A Review of Stress Corrosion Cracking of High-Level Nuclear Waste

CNWRA 92

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Prepared for

Nuclear Regulatory CommissionContract NRC-02-88-005

Prepared by

Center for Nuclear Waste Regulatory AnalysesSan Antonio, Texas

August 1992t

I.2 T199209150C008A P ev iew of Stress CorrosionCrack 1ring of High-Level WasteC';- d -airner Materials-I-I-CNWRA'-, - 21

Property of CNWRA 92-021CNWRA Library

A REVIEW OFSTRESS CORROSION CRACKING

OF HIGH-LEVEL NUCLEAR WASTECONTAINER MATERIALS - I

Prepared for

Nuclear Regulatory CommissionContract NRC-02-88-005

Prepared by

Gustavo CragnolinoNarasi Sridhar

Center for Nuclear Waste Regulatory AnalysesSan Antonio, Texas

August 1992

Number

CNWRA 91-004

CNWRA 91-008

PREVIOUS REPORTS IN SERIES

Name

A Review of Localized Corrosion of High-Level Nuclear WasteContainer Materials - I

Hydrogen Absorption and Embrittlement of Candidate ContainerMaterials

Date Issued

April 1991

June 1991

ii

ABSTRACT

In this report, the literature on stress corrosion cracking (SCC) of the candidate container materials forthe proposed Yucca Mountain repository site, including Fe-Cr-Ni alloys, such as Types 304L and 316Lstainless steels and alloy 825, and copper-based alloys, such as CDA 102 (oxygen-free, high conductivitycopper), CDA 613 (Cu-7AI-3Fe), and CDA 715 (Cu-3ONi), is reviewed. The advantages and limitationsof different stress corrosion cracking (SCC) test methods are discussed in terms of their suitability fordetermining parameters that can be used for long-term prediction of SCC resistance performance,followed by a detailed review of all the investigations related to the Yucca Mountain project conductedby Lawrence Livermore National Laboratory (LLNL) and its subcontractors on behalf of the DOE andby Cortest Columbus Technologies (CCT) on behalf of the NRC. Finally, results available in the generalliterature for the materials considered are examined, using a phenomenological approach which takes intoconsideration mainly the role of environmental variables on SCC and the adequacy of the concept of acritical potential for SCC as a bounding parameter for performance assessment.

iii

CONTENTS

Section Page

FIGURES .......................TABLES .......................ACKNOWLEDGMENTS.EXECUTIVE SUMMARY ............

. . . . . . . . .

. . . . . . . . .

. . . . . . . . .

. . . . . . . . .

viixxi

xii

1

22.12.22.2.12.2.22.2.32.2.4

33.1

INTRODUCTION ........... ................................ 1-1

STRESS CORROSION CRACKING TEST METHODS AND PARAMETERS ..... 2-1VARIABLES AND PARAMETERS ................................ 2-1STRESS CORROSION CRACKING TEST METHODS ................... 2-2Slow Strain Rate Test ...Constant Deflection TestConstant Load Test .

Fracture Mechanics Test .

2-32-52-72-8

STRESS CORROSION CRACKING Op Fe-Ni-Cr-Mn ATLT.OYSREVIEW OF THE DOE/NRC HLW RESEARCH *vs . .*.

3.1.1 Materials ...............................3.1.2 Environmental Considerations ..................3.1.3 Fe-Ni-Cr-Mo Alloys ........................3.1.3.1 Slow Strain Rate Tests ......................3.1.3.2 Constant Deflection Tests .....................3.1.4 Other U.S. HLW Investigations .................3.1.5 European HLW Investigations ..................3.2 REVIEW OF THE GENERAL LITERATURE.3.2.1 Effect of Environmental Factors .................3.2.1.1 Effects of Chloride and pH ....................3.2.1.2 Effect of Temperature .......................3.2.1.3 Effect of Radiation.........................3.2.1.4 Effect of Potential .........................3.2.2 Effect of Mechanical Factors ...................3.2.2.1 Effect of Stress and Strain Rate .................3.2.2.2 Effect of Stress Intensity .....................3.2.3 Effect of Metallurgical Factors .................3.2.3.1 Effect of Chemical Composition ................3.2.3.2 Effect of Microstructure .....................3.2.3.3 Effect of Cold Work and Surface Preparation.

.... . . . . . . . . ...................................................................................................................................................................................................................................................................................

..... .. 3-1..... .. 3-1..... .. 3-1..... .. 3-3..... .. 3-9..... .. 3-9.... 3-14

... 3-20.... 3-20.... 3-21.... 3-21.... 3-23.... 3-25.... 3-27.... 3-29.... 3-37.... 3-37

3-383-413-41

.... 3-45.... 3-46

..... .. 4-1..... .. 4-1..... ..4-4..... ..4-4

44.14.24.2.1

SCC OF Cu AND Cu-BASED ALLOYS ......REVIEW OF THE DOE/NRC HLW RESEARCHREVIEW OF THE GENERAL LITERATURE . .Effect of Environmental Factors ............

iv

CONTENTS (Cont'd)

Section Page

4.2.1.1 Effect of Nitrate and Nitrite .4.-.................................. 444.2.1.2 Effect of Other Anionic Species ................................. . 4-74.2.2 Effect of Alloy Composition ..................................... 4-8

5 SUMMARY AND RECOMMENDATIONS .......... ................. 5-1

6 NOMENCLATURE ...................................... 6-1

7 REFERENCES . ...................................... 7-1

v

FIGURES

Figure Page

2-1 Specimen configuration for the wick-test (Dana, 1956) ..................... 2-6

2-2 Schematic of the apparatus for constant load SCC test with heated tensile specimen.(a) Side view; (b) Top view. (Sato et al., 1990) .2-2

3-1 Concentration changes in J-13 water during reaction with crushed tuff and stainlesssteels corrosion samples at 1000 C. Data at various time periods are from independenttest cells. (Abraham, 1986). 3-6

3-2 Concentration changes in lOX J-13 water (solution where all species are present to tentimes their concentration in J-13 water) during reaction with crushed tuff and stainlesssteels corrosion samples at 100TC. Data at various time periods are from independenttest cells. (Abraham, 1986) .3-7

3-3 Relative cracking susceptibility of six candidate container materials in simulated J-13well water at 930 C. Nominal strain rate: 1 x IC- per second. (Maiya et al., 1990) . . 3-13

3-4 Corrosion potentials of Types 304L stainless steel and alloy 825 in long-term boildowntests in aerated, simulated J-13 water (Beavers, 1992) .3-19

3-5 Effects of oxygen and chloride concentration on the SCC of austenitic stainless steelsin high-temperature (250-350'C) water (Gordon, 1980) .3-24

3-6 Effects of temperature and chloride concentration on the SCC of solution-annealed Type304 stainless steel in aerated NaCI solutions at various pHs (Truman, 1977) .3-25

3-7 Effect of chloride concentration on the SCC of solution annealed AISI 304 stainlesssteel in NaCl solutions at 100IC using the wick test. Note the distribution of failuretimes for the cracked specimens. (Warren, 1960) .3-27

3-8 Effect of temperature on the SCC of solution annealed AISI 304 stainless steel in100 ppm Cl- solution using the wick test (Warren, 1960) .3-28

3-9 Combined effects of temperature and chloride concentration on the SCC of solutionannealed AISI 304 stainless steel using the wick test (Warren, 1960) .3-29

3-10 Effect of the metallic cation on the SCC of solution annealed AISI 304 stainlesssteel in 100 ppm Cl- solutions at 1000 C using the wick test (Warren, 1960) .3-30

3-11 Ranges of potential and temperature over which SCC of solution annealed Type 316stainless steel occurs as determined in slow strain rate tests conducted in watercontaining 5 ppm Cl- (Yang, 1992) .3-32

vi

FIGURES (Cont'd)

Figure Page

3-12 Effects of temperature and applied potential on the strain for crack initiation forboth solution annealed and sensitized Type 316 stainless steel in water containing5 ppm Cl- as determined in slow strain rate tests (Yang, 1992) ...... . . . . . . . . . 3-33

3-13 Effect of potential on the time to failure of sensitized (15 hours at 650'C) Type 304stainless steel at various stress levels in 10WM NaCI solution at 1000C(Herbsleb, 1980). 3-34

3-14 Determination of crevice corrosion repassivation potentials (ER) for Type 316 stainlesssteel in dilute chloride solutions at 80'C (Tamaki, 1990) .3-35

3-15 Crack growth rate as a function of stress intensity (K,) for creviced, tapered, double-cantilever-beam specimen of Type 316 stainless steel showing results for twochloride concentrations at a potential just above the crevice corrosion repassivationpotential (Tamaki, 1990) .3-36

3-16 Ranges of chloride concentration and potential for the SCC of creviced specimensof Type 316 stainless steel at 80'C (Tamaki, 1990 as summarized by Newman,1990). 3-37

3-17 Effects of pH and strain rate defining domains of pitting and SCC for solutionannealed Type 304 stainless steel in 5 M NaCl solution at 1 10'C (Mancia, 1988) . . . . 3-39

3-18 Effect of stress intensity on the crack growth rate of solution annealed Type 304Lstainless steel exposed to 42% MgCl2 solution at 130'C and to 22 % NaClsolution at 105 0C (Speidel, 1981) .3-40

3-19 Summary of the effect of the elements in the periodic table on the SCC resistanceof austenitic stainless steel in chloride solutions (Sedriks, 1979) .3-42

3-20 Effect of the nickel content on the maximum crack velocity (Stage 2) forFe-Cr-Ni alloys in aerated 22% NaCl solution at 105'C. Arrows indicate nocrack growth. (Speidel, 1981). 3-43

3-21 Effect of the nickel content on the threshold stress intensity (KL) for the SCC ofFe-Cr-Ni alloys in aerated 22% NaCl solution at 105'C (Speidel, 1981). 3-44

3-22 Effect of the molybdenum content on the threshold stress intensity (Kb,,) for austeniticFe-Cr-Ni-Mo alloys in aerated 22% NaCl solution at 105'C (Speidel, 1981) .3-45

4-1 Corrosion potentials of CDA-102 and CDA-715 in long-term boildown tests in aerated,simulated J-13 water (Beavers, 1992) .4-3

vii

FIGURES (Cont'd)

Figure Page

4-2 Effect of the nickel content on the ultimate tensile strength (UTS) in air and SCCthreshold stress (oh) for copper-nickel alloys in ammonia atmosphere (Speidel,unpublished).. . .. 4-10

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TABLES

Number Page

3-1 Nominal chemical composition of current DOE candidate and selected alternatecontainer alloys .............................................. 3-2

3-2 Chemical composition ranges of J-13 well water (Glassley, 1990) ............... 3-3

3-3 Analysis of solutions from U-bend stress corrosion cracking tests in aerated J-13water under y-irradiation. Fresh J-13 water was added at the end of each period.(Westerman, 1987). 3-5

3-4 Summary of slow strain rate tests on Type 304L stainless steel at 900 C in variousenvironments. Nominal strain rate: 1 x 104/sec. (Beavers, 1989b) .3-10

3-5 Summary of slow strain rate tests on Type 304L stainless steel at 90'C in variousenvironments. Nominal strain rate: 1 x 104/sec. Specimens were shorter andsmaller diameter than in Table 34. (Beavers, 1989b) .3-11

3-6 Results of constant deflection tests on various austenitic alloys in J-13 water(Westerman, 1987; Juhas, 1984) .3-15

3-7 Effect of nitrate concentration on corrosion potential of Type 316L stainless steelin simulated J-13 Water at 950C. Solutions were deaerated by argon.(Cragnolino, 1991) .3-17

3-8 Changes in the concentrations of various species during the long-term boildowntests with simulated J-13 water (Beavers, 1992a) .3-18

3-9 Stress corrosion cracking test results on 4-point bend specimens exposed to J-13water at 900 C. Total exposure time: 9000 Hours. (McCright, 1985) .3-21

3-10 List of liquid environments known to cause stress corrosion cracking in austeniticstainless steel and Ni-based alloys .3-22

3-11 Stress corrosion threshold stresses of solution annealed austenitic stainlesssteels exposed to concentrated solutions of MgC12 in water at 145°C-150'C.(Speidel, unpublished a) .3-38

4-1 Results of slow strain rate tests for CDA 102 in nitrite solutions. Nominalstrain rate: 1 x 104 /sec. (Beavers, 1990c) .............. .. .. .. ... .. .. . 4-2

4-2 Environments that promote SCC of copper-based alloys ......... . . . . . . . . .. . . 4-5

ix

TABLES (Cont'd)

Number Page

4-3 The corrosion potential, ratio of fracture stress in solution to that in air, averagecrack velocity, surface appearance of the fractured specimen, and the cracking modeof admiralty brass in various solutions at 250C. (Kawashima et al., 1979) ..... . . . . 4-6

4-4 Nominal transgranular crack velocities for copper in 1 M sodium nitrite, pH 9, testedunder slow strain rate at 250C ...... .............. 4-8

4-5 Results of slow strain rate tests on copper monocrystals in 0.1 M sodium acetate(Cassagne et al., 1990).. ........... . 4-9

x

ACKNOWLEDGMENTS

This report was prepared to document work performed by the Center for Nuclear Waste RegulatoryAnalyses (CNWRA) for the U.S. Nuclear Regulatory Commission under Contract NRC-02-88-005. Theactivities reported here were performed on behalf of the NRC Office of Nuclear Regulatory Research,Division of Regulatory Applications. The report is an independent product of the CNWRA and does notnecessarily reflect the views or regulatory position of the NRC.

The authors gratefully acknowledge the technical review conducted by other members of the CNWRAand by Stephen J. Lukezich from the Institute staff. Appreciation is due to Bonnie L. Garcia for herassistance in the preparation of this report.

xi

EXECUTIVE SUMMARY

This report is a review of the literature on stress corrosion cracking (SCC) of the candidate materialsidentified in the Site Characterization Plan (SCP) by the Department of Energy (DOE) for theconstruction of high-level waste (HLW) containers in the proposed repository site at Yucca Mountain.These materials are Fe-Cr-Ni alloys, including Types 304L and 316L stainless steels and alloy 825, andcopper-based alloys, such as CDA 102 (oxygen-free, high conductivity copper), CDA 613 (Cu-7AI-3Fe)and CDA 715 (Cu-3ONi).

The effect of environmental variables on SCC is discussed in general terms to emphasize the need foridentifying experimental parameters that can be used as bounding parameters for the long-term predictionof material degradation due to SCC. It is suggested that a critical potential for SCC may exist for therange of environmental factors of interest, including the concentration of aggressive and inhibitingspecies, pH, temperature, and other variables, under the anticipated range of repository conditions. Asan example, the existence of a critical potential has been demonstrated for the transgranular stresscorrosion cracking (TGSCC) of solution-annealed Type 304 stainless steel in boiling, concentrated MgCI2solutions. The critical potential concept was recently applied to the intergranular stress corrosioncracking (IGSCC) of sensitized Type 304 stainless steels in high-temperature aqueous environmentscharacteristic of recirculating lines in boiling water reactors (BWRs), and to the irradiation-assisted stresscorrosion cracking (IASCC) of austenitic stainless steel, irradiated to high neutron fluence levels in thecore of BWRs. An alloy with a sufficiently high critical potential for SCC in an anticipate repositoryenvironment could be a suitable material if the predicted corrosion potential remains lower than thecritical potential for a comprehensive range of environmental conditions.

Different SCC test methods are evaluated by comparing their advantages and limitations for conductingaccelerated tests and for determining suitable bounding parameters for life prediction. The test techniquesare classified according to the stressing method as slow strain rate test (SSRT), constant deflection test,constant load test, and fracture mechanics test. The advantages of the SSRT as an appropriate test fordetermining environmental conditions, including electrode potential ranges, in which a given material canbe susceptible to SCC are emphasized. Due to uncertainties in the stress conditions, the use of crackgrowth rate data obtained with this technique may be considered unreliable. However, the technique isuseful for comparison or ranking purposes. The advantages of constant deflection tests are mainly interms of simplicity and possible adaptation for experiments under heat transfer conditions. This techniquealso requires, in order to reduce uncertainties, a large number of simultaneous tests, even under wellcontrolled electrochemical conditions. Constant load tests can be used to obtain the threshold stress forSCC (ah). However, the threshold stress is not appropriate as a bounding parameter due to theunpredictable influence of flaws, notches, etc. on its value. Test methods using a linear elastic fracturemechanics (LEFM) approach are reviewed concluding that the threshold stress intensity for SCC (K,,)could be a useful bounding parameter for assessing SCC resistance of different candidate materials. Theirmain limitation is the extremely long test time that may be required to obtain reliable values of K,particularly for materials with a relatively high SCC resistance.

A brief description of the material selection process conducted by Lawrence Livermore NationalLaboratory (LLNL) is presented. The possible variations in the composition of the environment incontact with container materials are then discussed taking into consideration rock-water interactions as

xii

affected by temperature, changes in occluded regions such as crevices, modifications due to wet/drycycles, and the effect of radiolysis.

The investigations conducted and sponsored by LLNL, as well as those conducted by Cortest ColumbusTechnologies (CCT), are reviewed in detail considering separately the classes of Fe-Ni-Cr alloys andcopper-based alloys. Investigations conducted within HLW programs in some European countries arebriefly considered in the review. A comprehensive review of the general literature is also included, inwhich the effect of environmental, mechanical, and metallurgical factors are separately discussed.

From the data available to date on the Fe-Cr-Ni alloys it is difficult to determine if SCC is a viabledegradation mode, especially for the high nickel alloys such as alloy 825. The chloride ion, alwayspresent in groundwaters, is clearly identified as the main aggressive species for inducing SCC ofFe-Cr-Ni alloys. Types 304L and 316L stainless steels are both prone to cracking at relatively lowconcentrations and temperatures, particularly under heat transfer conditions and in the presence ofcrevices.

SCC does not appear to be a viable degradation process for the copper-based alloys, unless highconcentrations of the nitrite ion (which can be formed by y-radiolysis of humid air) are present atrelatively high potentials. These conditions do not seem to be highly probable in the repositoryenvironment.

However, there are many uncertainties that should be investigated. There is limited information on theSCC susceptibility of alloy 825 for a wide range of environmental variations, including the effect ofpotential. In the case of copper-base alloys there is a lack of information on the combined role ofbicarbonate, chloride, and sulfate. Other sources of uncertainties, such as the statistical significance ofthe number of specimens tested, variations in surface conditions between laboratory specimens andfabricated components, and limitations of test methods are also discussed. These issues have a directbearing on the long term performance evaluation of the selected materials.

One of the main conclusions of this review is that the concept of a critical potential for cracking inrelation to the repassivation potential for crevice corrosion of Fe-Ni-Cr-Mo alloys should be investigatedto determine its usefulness as a bounding parameter for performance assessment. The role of the mainspecies present in groundwater should be carefully addressed, as well as the relative importance of pH,temperature, and heat transfer effects, in conjunction with an appropriate prediction of the evolution ofthe repository environment with time.

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1 INTRODUCTION

The Nuclear Regulatory Commission (NRC) regulation 10 CFR 60.113 requires waste packages toprovide substantially complete containment of radionuclides for a minimum period of 300-1000 years.Arising from this requirement is the need for the Department of Energy (DOE) to demonstrate, throughproper selection, design, testing, and analyses, the long-term performance of waste packages. In orderto evaluate DOE's resolution of these technical issues, NRC must develop an understanding of theimportant factors that affect long-term performance of waste package materials and components. Otherconsiderations are the suitability and limitations of various test methods used to demonstrate performance,as well as the influence of factors not addressed by DOE that may affect the performance of wastepackages. The Integrated Waste Package Experiments (IWPE) research program at the Center forNuclear Waste Regulatory Analyses (CNWRA) supports the NRC in attaining this understanding througha program plan divided into five interconnected tasks:

Task 1: CorrosionTask 2: Stress Corrosion CrackingTask 3: Materials StabilityTask 4: Microbiologically Influenced CorrosionTask 5: Other Degradation Modes

The present report is part of the activities conducted under Task 2, being focused on reviewing theliterature on stress corrosion cracking (SCC) of the classes of candidate container materials selected byDOE for the proposed repository site at Yucca Mountain. This review is limited essentially to the classesof candidate materials already selected, namely austenitic Fe-Ni-Cr-Mo alloys and Cu-based alloys, eventhough the conceptual design of the waste package is currently under revision and changes in the list ofcandidate container materials are expected. Since this is the first of the two reviews on the same subjectplanned under the IWPE project, it is anticipated that the second one, scheduled for a later date, willaddress the stress corrosion cracking of alternate container materials and designs.

SCC is defined as a phenomenon by which ductile metals and alloys fail in a brittle manner through theinitiation and propagation of cracks resulting from the combined action of a sustained tensile stress(applied and/or residual) and a specific corrosive environment. In the context of this review, theenvironments of interest are aqueous and can be condensed layers of moisture or bulk solutions. In recentyears, a tendency has evolved to consider other forms of failure under the simultaneous presence of stressand a corrosive environment, such as corrosion-fatigue and hydrogen embrittlement, as parts of aspectrum of environmentally assisted cracking. However, despite certain similarities, these phenomenaexhibit substantial differences, thereby justifying a separate discussion. Whereas corrosion-fatigue is notconsidered a plausible failure process under the anticipated repository conditions (Manaktala, 1990), thepossibility of hydrogen embrittlement has been previously reviewed (Sridhar, 1991) for the currentcandidate materials.

SCC is one of the most insidious forms of metal failure because it usually occurs in metal and alloys thatare extremely resistant to uniform corrosion as a result of the formation of a film on the metal surfacethat slows down the corrosion rate. These films may be passivating layers, tarnish films, or dealloyedlayers. The subject has been extensively covered in books and specialized conferences (Logan, 1966;Staehle, 1969; Scully, 1971; Staehle, 1977; Gangloff, 1984; Gangloff, 1990). A large number of failures

1-1

have been observed in many industries, including the nuclear power industry in which the materialselection process is generally rigorous as a consequence of stringent safety considerations. It isinteresting to note that in a 1973 Conference, Bush and Dillon (1977) reviewed 88 cases of SCC innuclear power plants. A survey of this nature is almost an impossible task today as a result of theextremely large number of service failures. However, many alloys have performed extraordinarily wellunder very aggressive environmental conditions, indicating that a careful material selection and designprocess and an adequate knowledge of the anticipated environmental and mechanical conditions can reducethe risk of SCC failures to a reasonable minimum.

It is important to emphasize that the understanding of SCC is largely phenomenological. No satisfactorygeneralized theory exists to explain the behavior of all metal/environment combinations that lead to thiscomplex form of coupled environmental/mechanical damage. Although a variety of mechanistic modelshave been postulated and some mathematical models exist, none of them can be used with confidence forthe prediction of SCC without the support of additional experimental data.

The sources of the experimental data discussed in this review include published reports of DOElaboratories, various reports and papers of NRC-sponsored research at Cortest Columbus Technologies(CCT), published reports of research on high-level nuclear waste container corrosion from othercountries, and selected publications from the open literature.

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2 STRESS CORROSION CRACKING TESTMETHODS AND PARAMETERS

2.1 VARIABLES AND PARAMETERS

A critical issue in the design of high-level nuclear waste containers is the selection of the mostsuitable material and, related to this, the prediction of its performance over a very extended period.While the preliminary selection of a candidate material in terms of resistance to corrosion and stresscorrosion cracking (SCC) is essentially based on the experience accumulated in previous serviceapplications, the prediction of its long-term performance requires a joint experimental and modelingapproach.

Although mechanistic understanding and modeling of SCC have experienced significantdevelopments in recent years, suitable experimentation is still essential to define the range ofenvironmental and mechanical conditions in which a given material with a particular microstructure maybecome susceptible to SCC. The approach traditionally adopted for material evaluation and testing inother applications, such as in the electric power generating industry, of simulating as closely as possiblethe operating environmental conditions cannot be safely applied to high level-waste disposal for tworeasons. First, the wide range of possible changes in the physical and chemical properties of theenvironment with time are not completely characterized and second, the use of accelerated tests isrequired as a consequence of the extremely long life expected for the containers. The most fruitfulapproach is based on testing the individual and combined effects of different environment componentsthrough suitable accelerated tests, where the main environmental variables (i.e., temperature,concentration of aggressive and inhibiting species, pH, redox potential, etc.) are properly controlled andenhanced, in order to yield relevant parameters that can be used for incremental extrapolation andbounding models. In particular, the impact of upsetting environmental conditions must be properlyevaluated, as the experience of unexpected failures by SCC in many industrial applications teaches overand over again.

During the 1960's, it was clearly recognized in laboratory experiments that one of the mostinfluential factors in the propensity of a given material to fail as a result of the initiation and growth ofstress corrosion cracks is the electrode potential. In particular, it was established through potentiostaticSCC experiments that transgranular cracking of annealed AISI 304 stainless steel in hot, concentratedchloride solutions (e.g., boiling MgCI2) occurs only above a critical potential (Hines, 1958; Barnartt,1961; Smialowski, 1967; Uhlig, 1969; Brennert, 1960). As noted by Lee and Uhlig (1970), this potential(- -0.15 VSHE) was found to be slightly dependent on applied stress, chloride concentration, andtemperature, for boiling (100 to 1540C) concentrated solutions. They pointed out that a necessaryrequirement for cracking to occur under open circuit conditions is that the corrosion potential must lieabove that critical potential. Moreover, by appropriate modification of the environment through theaddition of sodium acetate or nitrate, the critical potential can be shifted to more noble values than thecorrosion potential, and, therefore, SCC can be inhibited (Uhlig, 1969). Although SCC testing ofaustenitic stainless steels in boiling MgCl 2 solutions is currently not considered an appropriate simulationof SCC failures occurring in dilute chloride-containing solutions, the concept of a critical potential forcracking deserves careful examination as a potentially suitable bounding parameter.

2-1

In recent years, it has been demonstrated that intergranular stress corrosion cracking (IGSCC)of sensitized AISI 304 stainless steel in high-temperature aqueous solutions occurs only above a certaincritical potential which depends on the anionic composition of the solution, pH, and temperature(Cragnolino, 1984; Cragnolino, 1985; Govindarajan, 1987). IGSCC has been a recurrent problem inrecirculating lines of boiling water reactors (BWRs) because under normal operating conditions, thecorrosion potential of the steel in the oxygenated reactor water is generally above that critical potential.The current approach to avoid the occurrence of this type of failure by merely changing the environmentalconditions lies in the reduction of the corrosion potential through the injection of hydrogen into therecirculation system. Under typical operating conditions in terms of temperature (288°C), conductivity(< 0.1 tS/cm), etc., the critical potential is about -0.230 VSHE (Kass, 1986), and the corrosion potentialcan be decreased to values lower than that with controlled additions of hydrogen gas to suppress IGSCC.Furthermore, it has been shown (Cragnolino, 1984; Macdonald, 1986; Indig, 1986) that propagatingcracks can be arrested by decreasing the potential below the critical value for a particular solutioncomposition.

Even more recently, Indig et al. (1992) have extended the concept of a critical potential toirradiation-assisted SCC (IASCC), which is the failure of austenitic stainless steel components irradiatedto high neutron fluence levels in the core of BWRs (Nelson, 1992).

These examples illustrate the importance of evaluating this concept of a critical potential as abounding parameter for the SCC resistance of a container material. It is apparent that an alloy with ahigh critical potential for SCC in an anticipated repository environment could be a suitable choice if thepredicted corrosion potential remains lower than that critical value for a comprehensive range ofenvironmental conditions.

Although it is well established that temperature is a major influential variable on SCCsusceptibility for many alloy/environment combinations, it is by no means clear that a critical temperaturefor SCC exists, as in the case already discussed (Cragnolino, 1991a) of pitting and crevice corrosion.In general, temperature exerts a monotonic effect on SCC-related parameters, such as critical potential,crack growth rate, etc. Although these parameters may exhibit a minimum (or maximum) at someintermediate temperature, it is difficult to anticipate an abrupt change or discontinuity in the effect of thisvariable.

Another important group of parameters that must reach some critical value for SCC to occurare those related to mechanical variables such as stress and/or strain. These parameters, such as thethreshold stress intensity factor, could be extremely useful as bounding parameters. However, it shouldbe noted that they are strictly valid for a material/environment combination rather than for a materialalone, and, therefore, a knowledge of the evolution of the environment with time is also required. Theyare discussed in the following sections in relation to different test methods for SCC, since these methodsare classified according to the nature of the stressing system.

2.2 STRESS CORROSION CRACKING TEST METHODS

As noted above, SCC test methods can be classified according to the stressing system adoptedfor testing. In many circumstances, a particular test method is adopted in order to reproduce or simulatethe stressing conditions present in a given component or structure. In other cases, what is required isa simple method for screening or ranking a group of alloys in a particular environmental condition.

2-2

Finally, an appropriate test or modifications of an existing test method can be adopted for a mechanisticstudy. The subject has been well covered, emphasizing either practical or fundamental aspects of thevarious methods, in several publications and books. A recently published book (Sedriks, 1989) offersa good overview of different test methods. For that reason, only a brief summary is presented here,mainly with the purpose of discussing the effect of test acceleration by increasing the severity of testvariables on parameters that may be relevant for long-term life prediction.

2.2.1 Slow Strain Rate Test

The slow strain rate test (SSRT), also known as constant extension rate test (CERT), involvesthe application of a relatively low rate of extension (usually ranging from 10 to 104 mm/sec) to anuniaxially loaded tensile specimen immersed in a corrosive environment (Parkins, 1979). The purposeof the test is to produce a fracture surface and, eventually, secondary cracks, which are fractographicallyand metallographically indistinguishable from those produced in the same environment under constant loador constant deflection conditions (Parkins, 1976).

Although this test method was introduced by Parkins in the late 1960's and became very popularin the last decade, only recently have standard procedures been developed. The InternationalOrganization for Standardization has adopted one that summarizes general guidelines (InternationalOrganization for Standardization, 1989).

One of the most useful applications of this test method is the identification of environmentalconditions (solution composition, pH, temperature, etc.) in which a given material is susceptible to SCC.In particular, for a given environmental condition it is relatively simple to determine very accurately acritical potential for SCC, or a potential range within which SCC occurs, using a limited number of tests.A large number of parameters, such as time to failure, elongation to failure, maximum load or stress,fracture stress, reduction in area, and the ratio of any of these parameters to those obtained in an inertenvironment (air, oil, glycerine, etc.), can be used to assess SCC susceptibility (Sedriks, 1989).Combinations of these parameters, in the form of particular indexes, have also been used (Clarke, 1978;Hishida, 1977). In many circumstances, a particular index or parameter is selected to establish acorrelation or dependence with another parameter associated with a different but related phenomenon (i.e,degree of sensitization, as measured by a particular electrochemical technique, or grain boundarysegregation, as measured by a surface analytical technique). In certain cases, appropriate parameters arederived from metallographic or fractographic examination of the failed specimens. For example, theFracture Area Ratio (FAR), obtained as a ratio of the area in the fracture surface exhibiting SCC to thetotal fracture surface area was found to be more reliable for assessing the effect of heat treatment on theIGSCC of alloy 600 in hot caustic solutions than other mechanical parameters (Lee, 1985).

SSRT is commonly considered a very severe test because crack initiation occurs readily evenin smooth, plain tensile specimens. However, this is not always the case. It has been shown by usinginterrupted tests (Andresen, 1982; Lin, 1981) that crack initiation may take place at relatively high strainvalues for sensitized Type 304 stainless steel in high-temperature aqueous solutions. It appears that forcertain metal/environment conditions in which the concentration of an aggressive species is very low, aspecific environment is required to build up locally to induce crack initiation. This may demand a timethat extends far beyond the total test time in a slow strain rate test (usually around a week), or thedevelopment of occluded cell conditions typical of crevice corrosion which are not easily established ina plain tensile specimen. For some alloy/environment systems, such as alloy 600 in high-temperature,

2-3

deaerated water, it is very difficult to reproduce the IGSCC of steam generator tubing observed inpressurized water reactors (PWRs) using SSRTs, even at strain rates as low as 10-'/sec. Reproduciblefailures were obtained (Totsuka, 1986) only by using split-tube, tensile specimens with a pressed hump(high levels of localized cold work). Recently, the results of a survey regarding the application of SSRTfor prediction of SCC in systems typical of the chemical process industry were critically analyzed(Beavers, 1992a). It was concluded that anomalous results, as compared to those obtained using otherSCC test methods, could be attributed to inadequate potentiostatic or chemical (redox couples) controlof the potential in some cases, and in others to an excessively high extension rate. However, it shouldbe noted that it is almost impractical to conduct SSRTs at initial strain rates lower than 1.Ox1O-'/sec (asapparently required in some alloy/environment systems) because it ties up relatively expensive equipmentfor a prolonged period without assurance of obtaining relevant data. As noted above, probably a criticalfactor is the initiation time or the development of conditions for crack initiation. This may require anextended pre-exposure time, preceding the onset of straining.

In addition to these parameters discussed, a more fundamental quantity required to characterizeSCC of an alloy/environment system is the crack growth rate and its dependence on metallurgical,mechanical, and environmental variables. To estimate the crack growth rate in a SSRT, the length of thelongest crack measured in a metallographic cross section of the failed specimen is usually divided by thetest time. Otherwise, through a scanning electron microscope (SEM) fractograph, the extension of thecrack on the fracture surface can be obtained for the same purpose. However, this approach may notyield a true average or maximum crack growth rate because the initiation time may represent an importantfraction of the test time. Nevertheless, the most important objection to the use of crack growth ratesmeasured in SSRTs is the lack of a precise definition of the stress conditions at the crack tip.

Several authors have used notched specimens in order to measure more accurately the rate ofcrack propagation and its dependence on the extension rate because multiple crack initiation sites can beavoided. Takano et al. (1977; 1981) used circumferentially notched, cylindrical specimens to study theeffect of extension rate on the SCC of Type 304 stainless steel in boiling MgC12 and of Cu-30%Zn inammoniacal solution. The same specimen geometry was used by Ljungberg et al. (1988), in conjunctionwith interrupted tests conducted at very slow extension rates (nominally 10-'/sec), to study the IGSCCof several alloys in high-temperature aqueous environments. The purpose, in this case, was to inducegrowth of cracks during a well defined and relatively short time span in order to calculate an averagecrack growth rate from fractographic measurements once the specimen was brought to failure in air, afterbeing removed from the testing environment. Scully and Adepoju (1978) have measured crack growthrates optically as a function of extension rate for titanium alloys in chloride-containing solutions usingflat, one-sided notched specimens. Parkins (1979) has used precracked beam specimens of carbon steelin a CO3

2 /HCO; solution, and observed that under constant deflection rate conditions a thresholddeflection rate exists above which the crack growth rate increases with increasing deflection rate.

These observations indicate that crack growth rates as measured in SSRTs are dependent on thestrain rate. In general, the crack growth rate increases with increasing strain rate until a strain rate isreached in which the duration of the test is too short for crack initiation to occur and the specimen failsas a result of ductile fracture. However, it should be noted that only the initial strain rate is well defined.As soon as cracks initiate, and particularly when multiple cracks are simultaneously growing, the simplerelationship between the crosshead speed of the tensile machine and the strain rate on smooth, plainspecimens is no longer valid. On the other hand, for notched specimens, the strain rate at the notch tipis many times higher than the nominal strain rate. Therefore, the variable usually reported for notchedspecimens is the crosshead speed.

2-4

Variations in the strain rate affect the fracture morphology also. Several authors have observedtransitions from intergranular to transgranular or vice-versa by altering the strain rate without changingany other variable.

In summary, it can be concluded that the SSRT method is an excellent technique for definingenvironmental conditions conducive to SCC and for comparing the relative susceptibilities to SCC ofdifferent materials or different thermomechanical treatments for a given material. However, theapplication of crack growth rate data obtained with this technique to the prediction of actual in-servicebehavior can be considered unreliable. Nevertheless, the technique could be extremely useful for definingcritical potentials as bounding parameters for SCC once the evolution of the plausible environmentsexpected at the repository site is established.

2.2.2 Constant Deflection Test

In constant deflection (or strain) tests, stressing is accomplished by bending a beam machinedfrom a sheet, plate, or bar using an appropriate jig. The simplest type of smooth, plain specimen forSCC tests is the U-bend. It consists of a strip specimen bent to approximately 1800 around apredetermined radius and maintained in a plastically deformed condition during testing. The stress ofinterest is the circumferential stress in the outer fiber, which is not uniform across the width of thespecimen. The applied strain can be calculated with a simple equation

e= /2r fort<<r (2-1)

where t is the specimen thickness and r is the radius of curvature. An approximate value of the outerfiber stress can be obtained from the stress-strain curve of the test material by using the value of straingiven by Eq. (2-1). Details of the test can be found in a ASTM standard (ASTM, 1991a). Sedriks(1989) has discussed the advantages and limitations of this simple test method.

In recent years, the double U-bend specimen has been extensively used for SCC testing in high-temperature aqueous environments typical of the nuclear power industry. This type of specimen consistsof two U-bend members held together forming a crevice between the inner and outer members. Withthis configuration, it is relatively easy to initiate cracking in the outer surface of the inner U-bend as aresult of the build-up of a localized and more aggressive environment in the crevice.

The simplicity of this type of specimen, as well as its low cost, makes this test methodappropriate for the simultaneous evaluation of different materials or the same material at different appliedpotentials in a given environment, using a single test vessel.

As shown in Figure 2-1, the U-bend configuration has also been adopted in the "wick test"(Dana, 1956). This test was developed to simulate a wet/dry cycle, in which aggressive species presentin the solution at low concentrations (i.e., chloride ions) are transported through the glass wool andconcentrated in the outer surface of an electrically heated U-bend specimen. Although the actual chlorideconcentration is not measured, it is useful as an accelerated test to simulate conditions present in heat-transfer surfaces.

In other types of constant deflection tests, the strain applied is maintained within the elasticrange. For this purpose, bent beams are used in which the stress of interest is the longitudinal tensile

2-5

Electric

Insulation---

C,t

(

X WI~~

Specimen

Solution

Figure 2-1. Specimen configuration for the wick-test (Dana, 1956)

stress on the convex surface. It is obvious that a stress gradient exists through the thickness, but also thelongitudinal stress can vary across the width of the beam depending on the width-to-thickness ratio.There are expressions for the stress for 2-, 3-, and 4-point bending in which the center-to-edge stressvariations are not considered. All these configurations are extensively used for SCC testing, usuallyfollowing ASTM standards (ASTM, 199 lb). One of the limitations is that crevice corrosion may developin an uncontrolled manner in the point of contact between the specimen and the holder. Galvaniccorrosion as a result of contact with a different metallic material should be avoided.

A variation of the previous approach is the crevice bent-beam test (Akashi, 1988), in which afully supported bent-beam specimen containing a carbon fiber packing between the specimen and theholder is used in order to create a well-defined crevice. This type of specimen has been widely used toaccelerate the initiation of cracks in alloys tested in high-temperature aqueous environments. It appears

2-6

that initiation and growth of cracks are accelerated by the combined effect of the crevice geometry andthe galvanic coupling with the carbon fiber surface.

These test methods can be used in conjunction with monitoring techniques, such as acousticemission or electrochemical noise, to determine the initiation time for cracks. Eventually, it will benecessary to evaluate the accuracy and reproducibility of these monitoring techniques to follow theinitiation and propagation of cracks. While acoustic emission is a well-established technique for SCCmonitoring, at least in some alloy/environment systems in which mechanical cleavage events occur, theviability of electrochemical noise is still a matter of controversy. The application of the "wick-test' toreproduce the wet/dry cycle expected under certain conditions in the repository environment seemspossible, probably in combination with a better defined crevice configuration, as in the crevice bent-beamspecimen. A method should be developed for measurement of the aggressive anion concentration on themetal surface.

Although there are other constant deflection tests, based in the use of C-rings, O-rings andcapsules, these are not considered appropriate for the testing required in the evaluation of materials forhigh-level waste (HLW) containers. They are commonly used for testing other product forms, such astubing.

In summary, both U-bends and bent-beam specimens could be useful for testing candidatecontainer materials. Since they are relatively inexpensive and easy to contain in a testing vessel orautoclave, a large number of specimens can be tested simultaneously. They are suitable for long-termtesting, and different schemes for the periodic removal of specimens for optical inspection can be easilydesigned. Monitoring techniques can be used to determine crack initiation It should be noted that inthese types of specimens the stress relaxes with the initiation and propagation of cracks. Therefore, thesetest methods (with the application of elastic or plastic strain) are useful for evaluating the effect ofresidual stresses.

2.2.3 Constant Load Test

In these type of tests, the specimens are uniaxially loaded by means of a dead weight, generallyusing a lever device. The advantage of this test method is that the tensile stress increases with thepropagation of cracks; hence, there is no possibility of crack arrest as a result of stress relaxation as itmay happen in constant deflection tests. There is a standard practice for preparation of specimens(ASTM, 1991c). The dependent variable generally used to assess the susceptibility to SCC is the timeto failure, which is plotted as a function of the initial applied stress. For many alloy/environmentsystems, a threshold stress (ah) can be determined. In general, this threshold stress is higher than theyield stress (ay) of the material. However, in many cases the ratio ahlay can reach very low values (i.e.,0.1), which is an indication of a material with extreme susceptibility to SCC. The ratio can be used forcomparing different materials on a normalized basis. From the point of view of SCC resistance, a,,, canbe used as a bounding parameter in terms of applied and/or residual stresses for a givenalloy/environment combination. Since the parameter is measured over a limited test period assuming anasymptotic extrapolation, a safety factor should be considered. However, the main limitation of thisapproach is that the presence of flaws, notches, or other minor defects on the metal surface in fabricatedcomponents may enhance the stress level locally, as discussed below, leading to failure even in thepresence of nominally applied stresses lower than the threshold value determined in laboratory tests. Thevalidity for long-term prediction of this extrapolation could be questionable.

2-7

Sometimes, large and cumbersome frames are necessary to attain the applied stresses requiredwith specimens of the size normally tested. An alternative is the use of spring loaded frames. Proofrings have been used extensively to perform constant load tests for detecting susceptibility to crackingin aqueous H2S environments, according to the NACE Standard Tests TM-01-77. Test procedures forthese configurations are described in a standard practice (ASTM, 1991c)

Since in many alloys/environment systems, the initiation time (t.) represents a significant fractionof the time to failure (tf), there is a strong incentive for measuring that time. This is a particularlyimportant issue in the case of container materials, because if a crack is initiated, unless a crack arrestmechanism can be postulated, minimum detectable crack growth rates which are in the order of l102 to10.11 m/s will lead to failure within the containment period. The initiation time can be measured throughinterrupted tests accompanied with optical examination or through a continuous monitoring approach.As noted above for constant deflection tests, electrochemical noise or acoustic emission may be suitabletechniques, but in the case of constant loaded specimens the elongation can also be measured as a functionof time. This approach has been used in the past with limited success through the application of opticalmeasurement techniques or displacement transducers (Wilde, 1969). A higher resolution level usingmodern methods such as video or high resolution digital extensometers (Brinton Ferguson, 1992) canfacilitate the use of this technique.

Measurements of potential in galvanostatic or open circuit tests and current in potentiostatic testshave been used in the past for monitoring crack initiation and growth in constant load tests. Thesemeasurements were traditionally based in recording average potential or current values rather than beingconcerned with transients and fluctuations, as is the case in recent electrochemical noise analyses usingzero resistance ammeters (Stewart, 1990). In any case, without an independent verification of crackadvance through other techniques, indirect electrochemical methods rely heavily on interpreting currentincreases and potential drops as signs of crack propagation. While this is generally valid, crack growthmay not be accompanied by a detectable variation of potential or current in the presence of a largepassive/active area ratio.

Simultaneous monitoring of elongation and potential has been used by Sato et al. (1990) for thestudy of SCC of Type 304 stainless steel in dilute, neutral chloride solutions under heat transferconditions. A schematic of the apparatus, is shown in Figure 2-2. The horizontal set-up of the tensilespecimen allowed the test solution to be dropped onto the specimen surface via the salt bridge. Thespecimen temperature can be controlled from 25 to 200'C, simulating a wet/dry cycle.

In summary, constant load tests are useful for testing candidate container materials because athreshold stress below which SCC does not occur can be measured. However, the risk exists that theselected test time can be shorter than the initiation time, and, hence, the measured threshold stress wouldbe higher than the true value. This limitation and the existence of stress concentration effects in realstructures makes ah not adequate as a bounding parameter for SCC.

2.2.4 Fracture Mechanics Test

Although precracked or notched specimens have been used for a long time in SCC testing, theintroduction of linear elastic fracture mechanics (LEFM) concepts led to a significant advance in termsof obtaining quantitative data for design purposes. Following the classical review of Brown (1968), alarge database has been accumulated for many alloy/environment systems. The basic concept in LEFM

2-8

Specimen -

Electric SolutionHeater

(a) (b)

Figure 2-2. Schematic of the apparatus for constant load SCC test with heatedtensile specimen. (a) Side view; (b) Top view. (Sato et al., 1990)

is that the stress field ahead of a sharp crack can be characterized in terms of the stress intensity factor(K) which is a function of both the nominal stress (a) and the size of the crack (a). In the case of purelytensile loading, the relative movement of the two crack surfaces corresponds to the opening mode (ModeI), and the associated K is defined as K,.

By using precracked specimens of well-defined geometries, crack growth rates (da/dt) can beplotted as a function of K, for different environmental conditions, including the effect of potential. It isbeyond the scope of this review to discuss different specimen geometries, as well as various approachesto this type of testing (i.e., constant load versus constant displacement). The subject is well covered inseveral books and adequately summarized by Sedriks (1989). As discussed in Section 2.2.3,measurement of crack growth rates are not relevant for performance assessment of container materials,but a useful bounding parameter can be defined from da/dt versus K, curves. This is the critical stressintensity factor for SCC (K1,<,), which corresponds to the minimum K1 value below which a crack willnot grow. Use of Ku. is better than the use of oa because the existence of flaws as stress raisers is takeninto account. It should be noted, however, that a high level of resolution in nondestructive examinationof fabricated components is required to identify flaws of small size, so that K1, values can be used withconfidence.

2-9

Although an approach based on crack propagation is not valid as a criterion for performanceof containers, laboratory or field measurements of crack growth rates could be useful to examine crackarrest processes, as related to changes in redox potential or other environmental variations. Formeasurements of K1,,,, it is worthwhile to note that prolonged exposures may be required to initiate crackgrowth. The existence of a notch, usually followed by a fatigue precrack, does not eliminate the initiationtime which, in many cases, seems to be related to the development of appropriate electrochemicalconditions at the crack tip.

For the measurement of crack extension, there are methods based on visual observation andcrack opening displacement (COD) monitoring using a linear voltage displacement transducer (LVDT).In recent years, for measurements in high-temperature, high-pressure autoclave systems, voltage droptechniques have been developed (Prater, 1981; 1984). A DC reversing current signal is used to avoidthe electrochemical polarization of the specimen and thermoelectric effects. For the same application,modified wedge-opening-loading (WOL) specimens loaded by a bolt, internally instrumented with a straingage, have also been used (Chung, 1985a; 1985b).

An additional point should be emphasized. It was originally assumed that LEFM is onlyapplicable to SCC testing of high strength (high ay) materials in order to maintain plane strain conditionsover the width of the specimen. This is a requirement for the validity of the equations expressing K, interms of the crack size and is related to the size of the plastic zone (ry) ahead of the crack tip, as definedby

ry = A (K11a 2(2-2)

where A is a geometric constant. For plane strain conditions, the size of the plastic zone should besignificantly smaller than the specimen thickness. In the presence of an environment that promotes SCC,that condition is generally satisfied even for ductile (low ay) materials because crack growth occurs at lowK, values, well below K1., which is the fracture toughness.

In recent years, elastic-plastic fracture mechanics has been developed to deal with fracture ofductile materials. Attempts have been made to extend this concept to the analysis of SCC, but there isno consensus regarding the advantages of this approach.

In summary, fracture mechanics techniques can be applied for the study of SCC of candidatecontainer materials. However, from a performance assessment perspective, the main objective is toestablish the validity of K,. as a bounding parameter. Although this approach has been extremelysuccessful for dealing with engineered structures in which flaws and defects that act as preexisting cracksare present, it should be noted that an accurate determination of K1. requires extensive testing time.Only when da/dt exhibits a very strong dependence on K, at low K, values is a reliable determination ofK1.,c possible. In practice, however, K1,C is related to the lowest measurable crack growth rate. For adesign wall thickness of 1.25 cm and a minimum life time of 1000 years, the maximum acceptable crackgrowth rate is 4 x 10"l m/s. This extremely low crack growth rate requires extended test time (- 1year) to be able to detect a crack advance of 10 Am, which is a reasonable detection limit.

2-10

3 STRESS CORROSION CRACKING OF Fe-Ni-Cr-Mo ALLOYS

3.1 REVIEW OF THE DOE/NRC HLW RESEARCH

3.1.1 Materials

In his "Annotated History of Container Candidate Material Selection" McCright (1988)described the evolution of not only the candidate materials, but also the conceptual and alternate designsof waste packages for disposal of spent fuel and high-level reprocessed waste. Although the conceptualdesigns are still evolving (Harrison-Giesler, 1991), it is necessary to consider briefly the history of thewaste package design in order to better understand the context within which the experimentalinvestigations have been performed by DOE in the past few years.

In the years between 1981 and 1983, the main interest was in the saturated zone of the YuccaMountain site for the repository horizon; hence, thick, self-shielded containers, made of carbon and low-alloy steels, were the main considerations. Limited testing of Type 304L stainless steel, mainly to studythe effect of gamma radiation, was started in December of 1982 at Pacific Northwest Laboratories (PNL).

By February 1983, the decision was made to study the unsaturated zone as the potentialrepository horizon. In October of 1983, a survey of 31 alloys as alternatives to Type 304L stainless steelwas published (Russell, 1984). This list consisted of a wide variety of alloy classes including stainlesssteels, Ni-base alloys, Ti, Zr, and Cu-base alloys. The list was later shortened to 17 alloys (McCright,1988), while retaining the same diversity of alloy classes. Out of this list, four alloys that obtained thehighest ranking in the combined consideration of cost, corrosion resistance, mechanical properties, andweldability were selected (Types 304L, 316L, and 321 stainless steels, and alloy 825). In mid-1986,testing of Type 321 stainless steel was ended since this material did not offer a great advantage over Type316L stainless steel. Some testing of carbon steel continued at PNL because it was still being consideredas the only candidate for borehole liner. Inclusion of copper alloys was requested early in 1984 at theurging of Congress, and investigations on three copper alloys began in October of 1984. The nominalcompositions of the six candidate alloys are given in Table 3-1.

Initially, the pour canister, made of 304L Type stainless steel, was visualized as the onlycontainer for reprocessed HLW, but, after late 1984, a decision was made to use an overpack. The pourcanister is subjected to thermal cycling during the glass pouring and resistance welding operations thatmay sensitize the material to intergranular corrosion. Because the pour canister was designed to bedirectly exposed to the groundwater environment, the effects of both high-temperature and low-temperature sensitization were examined in great detail initially. In June 1986, a project-wide stop-workorder was issued to better identify quality assurance needs of various aspects of the work, and most ofthe experimental work was terminated by the end of 1986. The draft Site Characterization Plan (SCP)was completed in mid-1986.

The first six alloys shown in Table 3-1 are still considered to be the reference list of alloys(McCright, 1990), and, hence, will be the focus of this report. After the issuing of the SCP, manyalternate materials (e.g., carbon steel, Ti-grade 12, Hastelloy alloy C4, A1203) and conceptual designs(e.g., self-shielded, bimetallic, composite containers) have been proposed (Gdowski, 1991; Harrison-Giesler, 1991). However, no significant work has been performed by DOE in characterizing the stresscorrosion cracking (SCC) resistance of the alternate candidate materials in a tuff repository. Information

3-1

- - -M- - - -- - - - m m M M M M M M

Table 3-1. Nominal chemical composition of current DOE candidate and selected alternate container alloys

Composition, Wt. Percent

Alloy UNS No. C, max. Cr Cu moFe t MoP Ni Others

304L SS S30403 0.03 19.0 Bal. - 10.0 Mn: 2.0 max.S: 0.03 max.P: 0.045 max.

316L SS S31603 0.03 17.0 - Bal. 2.5 12.0 Mn: 2.0 max.S: 0.03 max.P: 0.045 max.

Alloy 825 N08825 0.05 21.5 2.0 29.0 3.0 42.0 Ti: 1.0S: 0.03 max.

_ _ _ _ _ _ _________ ________ ________ M n: 1.0 m ax.

CDA 102 C10200 - - 99.95 min. - - - -

CDA 715 C71500 Bal. 0.7 31 -

CDA 613 C61300 - - Bal. 2.5 - - Al: 6.8__________ _________Sn: 0.35

Alloy C-4 N06455 0.01 16.0 - 3 max. 15.5 Bal. Ti: 0.7 max.

Ti Grade 12 R53400 0.08 - 0.30 0.30 0.75 Ti: Bal.

wt�)

on SCC of some of these alternate materials, mainly available from sources outside the DOE high-levelwaste (HLW) program, will be the focus of a separate report. The nominal compositions of two of theproposed alternate candidate materials are also shown in Table 3-1.

3.1.2 Environmental Considerations

It is well known that water extracted from the J-13 well, located in the vicinity of the YuccaMountain site, has been used as the reference groundwater for experimental studies (Glassley, 1986),assuming that its composition is close to that of the vadose water in the unsaturated zone in the TopopahSpring tuff. However, the J-13 well water has a range of composition for various species (Glassley,1990). As shown in Table 3-2, J-13 water is a neutral pH water in which the prevailing ionic speciesare HCO3 and Na'. Other anions, such as Cl-, F, S04

2 and NO;, are present at lower concentrations.The silicon content of the water, in the form of silicic acid, is relatively high. Waters from other wellsin the vicinity of Yucca Mountain have similar chemical composition, but differences in pH and ionconcentration ratios are considered to be significant (Kerrisk, 1987).

Table 3-2. Chemical composition ranges of J-13 well water (Glassley, 1990)

Species mg/L mMoles/L l

Li+ 0.04-0.17 0.006-0.024

Na' 42-50 1.83-2.17

| K+ 3.7-6.6 0.10-0.17

Mg2 + 1.7-2.5 0.07-0.10

Ca2+ 11.5-15.0 0.29-0.37

Sr2+ 0.02-0.1 0.00024.001

Fe3' < 0.01-0.16 < 0.0002-0.003

Al3+ 0.008-0.11 0.0003-0.004

Si(SiO 2) 26.6-31.9 0.95-1.14

NO3- 6.8-10.1 0.113-0.168

F 1.7-2.7 0.029-0.135

Cy- 6.3-8.4 0.1784.237

HCO; 118-143 1.93-2.34

S042- 17-21 0.18-0.22

pH 6.8-8.3

3-3

Of greater importance to SCC and, indeed, to all localized corrosion processes, is the evolutionof the environmental composition as a result of the emplacement of the waste packages in the repository.These changes in the environmental composition can occur by the following processes:

* Rock-water interactions over long-periods of time and modified by the increase intemperature:

Modeling efforts in this area (Murphy, 1990) have shown that as the temperature increases,the bicarbonate concentration will decrease due to CO2 volatilization resulting in aconcomitant increase in pH. In many of the reported SCC experiments, J-13 waterconditioned by contact with crushed tuff at the test temperature has been used. It can bepresumed, however, that the concentration should be extremely dependent on themineralogical composition of the rock and the temperature and also, unless equilibriumconditions were rapidly established, on the duration of the conditioning process and thevolume ratio of the rock to the water. The evolution of J-13 water in contact with crushedtuff at 50'C and 90'C and under -y-irradiation is shown in Table 3-3 (Westerman, 1987).In Westerman's experiments, the autoclaves were continuously purged with air and wereopen to the atmosphere via 9 meters of outlet tubing. After each period of testing, as shownin the first column of Table 3-3, the solution was replaced with fresh J-13 water. Thecrushed rock was replaced after each test period up to 10 months and not thereafter.Although some of the changes in the concentration of species such as NO2- and NO3- maybe the result of radiolysis, changes in the concentration of other species such as C[, HCO3-;and SO42- are most probably due to interaction of the solution with rock and specimens.Unfortunately, no systematic variations in the concentrations of any species is discernible.Since no pretreatment of the rocks is reported, it is not known whether the increase in someionic species originated from rock heterogeneities or incidental impurities in the rock. ThepH at the end of each period is higher than the pH of 7.1 to 7.6 measured for the fresh J-13water, in agreement with the model predictions. The Cl- and SO,2- concentrations aregenerally much higher than found in the initial J-13 water that was used by the investigators.These results are corroborated by the measurement of high conductivity at the end of eachperiod. Westerman et al. (1987) report that the measurements of high ionic concentrationsare reproducible.

A further confirmation of the increase in concentration of ionic species due to rock-waterinteractions has been presented by Abraham et al. (1986). At boiling temperatures (100TC),it has been shown (Abraham, 1986) that the concentration of several anions and cations, suchas SO4

2', NO3;, Cl-, NaI, K+ and Ca2', increases significantly with time when synthetic J-13water was heated in the presence of crushed tuff, as shown in Figures 3-1 and 3-2. Detailedsolution analyses were carried out after various time periods of testing. Interestingly,reaction of boiling distilled water with tuff produced an initial increase in chloride to about160 ppm, which, upon further treatment with new distilled water, decreased to a steadyvalue of about 20 ppm, which is more than three times the nominal concentration reportedfor J-13 water.

In contrast to the above experiments, after reviewing rock-water interaction experimentsperformed at Lawrence Livermore National Laboratory (LLNL) (Knauss, 1985; Oversby,1985), Glassley (1986) indicated that no significant changes from the initial J-13 values werefound for chloride, fluoride, nitrate, and sulfate. These experiments were conducted in

3-4

MM mMM M" - M M -M M M - -

Table 3-3. Analysis of solutions from U-bend stress corrosion cracking tests in aerated J-13 water under -y-irradiation.Fresh J-13 water was added at the end of each period. (Westerman, 1987)

Concentration, mg/LT _me, _Cnd., IOS04CCa NaISiINH 3months pH /AS/CM | Cl NO, O HCO, 0 aN |S ~

50'C Test

3 7.7 780 2.2 51 5 42 190 100 42 110 25 -

5 7.3 5400 7.0 420 570 670 210 970 530 590 47 4.77 7.7 6800 7.2 470 680 770 180 1130 560 660 40 1110 7.7 2520 2.5 119 106 201 142 802 278 218 37 2.016 7.8 1926 <0.1 110 24 189 100 510 140 200 26 0.4

24 8.1 2380 <0.1 37 37 323 34.3 806 210 275 37 0.03

25 8.1 1000 <0.1 25 < 1 125 56.3 219 98 97 28 1

900C Test I

3 9.2 1240 1.5 140 2 1 24 300 23 250 66 -5 8.2 1085 0.6 14 11 23 65 88 44 27 28 1.4

7 8.0 8450 12 600 1700 110 230 1400 820 950 58 0.3

10 8.3 123 0.2 0.3 3 10 34 8 22 3 22 -

14 8.9 115 0.2 0.6 3 7 33 12 14 7 34 <0.1

23 9.0 601 <0.1 27 43 <1 47 138 18 101 28 0.04

W.

(!f

900

800 -

700-

600

500

200-

3 6 9 12

Test Time (months)

Figure 3-1. Concentration changes in J-13 water during reaction with crushed tuffand stainless steels corrosion samples at 100'C. Data at various time periods arefrom independent test cells. (Abraham, 1986)

essentially closed systems. In the results reported by Oversby (1985), the pH increased toabout 9 after 70 days at 150TC, whereas, a slight decrease in the dissolved bicarbonatecontent was observed. This was attributed to the gradual exsolution of CO2 from thesolution through the pores in the polytetrafluoroethylene (PTFE) liner. Knauss et al. (1985)used an impermeable gold-bag; and, hence, did not observe any increase in pH during thecourse of their experiments. A slight decrease in the pH (measured at 250C) was noted andattributed to the precipitation of calcium and magnesium carbonates which have retrogradesolubility. The most significant compositional change was due to the dissolution of siliconfrom the tuff.

* Changes due to occluded regions such as crevices:

It has been well established that concentration of anionic species, such as chloride, canincrease greatly in crevice regions (Turnbull, 1983; Alavi, 1987; Luo, 1992), while pH candecrease due to hydrolysis of cationic species such as Fe2+, Cri+, and Ni2+. Additionally,

3-6

(3~~~~~~~~~~~~0

200W3 6 9 12

Test Time (months)

Figure 3-2. Concentration changes in lOX J-13 water (solution where all species arepresent to ten times their concentration in J-13 water) during reaction with crushedtuff and stainless steels corrosion samples at 100'C. Data at various time periods arefrom independent test cells. (Abraham, 1986)

highly reducing conditions may be created in occluded regions due to the rapid depletion ofoxygen and slow rediffusion from the bulk. The radiolysis reactions in acidic, reducingsolutions may be significantly different from those in the bulk oxidizing solutions (Spinks,1990).

Changes due to repeated/episodic evaporation and re-wetting of container:

This process has been poorly characterized. Experimental measurements of Beavers et al.(1992b) have shown increases in anionic concentrations proportional to the number ofboildown and refill cycles. However, the relationship of these procedures to actualconditions in the repository is not clear. Westerman et al. (1987) performed boildown testsin autoclaves at 200(C and 1000 psig with seven days in liquid and one day under 'dry'conditions. The "dry' period was achieved by reducing the pressure. The reportedconcentration of chloride and sulfate after 15 and 50 boildown cycles were higher than initial

3-7

values, but not high enough in proportion to the number of boildown cycles if completeevaporation is assumed.

Changes due to radiolysis:

The gamma dose rate from spent fuel containers is expected to be as high as 0.1 Mrad/hrand, typically, about 0.01 Mrad/hr (Reed, 1991). The dose rate from glass waste isexpected to be somewhat less, about 0.005 Mrad/hr (van Konynenburg, 1986). Some of theearly field evidences of the effects of radiolysis on corrosion and SCC were observed in theSpent Fuel Test-Climax Program (Patrick, 1986). Corrosion of the carbon steel linersexposed to spent-fuel containers and cracking of Ni-Fe extensometer connecting rods werenoted. In the latter case, laboratory simulation of SCC of Ni-Fe specimens in groundwaterenvironments was achieved only upon addition of an oxidizing salt, CuCl2. In addition togalvanic effects due to coupling with copper components, it is possible that radiolysis of thegroundwater environment increased the concentration of oxidizing species; thus raising thecorrosion potential high enough to cause corrosion and cracking.

Experimental investigations of radiation chemical and electrochemical effects have beenreported by van Konynenburg (van Konynenburg, 1986; Glass, 1986; Kim, 1987; Reed,1990, 1991). Van Konynenburg (1986) interpreted the experimental results related toreaction of various types of glasses with water and water reacted with tuff. The majorspecies present as a result of y-radiation of water in contact with air were nitrate and nitrite.The ratio of nitrite to nitrate was dictated by the presence of catalytic surfaces such as tuffand ionic species in the water such as bicarbonate. In the absence of tuff rock and ionicspecies in water, the nitrite/nitrate ratio was close to 1. This is presumably due to theoxidation of nitrite to nitrate by the hydrogen peroxide that is formed in the water as a resultof radiolysis. In the presence of catalysts that promote decomposition of hydrogen peroxide,the nitrite concentration increased. While the hydrogen peroxide concentration was notreported, it was speculated that it would have been quite low. Reed et al. (1990, 1991)reported that in moist-air systems (at relative humidities of 15 percent at 90'C), nitrites andnitrates were the main species as a result of radiolysis. Copper specimens under theseconditions formed hydrated cupric nitrates. In the 100 percent relative humidityenvironment, only Cu20 and CuO were found on copper specimens. Irradiation experimentsby Yunker (1990) on various copper alloys under 100 percent relative humidity conditionsat 950 C have resulted mainly in the formation of CuO. Yunker (1990) reported estimatedconcentrations of nitrite in the gas phase of moist air mixtures exposed to irradiation but didnot report any observation of nitrate corrosion products on copper specimens. ColloidalFe(III) compounds, thought to be originated from the reaction of stainless steel vessels withthe moist environment, have also been reported in these investigations (van Konynenburg,1986). In situations where the aqueous phase is predominant (Glass, 1986) or where theaqueous phase is in contact with an inert gas such as argon (Kim, 1987), formation of H202and °2 in the solution are more likely, although only indirect evidence for this has beenprovided (Spinks, 1990). Unfortunately, the effect of radiolysis on environments withinoccluded regions, such as crevices and cracks, have not been studied systematically.

3-8

3.1.3 Fe-Ni-Cr-Mo Alloys

3.1.3.1 Slow Strain Rate Tests

The literature up to 1986 on SCC pertaining to tuff repository environment has been reviewedby Beavers and Thompson (1990a). However, some of this material is included in the current reviewfor completeness and for comparison to later information.

Early studies, using slow strain rate testing, were conducted on annealed and annealed plussensitized Types 304, 304L, and 316L stainless steels (McCright, 1983; Juhas, 1984; Pitman, 1986)under unirradiated conditions. The sensitization treatment was different for the three alloys. Sensitizationwas characterized by using the electrochemical potentiokinetic reactivation (EPR) technique. The EPRtest results indicated that the sensitization treatments adopted for the SCC tests were quite severe forTypes 304 and 304L stainless steels. The sensitization treatment adopted for Type 316L stainless steelwas at too low a temperature (250'C), and, hence, did not induce any sensitization. The tests wereconducted at 150'C in a pressurized autoclave using J-13 water that was sparged with air in a separatevessel and then continuously pumped through the test vessel. The bottom of the test vessels containedcrushed tuff which was contacted by the test solution before contacting the specimens. SCC was observedonly on solution-annealed plus sensitized specimens of Type 304 stainless steel. Cracking wasintergranular. SCC was observed only on the specimens strained at the nominal strain rates of 1 X 1Pand 2 x 10' per second. Higher strain rates did not induce cracking, and cracking was more severe atthe lowest strain rate as expected. No SCC was observed on Type 304L stainless steel regardless of themetallurgical treatment or strain rate. Further slow strain rate testing on Types 304 and 316L stainlesssteels were conducted at 950C in the same environment (Westerman, 1987). No SCC was observed.However, no conclusions regarding temperature effects can be made from these results because thesensitization treatments used in these tests did not produce any sensitization on either alloy.

Slow strain rate tests (SSRTs) were performed on Type 304L stainless steel in the mill-annealedcondition by Beavers et al. (1989b). The results are summarized in Tables 3-4 and 3-5. No SCC wasobserved in this alloy at 90'C. It can also be seen that acidification (using C0 2), increasing the corrosionpotential (using H202), applying an anodic potential, or increasing the chloride concentration to 10,000ppm did not have any effect on the SCC susceptibility of the alloy. The use of aeration or deaeration inthe 10,000 ppm chloride solutions was not reported. In the 1000 ppm chloride solution at an appliedanodic potential of 200 mVsce, drops in ductility and time to failure were noted. However, this wasattributed to the effects of pitting rather than SCC. Based on other published data (Mancia, 1988), theauthors speculated that the strain rate adopted (1 x 0l per second) may have been too high. Becauseof the lack of observed SCC on Type 304L stainless steel, an alloy quite susceptible to SCC even indilute chloride environments, these investigators abandoned further slow strain rate testing on othercandidate alloys which are expected to be more resistant to cracking.

A single SSRT at a nominal strain rate of 1 x 10' per second was performed by Beavers et al.(1987) on a Type 304L stainless steel specimen in simulated J-13 water. The temperature of the solutionwas 175°C, and the specimen was heated from the inside by a bayonette heater. The heat output of thebayonette heater was not reported. The solution was ultrasonically atomized under partially aeratedconditions. No SCC was observed.

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Table 34. Summary of slow strain rate tests on Type 304L stainless steel at 90'C invarious environments. Nominal strain rate: 1 x 104/sec. Solutions were exposed toambient air unless otherwise specified. (Beavers, 1989b)

iE _Time tol E.w, Failure, Elong., RA,

Environment Phase mVSCE Hours % %Oil L _ 61 20.3 74.6

Oil L - 63 20.5 71.3

Simulated J-13 L -136 73 24.1 74.6Simulated J-13 L -113 71 23.7 72.2Simulated J-13 V - 73 24.3 73.8

Simulated J-13 + C02(a) L - 73 24.2 74.6

Simulated J-13 + CO2 L -55 75 26.8 76.2Simulated J-13 + CO2 V - 80 25.9 74.6Simulated J-13 + CO2 V - 80 27.6 75.4

Simulated J-13 without HCO3; L 32 81 27.0 73.8J-13 without HCO3-, + H202(-) L 450 81 26.5 71.2

LV(a)(b)

- Liquid- Vapor- Pure CO2 sparged through solution at 10 ml/min.- 15%H202 dripped in cell at 0.02 ml/hour.

Additional SSRTs at nominal strain rates of 104 and 10-7 per second were carried out by Beaversand Durr (1992c). These indicate that on Type 304L stainless steel, transgranular stress corrosioncracking (TGSCC) was observed only at a strain rate of 10-7 per second when H202 was added to 1000ppm chloride solution. At the higher strain rate, TGSCC was observed in the 100,000 ppm chloridesolution without H202 and 10,000 ppm chloride solution to which H202 was added.

In summary, SCC in the SSRTs has been observed only on severely sensitized Type 304stainless steel in aerated J-13 water at 150°C, and at nominal strain rates of 1 x 104 per second, orlower. Cracking has not been observed on Type 304L stainless steels in the unsensitized condition atchloride levels as high as 10,000 ppm. In contrast, there is ample evidence of SCC of Type 304Lstainless steel in service at much lower chloride concentrations as discussed in Section 3.2. Presence ofH202 increased the susceptibility to TGSCC of Type 304L stainless steel in environments containing10,000 ppm chloride.

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Table 3-5. Summary of slow strain rate tests on Type 304L stainless steel at 90'C invarious environments. Nominal strain rate: 1 x 10 4/sec. Specimens were shorter andsmaller diameter than in Table 34. Solutions were exposedotherwise specified. (Beavers, 1989b)

to ambient air unless

T Time toI E. Failure Elong., RA,

l ~~~Environment Phase MVSMz Hours % %

Oil L _ 106 40.0 77.2

Oil L - 106 38.9 77.0

Simulated J-13 without HCO; L -92 108 36.6 75.5

J-13 without HCO3-, + H202(b) L +337 109 36.6 78.1

Simulated J-13+ 1000 ppm Cl L +200(applied) 88 41.6 68.4

Simulated J-13+ 1000 ppm Cl L +200(applied) 82 28.2 66.9

Simulated J-13+ 1000 ppm Cl L +5OLA/cm 2l

(applied) 106 38.4 77.2

Simulated J-13 + 1000 ppm Cl L +50A/cm2l

(applied) 112 38.8 74.6

Simulated J-13+CO2 +H2 02(b) V _ 122 41.4 78.7

Simulated J-13 +CO2+H2 02(b V _ 128 42.8 78.9

Simulated J-13 + 10,000 ppm Cl W L _ 102 36.2 84.0

Simulated J-13 + 10,000 ppm Cl- (W) L _ 102 32.2 76.1

LV(a)(b)(c)

LiquidVaporPure CO2 sparged through solution at 10 ml/min.15% H202 dripped in cell at 0.02 ml/hour100°C

Maiya et al. (1990) investigated the SCC susceptibility of mill-annealed Types 304L and 316Lstainless steels and alloy 825 in simulated J-13 water at 93°C. The environment was aerated with amixture of 12 % CO2 - 20% 02 - 68% N2 at a slight excess pressure of 5 psig. Interrupted SSRTs wereused where the specimens were strained in simulated J-13 solution at nominal strain rates of 1IV and 10'per second up to a predetermined strain, then pulled to failure in liquid nitrogen. The specimens werecylindrical with two through-holes in the gage section transverse to the stressing direction. In some

3-11

specimens, pins of the same alloy were inserted through the holes to create a tight crevice. The holeswere located such that one was immersed in the solution at all times while the other hole was immersedin the solution part of the time, creating an alternating wet and dry condition. At the end of the tests,the specimens were sectioned longitudinally and examined for cracks in the holes through a scanningelectron microscope (SEM) and optical metallograph.

All the alloys showed micro-cracks which were, at most, 10 to 20 Um deep. The authorsindicated that the cracks on the Type 304L stainless steel were more open than on others. No attemptwas made to determine the mode of cracking. No difference was noted between the creviced andnoncreviced specimens, and between completely immersed and wet-dry conditions. No significantdifference was noted between the high and low strain rates. In order to quantify the crackingsusceptibility, the authors calculated the stress and strain ratios defined as:

Stress Ratio, SR = V'. - (C /3) (3-2)O~'- (ay / 3))

where or. and j,, denote the stresses in the environment and air, respectively, at a given plastic strain,e. The yield strength (a,) of the alloys is measured in air at 230 C, at a strain rate of 104 per second.The stress ratio was calculated at various nominal plastic strains. The strain ratio was calculated atvarious nominal stress values as shown below:

Strain Ratio = (3-3)

The calculated strain ratio was then plotted as a function of plastic strain at the various stress values usingthe stress-strain curves for air. These ratios are shown in Figure 3-3(a) and (b) for the various alloystested. From these figures, it is apparent that Types 304L and 316L stainless steels exhibited a greaterextent of SCC (lower stress and strain ratios) than alloy 825 and the Cu-base alloys. However, thisconclusion must be treated with some caution for the following reasons:

* SEM and optical metallography revealed that all alloys exhibited cracking and that crackingwas quite minor. The evidence for these microcracks to be stress corrosion cracks is notconclusive as no determination of the mode of cracking was made. The authors did notreport any microscopic observation of the specimen tested in air.

* The stress and strain ratios are determined by a variety of factors in addition to SCC. Theauthors mention that the tests in air were performed at a different strain rate than those insimulated J-13 solution. They did not report the temperature of the air tests. In the caseof austenitic stainless steels such as Type 304L, plastic deformation will introduce formationof e and a martensite at various strain levels. The deformation-induced phase transformationdecreases at higher temperatures and is not observed at temperatures above the Mdtemperature. The Mo, temperature (temperature at which 50 percent of the austenite istransformed to martensite at a true strain of 30 percent) varies from about 8 to 30'C forType 304 stainless steel (Hannula, 1984). Thus, if the air test was performed at a lower

3-12

1.4

1.2

1.2la

U)a)

ci,

0.8

0.6

'L I I I I i I

- 'At 4 -A A A la+ ^-* * xOxOXO I 4f ft x'xOXI O X x -

0 0 00 0 0 00 00000 -

- Cu-7AI

_ X Cu

- + Cu-30Ni -_ 0 Incooy 825_0 316L_ - 304L

I

0.4

0.2

00.5 1 1.5 2 2.5 3 3.5 4

Plastic Strain (%)(a)

1.4

1.2

1

I0.8

0.6

0.4 5

0.2 -

0.5 1 1.5 2 2.5 3 3.5 4

Plastic Strain (%/)(b)

Figure 3-3. Relative cracking susceptibility of six candidate container materials insimulated J-13 well water at 93 0C. Nominal strain rate: 1 x 10O per second.(Maiya et al., 1990)

temperature and faster strain rates, more martensite is likely to have formed resulting ingreater stress for a given plastic strain. This may explain the lower stress and strain ratiosfor Type 304L stainless steel. The stability of the austenite phase increases in the order:Type 304L < Type 316L < Alloy 825. Type 316L stainless steel is more stable than Type304L stainless, but has been shown to form deformation-induced martensite at lowtemperatures or high strains. Alloy 825, being a Ni-base alloy, does not undergo martensitictransformation upon plastic deformation. The copper-base alloys have a stable face-centeredcubic structure, hence, they are not expected to show a significant dependance ofdeformation behavior with temperature at these low temperatures.

3-13

3.1.3.2 Constant Deflection Tests

Results of various constant deflection tests have been compiled in Table 3-6 (Juhas, 1984;Pitman, 1986; Westerman, 1987). Some of these tests were conducted in the presence of -y-radiation,generated by a Co-60 source. Unfortunately, several factors conspire to cloud any attempt to interpretthese results purely in terms of the effect of irradiation:

* The test techniques for the irradiated (U-bend) and the unirradiated (4-point bend) weredifferent. The U-bend tests were likely to have produced greater strains as well as straingradients. Additionally, the U-bend tests were conducted for a longer test period with thestrain being increased by an unknown amount after each time period (shown in Table 3-3).The orientation of the principal stress with respect to the rolling direction was not specifiedin either test program.

* Equally important are the differences in the metallurgical condition between the irradiatedand unirradiated specimens. The unirradiated specimens were either cold-worked, welded,and then aged at 700'C, or cold-worked, aged at 700'C, and then welded. The irradiatedspecimens were in the solution-annealed and solution-annealed plus aged (at 6000 C)condition. While it is likely that both types of test specimens in the case of Type 304stainless steel were severely sensitized, the cold-worked materials probably exhibited agreater degree of sensitization (Cihal, 1984; Advani, 1991).

* The test temperatures used in the irradiation tests were lower than the unirradiated tests.

The y-irradiation tests showed that sensitized specimens cracked more readily than solution-annealed specimens. Types 304 and 304L stainless steels exhibited similar resistance to SCC in thesolution annealed condition, but the low-carbon alloy was better in the solution annealed plus sensitizedcondition. In the solution annealed condition, only transgranular cracking was noted in both 304 and304L specimens. As expected, Type 304 stainless steel exhibited intergranular cracking in the sensitizedcondition. However, a "sensitized" 304L specimen exhibited predominantly transgranular cracking,probably due to the decrease in Cr-depletion as a result of overaging. This is apparent from the resultsof EPR tests. Unfortunately, only one specimen was examined.

Generally, the susceptibility to cracking was greater in the vapor or vapor plus rock phases thanin the liquid phase, and more incidence of cracking was observed at 90'C than at 50'C. Surface X-rayphoto spectroscopy (XPS) analysis did not indicate any concentration enhancement of chloride in thevapor phase. It is possible that the corrosion potential was higher in the vapor space due to easier accessof oxygen through the thin liquid film formed on the specimens or due to differences in radiolytic speciesbetween liquid and vapor phases. As a result of radiolysis, NO; and NOi species can be expected inthe moist vapor phases; whereas, H202 may be present in the liquid phase. Nitrate and nitrite mayincrease the corrosion potential only above approximately 10,000 ppm as shown in Table 3-7. On theother hand, nitrate and nitrite have been shown to inhibit localized corrosion considerably (Cragnolino,1991b) and may have a similar effect on SCC. Hydrogen peroxide, on the other hand, can also increasethe corrosion potential considerably (Table 3-4). Hence, it is difficult to conclude, on the basis of thedata in Table 3-6, that radiolysis was an important factor in the greater cracking susceptibility in thevapor space. The changes in solution chemistry during the course of the irradiation tests were discussedin Section 3.1.2 and are shown in Table 3-3. Since the ionic concentrations increased during the courseof the tests, it is difficult to determine the specific environmental factors that lead to cracking in thesetests.

3-14

mm - - - - - - - - -W m W

Table 3-6. Results of constant deflection tests on various austenitic alloys in J-13 water (Westerman, 1987; Juhas, 1984)

W

ItA

No. Cracked/No. Tested

J-13 WaterJ-13 Water ° J-13 Water °a1000C

50°C 90°C No6 x 105 Rad/h 3 x 105 Rad/h Irradiation

| Test Type Alloy Condition RL RV_ _ V _ _RL |_RV | V VU-bend 304 Solution Annealed 0/2 0/2 0/2 0/2 0/2 1/2(23) _ -

U-bend 304 S. A. + 600°C/24 Hrs 1/2(24) 2/2(7,25) 2/2(1,3) 2/2(1,1) 1/2(3) 2/2(3,5)

U-bend 304L Solution Annealed 0/2 0/2 0/2 0/2 0/2 1/2(23) _

U-bend 304L S. A. + 6000C/24 Hrs 0/2 0/2 0/2 0/2 1/2(14) 2/2(10,14) - -

4-point bend 304 CW20% + GMAW,90% Y.S. ° 2pass + 700 C/8hrs - - - - _ - 0/9 0/9) |

4-point bend 304 CW20% + GMAW,90% Y.S. 2pass _ - - - - - 0/3 0/34-point bend 304L CW20% + 700 C/890% Y.S. hrs + GMAW, 2pass - - - - - - 0/9 ) 0/9

4-point bend 304L CW20% + GMAW,90% Y.S. 2pass - - - - - - 0/3 (c) 0/3

4-point bend 316L CW20% + 700 C/890% Y.S. hrs + GMAW, 2pass - - - - - - 0/9 °) 0/9

4-point bend 316L CW20% + GMAW,90% Y.S. 2pass - - - - - - 0/3 (c) 0/3

(continued on next page)

m m M M - - M - - - -M

Table 3-6. Results of constant deflection tests on various austenitic alloys in J-13 water (Westerman, 1987; Juhas, 1984)(Cont'd)

No. Cracked/No. Tested

J-13 WaterJ-13 Water ()J-13 Water ()1000 c

500C 900C No6 x 105 Rad/h 3 x 105 Rad/h Irradiation

Test Type Alloy Condition RL RV_ V RL7 RV _V L V

4-point bend 321 CW20% + 700 C/890% Y.S. hrs + GMAW, 2pass- 0/9 e 0/9

4-point bend 321 CW20% + GMAW,90% Y.S. 2pass 0 _ _ - - 0/3(c) 0/3

WI-

07~ RLRVLV(a)(b)(c)(d)

- Rock + Liquid- Rock + Vapor- Liquid- Vapor- Test Time = 25 months- Test Time = 4016 hours- Test Time = 2000 hours- Y.S. - Yield strength; Stressing and welding directions not reported

( ) - The numbers in parentheses indicate time to failure in months.

Table 3-7. Effect of nitrate concentration on corrosion potential of Type 316Lstainless steel in simulated J-13 Water at 950 C. Solutions were deaerated by argon.(Cragnolino, 1991)

Nitrate Concentration,mMoles/liter E., mVSCE

0.16 -641 + 86.9

16.0 -721 ± 79.8

160.0 -495

Pitman et al. (1986) reported the results of boildown tests in unirradiated conditions onsensitized (presumably from mill-annealed materials) Type 304 and 304L stainless steel specimens. Testswere conducted in 200'C/J-13 water that was initially aerated. Once every seven days, the pressure wasreduced to boil-off water for 24 hours, then refilled by air-saturated J-13 water. In this test, theautoclaves were "filled to the top," so it is difficult to know what the dissolved oxygen content was attemperature. Cracking of sensitized 304 stainless steel was reported after 4 months and 15 boildowns.No failure was observed on 304L specimens even after 50 boildowns (12 months of testing). Asmentioned in Section 3.1.2, the chloride concentration even after 50 boildowns was only 90 ppm, lowerthan anticipated assuming complete drying after each cycle.

Beavers et al. (1987) performed U-bend tests on Type 304L stainless steel specimens. In onetest, the specimens were wrapped around a 1000 watt heater and placed in contact with glass wool, partof which was immersed in J-13 water. The solution was wicked by the glass wool onto the heatedspecimen surface as in the "wick-test" described in Section 2.2.2. No cracking was noted in 403 hoursof exposure. In two other tests, the specimens were heated to 150'C, and J-13 water was dripped ontothem at a rate of 1 drop per 80 seconds. The specimen remained wet for about 50 seconds. In thiswet/dry test, no cracking was observed for up to 3000 hours. It is interesting to compare the result ofthe single wick test to those reported by Warren (1960) on Type 304 stainless steel. In wick tests in a10 ppm chloride solution (corresponding roughly to J-13 water) at 100'C, Warren reported that roughly10 percent of the specimens cracked in less than 500 hours. This indicates that single specimens are notsufficient to study SCC in short-time tests.

Beavers et al. (1992b) reported results of long-term U-bend tests on Type 304L stainless steeland alloy 825 in boildown tests. In these tests, the specimens were immersed in simulated J-13 waterat 90°C while being purged by air in an open-loop system. The rate of purging and condensation ofoutgoing vapors was adjusted such that complete evaporation occurred once a week. Fresh solution wasadded and the tests continued for a total of 13,400 hours. The solution analyses at the end of this timeperiod indicated that the concentration of species such as chloride corresponded to the expected increasethrough total boildown and rewetting cycles. As an example, the solution composition for the boildowntests on alloy 825 is shown in Table 3-8. The corrosion potentials of unstressed specimens exposedsimultaneously are shown in Figure 34. No SCC was observed on any of the specimens at the end of

3-17

Table 3-8. Changes in the concentrations of various speciesboildown tests with simulated J-13 water (Beavers, 1992a)

during the long-term

r Initial Concentration, Final Concentration,Environmental Species mg/i mg/I

Nitrate 12.4 1169

Sulfate 23.5 1975

Chloride 5.5 750

Fluoride 1.8 144

Calcium 10.0 2

Magnesium 1.5 < 1

Sodium 42 4450

Potassium 6.3 380

Silicon (soluble) 27.5 1170

Alkalinity, HCO3 102.0 0

Alkalinity, Phenolphthalein 102.0 4925

Alkalinity, Total 102.0 7150

| pH7.0 ± 0.2 10.68

this time period. It is possible that the precipitation of calcium carbonate on the sample surfacesminimized access of solution to specimen surfaces. The increase in pH also may have assisted inmaintaining the passivation of the alloy.

Additional tests were carried out by Beavers and Durr (1992c) on Type 304L stainless steel andalloy 825 U-bend specimens in simulated J-13 water at 90'C containing 1000, 10,000, and 100,000 ppmof chloride added as NaCl, and 10,000 ppm chloride as CaCl 2. These tests were carried out with andwithout the addition of H202 (200 ppm daily). TGSCC was observed on Type 304L stainless steel in the100,000 ppm chloride solution without H202 and was also observed in the less concentrated chloridesolution if H202 was added. Cracking occurred mainly in the vapor space. No cracking was observedon any of the alloy 825 specimens. Pitting and crevice corrosion were detected on all the specimens ofboth alloys.

Abraham et al. (1986) used notched C-ring tests on seamless tubing (0.75" - 0.84" 0. D. x0.125" - 0.109" Wall) to study stress corrosion. The specimens were tested in the solution-annealed(mill-annealed?) and solution-annealed plus sensitized (600'C/100hours) conditions. The specimens were

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0

-100

E -200

(5

g -300

-400LEGEND

0 - 0 8250 - 0 304L

-500 . . I . . | . | . I . . I . . I . . I . . . . . ,_0 1500 3000 4500 6000 7500 9000 10500 12000 13500 15000

Test Time, Hours

Figure 3-4. Corrosion potentials of Types 304L stainless steel and alloy 825 inlong-term boildown tests In aerated, simulated J-13 water (Beavers, 1992)

stressed to 90 percent of the elastic limit determined on the C-ring specimens in the un-notched condition.Simulated J-13 and loX J-13 solutions were used along with crushed tuff. Specimens were tested bothin the immersed and vapor phase conditions without any intentional aeration. A condenser was usedwithout any trap; and, hence, some steady state aeration of the solution probably occurred. Tests weredone independently for 3, 6, and 12 months. Oxygen concentration was measured both in the steamphase and in the water phase after cooling to room temperature. It was found to be lower than theoxygen partial pressure in air or the expected solubility of oxygen in water, indicating either corrosion-induced consumption or rock-water interaction. Microcracking, ranging from 7&m to 65gLm, wasobserved under SEM on many specimens of Type 304L stainless steel, whereas, alloy 825 exhibited onlyone case of cracking. The grain size of Type 304L stainless steel, shown in the publishedphotomicrographs of Abraham et al. (1986) ranges from 50 to 100rm, although the surface grains appearto be smaller. No trend with respect to material condition or environment was obvious. Cracking wasnot always found at the notch tip. In some specimens, much pitting is evident and some of the cracksmay also have been pits that were expanded or coalesced by applied strain. The differences between theoutcome of these tests and those of Westerman et al. (1987) and Beavers et al. (1992a) may be attributedto the level of observation. Abraham et al. observed the specimens under the SEM at magnifications

3-19

ranging from 1000 to 2000X; whereas, the other investigators examined the specimens under astereoscope, probably at magnifications of about lOOX. The differences in the results can also beattributed to differences in the material conditions (seamless tubing versus plate). Unfortunately, notenough microstructural details have been reported in any of these test programs. The effect of the notchis not obvious since cracking was not always found at the notch tip.

Park et al. (1991) performed low frequency fatigue tests on Type 304L and 316L stainlesssteels, and alloy 825 in simulated J-13 water, at 930C under aerated (02 + N2 + C0 2) conditions at lowflow velocity. Compact tension specimens were side-grooved to prevent crack branching and deviation.The specimens were loaded to a variety of R-values (R = minimum stress intensity factor/maximumstress intensity factor) ranging from 1 to 0.5. They then compared the crack growth rate in the J-13water to crack growth of austenitic stainless steels in air at 930 C predicted by ASME Section XIcorrelation. They concluded that essentially no acceleration of cyclic crack growth occurred due to J-13water, even at frequencies as low as 10-' Hz.

3.1.4 Other U.S. HLW Investigations

In another investigation related to tuff repository, McCright and Weiss (1985) studied the SCCon three grades of steels-Nodular cast iron, 9Cr-lMo steel, and 1020 plain carbon steel. These areproposed to be used as partial borehole liners. Longitudinally welded specimens were tested in the4-point bent beam configuration at an applied stress of 90 percent of the yield strength of the base metal.Specimens were simultaneously exposed to J-13 water at 90'C for up to 9000 hours. The results areshown in Table 3-9. Cracking was intergranular and along the welds. The failures of the 9Cr-lMo steeland the nodular cast iron were attributed to hydrogen embrittlement, due to a combination of cathodicpolarization of the weld through galvanic contact with the base metal and the presence of martensiteformed in the weld as a result of rapid quenching. Weld cracking of a carbon steel (composition was notreported) was also observed in the Spent Fuel-Climax test program (Patrick, 1986), where it was usedas a full liner. Cracks were observed in the circumferential single-bevel, groove welds between thebottom plate and liner wall. However, cracking appears to have initiated from the inside of the linerwhere there was incomplete penetration of the weld. The cause of cracking may not be related tointeractions with the groundwater, but may be due to other factors such as prior hydrogen penetrationduring fabrication or presence of moisture in crevices. In the same test program, cracking ofSuperinvar&1 (Fe-32%Ni-5SCo) connecting rods was experienced and attributed to SCC under galvaniccoupling with copper in contact with groundwater at 45-550 C. Similar morphology of cracks wasreproduced in laboratory tests of this alloy in groundwater solutions containing CuCl2.

3.1.5 European HLW Investigations

The results by various European HLW programs through 1982 have been summarized byAccary (1985). Of interest are the investigations of the French program performed by the Commissariata l'Energie Atomique (CEA), the German program performed by Kernforschungszentrum KarlsruheGmbH (KfK), and the Belgian program performed by the Centre d'Etudes Nucleaires (CEN/SCK), sincethese programs examined SCC of Ni-base and Ti alloys.

1 - Superinvar is a trade name of Imphy S.A.

3-20

Table 3-9. Stress corrosion cracking test results on 4-point bend specimens exposed toJ-13 water at 90'C. Total exposure time: 9000 Hours. (McCright, 1985)

Material I Observation

Nodular Cast Iron 1 out of 9 cracked

9Cr-lMo Steel 6 out of 9 cracked, IG cracking

1020 Plain Carbon Steel No cracking

The French investigation focused on alloys 825, C-276, and 625. Researchers used a simulatedgranitic water that was enriched in chlorides to 534 mMoles/liter and acidified to simulate crevice sites.The tests consisted of SSRTs (strain rate unreported) at 80'C. It was reported that alloy 825 did notsuffer any SCC. Alloy 625 exhibited SCC in the pH range of 1.2 to 3. These results are surprisingbecause alloy 625 with a higher Ni content is expected to be more resistant than alloy 825. It is possiblethat heavy general corrosion of alloy 825 in the acidic environment precluded SCC. Another factor isthe microstructural condition of the alloys which was not reported.

In the German program, U-bend specimens of alloy C-4 were tested along with those ofTi-0.2% Pd alloy in a brine solution consisting of 27%MgCl 2 + 4.7%KCI + 1.4% MgSO 4 + 1.4%NaClat 90 and 170'C for 112 days. No SCC was observed. However, more recent investigations (Smailos,1992) have concentrated more on the corrosion resistance of alloy C-4 weld overlay over carbon steelthan on SCC.

In the Belgian program, U-bend tests were conducted on alloy C-4 at 50'C in a claygroundwater that consisted of 35 mg/liter of chlorides at a pH of 7.4. No cracking was noted in 28,496hours. In comparison, Type 316L stainless steel failed after 1050 hours.

3.2 REVIEW OF THE GENERAL LITERATURE

3.2.1 Effect of Envirommental Factors

One of the major influential factors on SCC cracking of ductile alloys is the physicochemicalproperties of the environment. It has often been stated that the environmental requirements for SCC arehighly specific for a given alloy system. However, the number of environments known to induce SCCin a given alloy has increased substantially during the last three decades and the concept of solutionspecificity is not as firmly based as before. As an example, Table 3-10 shows the environments knownto induce SCC of austenitic stainless steels in 1950, as compared to those reported in 1990. It shouldbe emphasized, however, that SCC occurs in these environments under specific and interrelated conditionsof solution composition, potential, temperature, pH, etc. Also, an important aspect to note is that SCCof mill-annealed stainless steels, occurred mainly in the first four general environments listed in Table3-10. These are acid chlorides, neutral chlorides in the presence of oxygen, concentrated hydroxides,

3-21

Table 3-10. List of liquid environments known to cause stresscorrosion cracking in austenitic stainless steel

Up to | Environments

1950 Acid Chlorides, MgCl2, CaCI2, ZnC12

Neutral Chlorides, NaCi + 02 (or H2 02 )

1990 Hydroxides, OR

Acidified Chlorides, NaCI + H2SO4

Fluorides, F

Bromides, Br,

Water + 02

Thiocyanates, SCN-

Thiosulfates, S2032

Tetrathionates, S4062l

Polythionic Acids, H2SXOl

Sulfurous Acid, SO2 + HSO3l

Hydrogen Sulfide, H2S + 02

and chlorides in the presence of relatively concentrated sulfuric acid. In most of the remainingenvironments of that table, a necessary precondition for SCC to occur is the existence of a Cr-depletedregion along grain boundaries, often a result of alloy sensitization during stress-relief heat treatmentsand/or welding procedures.

As can be inferred from the possible repository environments described in Section 3.1.2, whenconsidering the SCC of austenitic Fe-Cr-Ni alloys, the environments of interest are those containingchloride as the predominant anion. In the proposed repository site at Yucca Mountain, hydroxideconcentrations higher than 10- mole/liter, corresponding to a maximum pH of 10, are not expected underany anticipated conditions. These concentrations are several orders of magnitude lower than those foundunder heat transfer conditions in creviced areas of nuclear steam generators. It is well known that undersuch aggressive conditions, Fe-Cr-Ni alloys including austenitic stainless steels and nickel-base alloys areextremely susceptible to SCC (Macdonald, 1989). The minimum OH- concentration required for SCCof stainless steels seems to be about 2.5 x 10-3 mole/liter (Speidel, unpublished a; Latanision, 1969).Also, as a result of the oxidizing conditions prevailing in the repository, the presence of metastable sulfuroxyanions, such as thiosulfate and tetrathionate, and reduced sulfur compounds, such as hydrogen sulfide,

3-22

which are all strong promoters of intergranular stress corrosion cracking (IGSCC) in sensitized alloys,is excluded.

In the following subsections, specific environmental factors affecting the chloride-induced SCCof Fe-Cr-Ni alloys are discussed separately while taking into consideration their interdependence. Thediscussion is necessarily confined to the austenitic stainless steels, because alloys containing more than25 % Ni appear to be far less susceptible to SCC within the ranges of concentration, temperature, and pHthat are covered in the reviewed studies. The literature on the SCC of Ni-base alloys has been recentlyreviewed (Sridhar, 1992b).

3.2.1.1 Effects of Chloride and pH

Since transgranular cracking of Type 304 stainless steel was discovered in the 1930's, extensivetesting has been performed in boiling, concentrated MgCk2 solutions because in this environment manyservice failures were easily reproduced at the open-circuit potential. Because of the concern with thevalidity of using tests carried out in MgC12 solutions for predicting the behavior of materials in moredilute solutions, such as those frequently found in service applications, in recent years studies have beenconducted in dilute, neutral chloride solutions. Figure 3-5 shows a quite extensive compilation of SCCdata for austenitic stainless steels, including Type 304, 304L, 316, 316L and 347 (Gordon, 1980).Gordon updated previous plots of cracking/no-cracking domains (Williams, 1957; Hfibner, 1974) withmore recent results, in terms of chloride and oxygen concentrations, and for temperatures in the rangeof 250 to 350'C. The importance of oxygen concentration and, hence, potential, on the SCCsusceptibility at a given chloride concentration is clearly seen. The low and intermediate boundariescorrespond to the IGSCC of sensitized material and the TGSCC of nonsensitized material, respectively,in liquid water under conditions typical of nuclear water reactors. The upper boundary, correspondingto TGSCC of nonsensitized material (Williams, 1957), defines the oxygen/chloride combinations abovewhich cracking occurs at 250'C in steam, under intermittent wetting by contact with an alkaline-phosphate treated water. These conditions, in terms of temperature and aqueous phases present, are moreclosely related to those expected in the repository, if liquid water becomes available to the metal surfaceduring the initial part of the containment period. It is seen in Figure 3-5, that at a dissolved oxygenconcentration of 1 ppm, which is typical of a slightly oxidizing environment, TGSCC may occur atchloride concentrations as low as 6 ppm (which is the approximate chloride content of J-13 well water).

At lower temperatures, below the boiling point of water at atmospheric pressure, domains ofcracking/no-cracking for solution annealed Type 304 stainless steel in aerated NaCl solutions, with thepH adjusted by the addition of HCI or NaOH, were determined by Truman (1977) as a function ofchloride concentration. A maximum test period of 565 days was adopted for those specimens that didnot exhibit visible signs of cracking in a shorter time. As summarized in Figure 3-6, TGSCC wasobserved at chloride concentrations above 100 ppm in aerated solutions. The interrelated influence ofchloride concentration, temperature, and pH on TGSCC is apparent from this figure. By decreasing thepH from neutral to pH 2, the cracking/no-cracking boundary is lowered by more than 20'C. No TGSCCwas observed in neutral solutions below 80'C, but the temperature threshold decreased to approximately50'C at pH 2. Within the chloride concentration range used in this study (102 to 105 ppm), Truman didnot find a strong effect of the cationic species used, but a slight increase in the severity of cracking inthe order Na' > Mg2" > Ca2l > Zn2'. This is not usually the case, but it should be noted that inthese tests the initial pH of all solutions was adjusted to 7.

3-23

1000

100- V V

1 TGSCC-Annealed

8

0 BWR c11111 17 D V's 0+v co 1 so

8 0.1 _ IGSCC-Sensitized 0A2 * wtr) 0a a _ \ TGSCC-Annealed 0

r ' f t (Water). P7WR7 fttf ft V

0.01 _ 7 vsce Ar"b* *x

* NOS CC Pautially Sestized0 NO SOC Sapsitized

No NSCCSco

0.001 I I I l -0.01 0.050.1 0.2 0.5 1 10 100 1,000 10,000

Chloride Concentration, ppm

Figure 3-5. Effects of oxygen and chloride concentration on the SCC of austeniticstainless steels in high temperature (250-350'C) water (Gordon, 1980)

As shown in Figure 3-6, the pH of the bulk environment, at least at pHs below 7, has a strongeffect on the SCC susceptibility. The time to failure or the time for crack initiation decreases withdecreasing pH. Asphahani (1980) found that the addition of 1 % H3PO4 to a 4% (- 0.7 mole/liter) NaClsolution reaching a pH -2, leads to TGSCC of Types 304, 304L, 316, and 316L stainless steel at both141 and 820C. Similar results were observed in 0.8% (-0.14 mole/liter) NaCI solution with the additionof 0.2% H 3PO4, 0.5% CH3COOH (acetic acid) or HC1 in a concentration enough to reach pH 2.2 at141 'C. Only in the case of the acetic acid containing solution was the TGSCC accompanied by pittingcorrosion. All these results, even in relatively dilute solutions, explain why chloride salts of cations thatundergo hydrolysis, generating acid solutions, such as Mg2" and Ca2+, are more severe promoters ofTGSCC than neutral chlorides. On the other hand, an alkaline pH in the bulk environment tends toinhibit the initiation and propagation of cracks, leading to improved cracking resistance.

3-24

b.p.

8000

X 60_Q_E

I- No SCC40

AISI 304 Stainless Steel

20 Aerated NaCI SolutionI I l l

102 103 104 105

Cl- Concentrationppm

Figure 3-6. Effects of temperature and chloride concentration on the SCC ofsolution-annealed Type 304 stainless steel in aerated NaCl solutions at various pHs(Truman, 1977)

3.2.1.2 Effect of Temperature

Temperature is one of the most influential variables in the TGSCC of austenitic stainless steel,which generally occurs at temperatures well above room temperature. As noted by Sedriks (1979), thetraditional engineering viewpoint, based on practical experience, is that chloride cracking does not occurat temperatures below 60TC. However, many cases of cracking have been reported at lowertemperatures, close to ambient temperatures, generally under poorly defined conditions. In laboratorytests, it is difficult to detect TGSCC even at temperatures close to 100IC because very long initiationtimes may be involved. No SCC of Type 304 stainless steel was observed by Bernhardsson (1984) attemperatures below 1000C in refreshed solutions containing up to 10' ppm chloride and oxygen contentsranging from 4.6 to 10 ppm with pH varying from 4.5 to 7.1. Tests were conducted using constant strainmethods, but were interrupted after 1000 hours (42 days) when no signs of cracking were detected.Probably the testing time was not prolonged enough to initiate cracks, as a comparison with the data

3-25

shown in Figure 3-6 would suggest. Another tentative explanation is that these tests were conductedunder open circuit potential conditions in slowly flowing solutions, and, therefore, if the corrosionpotential lies outside a likely potential range for cracking, no SCC will occur.

A relevant observation for the behavior of container materials is that under heat transferconditions, particularly in crevices or regions of very limited fluid flow, the concentration of chloride ionscould be many times higher than in the bulk environment. This situation will tend to promote SCC inan environment that otherwise could be considered safe. In many cases, TGSCC has been observed inouter surfaces of stainless steel columns, pipes, and towers exposed to the atmosphere. This is the so-called external SCC (ESCC) and it has been detected frequently under thermal insulation. Precisely thistype of problem led to the development of the wick-test, which was described in Section 2.2.2. Asshown in Figure 3-7, SCC of Type 304 stainless steel can occur at a very low chloride concentration inthe bulk environment (10 ppm), under heat transfer conditions at 100'C (Warren, 1960). The set ofcurves illustrates clearly the broad distribution of failure times, showing also that the median time tofailure is significantly reduced with increasing chloride concentration. Figure 3-8 shows that for achloride concentration of 100 ppm, the median time to failure increases with decreasing temperature, butTGSCC can occur even at 60'C. The median time to failure in this case is close to six months,suggesting that cracking may also occur at lower temperatures if specimens were exposed longer. Thecombined effects of chloride concentration and temperature under heat transfer conditions is clearlyillustrated in Figure 3-9, where it is seen that cracking occurs even at 40'C in the presence of 1800 ppmchloride in the bulk environment. The effect of different cationic species on the cumulative percent ofcracked specimens is shown in Figure 3-10 for 100 ppm chloride solutions at 1000 C. This figure clearlyillustrates the accelerating effect that can be expected under heat transfer conditions in the presence ofacidic chlorides even at a relatively low chloride concentration in the bulk environment. One of thelimitations of these results is that no attempt was made to determine the chloride concentration in contactwith the metal surface. It should be noted, however, that due to the nature of the testing method,significant variations could be expected in different tests, not only in the chloride concentration, but alsoin the corrosion potential of the U-bend specimens. The wide dispersion of failure times for a givenexperimental condition is most likely a result of these variations.

Sato et al. (1990) have analyzed using a Weibull distribution the time to failure data for the SCCof Type 304 stainless steel in 0.5 M and 0.05 M NaCl solutions. The data were obtained in constant loadtests using the apparatus shown in Figure 2-2 (metal surface at 1000C) to simulate a wet/dry cycle.

TGSCC of mill-annealed Type 304 and 316L stainless steel has been observed underatmospheric conditions, even at room temperature, at various degrees of relative humidity in the presenceof deposited chloride films (Shoji, 1989). Films were developed by depositing drops of 0.5 M solutionsof NaCl, CaCI2, ZnCl2 and MgCl2, and simulated sea water on U-bend specimens. The specimens werethen exposed to moist air with various levels of controlled humidity. At room temperature, TGSCC wasobserved only in the presence of the acid chlorides and within a very restrictive range of relative humidity(10 to 30%) after prolonged exposure times (24 months). However, at 50 and 70'C crack initiationoccurred rapidly and crack growth rates were faster. Even at room temperature, crack growth ratesranging from 0.2 to 1.0 mm/year were detected, depending on steel composition, relative humidity, andthe cation of the chloride salt. These results illustrate clearly the risk of cracking for potentiallysusceptible materials even at ambient temperatures.

In recent years, a rash of failures of stressed components made of Type 304 and 316 stainlesssteels in the atmosphere of indoor swimming pools led to additional studies on TGSCC at low

3-26

100

'E80X~ 1 / / 00 ppm Cl-

75 60E

ppm Cl1.E 40

0)Cn

0u < RAISI 304 Stainless SteelSCC-Wick Test

20 NaCI Solution, 1000C

1 Day 1 Week1 Month 6 Month

010 100 1000 10,000

Exposure Time, Hours

Figure 3-7. Effect of chloride concentration on the SCC of solution annealed AISI304 stainless steel In NaCI solutions at 1000C using the wick test. Note thedistribution of failure times for the cracked specimens. (Warren, 1960)

temperatures (Herbsleb, 1989; Oldfield, 1990, 1991; Dillon 1990). In these studies, it was postulatedthat local increases in the chloride concentration coupled with the development of very acidic conditions,through the formation of thin liquid films, are the conditions for cracking to occur. The role of oxidants,such as chlorine or even oxygen, has not been sufficiently investigated.

3.2.1.3 Effect of Radiation

Unfortunately, there have been no studies on the effect of y-radiation on the TGSCC of mill-annealed stainless steels under conditions similar to those expected in the repository with the exceptionof the experimental results reported in Section 3.1.3.2. There are, however, several publications (Fujita,1981; Saito, 1990; Nakata, 1992), dealing with the IGSCC of sensitized Type 304 stainless steel underconditions typical of boiling water reactors (BWRs). It was found (Fujita, 1981; Saito, 1990) thaty-radiation produced by a Co-60 source enhanced the susceptibility to IGSCC under open circuit

3-27

100

AISI 304 Stainless SteelSCC Wick-Test100 ppm or-

'E80

a. L °1 OOI

10 100 0

CD~ ~ ~ ~~~~~~~8 0

~~~~~ 60 ~ ~ ~ ~ ~ ~ ~ ~ 60


Figure 3-8. Effect of temperature on the SCC of solution annealed AISI 304stainless steel in 100 ppm Cl solution using the wick test (Warren, 1960)

conditions at 288°C in the presence of relatively high oxygen contents (0.2 to 8 ppm), as evaluated bySSRTs. On the contrary, under reducing conditions due to the addition of hydrogen to the water,irradiation suppresses the tendency to IGSCC. This effect was found to be particularly evident in thepresence of Na2SO4, an impurity that enhances the susceptibility to IGSCC with respect to that in purewater. These results were interpreted in terms of the effect of -y-radiation on the corrosion potential ofthe alloy, because the potential increases in the oxidizing environment as a result of the formation of H202while it decreases in the reducing environment.

Similar effects were observed more recently (Nakata, 1992) in crack propagation studies usingcompact tension specimens. Crack growth rates were accelerated by 'y-radiation under oxidizingconditions but were substantially reduced by hydrogen injection into the water in the presence of--radiation. It is obvious that these results are not directly applicable to the container materials and

conditions expected in the repository, but illustrate clearly the complex interdependence of radiation andthe chemical composition of the environment and their joint effect on the corrosion potential.

3-28

o. 8 100 ppm Cl0

Ii 4000 0,60 1000 ppm0 10,00E~~~~0 1800mpmCo

0

E 40

CU

20

1 Day 1Week 1 Month 6 Month

10 100 1000 10,000


Figure 3-9. Combined effects of temperature and chloride concentration on theSCC of solution annealed AISI 304 stainless steel using the wick test (Warren,1960)

3.2.1.4 Effect of Potential

The effect of potential on the TGSCC of Types 304 and 316 stainless steels has been studiedextensively at temperatures ranging from 80 to 140'C, but mainly in highly concentrated, acid chloridesolutions such as MgC12 and CaC12. More recently, studies have been conducted using Type 304 stainlesssteel in concentrated (2.5 to 11.8 mole/liter) LiCO solutions in which the pH is close to 7.0 (Shamakian,1979; Shamakian, 1980; Duffo, 1988). In these solutions, TGSCC occurs above a critical potentialwhich, at 105'C, decreases from -0.13 to -0.17 VSH with an increase in the chloride concentration from2.5 to 11.8 mole/liter (Shamakian, 1980). Since the corrosion potential is almost constant in thatconcentration range, reaching a value of approximately -0.11 V.. after 20 hours exposure, TGSCCoccurs with increasing severity and, therefore, at shorter failure times with increasing chlorideconcentration. It was also found that the critical potential for SCC and the corrosion potential areindependent of temperature in the 80 to 1400C range. Duffo et al. (1988) also found that TGSCC occurs

3-29

100

>. 80 MgCI 2(D

0.. CaCI 2(D

Cn

0

X ll / ~~~~~~~AISI 304 S5r _ || / ~~~~~~ScC Wick-Test_ 20 ll| ~~~~~~100 ppm Cl-, 10000

0 10 100 1000


Figure 3-10. Effect of the metallic cation on the SCC of solution annealed AISI304 stainless steel in 100 ppm cr solutions at 1000C using the wick test (Warren,1960)

at potentials above -0.15 VSH8, but they observed IGSCC at potentials below this value. It should be notedthat Shamakian et al. (1980) used constant load tests at a nominal stress ranging from 60 to 100 percentof the yield strength, whereas, Duffo et al. (1988) performed SSRTs. The latter authors also noted thatat constant potential and temperature, the crack propagation rate increases with increasing LiCIconcentration and with increasing temperature at constant LiCI concentration and potential. Crack growthrate varied from approximately 10- to 1r' m/s over all the range of concentrations, temperatures, andpotentials covered in their study.

As reviewed by Cragnolino (1982), several investigators reported the occurrence of partiallyor wholly intergranular cracking on annealed, nonsensitized austenitic stainless steels under some specificconditions of stress, strain rate, and potential in concentrated boiling solutions of MgCI. and in dilutesolutions of NaCI at room temperature, acidified by the presence of relatively high concentrations ofH2SO4. It can be concluded that in annealed austenitic stainless steels, even though the predominant

3-30

cracking mode is transgranular, IGSCC can also occur. It is not currently known if IGSCC can occur atpotentials lower than the critical potential associated with TGSCC.

Data on the TGSCC of austenitic stainless steel in high-temperature, dilute chloride solutionshas been published in recent years (Conglenton, 1985, 1990a, 1990b; Yang, 1992). For solutionannealed Type 316 stainless steel in solutions containing 5 ppm Ct , TGSCC occurs in two regimes ofpotential in the temperature range of 100 to 300'C, as shown in Figure 3-11 (Conglenton, 1990b). Onthe other hand, IGSCC occurs only above a critical potential in the sensitized condition (Yang, 1992),in a similar fashion as found by several authors (Andresen, 1980; Lin, 1981) for sensitized Type 304stainless steel. Cracks in the low potential regime were found to grow at a significantly lower crackpropagation rate, and, as clearly shown in Figure 3-10, the potential required for crack initiation is veryhigh at 100'C. Yang et al. (Yang, 1992) found that the strain needed to initiate cracks increasesmarkedly with decreasing potential for the annealed material, as shown in Figure 3-12. Also, the strainat which cracks initiate is extremely high at 150'C, regardless of the potential. Cracking initiates at thelowest strain at temperatures around 250'C.

Apparently, the difficulties associated with the initiation of cracks in neutral chloride solutionsat temperatures lower than 1000C are reflected in the scarce number of papers dealing with the effect ofpotential under such conditions. Several authors have noted that under certain conditions, such as lessconcentrated MgCl2 solutions or at temperatures lower than the boiling point, cracks seem to nucleatefrom pits. This is also the case in most of the data points shown in Figure 3-13 (Herbsleb, 1980). Theeffect of potential in the SCC susceptibility of Type 304 stainless steel in a very diluted NaCI solutionat 1000C is seen, as well as the coincidence between the critical potential for cracking and the pittingpotential. Unfortunately, although TGSCC was observed, IGSCC predominated owing to the sensitizedcondition of the alloy microstructure. These issues concerning the relationship between pitting corrosionand SCC were reviewed and discussed at length by Szklarska-Smialowska (1986). Important points toemphasize are: i) pit bottoms are highly enriched in chloride, generally with the associated formation ofsalt films, and ii) the pH inside pits in Fe-Cr-Ni alloys is very low, usually lower than 1.0. The existenceof these aggressive, depassivating conditions, highly localized in a microscopic area, seems to establisha bridge between the SCC behavior in highly concentrated and acidic chloride solutions and that in diluteand neutral chloride environments. In the latter, a necessary condition for crack initiation is thedevelopment of a localized environment. The generation of an aggressive microenvironment inside pitsis of greater importance to initiation of SCC than the purely mechanical effect of stress concentration.

As discussed in detail in a previous review on localized corrosion (Cragnolino, 1991a), theseparticular environmental requirements for SCC to occur are attained much more easily in a geometricallycreviced area. In a series of papers, Tsujikawa and coworkers (Tsujikawa, 1985; Tamaki, 1990) haveaddressed that issue. Figure 3-14 shows the repassivation potentials (ER) for crevice corrosion for Type316 stainless steel in dilute NaCl solutions of various concentrations at 80'C (Tamaki, 1990). As acomplement of these data, Figure 3-15 shows the crack growth rate measured as a function of the stressintensity factor (K,) for creviced, tapered double-cantilever-beam specimens at a potential just 10 mVhigher than ER for two chloride concentrations. It should be noted that 0.03% NaCI represents anextremely dilute concentration (0.005 mole/liter). It is clearly seen that crack initiation and propagationis possible at 80'C in a dilute chloride solution, however, the presence of an active crevice is a necessarycondition for SCC to occur. A summary of the results showing the effect of potential and NaCIconcentration on crevice corrosion and SCC is shown in Figure 3-16. It is clearly seen that no crackpropagation occurs below the critical potential, which, in this case, has been identified as therepassivation potential for crevice corrosion. This can be interpreted as the minimum potential at which

3-31

B

0.41-

DCracking

AISI 316 Stainless Steel5 ppm cr

A

0.21-

wW

C

0a-

0 -

-0.2 -

-0.41-

CNo Cracking

I I II

0 100 200 300

Temperature, OC

Figure 3-11. Ranges of potential and temperature over which SCC of solution annealedType 316 stainless steel occurs as determined in slow strain rate tests conducted inwater containing 5 ppm Cr (Yang, 1992)

3-32

0-~~~~~I

la 30 _"

20 _ \ 4,

O~~~~~~~ _10

0

150 200 250 300

Temperature,OC

Figure 3-12. Effects of temperature and applied potential on the strain for crackinitiation for both solution annealed and sensitized Type 316 stainless steel in watercontaining 5 ppm cr as determined In slow strain rate tests (Yang, 1992)

3-33

I # I I ;&# AISI 30 4 Stainless Steel103 ~ If At Army Sensitized 15 hr at 6500C

10-4 M NaCIT =1000C

0

LU 102 ; °0/;'.5

0) A0.75.E 0 1.25

_ 1.75

-* *AOVPitting 0o Intergranular Cracks

OTransgranular Cracks

O Inter and Transgranular Cracks

10 l l I l l l l0 0.2 0.4 0.6 0.8 1.0

Potential, VSHE

Figure 3-13. Effect of potential on the time to failure of sensitized (15 hours at6500C) Type 304 stainless steel at various stress levels in 10"M NaCl solution at1000C (Herbsleb, 1980)

localized acidification can be maintained within the crack enclave (Newman, 1990). In other words, itrepresents the corrosion potential corrected for ohmic potential drop in the crack or crevice environmentbecause, at a lower potential, a net cathodic current would consume the hydrogen ions.

Asphahani (1980) observed that TGSCC of Type 316L stainless steel in acidified chloridesolutions at 141'C can be suppressed by galvanic coupling to carbon steel, copper, nickel, alloy 400(70Ni-3OCu), and Type 304 stainless steel, indicating that galvanic coupling to more active metals broughtdown the potential of the Type 316L specimens below the critical potential for cracking.

These observations indicate that crack propagation does not occur at potentials below a certaincritical potential. As a result of the low ohmic drop which could be expected down the crack in thesesolutions, it appears that a propagating crack can be arrested by decreasing the potential, eitherpotentiostatically or by changing the redox potential of the environment, below this critical potential.

3-34

1 o2 U 0.03% NaCI* 0.3% NaCI

0 3% NaCI

1 0 If I

1V

I I-f--15gA

-1 O0^ li [t.1 -19 9A |

-0.25 -. 20 -0 .15 -0. .10 -0 .05

External Potential, VSCE

Figure 3-14. Determination of crevice corrosion repassivation potentials (ER) for Type 316stainless steel in dilute chloride solutions at 80°C (Tamaki, 1990)

3-35

AISI 316 Stainless SteelER + 0.01 V, 800C0 3% NaCI 10-8o 0.03% NaCI

3 / 10-2 nE

13 1 1

o 0Z

~ 10-4 0

10-10

10-4 ' I10 20 30

Stress Intensity, kgf/mm3/2

Figure 3-15. Crack growth rate as a function of stress intensity (K,) for creviced, tapered,double-cantilever-beam specimen of Type 316 stainless steel showing results for twochloride concentrations at a potential just above the crevice corrosion repassivationpotential (Tamaki, 1990)

3-36

AISI 316 Stainless SteelK1 = 28 kgf/mm3' 2 , 80°C

0

.98I

80

3.0

0.3

. .CREVICECORROSION

0 CREVICECORROSION

NO LOCALIZEDCORROSION

0 00.03 00 0 N 0 a3 0 .

, 1I,. . . .I I . . . I - - aI . . . . .. . . . . .. . . . .I I I a I I I I -

-0.3 -0.2 -0.1 0 0.1 0.2

PotentIal, VSCE

Figure 3-16. Ranges of chloride concentration and potential for the SCC ofcreviced specimens of Type 316 stainless steel at 80'C (Tamaki, 1990 assummarized by Newman, 1990)

Although simple, this concept has not been experimentally demonstrated as can be judged from theliterature reviewed.

3.2.2 Effect of Mechanical Factors

As in other alloy/environment systems, the SCC of Fe-Cr-Ni alloys is strongly influenced bymechanical factors, such as stress, strain, and strain rate. As discussed in Section 2.2, these variablescan be independently controlled in different types of SCC tests. Some relevant experimental results willbe discussed in the following subsections.

3.2.2.1 Effect of Stress and Strain Rate

By using constant load tests, the influence of applied stress on the time to failure for smoothspecimens has been determined in many studies, as reviewed by Speidel (unpublished a). Althoughthreshold stress values have been determined, they cannot be used with confidence for design purposesbecause they are extremely dependent on surface preparation and other testing conditions. A carefulselection of data reported by Speidel (unpublished a) and presented in Table 3-11 shows that the threshold

3-37

Table 3-11. Stress corrosion threshold stresses of solution annealed austenitic stainlesssteels exposed to concentrated solutions of MgCl2 in water at 1450C-150'C. (Speidel,unpublished a)

Threshold Stress (ath) Strength (ay)

Austenitic Stainless Steel r_ _ _ ksi MN/m2 ksi

Type 347 160 23.0 180 26

Type 316L 130 19.0 200 29

Type 316LN 100 14.5 230 33

Type 304 110 16.0 150 22

Type 302 120 17.4 -

values for several austenitic stainless steels in concentrated, boiling MgC12 are always lower than the yieldstrength. It should be noted that the data in Table 3-11 are confined to a specific solution and can beeasily altered by modifications in the environment.

The role of strain rate in the SCC of austenitic stainless steel is shown in Figure 3-17 (Mancia,1988). In this plot, regimes of pitting and SCC for type 304 stainless steel in boiling (110'C) SM NaCIsolutions at various pHs are shown as a function of strain rate for specimens tested at the corrosionpotential. It is apparent that strain rates as low as 2.2 x 10-7/sec are required to induce cracking inneutral solutions. The steady state corrosion potentials observed in these tests were close to -0.38 Vsm(-0.14 VSHE), regardless of the solution pH. The authors identified this potential as the pitting potentialof Type 304 stainless steel in 5 M NaCI solutions. On the other hand, it should be noted that Duffo etal. (Duffo, 1988) have observed TGSCC in 6.5 M LiCI solution at 80 and 105'C at potentials higher than-0.15 VSHE using a nominal strain rate one order of magnitude faster (1.2 x 104/sec). These results seemto indicate that LiCI is a more aggressive electrolyte than NaCi for causing SCC of Type 304. However,more work is needed to confirm this observation.

3.2.2.2 Effect of Stress Intensity

Many authors have studied the effect of the stress intensity, K1, on the crack growth rate ofaustenitic stainless steel in chloride containing solutions. A good example is shown in Figure 3-15 forType 316 stainless steels under potential controlled conditions. Figure 3-18 shows crack velocity versusK, curves for Type 304L stainless steel under open circuit conditions in two chloride solutions (Speidel,1981). The crack growth rate at the plateau corresponding to Stage 2 is almost one order of magnitudehigher in MgC12 than that in NaCI. In addition, the threshold stress intensity, K 1,.. is significantly lowerin the acid chloride. It should be noted, however, that there are significant differences in concentrationand temperature between both solutions. In the NaCI solution, the concentration of chloride ions is 4.9mole/kg H20, whereas, in the MgC12 reaches 15.2 mole/kg H20. On the other hand, the crack growthrates in neutral NaCI solutions appear to be very similar for Type 304 and 316 stainless steel regardless

3-38

1

3

U U

I II I

I I

I I

I II II I

- 0 Us* l

- I

Localized *Corrosion I

..

_ : :~~~~~~~.

00I~..

......

SCC

............

U

I0.

5

7

Pitting + SCC

U

1 1

10-5 10-6 10-7 10-8

Strain Rate, S-1

Figure 3-17. Effects of pH and strain rate defining domains of pitting and SCC forsolution annealed Type 304 stainless steel in 5 M NaCI solution at 110'C (Mancia,1988)

of concentration and temperature, as seen by comparing the data shown in Figure 3-15 with that in Figure3-18. However, KI, seems to be abnormally low in Figure 3-15.

Threshold stress intensity values of 12 MN * m-" and 14 MN* m-32 were reported by Eremiasand Marichev (1979) for Fe-l8Cr-lONi-0.5Ti stainless steel in 46% LiCI at 1050C and 44.5% MgCI2at 115'C, respectively. Lefakis and Rostoker (1977) determined the K. values for Type 304 stainlesssteel in boiling (144°C) 42% MgCl2 and in the vapor phase above the boiling solution. They found avalue close to 10 MN * m3W in the liquid phase, that decreases to 1.1 MN * mnr in the vapor phase.This significant decrease in K, could be attributed to the predominant presence of HCI in the gaseousphase as a consequence of the decomposition of MgCl 2 by boiling, as noted by Duffo et al. (Duffo,1988). Although this is an important observation, it may not be relevant for a container material becauseit is doubtful that a strongly acidic gaseous phase can be formed under the conditions present in therepository.

3-39

ia- 5

AISI 304L Stainless Steel

* 42% MgCI2 Solution, 1300C

V 22% NaCI Solution, 1050C10-6 I-

0

E

U

cu

h.6.

c0.-

-V0

o

1-

00

a,

C,)ut)

I-

1a- 7 _-

*-�Ik.P PI

10-8 I- I

W W W1E�

10-9

10.10

'I TTI10o-1 1 I

~K:SCC_ .~~~~~~~

aII

ISCC

10-12 I I Ii

0 10 20 30 40 50 60Stress Intensity, MN-m-1/2

Figure 3-18. Effect of stress intensity on the crack growth rate of solution annealedstainless steel exposed to 42% MgCI2 solution at 130°C anc to 22% NaCI solution at(Speidel, 1981)

Type 304L1050C

3-40

3.2.3 Effect of Metallurgical Factors

The influence of alloy composition and the effect of thermomechanical treatments on thechloride cracking are the subject of the following subsections.

3.2.3.1 Effect of Chemical Composition

For many years, there has been extensive research to investigate the effects of alloying additionson the SCC resistance of austenitic Fe-Cr-Ni alloys, mainly by testing in boiling MgCl2 solutions.Sedriks (1979) has summarized this information as shown in Figure 3-19, following the broad andthorough reviews of Latanision and Stahele (1969) and Theus and Stahele (1977). It is evident that manyadditions are detrimental to chloride cracking. However, several alloying elements are beneficial, suchas nickel and silicon.

One of the most important factors controlling the susceptibility to SCC of austenitic Fe-Cr-Nialloys in chloride containing solutions is the nickel content. This was established many years ago byCopson (1959) using wire specimens immersed in boiling 45% MgCl2 under constant load conditions.Alloys containing above 45% Ni did not crack in 30 days of testing. Speidel (1981) has used a fracturemechanics approach to measure the crack growth rate and K, for a variety of commercial alloys inboiling (105'C) 22% NaCI solutions. Tests were conducted in fully aerated solutions under open circuitconditions. Figure 3-20 shows that crack growth rates in the stress intensity independent region (Stage2) are too low to be measurable for alloys with nickel contents above 35 percent. The limit of detectionwas 3 x 10-" m/s. On the other hand, the highest crack propagation rates were found for chemicalcompositions similar to Type 304 stainless steel (approximately 8 to 10% Ni). The effect of nickelcontent on the threshold stress intensity factor, Ki., is shown in Figure 3-21. It is clearly seen thatalloys containing more than 30% Ni exhibited extremely high K,, values, and a minimum is observedat the same nickel contents at which a maximum appears in the crack growth rate versus nickel contentplot (Figure 3-19).

Whereas Figure 3-18 indicates that the effect of molybdenum is variable, Figure 3-22 clearlyshows that Mo also has a strong beneficial effect on the SCC resistance of Fe-Cr-Ni alloys in 22 % NaCIsolutions at the boiling temperature. Alloys with around 5% Mo are fully resistant to SCC in thisconcentrated chloride environment. The apparent discrepancy arises from the fact that in MgCl 2solutions, molybdenum additions first decrease and then increase the resistance to chloride cracking, witha maximum susceptibility at about 1.5% Mo (Sedriks, 1979), while no such pattern is evident inconcentrated NaCl solutions where a continuous increase in SCC resistance with molybdenum content isobserved.

The information presented above indicates that alloys with chemical compositions similar toalloy 825 are extremely resistant to SCC in concentrated, neutral chloride solutions as a result of thecombined beneficial effect of nickel and molybdenum. It should be emphasized, however, that thisconclusion is strictly valid for the material in the solution annealed condition when exposed to a well-controlled environment. Asphahani (1980) reported that alloy 825, as well as other high nickel alloys(with Ni > 25% and Mo > 2.5%), were resistant to SCC in acidified chloride solutions up to 141'C,under the same conditions (30 days test) in which Type 304, :104L, 316, and 316L stainless steel werefound to be extremely susceptible. On the other hand, there ar . references (Staehle et al, 1970; Chiang,1985) reporting TGSCC of alloy 825 in concentrated, boiling MgCl2 solutions, as well as in 100 ppm

3-41

IIA VIA VIA

I I%/II I0 1 lI,

IIIA

B

V

AlV

IVA VA

C N

V P

Si _

* X1%/R X/P X/IQ xi II CaI v Iv V Iv vILD V1IL 1 VI IDI IIID' *ITi V Cr Mn Fe Co Ni Cu Zn Ga Ge AsX X V V Base V * * * X

Zr Cb Mo Tc Ru Rh Pd Ag Cd In Sn SbX X - - -X X X - * V XHf Ta W Re Os Ir Pt Au Hg TI Pb Bi… -X X X I X X _ _ X X

Segment of the Periodic Table of Elements R.E. R.E.* Beneficial, V Variable, [ No Effect

X Detrimental, Not Investigated F-1

Figure 3-19. Summary of the effect of the elements in the periodic table on theSCC resistance of austenitic stainless steel in chloride solutions (Sedriks, 1979)

chloride solutions at 300C in the presence of 50 ppm oxygen (Boyd, 1972). Kolts (1982) showed thatalloys with nickel contents higher than 35% can undergo SCC in deaerated, concentrated NaCI, CaC12,and MgCI 2 solutions at temperatures higher than 141TC. The temperature at which cracking wasobserved increases with increasing nickel content. Undoubtedly, significant increases in temperature anddecreases in pH can make Fe-Cr-Ni alloys with nickel contents about 40 percent susceptible to chloridecracking. The same decrease in the resistance to cracking is observed in chloride-containing solutionsin the presence of H2S and elemental sulfur (Sridhar, 19S,2b), probably due to the formation ofpolysulfides that act as very active depassivating agents for Ni-based alloys. These types of sulfurcompounds are not expected to be present under the milt, oxidizing conditions prevailing in therepository, where the predominant sulfur compound seems to be sulfate.

3-42

10-7

1 0-8

10 9

10-11-

15.5 - 21 %Cr

025 - 30% Cr

-12 I s0 20 40 60 80

Ni Content, %

Figure 3-20. Effect of the nickel content on the nmaximum crack velocity (Stage 2) forFe-Cr-NI alloys in aerated 22%Z NaCI solution at 1 05°C. Arrows indicate no crackgrowvth. (Speidel, 1981)

3-43

80

on C60

C,

CII20

Figure 3-21. Effect of the nickel content on the threshold stress intensity (Kim for theSCC of Fe-Cr-Ni alloys in aerated 22%t NaCI solution at IO5°C (Speidel, 1981)

344

60Austenitic Stainless SteelsAerated 22% NaCI Solution, 1 05C

cm 50

CL ~~~~~~~~~~904L

E3 0 40

317

(X30

5 20 316L

Crand Ni Range:C-) I 15.5 - 21% Cr

10 13.0- 24.5% Ni

o~~~

0 1 2 3 4 5 6

Molybdenum Content, wt-%

Figure 3-22. Effect of the molybdenum content on the threshold stress intensity (Kay)for austenitic Fe-Cr-Ni-Mo alloys in aerated 22%* NaCl solution at 105'C (Speidel,

3.2.3.2 Effect of Microstructure

All the Fe-Cr-Ni alloys of interest are single phase, solution strengthened alloys in which themajor microstructural modification that can be anticipated is the precipitation of chromium carbides tograin boundaries under certain heat treatments. Although Type 304L and 316L stainless steel have lowcarbon contents (< 0.04%), prolonged heat treatment in the 550-800°C temperature regime may leadto the precipitation of CrzC6 along grain boundaries and the associated chromium depletion. Althoughalloy 825 is stabilized with titanium, it can exhibit sensitization due to chromium depletion at precipitatedcarbides (Raymond, 1968; Brown, 1969). In commercially pioduced material, sensitization is reducedby mill annealing at a relatively low temperature, i.e. 940'C (S adriks, 1982), such that titanium carbidesare formed preferentially and grain boundary precipitation of chromium carbides is minimized.

3-45

The main effect of a sensitized microstructure is a greater tendency to IGSCC in chloridecontaining solutions. This is well documented for the austenitic stainless steels (Sedriks, 1979; Kowaka,1990), but it can also occur in alloy 825. No data on the effect of sensitization is available in the openliterature for this alloy.

3.2.3.3 Effect of Cold Work and Surface Preparation

For both Types 304 and 316 stainless steel, it has been shown that cold work decreases furtherthe resistance to SCC in boiling 45% MgCI2 solutions (Kowaka, 1990). However, when the degree ofcold work is above 20 percent there is a significant recovery in the resistance to SCC in the case of Type304 stainless steel, probably as a result of the formation of ci-martensite which is not the case of Type316. As noted by Sedriks (1979), the complex influence of various surface treatments, which canproduce variations in local work hardening, induce martensite transformation, generate residual stresses,and introduce embedded particles from abrasives or machining tools, are impossible to quantify. It isusually considered that shot peening, which introduces compressive surface stresses, is beneficial becauseit delays or suppresses crack initiation.

For predictive laboratory evaluations, it is important to study the effect of surface conditionsto be used in service. This is of particular concern in the case of some Ni-based alloys, such as alloy825 and Hastelloy C-22, because the surface layer can be impoverished in chromium as a result of milldescaling operations. Crack initiation may be then facilitated by surface Cr-depletion (Place, 1991).

3-46

4 SCC OF Cu AND Cu-BA'SED ALLOYS

4.1 REVIEW OF THE DOE/NRC HLW RESEARCH

As mentioned before, inclusion of copper base alloys in various test programs began in Octoberof 1984. Hence, most of the experimental investigations, whil-h were started prior to this and reportedsubsequently, did not include copper base alloys. The results of long-term corrosion tests (13 to 16months) on Cu-base alloys in irradiated environments from FY 85 through FY 86 have been reported byYunker (1990). The environments consisted of exposure to gaseous air-water mixtures at 95 and 1500 Cand aqueous J-13 water at 950 C. In the latter case, the specimens were partially immersed. The gammairradiation dose rate was 1.9 x 103 to 4.9 x 105 rads/hour. Specimens of CDA-102, CDA-715, andCDA-613 (1.5 mm thick), which were bent into the form of a tear-drop and welded using the GasTungsten Arc process at the foot of the bend, were tested for stress corrosion cracking (SCC). In somecases, the specimens were "pre-pitted" electrochemically by polarizing them in simulated J-13 water atan unspecified temperature and at various anodic potentials. In this environment, it must be noted thatthe corrosion mode of CDA-102 and CDA-715 is uniform, unless the temperature is 950C (Sridhar,1992a). Indeed, in the micrographs of pre-pitted specimens shown by Yunker (1990), a uniform typeof corrosion with broad surface undulations is the predominant form of corrosion. No SCC of anyspecimen was observed in this test program up to a test time of 16 months for the 150'C tests and 13months for the 950C tests. No significant changes in the water chemistry were reported at the end ofvarious test periods.

Beavers and Durr (1990c) reported results of slow strain rate tests (SSRT) on CDA-102 andCDA-715 in various nitrite environments. Some of the results are shown in Table 4-1. It can be seenthat 0.005 M nitrite as sodium nitrite can cause transgranular stress corrosion cracking (TGSCC) ofCDA-102. Several other features are worthy of note in these :esults:

* Increase in temperature from 23°C to 95°C resulted in a shift in the potential regime inwhich cracking occurred from open-circuit to more positive potentials. The extent of thisshift is not known since the lowest applied potential at which cracking occurred at 95°C wasnot established. This parallels the behavior of tlis alloy in a bicarbonate system whereincreasing the temperature increases the passivity of the alloy.

* Addition of simulated J-13 water to the 0.005 M nitrite solution eliminated cracking evenat 23°C. It is not known whether a more anodic potential will cause cracking.

* The free corrosion potentials reported for CDA-102 in these tests are significantly higherthan those reported in other tests in aerated, simulated J-13 water (Figure 4-1). If thesedifferences are caused by the presence of oxidizing species such as NO2 , then experimentswhere N02- was added to model solutions simulating J-13 water should have exhibitedincreases in free corrosion potential. However, previous, partial factorial tests indicated thatnitrite did not have a significant effect on the free corrosion potential (Beavers, 1990b).

No SCC was observed on CDA-715 in 1 M NaNO. at 23°C at anodic potentials up to 127mVwE (Beavers, 1992c). However, the higher elongation and time to failure of the CDA-715 specimenscompared to CDA-102, even in an inert environment is puzzling. Typically, CDA-715 is expected to

4-1

Table 4-1. Results of slow strain rate tests for CDA 102 in nitrite solutions. Nominal strain rate:1 x 104/sec. (Beavers, 1990c)

= I CrackTime t o Velocity

Temnp., Potential, Failure, SCC x 10'1,Environment 0C mVscr hrs Observed nun/sec

Oil 90 _ 18.7 No _

Oil 90 _ 20.1 No _

IM NaNO2 23 (-13) * 23.4 TG -

1M NaNO2 23 (+8) 21.2 TG 1.97

IM NaNO2 90 (-2) 20.0 No -

1M NaNO2 90 (-17) 21.2 No -

1M NaNO2 23 +61 18.0 TG 8.02

1M NaNO2 23 +68 18.5 TG 5.26

IM NaNO2 90 +122 11.6 TG 20.40

1M NaNO2 90 +121 11.0 TG 12.60

IM NaNO2 + J-13 Water 90 (+6) 21.0 No -

IM NaNO2 + J-13 Water 90 (+15) 21.1 No -

0.005M NaNO2 23 (+35) 24.4 TG 1.02

0.005M NaNO2 23 (+12) 26.0 TG -

0.005M NaNO2 90 (+17) 20.5 No

0.005M NaNO2 90 (+13) 22.3 No -

0.005M NaNO2 23 +200 21.4 TG 5.84

0.005M NaNO2 23 +174 20.2 TG 5.09

0.005M NaNO2 90 +234 21.3 TG 0.52

0.005M NaNO2 90 +246 21.0 TG 1.19

0.005M NaNO2 + J-13 23 (+40) 25.0 No -

0.005M NaNO2 + J-13 23 (+29) 26.0 No

a -( ) Indicates potential under freely corroding condition

4-2

0

4100

E-200

Z

-300

-400LGND

* - A CDA 102SK - A CDA 715

-500 I I I IpI I pI0 1500 3000 4500 6000 7500 9000 10500 12000 13500 15000Test Time, Hours

Figure 4-1. Corrosion potentials of CDA-102 and CDA-715 in long-term boildowntests in aerated, simulated J-13 water (Beavers, 1992)

exhibit lower elongation than CDA-102. While the 1 M NaNO2 solution is far too concentrated to bea realistic simulation of radiolysis, the results of 0.005 M NaNO2 may be compared to the findings fromlong-term SCC tests under irradiation. The observation of cracking in the SSRTs may be a result of thegreater severity of the test, lower temperatures used, and higher potentials observed. If the highpotentials can only be obtained by the formation of H202 which is observed only in fully aqueousconditions, and if the fixation of nitrogen occurs only in air plus water vapor mixtures, then the resultsgenerated thus far suggest that SCC by nitrite or nitrate may not be a serious concern for Cu alloys.

Beavers et al. (1992b) also reported results of long-term boildown tests on U-bend specimensof CDA-102 and CDA-715 in simulated J-13 water. No SCC was observed in 13,400 hours.

Maiya et al. (1990) investigated the Cu-base alloys in the same manner as described for theFe-Ni-Cr-Mo alloys. The Stress and Strain Ratios (Figure 3-3) indicate no susceptibility to cracking.

4-3

However, small (10 to 20 ,um) cracks were noted in the scanning electron microscope (SEM). The modeand morphology of the cracks were not investigated.

4.2 REVIEW OF THE GENERAL LITERATURE

4.2.1 Effect of Environmental Factors

Copper and copper-base alloys are resistant to SCC in many environments, as revealed by theirsuccessful use in multiple applications over extended time periods. Nevertheless, they are susceptible tocracking in certain aqueous environments as shown in Table 4-2. Since the TGSCC of ca-brass in moistair contaminated with NH3 was identified as an important failure process at the beginning of this century,extensive studies in a large number of copper alloys have been conducted in the presence of ammonia orammoniacal species. As in the case of Fe-Cr-Ni alloys, however, the number of environments knownto induce SCC of copper alloys has increased significantly in the last 30 years. In particular, nitrogenoxyanions such as nitrate and nitrite have been identified as strong promoters of IGSCC and TGSCC incertain copper alloys. Under the mild oxidizing conditions prevailing in the Yucca Mountain proposedrepository site, the formation of nitrite and nitrate from humid air can be expected in the presence ofgamma radiation, as discussed in Section 3.1. Reed (1992b) reported the formation of traces of NH3 atlower oxygen concentrations, but the concentrations are probably too low to consider SCC in ammoniaenvironments a likely failure process under the repository conditions.

4.2.1.1 Effect of Nitrate and Nitrite

Brasses have not been selected as candidate container materials, among other factors, as a resultof their susceptibility to SCC in many aqueous environments. However, it is useful to discuss thedetrimental role of several anionic species on the basis of specific papers chosen from the extensiveliterature devoted to brasses as a guideline for evaluating potential risks of SCC failures in the candidatecopper alloys. Graf (1969) reported the occurrence of TGSCC of ca-brasses containing 20, 30, and 37at% Zn in cupric nitrate solutions, ranging in concentrations from 0.1 to 1.0 mole/liter, at 250C. TheSCC susceptibility decreased with the decrease in the Zn content, but no experiments were conducted inpure copper. A nominal stress corresponding to 80 percent of the yield strength was applied to thespecimens in open circuit tests. The failure times were found to be more than 100 times longer than insaturated solutions of Cu(NH3)42+ ions, representing approximately two orders of magnitude lower crackgrowth rates. Maximum susceptibility was observed for Cu-30Zn and Cu-37 Zn in 0.5 M Cu(NO3)2solution in which the pH was about 1.6. At concentrations above 1.5 mole/liter, intergranular corrosionbecame dominant as a result of the acidity of the solutions (pH 1.6).

A decade later Kawashima et al. (1979) found that admiralty brass CDA 443 (Cu-28Zn-lSn-0.04As) was more susceptible to TGSCC in nitrite than in nitrate solutions. These solutions wereprepared by using sodium salts at a concentration of 1 mole/liter and the pH adjusted to 8. SSRTs wereconducted under a controlled potential of 0.3 VSHE at 250 C. Many other anions were tested, as shownin Table 4-3 in which they are arranged from top to bottom in order of decreasing severity for promotingSCC. It can be seen that the estimated crack velocity was found to be more than one order of magnitudefaster in nitrite than in nitrate.

Pure copper (99.9 and 99.99%) was found to be extremely susceptible to TGSCC in nitritesolutions (Pednekar, 1979a). SSRTs conducted at the open circuit potential (-0.15 to 0.18 VSHE)

4-4

Table 4-2. Environments that promote SCC of copper-based alloys

Copper Cu-Al Cu-Ni Cu-Zn

Ammonia, NH3 Ammonia, NH3 Ammonia, NH3 Ammonia, NH3

Nitrite, NOi Steam, 300'C Steam, 300'C Nitrite, NOi

Acetate, CH3COOr _ Hg2(NO3)2 Nitrate, N03-

- Sulfate, SO42|

_ __- Acetate, CH3 COO

_ _ - Formate, HCOO-

_ _ - Tartrate

_ _ - Citrate

_ _ - Polyphosphates

_ _ _ Carbonates, C032- ?

revealed that the elongation to failure decreased from 55 percent in air to 25 percent in aerated 1 MNaNO 2. An average crack velocity of 1.0 x 10-' m/s was estimated, assuming that cracks were initiatedearly in the test. These results clearly demonstrated that ductile, pure metals are susceptible totransgranular, brittle failure in a specific environment that leads, as a necessary condition forenvironmental assisted cracking, to the formation of an oxide film (probably Cu2O) of appropriatethickness. Further work using copper single crystals in the same solution (Sieradzki, 1984) demonstrated,by measuring acoustic emission and electrochemical current transients, that crack advance isdiscontinuous, probably by fast, short-range cleavage events which are triggered by anodic processesassociated with the formation of a film. Average crack velocities of up to 3.0 x 10M8 m/s were measuredat 30'C under an applied potential of 0.24 VSHE using extension rates ranging from 7 x 10- to 3 X104 cm/sec. As in other face-centered-cubic (fcc) nonferrous metals and alloys, transgranular crackspropagate on { 110) planes in both copper and a-brass (Meletis, 1986).

Aaltonen et al. (1985) found that the SCC of oxygen-free high conductivity (OFHC) coppercould be intergranular or transgranular, depending on potential and temperature. At 25°C brittle,cleavage-like, TGSCC was observed in SSRTs in 0.3 M NaNO2 solution at potentials higher than thecorrosion potential (0.19 VSHE), whereas, mixed IGSCC/TGSCC was detected at potentials lower thanthat value. However, at 80'C, although TGSCC predominated at 0.34 VsHE, crack initiation wasintergranular. From anodic polarization curves conducted at 25°C, Aaltonen et al. (1985) concluded thatat 0.14 VSHE Cu+ is oxidized to Cu2 ' with the possible formation of Cu2O at intermediate potentials,while CuO is formed at 0.34 VSHB. This implies that TGSCC found at 0.19 VSHE could have been relatedto the formation of Cu2O. However, these conclusions are highly speculative without a more completeelectrochemical study.

4-5

- - - - - - m - m m ON

Table 4-3. The corrosion potential, ratio of fracture stress in solution to that in air, average crack velocity, surfaceappearance of the fractured specimen, and the cracking mode of admiralty brass in various solutions at 250C. (Kawashimaet al., 1979)

ll_| Surface CrackingAnion pH Eu, mVH a soln/a air Va r/s Appearance Mode

NO2 8. 1 80 0.09 2 x 10-' black film TC

NO3- 8.0 150 0.34 1 x HY reddish brown TC

CIO; 8.0 210 0.51 6 x 10-9 reddish brown TC

S04- 8.0 180 0.55 6 x 10-9 reddish brown TC

MoO 4- 10.2 100 0.68 4 x 104 bright yellow TC, mild

Cl- 8.4 20 0.84 4 x 10-10 black film TC, mild

W04- 9.4 80 0.89 2 x 10-9 bright yellow TC, mild

HP04- 8.1 120 0.93 - bright yellow NC

HCO; 8.0 80 0.95 - bright yellow NC

B407- 8.0 150 0.96 - bright yellow NC

CrO4- 8.7 40 0.97 - bright yellow NC

WTC = transgranular crackingNC = no cracking

6%,

Yu and Parkins (1987) conducted an extensive study on the SCC of ca-brass and pure copper(99.99%) in nitrite solutions at pH 9. TGSCC occurred predominantly in 1 M NaNO2 above a minimumpotential of -0.1 Vsm (0.14 VSHE) for copper and approximately 100 mV lower for brass. Crackingsusceptibility increased with potential and nitrite concentration, (0.01 to 1 mole/liter) as reflected in thedecreasing reduction in area (RA) and increasing crack velocity. At potentials above 0.2 VSc, generalcorrosion predominated over SCC. In the case of brass, it was found that crack velocity increases withincreasing strain rate (104 to 10- /sec) and temperature (20 to 100'C). No investigation of thesevariables was conducted on pure copper. Cassagne et al. (1990) studied the SCC of copper monocrystalsin 1 M NaNO2 solution at pH 9. Their results are comparable to those of Yu and Parkins, as shown inTable 4-4.

4.2.1.2 Effect of Other Anionic Species

As shown in Table 4-3, many other anions promote TGSCC of admiralty brass at anodicpotentials, including chloride and sulfate, which are found as components of the groundwater in theproximity of the Yucca Mountain site. The effect of sulfate was studied in great detail by Kawashimaet al. (1977). TGSCC was observed over a wide range of pHs (2 to 12) in 0.5 M Na2SO4 solutions withthe pH adjusted with either H2S04 or NaOH to the desired value. At pHs ranging from 4 to 9, theminimum potential for TGSCC was found to be 0.2 VSHE, which was only 20 mV higher than thecorrosion potential. In very acidic (pH < 2) and very basic (pH> 12) solutions, only general corrosionwas observed. In chloride-containing solutions, TGSCC was minor (Table 4-3), as it was also the casein the presence of two anions well known by their passivating action, such as molybdate (MoO4

2 -) andtungstate (WO4

2-). No cracking was observed in the presence of bicarbonate (HC03 ) (Kawashima, 1979),which is an important component of groundwater and has a profound effect on the localized corrosionof copper (Sridhar, 1992a).

As shown in Table 4-3, no SCC of admiralty brass was observed in alkaline orthophosphate(Na2HPO4) solutions. Rebak and Galvele (1989) confirmed this observation for a-brass. However, theyobserved that the alloy was susceptible to IGSCC in pyrophosphate (Na4P2Q7) and triphosphate (Na5P3O10)solutions. Crack velocity increased with increasing potential just above the corrosion potential. Theaddition of Cu2 1 ions to both solutions did not affect the crack velocity, even though it increases thecorrosion potential.

Some organic anions which, like NH3, are known to be strong complexing agents for copper(ll),also promote TGSCC of brasses. They are formate, acetate, and tartrate (Parkins, 1982). In an extensivestudy, in which potential and pH were varied for solutions containing these anions, Parkins and Holroyd(1982) found that IGSCC of 70/30 brass occurred within the potential-pH domain in which Cu2O is thethermodynamically stable solid phase. On the other hand, TGSCC was observed at potentials above thedomain of stability of Cu2O. SCC was also observed by Parkins and Holroyd (1982) in NaOH solutionsat pH 13. IGSCC was detected at low potentials in the Cu2O domain and TGSCC at higher potentials.

Very limited information on the effect of anions on SCC is available for the candidate containermaterials. Pure copper (99.9 and 99.999%) was found to be susceptible to SCC in cupric acetate solution(Escalante, 1971). Constant load tests were conducted in 0.05M Cu(CH3COO)2 solution under opencircuit conditions. It is important to note that IGSCC occurred only in the dark. No SCC was observedin specimens illuminated with 150 W incandescent lamps. Although the SCC of a pure metal wasconsidered, at that time, a unique observation, it can be argued that cracking occurred as a result ofimpurity segregation to grain boundaries.

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Table 4-4. Nominal transgranular crack velocities for copper in 1 M sodium nitrite,pH 9, tested under slow strain rate at 25'C.

Potential Nominal Crack Velocity, m/s

Cassagne, 1990 Yu and Parkins, 1987Monocrystals Polycrystals

VSE j VSHE Strain Rate = 104 s- Strain Rate = 1.8 x 104 s-

0.20 0.44 1.2 x 10-

0.10 0.34 1.2 x 10-l

0.05 0.29 3.3 x 10-9 1 x 10-l

0.00 0.24 1.9 x 10-9 9 X 10-9

-0.05 0.19 2.5 x 10-9

-0.07a 0.1 1.3 x 10-9

-0.10 0.14 _ 4.5 x 10-'°

-0.15 0.09 0 _

* - Corrosion Potential

More recently, TGSCC of copper monocrystals (99.99%) has been observed in sodium acetatesolution (Cassagne, 1990). As shown in Table 4-5, cracking occurred only in solutions of pH 5.5 and10.3 where an oxide film is formed. At pH 3, corresponding to the absence of a stable film on the metalsurface only ductile failure occurred. It appears that the minimum potential for cracking is just below 0.0Vsm (0.24 VSHE), and the crack velocity increases with potential.

Work by Aaltonen et al. (1985) showed the absence of SCC in OFHC-copper exposed to asimulated groundwater environment using a limited number of SSRT at an initial strain rate of 4.5 X10 7'/sec. Most of these tests were conducted at 80'C under aerated conditions. No SCC was observed,even in tests in which the chloride and sulfate concentrations were increased separately 100 times,corresponding to 7000 ppm chloride and 960 ppm sulfate, or simultaneously.

4.2.2 Effect of Alloy Composition

Most of the research work conducted on the SCC of copper base alloys has been performed inammonia environments, even though SCC has been reported in a variety of rural, coastal, and industrialenvironments. While pure copper and CDA 613 are susceptible to SCC in ammonia environments, itappears that CDA 715 is very resistant, as reflected in Figure 4-2. It is seen that only Cu-Ni alloys withnickel contents below 10 percent are susceptible to SCC in ammonia atmospheres, as indicated by thelower value of the threshold stress for SCC as compared to the ultimate tensile strength (UTS) in air(Speidel, unpublished b).

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M- -M m m M - M M M

Table 4-5. Results of slow strain rate tests on copper monocrystals in 0.1 M sodium acetate (Cassagne et al., 1990)

I-P

I Potential,2 1 Strain Rate, UTS, Elogation to Failure, Nominal CrackI pH VSCE sU ksi (MJPa) % Velocity, m/s Mode of Failure

l . ~~~ ~ ~~~~~~~~~~~AIR l- - 1 X 104 18.5(127.7) 35.2 _ ductile

________ ___________ ______________ O XIDE FILM

10.3 0.11 I x 10' 13.2(91.1) 15.5 1.6 x 10-9 TGSCC

5.5 0.05 8 x 104 1 5 (10 3 .5 )b 53.2 ductile

5.5 0.05 5 x 104 13.7(94.5)6 38.5 5.6 x 10-10 ductile + TGSCC

5.5 0.05 1 x 10' 8.5(58.7)p 8 4.6 x 10-10 TGSCC

5.5 0.11 1 X 104 10.2(70.4) 13 2 x 10-9 TGSCC

5.5 0.08 I x 104 - 5 x 10.30 TGSCC

5.5 0.04 I x 104 13.5(93.1) 32 4.7 x 10-1° TGSCC

5.5 0.00 I x 104 21.6(149) 35 2.5 x 10-10 ductile + TGSCC

No OXIDE

3 0.05 I x 104 17(117.3) 25 - ductile

3 0.01 I x 10' 14.3(98.7) 35 - ductile

3 -0.06 I x 104 23(158.7) 37 - ductile

3 -0.03 5 X 10' 15.9(109.7) 44 - ductile

3 -0.06 5 x 104 15.4(106.3) 37 - ductile

The corrosion potential for pH 3, 5.5, and 10.3 were respectively -0.04, -0.06, and -0.03 V,-, before the tests.

b Oriented crystals; the remainder are randomly oriented.

500

Fracture Strength of Copper-Nickel Alloys

o UTS in Air, Ambient Temperature

* SCC Threshold Stress in Ammonia Atmosphere

70

60

50

400 _-

w0_0

co(U

cI..

ca

112

U.

300

200

01 40 0e

(M

2Un

-1 30

- 20

100I

_ 10

0 I I I JAn0 10 20 30

Weight-Percent Nickel

Figure 4-2. Effect of the nickel content on the ultimate tensile strength (UTS) in airand SCC threshold stress (od) for copper-nickel alloys in ammonia atmosphere(Speidel, unpublished)

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No studies on the SCC susceptibility of CDA 613 and CDA 715 in the presence of anionstypical of groundwater compositions could be located in the open literature. However, it is worthwhileto note that CDA 706, which is a Cu-lONi alloy, was not found susceptible to cracking in 0.5 M Na2SO4solution over a wide range of pHs (1.2 to 13) when tested using SSRTs at 1 x 104/sec under open circuitconditions and at an applied potential of 0.2 VsHB (Pednekar, 1979). Under similar environmentalconditions, admiralty brass (CDA 443) was found to be extremely susceptible. Also, no SCC wasobserved in pH 7 Mattsson's solution (ammoniacal solution), 1 M NaNO2 and 1 M NaNO 3. These resultssuggest that Cu-Ni alloys are significantly more resistant to SCC than a-brasses and pure copper.

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h, -

5 SUMMARY AND RECOMMENDATIONS

The advantages and limitations of different stress corrosion cracking (SCQ) test techniques are brieflydiscussed in this review to guide the selection of the most appropriate experimental methods for assessingthe long-term performance of container materials. The different techniques are classified and describedaccording to the stressing system and the specimen geometry. Although no attempt was made to presenta comparative discussion of all specimen geometries used in SCC tests, particularly in tests using afracture mechanics approach, a clear distinction was established between plain, smooth specimens andnotched, precracked specimens. The validity of mechanical parameters used for assessing SCCsusceptibility, such as the threshold stress (ah) or the threshold stress intensity (K,,.), as boundingparameters for long-term extrapolation of material performance, is discussed. The presence of flaws,notches, or other minor defects on the metal surface, that may lead to failure even at low applied orresidual stresses, is the main limitation for using a nominal stress approach, as given by A. On the otherhand, the advantages of Kno as a bounding parameter are recognized, since a linear elastic fracturemechanics (LEFM) approach includes in a quantitative manner the effect of stress concentration inprecracked specimens. Another advantage is the easy generation of a severe microenvironment withinthe precrack enclave, which may facilitate crack initiation. It is noted, however, that an accuratedetermination of K,, necessarily requires a very extended test time that will increase with the resistanceof the material to SMC. An additional constraint, to avoid plane stress conditions, is the thicknessrequirements that generally accompany materials exhibiting high K,, values.

The observation, long ago reported, that the existence of a precrack does not necessarily eliminate theinduction or initiation time is emphasized. This is particularly important in the case of materialsmoderately resistant to SCC in relatively mild environmental conditions because the initiation time mayconstitute the predominant part of the total time to failure. Again, the implication is that, in these cases,very extended test times may be required to obtain a reliable assessment of the SCC performance of agiven material. Tests can be accelerated in these cases by increasing the value of some of the controllingvariables, such as temperature, concentration of aggressive species, etc. with the assumption, which hasto be demonstrated, that the mechanism of SCC is not altered.

The use of constant deflection tests is recommended with the purpose of studying, in some detail, crackinitiation and the influence of surface conditions, as affected by the presence of preformed oxide films,surface depletion of alloying elements, superficial work hardening, etc. The simplicity of the testspecimens could be appropriate for simulating heat transfer conditions in the presence of crevices. Thistype of tests can be used in conjunction with slow strain rate tests (SSRTs) to define environmentalconditions that lead to the occurrence of SMC. An important aspect to emphasize is that crack initiationcan be considered a stochastic process, particularly in the case in which there is no preferential path dueto chemical or microstructural inhomogeneities, and, hence, multiple specimens must be tested under thesame experimental conditions to obtain valid data.

The importance of the SSRT as a tool for identifying potential ranges or a critical potential for SCC isemphasized. This technique has been used successfully for determining critical potentials for theoccurrence of intergranular stress corrosion cracking (IGSCC) in sensitized austenitic stainless steels inhigh-temperature aqueous systems and in many other industrial applications. However, the use of crackvelocity data obtained in SSRTs for design or performance assessment is questioned because of thedifficulties involved in relating the crack velocity measured at a certain strain rate with that under the

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stress or strain conditions expected in the field. Also, the presence of multiple cracks in the gage sectionof tensile specimens alters the value of the applied strain rate. Attempts have been made (Ford, 1986;Lidbury, 1984) to relate the strain rate at the crack tip, which is considered a fundamental variable in thefilm rupture/slip dissolution model, with parameters such as the applied strain rate (Q) or the stressintensity (K,). This is based on the concept that, under constant load conditions, an actively growingcrack is experiencing a strain rate at its tip rather than purely static loading (Parkins, 1979; Ford 1986).The relationships are largely empirical and not confirmed by sufficient testing. Therefore, crack growthrates measured in precracked specimens using a fracture mechanics approach are considered more reliablethan those obtained in SSRTs. However, the crack velocities measured in slow strain rates are useful forcomparing materials and thermal treatments and for evaluating SCC susceptibility as a function ofinfluential environmental variables, such as temperature, pH, concentration of aggressive or inhibitingspecies, and potential.

As discussed in Section 2.4.4, the design approach to be taken for high-level waste (HLW) containersmust be based on propensity for crack initiation, rather than on crack propagation rate. This implies thatmaterials should be selected for the anticipated environmental conditions on the basis of relatively highcritical potentials for crack initiation, assuming that this concept can be demonstrated to be reliable. Theselected material must have a critical potential for cracking well above the predicted evolution of thecorrosion potential in the range of anticipated environments. Within this context, an issue that needsfurther evaluation is the concept of arresting an actively growing crack by decreasing the redox potentialof the environment, and, hence, the corrosion potential, below the critical potential for crack initiation.

As in the case of localized corrosion (Cragnolino, 1991a), besides the validity of using the criticalpotential for cracking as a bounding parameter, the issues of time and size extrapolation are crucial. Thevalidity of a test result beyond the test time depends critically on the application of a mechanistic model,and, therefore, parameters in the model should be related to those in the experiments. The test resultsshould also be extrapolatable in size, taking into consideration the size of each container and the largenumber of containers. If crack initiation is related to pitting or crevice corrosion in the environments ofinterest, methods of extreme value statistics can also be applied. In recent years, there have been manyattempts to use probabilistic approaches to SCC, and the problem of size scale-up should be consideredusing these methods.

Although many HLW test programs have been conducted under a variety of testing and environmentalconditions, it is still difficult to determine if SCC is indeed a viable mode of degradation for the candidatecontainer materials, especially for the high-nickel alloys such as alloy 825. Many questions remain tobe answered:

* Are the number of specimens tested in the various programs statistically significant incomparison to the time of exposure?

* What effects do environmental factors have, especially pertaining to the chemistry ofoccluded areas such as crevices? For example, Abraham et al. (1986) speculate that someof the cracks may have originated in crevices formed between the rock and the specimensurface.

* What effects do surface conditions have on SCC initiation? All the tests were performed onmachined surfaces; whereas, the surfaces of containers, especially the lateral surfaces, arenot likely to be machined. Mill-finished surfaces of some alloys may exhibit surface

5-2

'A

Cr-depletion (Cragnolino, 1990) which may lead to enhanced SCC susceptibility (Place,1991). This is not expected to be a significant problem for Types 304L and 316L stainlesssteels, but may be important for alloys 825 and C-4.

* Can regimes of electrode potential be defined for the candidate alloys in near-fieldenvironments wherein SCC will occur? This then can be coupled with measured orpredicted open-circuit potentials to predict cracking incidence. The investigations of Beaversand Durr (1990b) have attempted to do this using SSRTs. However, a similar approach isneeded using constant deflection tests.

* What are the relative susceptibilities to SCC of other candidate alloys in these tests? Whateffect does aging at temperatures of around 200'C during the "dry" initial period have uponsubsequent exposure to "wet" environment. Specifically, the Ni-based alloys such as alloyC-4 are known to undergo order-disorder transformation at these temperatures (Sridhar,1992b). The low-carbon stainless steels such as 304L and 316L have been reported (Strum,1988) to be susceptible to the growth of carbide precipitates in these temperature regimesaccompanied by Cr-depletion (i.e., low temperature sensitization).

* What effect does cold-work have? Cold-worked areas in dents, scratches, etc. areinevitable. The effect of cold work on SCC of the candidate Fe-Cr-Ni alloys has not beenclearly established, particularly for alloy 825. Additionally, cold-work has been shown topromote thermally-induced phase transformations, such as ordering and precipitation ofgrain-boundary carbides, that may be important in high-nickel alloys.

* Is the SSRT conservative? While many in the corrosion literature have shown this to be sofor a wide range of alloy-environment combinations (e.g., Parkins, 1984), more recentinvestigations (Beavers, 1992a) have indicated aspects of the technique that need furtherinvestigation. Again, duplicate or triplicate tests may not be statistically significant in lowchloride-containing environments.

The review of the general literature provides partial answers to some of these questions. It is apparentthat both Types 304L and 316L are significantly more susceptible to SCC in the presence of chloride thanalloy 825. Most of the data available refers to Type 304 stainless steel, and a minor proportion to Type316. However, the differences in the behavior of these alloys with respect to the low-carbon grades(Types 304L and 316L) are not significant for materials in the mill-annealed or solution-annealedcondition, as can be inferred from the results of Asphahani (1980). Although there are data (Sedriks,1979; Kowaka, 1990) indicating that Type 316 stainless steel is more resistant to cracking in chloridecontaining solutions than Type 304, the difference can be considered marginal for container applications.It is necessary, however, to quantify the margin of improvement obtained with alloy 825 and eventuallywith other potential candidate materials.

From the review of the literature, it can be concluded that phenomenological information and long-termtesting data is extremely limited in the case of alloy 825. This alloy is now being used more extensivelyin the oil and gas industry. Although not directly applicable, more information from field experience willbe available in the future. An experimental program on alloy 825 should examine in detail therelationship between crevice conditions and SCC initiation that is clearly valid for the austenitic stainlesssteels at relatively low chloride concentrations. It should be noted, for example, that the repassivationpotential for pitting corrosion of alloy 825 in 1000 ppm chloride-containing solutions at 950 C is about

5-3

100 mVscE (Cragnolino, 1991c). The repassivation potential for crevice corrosion is close to this value,and in the presence of H202, the corrosion potential reaches similar values, indicating that SCC may bepossible under such environmental conditions. The role of nitrate and silicates as potential SCC inhibitorsshould be studied, as well as the effect of heat transfer on the SCC susceptibility. The results for Type304 stainless steel suggest that the combined effects of heat transfer and the presence of crevices couldalso be important for nickel-base alloys.

The results in the literature do not clarify completely the role of strain rate in chloride SCC. Althoughsome results (Mancia, 1988) indicate that strain rates lower than 1 x 10-7/sec may be needed to detecttransgranular stress corrosion cracking (TGSCC) in austenitic stainless steel, it is possible that pre-exposure to the environment or other modifications in the surface conditions may lead to cracking,without discounting the important influence of potential, as discussed by Beavers (1992a).

The interrelated effect of several variables on SCC susceptibility, such as chloride concentration, pH,oxygen and hydrogen peroxide concentrations, potential, and temperature, requires a systematic approachfor testing a spectrum of plausible environmental conditions and a mechanistic understanding in order toselect the appropriate accelerating variables and their correct ranges of variation. This approach isapplicable to the nickel base alloys, such as alloy 825, and it may reveal the potential risk of SCC failuresin certain environmental conditions. The feasibility of these conditions should be evaluated to find outif SCC is a plausible degradation process for alloy 825 or other alternate candidate materials.

The use of SSRTs is recommended for determining environmental conditions and potential ranges forSCC of alloy 825. These tests must be coupled to constant deflection tests to evaluate the effect ofsurface conditions and long-term exposure on crack initiation. From the results of these short-term tests,environmental and surface conditions can be selected for long-term tests in which simulation of wet/drycycles should be incorporated. Constant-deflection tests, using U-bends, seem to be appropriate for long-term tests.

The results of the HLW SCC investigations on copper based alloys may be summarized as follows:

* SCC does not appear to be a major factor in these alloys except in the presence of highconcentrations of nitrites (and possibly nitrates) and relatively high potentials compared tothose caused by aeration alone. The combination of high nitrite concentrations and highpotentials does not seem to be a highly probable event in the repository.

* The effect of other environmental factors such as bicarbonate in combination with chlorideand sulfate needs to be investigated.

* While thermal effects on microstructural changes in terms of phase transformations are notanticipated for CDA-102 and CDA-715, the effects of grain-boundary S and P segregationdue to long-term thermal exposure have not been investigated.

The review of the general literature on copper alloys reveals that the bulk of the available informationis concentrated on ca-brass, probably due to its SCC susceptibility and extensive use in many industrialapplications. There is a long list of species, including neutral molecules and anions, that lead to SCC ofbrasses. However, most of these species (i.e., ammonia and several organic acids) are not expected tobe found under repository conditions in high enough concentrations to promote cracking. The list ofpotentially damaging species on the basis of the data available is much more limited for the copper-base

5-4

candidate container materials and essentially includes ammonia and nitrites for copper (CDA 102),ammonia and high-temperature steam for CDA 613, and only high-temperature steam for CDA 715.Although there is a mention in the literature of carbonates as cracking agent for a-brass, the results ofKawashima et al. (1979) contradict that information. Nevertheless, they have studied the effect ofbicarbonate within a very narrow range of conditions.

The SSRT method is appropriate and recommended for investigating SCC susceptibility of CDA-102 andCDA-715 in a range of chloride, sulfate, and bicarbonate concentrations at various applied potentials.For these copper-base alloys constant deflection test methods are also recommended to investigate theeffect of long-term exposure.

It can be concluded from this review that the concept of a critical potential for cracking and itsrelationship with the repassivation potential for crevice corrosion of Fe-Cr-Ni-Mo alloys should be fullyexplored to determine its usefulness as a bounding parameter for performance assessment. The role ofthe main species present in groundwater should be carefully addressed, as well as the relative importanceof pH, temperature, and heat transfer effects, in conjunction with an appropriate prediction of theevolution of the repository environment with time.

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6 NOMENCLATURE

i Strain ratef StraineP Plastic straina Stressath Threshold stress for SCCay Yield strengthBWR Boiling water reactorCERT Constant extension rate testCOD Crack opening displacementE., E. Corrosion potential or open circuit potentialEPR Electrochemical potentiokinetic reactivationER, E,.,. Repassivation potential for crevice corrosionHLW High-level wasteIASCC Irradiation-assisted stress corrosion crackingIGSCC Intergranular stress corrosion crackingK, Stress intensityK1, Critical stress intensity or fracture toughnessK1.cc Threshold stress intensity for SCCLEFM Linear elastic fracture mechanicsLVDT Linear voltage displacement transducerPWR Pressurized water reactorsSCC Stress corrosion crackingSEM Scanning electron microscopeSSRT Slow strain rate testtf Failure time for SCCTGSCC Transgranular stress corrosion cracking

Initiation time for SCCVSCB Volts in saturated calomel electrode scaleVSHE Volts in standard hydrogen electrode scaleWOL Wedge-opening-loading

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