and spent fuel in the various European countries. First, a screening of a wide range of materials was performed (based on literature studies, preliminary laboratory experiments, etc.) in order to decrease the number of potential candidate container materials. Then, more in-depth investigations were performed on a limited selection of candidate container materials (e.g. long-term immersion experiments, in situ experiments, electrochemical experiments, SCC-tests, etc.). These detailed experiments were aimed at determining the influence of important parameters on the long-term corrosion behaviour of the selected materials in environments relevant to disposal. The design of the experiments chosen depended on the class of materials that were being tested (i.e. actively or passively corroding metals), since the type of data that is required to predict the long-term corrosion behaviour depends on the choice of container material: for actively corroding materials (e.g. carbon steel), it has to be proven that the material exhibits an acceptable uniform corrosion rate, while for passive materials, the performance of the container is strongly influenced by its resistance to localised corrosion phenomena (pitting corrosion, crevice corrosion, SCC, etc.). Simultaneously, attempts have been undertaken to gather useful information from alternative research areas such as demonstration tests on a 1-1 scale, natural analogues, etc. Although modelling experience is more advanced in some countries than in others, the big challenge for all countries involved, in the future, remains to be able to predict the corrosion behaviour over a period of 1,000 years or more. This problem does not pose itself in any other industrial application.
240 B.2 POSSIBLE MODES OF CORROSION AFFECTING THE CONTAINER
Under disposal conditions, the overpack could be prone to one of the following modes of corrosion:
- general corrosion – aerated period. - general corrosion – anoxic period. - crevice corrosion. - pitting corrosion. - stress corrosion cracking (SCC). - intergranular corrosion – grain boundary attack. - galvanic corrosion. - microbially influenced corrosion (MIC). - hydrogen embrittlement. - radiation-induced corrosion. - stray current corrosion. - corrosion due to magnetic fields (electrical currents induced by magnetic fields). - atmospheric corrosion. - gaseous oxidation.
The two last of these (atmospheric corrosion and gaseous oxidation) only apply to the period of interim storage and to the disposal period before resaturation of the backfill.
TABLE 1 summarizes the factors affecting some of the above-listed types of corrosion.
TABLE 1. Factors affecting specific types of corrosion.
Type of corrosion Critical factors
Crevice corrosion geometry of crevice, size of cathodic area
residual stresses, applied load, size of surface defects, Stress corrosion cracking presence of stress concentrators, mechanical properties of the material
size of surface defects, presence of stress concentrators, Hydrogen embrittlement mechanical properties of the material, sub-surface defects
material combinations, relative areas, differential aeration Galvanic corrosion cells
Atmospheric corrosion relative humidity, concentration of atmospheric pollutants, (internal and external) air flow rates
241 Since the various national research programmes were not performed in parallel, there is a need to establish, for each national research programme separately, which corrosion modes are of primary importance for the service life of the container, based on the disposal concept, the container design, the type of host rock formation (i.e. the environmental composition) and the environmental conditions. This should be followed by an examination of whether the effect of each of these corrosion modes on the service life of the container has already been fully characterised or whether additional research is required.
One possible corrosion degradation mode that might increase in significance in the near future is microbially influenced corrosion (MIC). Metal corrosion can be accelerated by the action of micro-organisms by a number of different processes. They can produce a corrosive species as a by-product of their metabolic cycle, they may enhance the electron transfers involved in the electrochemical reactions, or they may be able to ionise the metal surface itself [324]. Three basic requirements have to be present simultaneously for MIC to occur: a source of organisms that is viable at the appropriate temperature, liquid water, and nutrients. Literature studies [325,326] and sampling studies [327,328] have concluded that microbial contamination of a geological repository, during construction and operation, is inevitable. There should, however, be an investigation of whether micro-organisms can maintain their activity in the disposal environment. In this respect, the risk of MIC affecting the integrity of the container strongly depends on the type of host rock formation and the type of candidate container material: • in rock salt, it is assumed that MIC will not pose a great threat because of the high salinity, high temperature (>200°C), and presence of radiation. • in clay and granite (with a bentonite type of backfill material), conditions in the near-field could turn out to be insufficiently extreme to ensure complete sterilisation. West et al. [329] concluded that pressure and temperature considerations alone are insufficiently intense to limit or control microbial activity in a geological repository.
Furthermore, the possibility of MIC occurring in a geological repository is strengthened by some observations in which certain micro-organisms were found to survive extreme conditions: • most sulphate reducing bacteria (SRB) grow ‘best’ at temperatures of between 10°C and 45°C according to the strain and its normal habitat. However, many strains exist that will grow up to at least 75°C and some SRB have even been found to survive temperatures at least up to the boiling point of water. Therefore, SRB could survive the extreme conditions (high T) occurring during the aerobic phase and become active during the subsequent anaerobic phase of disposal, being no longer constrained by lack of sulphate availability (sulphate will be formed during the aerobic phase by oxidation of the sulphurous compounds of the environment). • one species of sulphur oxidising bacteria (SOB) has been stated to remain active at pH 0.7, corresponding to more than 5% sulphuric acid [330].
To date, MIC in an underground repository environment has not been firmly identified from studies but the potential for bacterial activity in the oxic and/or anoxic phases of the disposal have been established. It will therefore be necessary to check the performance of candidate container materials for resistance to MIC through literature searches, or potentially by experimental work. Already, some preliminary studies
242 have started (and are still ongoing) to determine, for example, whether SRBs can maintain their activity in compacted bentonite [311].
243 B.3 PARAMETERS AFFECTING THE CORROSION BEHAVIOUR OF CANDIDATE CONTAINER MATERIALS
Several environmental, constructional (related to the repository concept and the container design) and metallurgical parameters can influence the lifetime of the container. These effects include:
- effect of corrosion products. - effect of pressure due to H2-production. - effect of welding on corrosion. - effect of concrete lining. - effect of γ-radiation. - effect of fabrication processes (e.g. thermal treatment, metal surface condition, residual stresses, welding, etc.). - effect of long-term metallurgical modifications (e.g. dealloying, segregation, creep, growth of grains, etc.).
TABLE 2 summarizes the various parameters that can affect corrosion.
TABLE 2. General factors affecting corrosion.
composition, microstructure, heat treatment, inclusions (size, composition, orientation), grain size, grain boundary Material composition, grain orientation, amount of hot or cold work, sensitisation, welding
Surface surface finish, surface treatment, electrochemical potential
temperature, composition, oxygen concentration, pH, concentration of anions, concentration of adsorbing organic Environment molecules, radiation effects, presence of microbial films, mass transfer limitations on oxygen (diffusion coefficients, flow patterns)
One of the more important parameters for public acceptance of the disposal concept is the influence of γ-radiation on the corrosion behaviour and corrosion mechanisms of the candidate container materials. To date, only a limited amount of R&D data has become available, partly because of the technical difficulties (e.g. safety issues) involved in performing experiments under radiation. Only in rock salt (German disposal system) and in granite (Swedish and Finnish disposal system), has a considerable amount of information become available: • the effect of γ-radiation (dose rate: 1-10³ Gy/h) on the uniform corrosion rate and the depth of localised attack (pitting and crevice corrosion) of the fine-grained carbon steel TStE 355, the cast steel GS 16Mn15, the nodular cast iron GGG 40.3, the Ni-resists D2 and D4, the Si-cast iron, the Ni-alloy Hastelloy C-4 and the Ti-alloy Ti 99.8-Pd in brines is summarized in part A (section A.4.1.1.2). • the effect of γ-radiation (dose rate: 10-10³ Gy/h) on the uniform corrosion rate of copper in air and water (1 mol/L Cl-) is summarized in part A (section A.4.3.1.3.5).
244 However, from the available information, most data were obtained in high dose radiation fields (≥ 1 × 10³ Gy/h). Except for rock salt environments, in which experiments have also been performed at lower dose rates (1 – 10³ Gy/h). However, at low dose rates (< 1 × 10³ Gy/h), a lot of questions still remain unanswered that can greatly influence the corrosion behaviour of the candidate container materials, such as:
- from the literature, it is evident that at high γ-radiation dose rates, the corrosion potential ECORR shifts upwards. But how large is this positive potential shift at low levels? - at low γ -radiation dose rates, it is ‘assumed’ that the values for ENP and EPP are stable. But this assumption has never been proven (or checked). - what is the lowest threshold value below which radiation will no longer influence corrosion (in this case, only threshold values are given for salt formations for carbon steel (see part A, section A.4.1.1.2.2.4) and Hastelloy C-4 (see part A, section A.4.1.1.2.3))? - what is the influence of long-term exposure to low dose rates?
Amongst researchers, a lot of controversy still exists on how γ-radiation can affect corrosion:
• some researchers [18,108,109,331] argue that only the corrosion potential is influenced by the presence of γ-radiation. Some data (apart from the data listed in part A) that supports this assumption is summarised in TABLE 3.
TABLE 3. Influence of γ -radiation on ECORR and ENP of candidate container materials.
Influence on Ecorr
- Material γ-dose rate [Cl ] pH Ecorr shift Ref. (Gy/h) (mg/L) (mV)
AISI 304L, AISI 316L 3.3 × 104 7 7.6 + 150-250 [108] AISI 304L 2 × 103 1,000 - + 300-400 [331] AISI 304L 2 × 103 300 - + 200-300 [109] AISI 304L 6.1 × 103 7 7.6 + 200 (initial)* [18] + 400 (after 1,000 years)*
* Mathematical modelling
Influence on Enp
- Material γ-dose rate [Cl ] pH Enp - Ecorr Ref. (Gy/h) (mg/L) (mV)
AISI 316L 3.3 × 104 650 - 429 [108] AISI 316L non-irradiated 650 - 197 [108]
245 The danger that can arise from an increase in the metal rest potential, resulting from irradiation, is that it can lead to a premature initiation of localised corrosion. Considering the importance of determining the accurate value of ECORR in order to evaluate the corrosion behaviour of stainless steel, the influence of radiation on this parameter should be fully understood. Other scientists claim that the pit nucleation potential can also be influenced by radiation.
One study [265], in which carbon steel and 316L stainless steel samples were exposed to a simulated cement porewater (0.1M KOH aqueous solution saturated with Ca(OH)2) deaerated with argon), concluded that radiolytic processes at the dose rate of 16 Gy/h would be insufficient to increase the corrosion potential of 316L stainless steel into the range where localised corrosion might occur. The only significant changes in potentials were those observed on oxygenating the system, but it was believed that this would probably also have taken place in the absence of radiation. The results are shown in FIGURE 1.
FIGURE 1. The effect of γ-radiation on the corrosion potential of carbon steel and 316L stainless - steel in 0.1M KOH saturated with Ca(OH)2 with and without oxygen and added Cl (gamma dose rate = 16 Gy/h) [265].
• it has also been found [108,109,332] that austenitic stainless steels become more susceptible to pitting under γ-radiation. Some data (apart from the data listed in part A) that illustrates the influence of γ-radiation on the pitting susceptibility of candidate container materials is shown in TABLE 4.
246 TABLE 4. Influence of γ-radiation on the pitting susceptibility of candidate container materials.
- Material γ-dose rate Media [Cl ] pH Max. dpit Ref. (Gy/h) (mg/L) (µm)
Low-alloyed steel non-irradiated MX-80 clay 70 8.0-8.2 160 [265] Low-alloyed steel 3.9 × 103 MX-80 clay 70 8.0-8.2 480 [265]
dpit : pit depth.
• γ -radiation can also contribute to a change in the chemical composition of the near-field environment: it is generally accepted that the imposition of a γ -radiation field generates a more oxidising nature of the aqueous environment due to the • radiolytic decomposition of water into e.g. H2O2 and OH . In the presence of a catalyst, such as iron, H2O2 will decompose into a.o. O2. The aqueous near-field environment could therefore act as a continuous source of oxygen, thereby extending the duration of the oxic period.
247 B.4 INTERNAL CORROSION
Besides the well-known corrosion problems affecting the exterior of waste containers, described in section B.2, there are some other possible causes of corrosion that are specific to the internal environment. For example, in the presence of a radiation field and after the underground repository has been sealed, the gaseous atmosphere inside containers may contain air and water vapour. It is known that irradiation of moist air produces nitric acid (HNO3), which is a potential corrodent of, for example, austenitic stainless steels [333,334]. The effect of temperature and HNO3 concentration on austenitic stainless steels has already extensively been investigated [335]. Increasing either or both these parameters raises the corrosion rate. One of the most prevalent problems with HNO3 is selective corrosion associated with chromium carbides precipitated around grain boundaries in the weld HAZ (Heat Affected Zone). Since welding will be used to seal the container, this could present a realistic problem, unless steps are taken to prevent sensitisation in the weld area. Also, the corrosion of austenitic stainless steels in HNO3 is accompanied by the formation of 6+ hexavalent chromium (Cr ), a species that increases the corrosivity of HNO3 solutions.
Some of the unanswered questions concerning this issue are listed below: • can nitric acid pose a serious threat to the integrity of the metallic container ? • have modelling efforts been undertaken in the past to calculate the maximum amount of nitric acid that can be produced inside an overpack ? • does the possibility exist of HNO3 being concentrated at local sites, such as crevices, even if only small amounts of nitric acid are produced (which seem harmless at first sight), thereby causing considerable damage such as SCC ? • can interactions between the waste form (vitrified HLW/spent fuel) and the metallic container influence its integrity (e.g. could fission products affect the corrosion of the container) ?
It is possible that the risk of HNO3-induced corrosion could be overcome by filling the space between the waste canister and the metallic container with inert gas (e.g. helium) or drying the air in the underground disposal site. However, the technical feasibility of these possible solutions should be explored in consultation with repository engineers.
248 B.5 INFORMATION FROM ARCHAEOLOGICAL ANALOGUES
In HLW/spent fuel disposal, the big challenge will be to predict the corrosion behaviour of candidate container materials (carbon steel, stainless steel, copper, etc.) for periods well exceeding the normal service life of constructional materials (30- 50 years). The prediction of this corrosion behaviour over periods of up to 1,000 years, or even longer, based only on relatively short-term experimental observations up to a maximum of 10 years (in situ experiments, immersion tests, etc.) and theoretical calculations, may be uncertain. Because it is difficult to simulate in a laboratory the time scales involved in the repositories, studying similar natural systems (i.e. natural analogues) that have been exposed to similar conditions over a long period of time (several hundreds to thousands of years) can provide useful information about corrosion kinetics and corrosion products, and can contribute to an improved confidence in predicting the future corrosion behaviour of waste container materials.
Some of the interesting information gathered from natural analogue studies, which is applicable to HLW/spent fuel disposal, is summarised below:
• natural analogues studies have shown the importance of oxygen availability on the corrosion rate. A hoard of nearly one million iron nails (approximately 10 to 40 cm long) was buried in a specially dug pit 5 m deep and covered with 3 m of compacted earth in Inchtuthil, in Pertshire (Scotland), in 87 AD [324,336-338]. The nails at the centre of the hoard showed minimal corrosion. It was found that the good state of preservation of these nails was attributed to the formation of an impermeable crust of corroded iron around the outside of the hoard, especially at the top. This solid iron oxide crust ensured virtually anoxic conditions inside the hoard, so that the amount of corrosion became progressively less as the centre of the hoard was approached. Corrosion of these internal nails was further inhibited by the presence of a coat of scale (produced by the high temperature production process) on most of them. Many nails were found in which the portion covered by a coherent coating of the scale was uncorroded, while the rest of the nail was badly attacked [324,337].
• the data from the archaeological analogues compare well with the data from the short-term laboratory experiments under anaerobic conditions. A detailed study of the long-term corrosion of iron was produced by Johnson and Francis [339] who documented information from over fourty iron archaeological artefacts. The corrosion rates deduced from iron archaeological artefacts are shown in FIGURE 2. Most of the artefacts came from a wide variety of environments, some of which were arid, and the range of estimated corrosion rates was still more than two orders of magnitude (10² - 104 µm per thousand years), with an average value (assuming a logarithmic distribution) of around 0.9 mm per thousand years. This is similar to the approximate value of 1 mm per thousand years obtained from more recent work [340], commissioned by ENRESA, the Spanish waste management agency. In this work, a wide range of modern chemical analytical techniques was used to study the corrosion of ancient iron artefacts. The results of an experimental programme, carried out by Romanoff of the US National Bureau of Standards and exposing different mild steels to a wide range of soils for a period of 9 years [341], indicated that the rate of corrosion decreased with time according to
249 parabolic kinetics. Extrapolating their results to a 1,000 year timescale gives a 1,000 year corrosion figure of about 0.5 mm for uniform corrosion and 2.5 mm for pitting corrosion. These values are not far from the average value of the Johnson and Francis data and the Enresa report [337]. The results of this extrapolation are added to FIGURE 2 (black squares).
FIGURE 2. Corrosion rates deduced from iron archaeological artefacts (after Johnson and Francis [339]) [337]. The data points shown as black squares are derived from extrapolation of data from short-term laboratory experiments.
Similarly, Dresselaers et al. [342] examined the corrosion of two grey cast iron shaft linings sunk into clay rocks after 60 and 90 years, which produced outer oxide layers of 300 and 350 µm thick respectively. In both cases the oxide had the structure of magnetite (Fe3O4), which suggests that oxidation was taking place at a low oxidation potential. Under this layer of magnetite there was a thicker layer where ‘internal’ oxidation had taken place preferentially at the graphite flakes that are a feature of this type cast iron. This layer was 600 µm and 700 µm thick after 60 and 90 years respectively.
The data from short-term laboratory experiments on copper samples also compared very well with the data from copper and copper-alloy artefacts, as illustrated on the basis of the following four archaeological analogue studies [324,337,339,343-345]: - calculations made by Tylecote [343], based on data from a wide range of copper and bronze artefacts retrieved from the seabed, suggest that a copper disposal container would need a thickness of between 10 to 20 mm to achieve a 1,000 year service life. - Tylecote [344] also carried out a study to investigate the effect of soil type on the corrosion of copper and bronze artefacts. He found that in mildly alkaline soil, and provided that the material was single phase, penetration of bronze was not expected to exceed 0.45 mm in 2,000 years.
250 - Johnson and Francis [339] calculated a corrosion rate of copper and copper alloys ranging between 0.006 to 4 mm per thousand years, based on a study of over 33 copper and copper alloy archaeological artefacts. The corrosion rate was estimated from the corrosion product thickness and the age of the artefact. - investigations of a 300-year old bronze (96% copper) cannon [345], recovered from the Swedish warship Kronan and which was partly buried in the clayey sediments on the sea bottom, revealed a total corrosion of about 50 µm over the 300-year period (corrosion rate ~0.15 µm/year) and that pitting was confined to inhomogeneities. The environmental conditions of the clay adjacent to the cannon are broadly representative of those expected in the Swedish KBS3 concept, which envisages the encapsulation of spent fuel in a 0.05 m thick copper canister surrounded by bentonite clay.
FIGURE 3 compares the data from the archaeological analogues with shorter-term experimental data [337]. The shorter-term data, shown in FIGURE 3, come from experiments exposing copper samples to soil for up to 14 years at different pHs. The least corrosive soils produced general corrosion rates between 0.05 and 2.5 µm/year. Pitting corrosion could be 20 times faster. More acidic soils could produce general corrosion rates as high as 35 µm/year [346]. To compare the shorter-term experimental data with the data from the archaeological artefacts, the upper and lower limits of the nominal 1,000 year corrosion values from Johnson and Francis have been used. Extrapolation (dashed lines), assuming parabolic corrosion, was adopted to create the upper and lower bounds. This may not be strictly applicable (especially for seawater where corrosion products can be eroded away), but if an artefact is to survive for hundreds of years, it is almost certainly necessary for the corrosion products to protect the metal to some degree. It is therefore convenient to assume parabolic corrosion as a means of connecting the short- and the long-term data [337].
• studies on archaeological artefacts have also shown that the general corrosion rates decrease with time of exposure [347]. For iron-based artefacts, one possible explanation could be that in the long-term, the amount of goethite increases and the amount of lepidocrocite decreases.
251 FIGURE 3. Corrosion of copper and bronze in soil, showing a comparison between short- and long-term data from various sources. The dashed lines show extrapolations through the 1,000 year (nominal) corrosion figures of Johnson and Francis, assuming parabolic oxidation [337].
From the above-mentioned examples, it is evident that archaeological analogue studies can provide useful information on general corrosion rates and pitting, the two processes that are believed to be most likely to cause degradation of the metallic container. However, some prudence is called for, since museum artefacts tend to focus on the best preserved items – it is quite possible that only a few samples survived while the rest was corroded away and this may therefore introduce a non- conservative bias in assessing corrosion rates. Furthermore, often the environmental conditions in which the artefacts were found either differ from the underground disposal conditions or they were not always well documented. Also, since many of the candidate container materials that will be used in underground repositories have been produced only in recent times (such as stainless steels and titanium alloys), they have no counterparts in nature or in archaeology. An intermediate solution to overcome these drawbacks might be the investigation of more recent ‘industrial artefacts’ such as the Chrysler Building (New York, USA) for the study of atmospheric corrosion of stainless steels. Yusa et al. [348] investigated the corrosion of buried gas and water pipes made from various steels and located in a clay-rich burial environment. They reported maximum corrosion rates of up to 10 µm/year and noted that the principal corrosion products were FeCO3 and iron oxyhydroxides. These results are in good agreement with the average corrosion data from Johnson and Francis.
A final remark is that a lot more information might be available from archaeological artefacts – information that was acquired without any thought for radioactive waste disposal – that could be useful in predicting the general corrosion rates and understanding the processes involved more accurately and therefore, a thorough literature search may be informative.
252 B.6 DETERMINING AN EVOLUTIANARY PATH: characterisation of the near-field environment as a function of time
A proper characterisation of the transient and changing environmental conditions with respect to for example temperature, pH, oxidants, reducing agents, changing concentration of chlorides, H2O2 due to radiolysis, reduced sulphur compounds, SRBs, etc. is necessary in order to establish the evolution of the corrosive near-field repository environment (chemically and physically):
• in the short term there may be microbial activity in the surrounding near-field environment to consider, with the possible formation of reduced sulphur compounds or even elemental sulphur and potentially sulphuric acid. Similar concerns are relevant to the period of exposure of the near-field environment to air oxidation where a variety of sulphur-containing species may co-exist, supporting a range of chemical reactions. • in the medium term there is the changing temperature at the container surface induced by radioactive decay of the waste. This temperature profile may, itself, induce changes in electrolyte chemical composition (e.g. chromatographic effect possibly raising the chloride concentration) and will alter the amount of water at the container surface. • in the medium and longer term there is the possibility of change in the environment caused by radiolysis (e.g. formation of hydrogen peroxide).
This is essential to identify the worst case conditions and to make any predictions of accumulated corrosion damage.
Once the evolutionary path has been determined, selected experiments can be elaborated, in environments and conditions that are predicted to prevail at different points along the evolutionary path, to verify corrosion performance and rates. This approach has been adopted by for example Japan [311]. These data can then be integrated to predict the long-term performance of the metallic container. However, one possible disadvantage of this approach may be the fact that the experiments, performed under the separate environmental conditions, are performed with freshly prepared specimens and therefore the ‘history’ of the specimens is not taken into account (e.g. corrosion products that are formed during the oxic phase or a change of the passive film properties during the oxic phase will affect the corrosion under anoxic conditions).
Another approach is to apply in situ monitoring techniques to differentiate between the effects of different phases of exposure on the corrosion. Examples include electrochemical techniques such as electrochemical impedance spectroscopy (EIS), linear polarisation resistance (LPR) or electrochemical noise (EN). All of these have the same characteristic, namely that they sum over all processes occurring at the sample surface (e.g. oxidation of metal, dissolution of metal cations, oxidation of solution ions on the surface, etc.). When the target is to specifically measure a very low corrosion rate as a function of time, possibly with changing environmental conditions, a possible technique could be the resistance wire method. In this technique a thin wire made of the material in question is exposed to the environment, and a direct current, I, is led through the wire and the resulting potential, U, is measured at the wire ends. The depth of corrosion
253 penetration (uniform or localised), dcorr, can be estimated from the resistance, R = U/I, according to the following formula [349]:
⎛ R ⎞ d = r ⎜1 − 0 ⎟ corr 0 ⎜ ⎟ ⎝ R ⎠ where r0 and R0 are the initial values of the radius and the resistance of the wire probe, and R is the resistance during corrosion. Differentiating this formula with respect to time gives us the possibility to estimate the corrosion rate (depth penetration rate) as
R 1 dR ∆ = 0.5 r 0 corr 0 R R dt
Another important feature of the resistance wire method is that it provides instantaneous corrosion rates. On the other hand, long-term immersion experiments, which are based on weight loss measurements, give an average corrosion rate over the entire immersion period.
Some of the difficulties that could arise in assessing the evolutionary path are: • what is the diffusion rate of oxygen in the saturated backfill/buffer material ? This parameter is needed to calculate the rate of supply of oxygen towards the metallic container. • is it possible that the presence of the γ-radiation field can cause a continuous supply of oxygen (due to the dissociation of the aqueous environment) ? Answering both questions may be essential in determining the length of the aeration period.
Determining the evolution of the near-field environment of the metallic container alone might not be sufficient since other factors might play a crucial role in the corrosion behaviour of the metallic container such as corrosion products which in some cases are known to act as a diffusion barrier.
254 B.7 APPROACH FOR MODELLING
The aim of this section is not to enumerate and describe all existing models, but rather to identify the different approaches existing today in long-term corrosion prediction.
With regard to their corrosion resistance, the metallic materials that are being considered as candidate container materials within the EU-countries generally belong to one of the following two main categories [350]:
• corrosion-allowance materials. Examples of these materials are carbon steel, low- alloy steels, cast irons and copper (in oxidising environments). These materials corrode actively under environmental conditions expected during geological disposal at a significant but relatively predictable uniform corrosion rate. The lifetime of these materials is attained by providing an appropriate wall thickness [350]. • corrosion-resistant materials. Examples of these materials are the passive materials (thermodynamically unstable but kinetically stable) such as austenitic stainless steels, Ni-Cr-Mo alloys and titanium alloys and the thermodynamically stable materials such as copper (in reducing environments and free of complexing agents). Passive materials spontaneously passivate in aqueous environments due to the formation of a protective oxide film, which considerably lowers the rate of general corrosion. The uniform corrosion rate is negligibly small as long as this passive layer, which is known to be stable over a relatively wide range of environments, remains intact. However, for these materials, the risk of localised corrosion, such as pitting and crevice corrosion, has to be taken into account because the protective film may break down locally [350].
The outcome of modeling for both types of material is quite different: for corrosion- allowance materials, the aim is to demonstrate that the general corrosion rate remains below an acceptable value, while for corrosion-resistant materials, the aim is to demonstrate, in a robust and convincing way, the absence of propagation of localised corrosion. Lifetime predictions are probably more difficult for corrosion- resistant materials than for corrosion-allowance materials.
B.7.1 Lifetime prediction for corrosion-allowance materials (carbon steel and low-alloy steel)
General corrosion is considered as the prevailing corrosion mechanism in the mean and the long-term. Actually, pitting or crevice corrosion, which occurs on the short- term, especially during the periods when oxygen is still present, tends to fade away because of the progress of general corrosion [350].
B.7.1.1 General corrosion
Different approaches have been used to predict and model the long-term corrosion of a steel container in underground repository conditions:
255 ½ the use of semi-empirical laws based on experimental data, taking into account the mass-balances and overestimating corrosion parameters, allows a realistic assessment of corrosion lifetime, which could be attained by providing an appropriate wall thickness. Containers of ‘reasonable’ thickness can thus be sized with these overestimating approaches [350].
A large number of experimental data already exists with respect to the general corrosion rate of carbon steel and low-alloyed steels in an argillaceous environment representative for an underground repository in clay. These data originate from several experimental programmes performed in various countries that are studying the feasibility of constructing a repository in clay and in granite (with a clay engineered barrier) [86]:
- AEA, United Kingdom [80,277,278,280,354]. - NAGRA, switzerland [355-358]. - SCK•CEN, Belgium [165,168,170,173,359,360]. - JNC (formerly PNC), Japan [311,361-363]. - ANDRA-CEA-EDF, France [364-367].
FIGURE 4 provides a compilation of the general corrosion rate data of carbon steel and low-alloyed steels in representative repository conditions.
FIGURE 4. Compilation of general corrosion rate data of carbon steel and low-alloyed steels in representative repository conditions [86,368].
256 Using these findings, Gras [353] derived the following laws to represent the general corrosion rate of carbon steel and low-alloyed steels under repository conditions:
−1340 T −1 • during the aerated phase: vCORR = 1.042 e (mm.year )
• during the deaerated phase: −1300 T −1 - for short periods: vCORR = 0.364 e (mm.year )
−1300 T −1 - for long periods (several years): vCORR = 0.162 e (mm.year )
½ some authors [278,351] have developed mechanistic models for the corrosion of carbon steel to be able to make long-term predictions of container lifetimes. The corrosion processes are modelled as a combination of electrochemical reactions at the carbon steel surface and mass transport of oxidants and corrosion products. These models do not take into account the effect of the corrosion products on lowering the corrosion rate. As can be seen in FIGURE 5, the average corrosion depth after 1 year of exposure estimated with such a model is five times larger than the experimental result and the difference between experimental results and estimated results increase with the time of exposure. This lack of agreement is attributed to the fact that the inhibition of corrosion by the build-up of corrosion products is neglected in the model [350].
FIGURE 5. Predicted corrosion penetration depth as a funtion of time at 80°C compared with experimental results [311,351].
More recently, a new model has been developed at ANDRA-CEA-EDF [369], taking into account the possibility of having a layer above the steel surface which
257 is responsible for the pseudo-passivity of the steel. This model considers a solid oxide, which adheres to the metal and through which the chemical species must diffuse to be able to react at the internal and external interfaces (FIGURE 6) [350].
FIGURE 6. Mechanistic modelling of corrosion processes, after Bataillon from CEA [369].
B.7.1.2 Localised corrosion
Two approaches, based on semi-empirical laws, to predict the long-term propagation kinetics of localised corrosion (pitting and crevice corrosion) in repository conditions have been developed [86]:
½ model based on pit propagation rates. It is generally accepted that the depth of pitting (P) varies with time t according to the following semi-empirical exponential law: P = k . t n
with k a constant dependent on the material and environmental conditions, whereas n is a numerical constant (mostly situated between 0.4 and 0.6).
This semi-empirical exponential law has been used frequently to predict the maximum depth of localised corrosion for carbon steel and low-alloyed steels under repository conditions, based on experimental data from long-term immersion tests and laboratory experiments performed in different environments and under various conditions (see Table 5) [86]:
• the American National Bureau of Standards (NBS) carried out a survey, between 1922 and 1952, with respect to the localised corrosion of carbon steel and cast iron tubing (six different materials) buried in 47 different soils. The results are given in terms of maximum pit depths as a function of soil characteristics and exposure time [276].
258 • The findings of Romanoff [276] have been reinterpreted by Mughabghad and Sullivan [370]. • Marsh et al. [278,280,282] have carried out long-term immersion tests on carbon steel coupons immersed for up to 30,000 hours at 90°C in chloride and -1 - carbonated solutions (0.1M NaHCO3, 1000 mg.L Cl , pH = 8.4). The measured maximum pit depths were analysed statistically using the distribution function of Gumbel. • Arnoux [373] performed an elaborate study on the cast iron circuits of the drinking water distribution network of Strasbourg. The maximum pit depths of localised attack were measured on 83 samples taken during system leakages. • Gras [39] formulated the following exponential equation, based on the statistical analysis of the entire set of above-mentioned data together resulting in average values for k and n:
P = 0.746 t 0.37
• a study performed by Zapp [372], however, concluded that the NBS findings cannot be used quantitatively. • Matsushimo [370] summarized pitting corrosion results of carbon steels observed between 1975 and 1999. The materials have been in service for a few years (ten years at most). Propagation rates were calculated by dividing the depth of attack by the exposure period.
TABLE 5. Prediction of maximum pit depths on carbon steel and low alloyed steel [86].
Conditions Equation Ref.
0.46 0.1 M NaHCO3; 1,000 mg/L Cl-; pH = 8.4; 90°C; 10,000 hours P = 8.35 ·t [278]
0.42 0.1 M NaHCO3; 1,000 mg/L Cl-, pH = 8.4; 90°C; 30,000 hours P = 7.02 ·t [280]
water; 65°C P = 0.01365 ·t0.584 [282]
50/50-bentonite/water mixture; 80°C; aerated, 180 days P = 9.4 ·t0.5 [362]
0.5 50/50-bentonite/water mixture; N2; 180 days P = 0.44 ·t [362]
0.45 solution cont. HCO3-, Cl-, NO2-; pH = 9.73; 50°C; 350 days P = 3.6 ·t [372]
According to [276], based on average values for k and n P = 0.746 ·t0.37 [39]
According to [276], good aeration of soil P = 0.94 ·t0.26
According to [276], average aeration of soil P = 0.94 ·t0.39 [370] According to [276], low aeration of soil P = 0.53 ·t0.44
According to [276], very low aeration of soil P = 0.53 ·t0.59
259 ½ model based on the pitting factor.
The models based on pit propagation rates are very conservative and usually lead to an overestimation of the importance of localised corrosion (pit depths) compared to general corrosion, especially in the long-term: localised corrosion, which occurs in the short-term, especially during the period when oxygen is still present, tends to fade away because of the progress of general corrosion.
Therefore, Gras [353] has proposed a new (and more pragmatic) approach, which was inspired by the method applied by JNC [311]. This approach is based on the ‘pitting factor’. The pitting factor F is defined as the ratio between the maximum depth of localised corrosion P and the average general corrosion depth X:
P F(t) = X
FIGURE 7 shows the relation between the average corrosion depth and the pitting factor for carbon steel and cast irons under various test conditions:
• Taniguchi et al. [311] used a set of corrosion data originating from (i) immersion tests that were conducted under atmospheric conditions in carbonate/chloride- containing aqueous solutions and (ii) a study performed by Romanoff [276] in which steel and iron specimens were exposed to 47 different soil types. These results are represented in FIGURE 7(a). Since the surface areas of the specimens used in the laboratory experiments were smaller than those used in the experiments in soils, the extreme value statistical analysis method was applied to the laboratory results in order to be able to compare them with those obtained in soil [311]. • Gras [353] used a set of corrosion data originating from experiments performed on carbon steel and cast iron specimens in a wide variety of environments (clay, granite or soil; distilled water, claywater, granitic water, cement-based water or seawater) and under various conditions (T: ambient - 90°C; non-aerated and aerated). These results are represented in FIGURE 7(b).
From both studies, it was observed (see FIGURE 7(a) and (b)) that almost all the data points are systematically located below an upper limit line described by an exponential law, indicating a decrease of the pitting factor with increasing average general corrosion depth. The exponential equation derived from FIGURE 7(a) is conservative compared with the findings shown in FIGURE 7(b). Gras proposed the following equation [86,353]:
F = 4.64 X −0.67 from which P = 4.64 X 0.33 with P and X (= vCORR.t) expressed in mm.
260
(a) (b)
FIGURE 7. Relationship between average corrosion depth and pitting factor for carbon steel and cast irons. (a) according to a study performed by Taniguchi et al. (JNC) [311]. (b) according to a study performed by Gras (CEA) [86,353].
B.7.2 Lifetime prediction for passive materials (stainless steel, Ni-Cr-Mo alloys)
The major corrosion mode for passive alloys exposed to the underground repository environment, which is near-neutral and mildly oxidising with the presence of aggressive halogen ions, is not general corrosion, but rather localised corrosion such as pitting, crevice corrosion and stress corrosion cracking, if the passive film is locally damaged and the metal is unable to restore this layer [374]. The initiation of localised corrosion is highly dependent on the redox conditions: in reducing environments, no localised corrosion is expected but oxidising conditions make pitting or crevice corrosion likely [350].
Several different approaches have been developed and applied to make corrosion failure life predicitons for passive materials:
½ deterministic approach. Because failures are rare events, there is generally insufficient data on any given system to derive reliable empirical models (based on the exact corrosion rate measured by laboratory and/or field tests) that capture the impact of all (or even some) of the important independent variables. The alternative prediction philosophy is a deterministic approach, in which predictions are made on the basis of mechanisms-based physical and chemical models whose outputs are constrained by the natural laws (conservation of mass and charge, Faraday’s Law) [375]. Macdonald [376,377] has developed the Point Defect Model (PDM) for the growth and breakdown of passive films on metal and alloy surfaces. Over the past decade, the PDM, together with the subsequent development of Damage Function Analysis (DFA), has become an interesting tool to predict the evolution of localised corrosion damage (pitting, crevice corrosion, stress corrosion cracking, corrosion fatigue, erosion-corrosion). The PDM has also been used to derive the dependence of the pitting potential on the voltage scan rate and to predict the
261 probability distribution function in the induction time for pitting [378]. Reliable assessment of the long lifetime in the repository, however, should be based on the exact corrosion rate or be supported by complementary data from a natural analogue [374].
½ stochastic approach. This approach describes corrosion life as the probability of occurrence. It has been widely accepted that localised corrosion is a stochastic process [379-383]: the sites where localised corrosion initiate, and subsequently propagate, occur at random and are therefore very difficult to predict. Due to the stochastic nature of localised corrosion, a scattering of corrosion data cannot be avoided, even under well-controlled laboratory conditions and therefore, it is imperative that the experimental data (such as pit nucleation potential) are analysed from a statistical point of view. Various types of probability distribution functions (pdf) [374,384] have frequently been used to analyse corrosion data, as illustrated in TABLE 6 [374].
TABLE 6. Probability distribution functions (pdf) observed in corrosion [374,384].
Probability distribution function Examples in corrosion
Normal distribution Pitting potential Log-normal distribution SCC1 failure time Poisson distribution Two dimensional distribution of pits Exponential distribution Induction time for pit generation, SCC and HE2 failure time Extreme value distribution Gumbel distribution Maximum pit depth Weibull distribution SCC failure time Generalised extreme value distribution Maximum pit depth, fatigue crack depth
1 SCC : stress corrosion cracking. 2 HE : hydrogen embrittlement.
Extreme value statistical analysis has become a well-established method for extrapolating corrosion data from small test coupons to the large surface areas of real structures. The analysis is based on the work of Gumbel [385]. Early applications of the technique to the analysis of corrosion data were described by Aziz [386] and Hawn [387], and there are numerous other examples in the literature including analysis of the corrosion of marine structures and pipelines [388-390], high temperature corrosion studies [391] and the corrosion of nuclear waste containers [392,393]. The technique is recommended in ASTM standards on pitting corrosion [394] and its application has been reviewed by Kowaka [395], Laycock [396], Turnbull [397] and Shibata [383]. Recently, a new and alternative methodology to the Gumbel approach was developed. This methodology consists of characterising and modelling both the morphology of pits and the statistical distribution of their depths from a limited inspection dataset. The heart of the data processing is based on the combination of two recently developed statistical methods that avoid making any choice about the type of the theoretical underlying parent distribution of pit depths: the Generalised Lambda Distribution (GLD) is used to model the distribution of pit
262 depths and the Bootstrap technique to determine a confidence interval for the maximum pit depth [398].
½ repassivation method. The basic idea behind this approach is that for a passive metal, there exists a critical potential for localised corrosion (pitting, crevice corrosion, stress corrosion cracking), such that in the potential domain below this critical value (or, less noble than), that particular form of localised corrosion will never take place. This approach involves the comparison of this critical potential to the free corrosion potential, ECORR, of the material concerned in an environment : in the domain where the free corrosion potential is lower than its critical potential for initiation of a certain form of localised corrosion, the possibility of that particular form of localised corrosion occurring is simply nil and the corrosion of the container will continue in a uniform manner at a low rate determined by the passive current density, i.e. the region of safe usage for that material, whereas if the free corrosion potential lies above (or, nobler to) the critical potential, there is always a possibility that kind of localised corrosion will initiate, i.e. the region in which that material should not be used [202]. In recent years, the protection (or repassivation) potential has been put forward as a conservative, lower-bound value for the critical potential [197-199,399,400]. This approach was developed by Cragnolino et al. [197-199,399,400] in the framework of the Yucca Mountain project evaluation and is also being used in Japan [200,202,401], France [350,402] and Belgium [175,184,188,189] as a criterion for the corrosion resistance of candidate passive alloys for high-level waste containers. In the framework of studies carried out in Japan, the USA and Belgium, numerous protection potential data have already been established.
The state of knowledge and the theoretical basis for making long-term extrapolations of the corrosion resistance of passive materials are still limited. Moreover, there is no natural or archaeological analogue of these alloys that can be used to gain confidence in the models used for long-term performance prediction. As remarked by the Nuclear Waste Technical Review Board, which is charged with performing an independent evaluation of the potential disposal at Yucca Mountain in Nevada, there is no experience with ‘metals protected by passivity for periods longer than 100 years’ [350,352].
The development of models to predict the lifetimes of nuclear waste containers is still incomplete. To be able to extrapolate corrosion phenomena up to thousands of years, a better understanding of the mechanisms and processes involved is necessary [350]. Furthermore, models predicting the failure life of the container should be coupled to models predicting the changes of the chemistry of the near-field environment due to, for example, radiolysis.
263 B.8 CEMENTITIOUS BACKFILL MATERIALS
B.8.1 Introduction of cracking due to variation in thermal conductivity of the various metallic and cementitious components
One potentially significant difference relative to the use of carbon steel reinforcement that is pertinent to the design of a nuclear waste repository concerns the thermal expansion coefficient of stainless steel compared to that of concrete and of carbon steel. The thermal expansion coefficients of ferritic steel and concrete are very close at 1.2×10-5 °C-1 and 1.0×10-5 °C-1 respectively whereas that of austenitic stainless steel is 1.8×10-5 °C-1. Thus compressive stresses can develop in stainless steel reinforced concrete on heating, leading, potentially, to defects and cracking in the adjacent concrete (under tensile stress) in the contact zone. Although there is no evidence that this has been a practical problem, it is clearly an important design parameter to be considered for the thermal phase of nuclear waste storage, which should be addressed by detailed finite element stress analysis.
B.8.2 Concrete spalling
Volume expansion associated with the corrosion of carbon steel can cause concrete spalling. If oxygen also has easy access due to high porosity or cracks and the passive film has broken down, general corrosion rates are likely to increase.
This phenomenon has been observed by Taylor et al. [261]. Carbon steel plates that were partly embedded in blocks of 3:1 PFA/OPC cement (to simulate corrosion of carbon steel at the air-steel-cement interface in a void between a waste container and the surrounding grout) caused the cement blocks to split open, even after 6 months, due to the replacement of alloy with corrosion products (volume expansion) generated under this condition. After cracking of the cement blocks, it was found that general corrosion became the dominant process and overwhelmed any localised corrosion that may have initiated.
264 ACKNOWLEDGEMENTS
The authors gratefully acknowledge their respective national authorities and institutions, and the European Commission for funding several of the projects which provided source information for this report. The authors would also like to express their thanks to G. Santarini (CEA/SACLAY, Gif Sur Yvette, France) who participated as an expert.
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