Alloy690inPWR

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Materials Reliability Program: Resistance to Primary Water Stress Corrosion Cracking

of Alloy 690 in Pressurized Water Reactors (MRP-258)

Materials Reliability Program: Resistance to Primary Water Stress Corrosion Cracking of Alloy 690 in Pressurized Water Reactors (MRP-258)

1019086

Final Report, August 2009

EPRI Project Manager K. Ahluwalia

ELECTRIC POWER RESEARCH INSTITUTE 3420 Hillview Avenue, Palo Alto, California 94304-1338 • PO Box 10412, Palo Alto, California 94303-0813 • USA


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ORGANIZATION(S) THAT PREPARED THIS DOCUMENT

Dr. John Hickling, Independent Technical Consultant

NOTE

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CITATIONS

This report was prepared by

Dr. John Hickling, Independent Technical Consultant Off Makariou III Street Prastio-Avdimou CY-4601 Cyprus

Principal Investigator J. Hickling

This report describes research sponsored by the Electric Power Research Institute (EPRI).

The report is a corporate document that should be cited in the literature in the following manner:

Materials Reliability Program: Resistance to Primary Water Stress Corrosion Cracking of Alloy 690 in Pressurized Water Reactors (MRP-258). EPRI, Palo Alto, CA: 2009. 1019086.

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PRODUCT DESCRIPTION

Wrought Alloy 600 and its weld metals (Alloy 182 and Alloy 82) were originally used in pressurized water reactors (PWRs) due to the material’s inherent resistance to general corrosion in a number of aggressive environments and because of a coefficient of thermal expansion that is very close to that of low alloy and carbon steel. Over the last 30 years, stress corrosion cracking in PWR primary water (PWSCC) has been observed in numerous Alloy 600 component items and associated welds, sometimes after relatively long incubation times. The occurrence of PWSCC has been responsible for significant downtime and replacement power costs.

Component repairs and replacements have generally used wrought Alloy 690 material and its compatible weld metals (Alloy 152 and Alloy 52 or 52M), which have been shown to be highly resistant to PWSCC in laboratory experiments and have been free from cracking in operating reactors over periods already up to nearly 20 years. The challenge is to attempt to quantify the longevity of these materials with respect to aging degradation by corrosion in order to provide a sound technical basis for the development of future inspection requirements for repaired or replaced component items. This document updates EPRI report 1009801 (MRP-111), published in 2004, and represents an extensive revision that takes into account recently obtained information on the PWSCC behavior of Alloy 690 base material. The performance of Alloy 152 and 52 weld metals is not considered here, but will be reported on separately at a later date.

Some consideration is also given to corrosion fatigue and low temperature crack propagation of thick-walled Alloy 690 material, since these topics were identified as knowledge gaps in the 2004 report (MRP-111).

Approach Building on the original MRP-111 report, numerous laboratory tests conducted over the last two decades that were performed with wrought Alloy 690 materials under various test conditions pertinent to corrosion resistance in PWR environments were reviewed, with the main focus on PWSCC. Wherever possible, the existing laboratory test data were evaluated to estimate the improvement factor of Alloy 690 relative to Alloy 600. In addition, Alloy 690 service experience in PWRs is reported to augment the laboratory findings.

Results and Findings It is concluded that wrought Alloy 690 is an acceptable and highly corrosion-resistant replacement material for Alloy 600 in PWRs, although further testing is still needed to examine some specific knowledge gaps that have emerged regarding PWSCC. These include the detrimental effect on resistance to crack growth of inhomogeneous cold work (particularly unidirectional cold rolling/tensile straining) and the possibility of enhanced cracking susceptibility in the heat-affected zone following welding. Relative improvement factors of 40 – 100 times versus Alloy 600 can now be derived for the initiation of cracking,

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but these numbers are clearly conservative, due to an absence of PWSCC in almost all Alloy 690 specimens within the test duration. Prototypical thick-walled Alloy 690 material (for example, as extruded piping for reactor pressure vessel top head penetrations) is extremely resistant to SCC growth from a pre-existing fatigue crack in simulated primary water, even under laboratory test conditions designed to maximize susceptibility. Measured crack growth rates to date are so low (< 5E-9 mm/s or 0.15 mm/year) as to be of no engineering significance and the relative factor of improvement for Alloy 690 CRDM material with respect to Alloy 600MA is thought to be well over 100 times. However, significant rates of intergranular cracking have been observed under constant load in some tests following the introduction of non-uniform cold work into Alloy 690 plate and bar materials that are not thought to be generally representative of plant components. The boundaries of this susceptibility are still being probed in order to confirm that such findings are unlikely to be relevant to long-term PWR operation.

No stress corrosion degradation of Alloy 690 has been observed in any replacement application to date. Service experience for inspected material exposed to PWR primary water ranges from approximately 10 to 20 years, depending upon the type of component considered.

Applications, Value, and Use This report should be of interest to utility engineers and scientists concerned with all aspects of Alloy 600 cracking in PWR primary water and especially to those involved in developing inspection regimes, making decisions on component replacement, and dealing with plant aging issues. It will be of direct value in obtaining regulatory acceptance for the solutions adopted by the industry to deal with increasing incidences of PWSCC that affect thick-walled Alloy 600 components in existing plants.

As part of an ongoing, comprehensive program involving utilities, reactor vendors, and engineering/research organizations, this report will also help to ensure that corrosion degradation of nickel-based alloys does not limit service life and that full benefit can be obtained from improved designs for both replacement components and new reactors.

EPRI Perspective The report describes the current knowledgebase on Alloy 690’s PWSCC resistance. The data show that Alloy 690 possesses significantly greater resistance to PWSCC than its predecessor material, Alloy 600. Although these results are encouraging and are expected to hold as additional research continues, there are factors such as cold work that can impair its superiority somewhat. On-going and planned research is focused on these factors to more clearly define Alloy 690 resistance to PWSCC and identify any vulnerabilities. Until this additional work is completed, the results presented in this report are for information only and should not be used for prediction of component life for components constructed from Alloy 690.

Keywords Alloy 600 Alloy 690 PWSCC Material degradation RPV penetrations

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ABSTRACT

Over the last 30 years, stress corrosion cracking (SCC) in pressurized water reactor (PWR) primary water (PWSCC) has been observed in numerous Alloy 600 component items and associated welds, sometimes after relatively long incubation times. Repairs and replacements have generally used wrought Alloy 690 material, which has been shown to be extremely resistant to PWSCC in laboratory experiments and has been free from detected SCC in operating reactors over periods already up to nearly 20 years. It is nevertheless prudent for the PWR industry to attempt to quantify the longevity of this material (and its welding alloys) with respect to aging degradation by corrosion in order to provide a sound technical basis for the development of future inspection requirements for repaired or replaced component items.

This document updates EPRI report 1009801 (MRP-111), published in 2004, and represents an extensive revision that takes into account recently obtained information on the PWSCC behavior of Alloy 690 base material. The performance of Alloy 152 and 52a weld metals is not considered here, but will be reported on separately at a later date.

Some consideration is also given to corrosion fatigue and low temperature crack propagation of thick-walled Alloy 690 material, since these topics were identified as knowledge gaps in the 2004 report.

Building on the original MRP-111 report, numerous laboratory tests conducted over the last 25 years that were performed with wrought Alloy 690 materials under various test conditions pertinent to corrosion resistance in PWR environments are reviewed, with the main focus on PWSCC. Wherever possible, the existing laboratory test data have been evaluated to estimate the improvement factor of Alloy 690 relative to Alloy 600. In addition, Alloy 690 service experience in PWRs is reported to augment the laboratory findings.

It is concluded that wrought Alloy 690 is an acceptable and highly corrosion-resistant replacement material for Alloy 600 in PWRs, although further testing is still needed to examine some specific knowledge gaps that have emerged regarding PWSCC. These include the detrimental effect on resistance to crack growth of inhomogeneous cold work (particularly unidirectional cold rolling/tensile straining) and the possibility of enhanced cracking susceptibility in the heat-affected zone following welding. Relative improvement factors of 40 to 100 times can now be derived for initiation of cracking, but these numbers are clearly conservative, due to an absence of PWSCC in almost all Alloy 690 specimens within the test duration. Prototypical thick-walled Alloy 690

a Note that Alloy 52 exists with both original and modified compositions, whereby the latter are often referred to

as Alloy 52M or Alloy 52(M). These terms are used interchangeably in the current report, which is focused on Alloy 690.

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material (for example, as extruded piping for reactor pressure vessel top head penetrations) is extremely resistant to SCC growth from a pre-existing fatigue crack in simulated primary water, even under laboratory test conditions designed to maximize susceptibility. Measured crack growth rates to date are so low (< 5E-9 mm/s or 0.15 mm/year) as to be of no engineering significance and the relative factor of improvement for Alloy 690 CRDM material with respect to Alloy 600MA is thought to be well over 100 times. However, significant rates of intergranular cracking have been observed under constant load in some tests following the introduction of non-uniform cold work into Alloy 690 plate and bar materials that are not thought to be generally representative of plant components. The boundaries of this susceptibility are still being probed in order to confirm that such findings are unlikely to be relevant to long-term PWR operation.

No stress corrosion degradation of Alloy 690 plate and bar materials has been observed in any replacement application to date. Service experience for inspected material exposed to PWR primary water ranges from approximately 10 to 20 years, depending upon the type of component considered.

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ACKNOWLEDGMENTS

The present report was prepared by Dr. John Hickling, acting as an Independent Technical Consultant to EPRI. It represents an extensive revision and update of EPRI report 1009801 (MRP-111), originally published in March 2004. That earlier report was prepared by Framatome ANP (now AREVA) and had the following principal investigators:

H. Xu and S. Fyfitch, Framatome ANP, Inc. 3315 Old Forest Road P.O. Box 10935 Lynchburg, VA 24506-0935

P. Scott Framatome ANP, SAS Tour AREVA 92084 Paris La Défense, France

M. Foucault Framatome ANP, SAS Porte Magenta – BP 181 71205 Le Creusot Cedex, France

R. Kilian and M. Winters, Framatome ANP, GmbH Freyeslebenstr. 1 – 91058 TGM – P.O. Box 3220 – 91050 Erlangen, Germany

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CONTENTS

1 INTRODUCTION ....................................................................................................................1-1 1.1 Background..................................................................................................................1-1 1.2 Purpose and Scope of Revising MRP-111 ..................................................................1-2 1.3 Causes of Alloy 600 PWSCC.......................................................................................1-4 1.4 References ..................................................................................................................1-5

2 ALLOY 690 PROPERTIES AND METALLURGY..................................................................2-1 2.1 Material Specifications.................................................................................................2-1

2.1.1 Typical PWR Specifications for Thin-Walled Alloy 690 SG Tubing.....................2-3 2.1.2 Specification and Manufacture of Thick-Walled Alloy 690 Components.............2-4

2.2 Phase Diagram of Alloy 690 ........................................................................................2-4 2.3 Carbon Solubility and Dynamic Strain Aging ...............................................................2-5 2.4 Intergranular Carbide Precipitation and Sensitization..................................................2-8 2.5 Effect of Elevated Temperature Exposure .................................................................2-12 2.6 References ................................................................................................................2-13

3 CORROSION BEHAVIOR OF ALLOY 690 APART FROM PWSCC ....................................3-1 3.1 General Corrosion Tests in Primary Water ..................................................................3-1

3.1.1 SG Tubing by Sedricks et al. 1979......................................................................3-1 3.1.2 SG Tubing by K. Smith et al. 1985......................................................................3-2 3.1.3 SG Tubing by Yonezawa et al. 1985...................................................................3-2 3.1.4 Esposito et al. 1991.............................................................................................3-2 3.1.5 Alloy Oxidation Studies related to the Mechanism of PWSCC ...........................3-3

3.2 Corrosion Fatigue Tests in Primary Water...................................................................3-4 3.3 Corrosion Behavior in Secondary Water....................................................................3-15 3.4 Low Temperature Crack Propagation (LTCP) ...........................................................3-17

3.4.1 Origins of the Phenomenon ..............................................................................3-17 3.4.2 Recent Studies to Assess the Possible Relevance of LTCP to PWRs .............3-18

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3.5 References ................................................................................................................3-23

4 PWSCC OF THIN-WALLED SG TUBING..............................................................................4-1 4.1 Laboratory Testing.......................................................................................................4-1

4.1.1 Early Studies .......................................................................................................4-2 4.1.2 Single U-Bend Test in Saturated Hydrogen Water with B/Li...............................4-2 4.1.3 CERT Tests in Hydrogenated Water with or without B/Li ...................................4-5 4.1.4 RUB Test in Hydrogenated Steam......................................................................4-6 4.1.5 Weibull and Weibayes Analyses of the Test Results ..........................................4-8 4.1.6 Improvement Factor by Weibull Analysis ..........................................................4-15 4.1.7 Improvement Factor with Minimum Alloy 600 Crack Time................................4-18

4.2 Field Experience ........................................................................................................4-21 4.3 References ................................................................................................................4-25

5 PWSCC OF THICK-WALLED ALLOY 690 MATERIAL ........................................................5-1 5.1 Laboratory Testing.......................................................................................................5-2

5.1.1 Crack Initiation Studies .......................................................................................5-3 5.1.1.1 Testing in Simulated Primary Water by Mitsubishi Heavy Industries (MHI) ......................................................................................................................5-3 5.1.1.2 Testing in Pure Supercritical Water at the University of Michigan ................5-8

5.1.2 Crack Growth Rate Studies...............................................................................5-10 5.1.2.1 Testing of Alloy 690 Material not directly related to PWR Components .....5-10

5.1.2.1.1 Feasibility Studies in Simulated Primary Water by General Electric Global Research (GE-GRC).....................................................................5-11 5.1.2.1.2 Investigations in Simulated Primary Water at Argonne National Laboratory (ANL)...................................................................................................5-15 5.1.2.1.3 Additional Studies in Simulated Primary Water by General Electric Global Research (GE-GRC) as part of the MRP Test Program ...............5-20 5.1.2.1.4 Investigations in Simulated Primary Water at the University of Tohoku in Japan....................................................................................................5-28 5.1.2.1.5 Investigations in Simulated Primary Water at the Institute of Nuclear Safety Systems (INSS) in Japan .............................................................5-29 5.1.2.1.6 Further Investigations in Simulated Primary Water in Japan ..............5-30 5.1.2.1.7 Investigations by Bechtel-Bettis in Deaerated High-Temperature Water ............................................................................................................5-31 5.1.2.1.8 Investigations by Westinghouse in Supercritical Water with Additions of Boron, Lithium and Hydrogen............................................................5-39

5.1.2.2 Testing of Alloy 690 CRDM Material without Deliberate Cold Working ......5-44

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5.1.2.2.1 Investigations in Simulated Primary Water at Argonne National Laboratory (ANL)...................................................................................................5-45 5.1.2.2.2 Investigations in Simulated Primary Water at Pacific Northwest National Laboratory (PNNL)..................................................................................5-48 5.1.2.2.3 Investigations in Simulated Primary Water at Studsvik in Sweden .....5-52 5.1.2.2.4 Investigations in Simulated Primary water at MHI in Japan ................5-52 5.1.2.2.5 Investigations by Westinghouse in Supercritical Water with Additions of Boron, Lithium and Hydrogen............................................................5-52

5.1.2.3 Testing of Alloy 690 CRDM Material after Deliberate Cold Working...........5-53 5.1.2.3.1 Investigations in Simulated Primary Water by General Electric Global Research (GE-GRC)..................................................................................5-53 5.1.2.3.2 Investigations in Simulated Primary Water at Pacific Northwest National Laboratory (PNNL)..................................................................................5-58 5.1.2.3.3 Investigations by Westinghouse in Supercritical Water with Additions of Boron, Lithium and Hydrogen............................................................5-59

5.1.2.4 Testing of Heat Affected Zones (HAZ) from Welding of Alloy 690 Material ....................................................................................................................5-60

5.1.2.4.1 Investigations in High-Temperature Water at KAPL............................5-60 5.1.2.4.2 Investigations in Simulated Primary Water at Argonne National Laboratory (ANL)...................................................................................................5-60 5.1.2.4.3 Investigations in Simulated Primary Water at Studsvik in Sweden .....5-61 5.1.2.4.4 Investigations in Simulated Primary Water at CIEMAT in Spain .........5-62 5.1.2.4.5 Investigations in Simulated Primary Water at Tohoku University in Japan ............................................................................................................5-62

5.2 Field Experience ........................................................................................................5-63 5.3 References ................................................................................................................5-63

6 DISCUSSION..........................................................................................................................6-1 6.1 Resistance of Alloy 690 to PWSCC.............................................................................6-1 6.2 Other Aspects of Alloy 690 Corrosion Behavior.........................................................6-10 6.3 References ................................................................................................................6-10

7 CONCLUSIONS .....................................................................................................................7-1

A TRANSLATED TABLE OF CONTENTS .............................................................................. A-1

日本語 (Japanese)................................................................................................................ A-2

한국어/조선말 (Korean)...................................................................................................... A-19

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LIST OF FIGURES

Figure 2-1 Carbon solubility diagram for Alloy 690 and Alloy 600 from [18] ..............................2-7 Figure 2-2 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690

Heat NX4459HG (0.06%C) from [18] .................................................................................2-9 Figure 2-3 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690

Heat NX9217H (0.01%C) from [18]..................................................................................2-10 Figure 2-4 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690

Heat NX9780H (0.01%C) from [18]..................................................................................2-10 Figure 2-5 High-temperature tensile properties of annealed Alloy 690. Data shown are a

composite of cold-and hot-worked products in the annealed condition and taken from [2] .............................................................................................................................2-12

Figure 3-1 Japanese fatigue data for Ni-base alloys in air at room temperature from [11] ........3-5 Figure 3-2 Japanese fatigue data for Ni-base alloys in simulated PWR water at 325°C

from [11] .............................................................................................................................3-5 Figure 3-3 Japanese data from [11] on relationship between calculated factor of

environmental fatigue enhancement (Fen) and strain rate for Ni alloys in simulated PWR water at 325°C ..........................................................................................................3-6

Figure 3-4 Japanese data from [11] on relationship between calculated factor of environmental fatigue enhancement (Fen) and temperature for various materials in simulated LWR water at 325°C ..........................................................................................3-7

Figure 3-5 Japanese data from [11] showing results of model predictions for the corrosion fatigue behavior of Ni-base alloys in simulated PWR water at 325°C compared with experiments ...............................................................................................3-8

Figure 3-6 Corrosion fatigue initiation data from [12] in high-temperature, deaerated water for high chromium weld alloys and Alloy 690 compared with stainless steels..........3-9

Figure 3-7 Approach to the analysis of environmental effects on cyclic crack growth developed at ANL and now applied to PWSCC testing of Alloy 690 and its weld alloys (from [13]) ..............................................................................................................3-10

Figure 3-8 No environmental enhancement of cyclic crack growth seen for thermally treated Alloy 690 material in either simulated PWR primary water or de-aerated pure water at 320°C [14], but under loading conditions not expected to favor EAC ........3-10

Figure 3-9 Appearance of significant environmental enhancement at lower rates of cyclic crack growth for unidirectionally cold-rolled Alloy 690 material (blue points) in simulated PWR primary water at 320°C (from [14]) .........................................................3-11

Figure 3-10 Japanese data from [16] indicating that a simulated PWR primary environment can enhance the fatigue crack growth rate of Alloys 600 and 690 by 5 to 10 times over a range of test conditions ......................................................................3-12

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Figure 3-11 Comparison of lines from a proposed (very conservative) Japanese model [16] for cyclic crack growth of Ni-base alloys in a PWR environment and experimental data that breach existing ASME curves......................................................3-13

Figure 3-12 Cyclic CGR behavior of Alloys 600 and 690 in Swedish studies [15] ...................3-14 Figure 3-13 Cyclic CGR threshold determined in simulated primary water for Alloy 690 in

Swedish studies [15] ........................................................................................................3-14 Figure 3-14 Classification scheme for categories of fracture resistance to LTCP after

Brown and Mills [36].........................................................................................................3-19 Figure 3-15 J-R curves determined by Brown and Mills for Alloy 690 in RT air and water

at various temperatures [36] ............................................................................................3-20 Figure 3-16 JIC and T values determined by Brown and Mills [36] for Alloy 690 in air and

water at various temperatures (values within bars indicate the dissolved hydrogen concentration) ..................................................................................................................3-21

Figure 3-17 Comparison of fractography from J-R testing under various conditions by Brown and Mills [36] of Alloy 690 (top) and Alloy 600 (bottom)........................................3-22

Figure 3-18 Results by Paraventi and Moshier [42] from J-R testing of Alloy 690 plate with additional, non-uniform cold work in 50°C water with varying contents of dissolved hydrogen ..........................................................................................................3-23

Figure 4-1 Vickers hardness number as a function of cold reduction % from [2]. Alloy 690 has a higher work-hardening rate than Alloy 600...............................................................4-5

Figure 4-2 Weibull plot from [2] of early Alloy 600 RUB results in primary water at 360°C generated in France. The Alloy 690 (three heats) Weibayes line assumes β = 5.0...........4-9

Figure 4-3 Weibull plot from [2] for RUB test results in deaerated water at 365°C reported by Norring et al. [17]. The Alloy 690 (many heats) Weibayes line assumes β = 5.0 ..............................................................................................................................4-10

Figure 4-4 Weibull plot from [2] for RUB test results by Norring et al. [17] on special production heats of Alloy 600 in deaerated water at 365°C. The Alloy 690 Weibayes line assumes β = 5.0 ........................................................................................................4-11

Figure 4-5 Weibull comparison from [2] for different tube diameters from RUB testing by Norring et al. [17] in deaerated water at 365°C. The Alloy 690 Weibayes line assumes β = 5.0...............................................................................................................4-11

Figure 4-6 Weibull plots from [2] of Japanese Alloy 600MA (one heat) CLT results at 340°C (644°F) in primary water. No failure was observed in any Alloy 600MA or Alloy 600TT CLT specimens tested at 320°C (608°F) and in the Alloy 690TT (one heat) CLT specimens tested at 360°C (680°F). The Weibayes lines for the unfailed specimens assume β = 5.0 ..............................................................................................4-12

Figure 4-7 Comparison from [2] of Japanese data for 20% pre-strained RUB specimens of Alloy 600MA or Alloy 600TT, tested in primary water at 320°C (608°F), with those for Alloy 690TT, tested under the same conditions, but at 360°C (680°F). The Alloy 690TT Weibayes line assumes β = 5.0 ............................................................................4-13

Figure 4-8 Weibull plot from [2] of RUB results at 360°C in primary water reported by Vaillant et al. [18]. The Alloy 600 RUB specimens were from four different heats in the MA and TT conditions. The Alloy 690 RUB specimens, also from four different heats in the MA and TT conditions, experienced no failure after up to 54,000 hours of exposure. The Alloy 690 Weibayes line assumes β = 5.0............................................4-14

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Figure 4-9 Weibull plot from [2] of SG mockups tested in deaerated water at 680°F by Framatome ANP, France. Alloy 690TT SG mockups experienced no failure after 100,000 hours of exposure. The Alloy 690 Weibayes line assumes β = 5.0....................4-14

Figure 4-10 Weibull θ and β factors for Alloy 600 tests listed in Table 4-3 ..............................4-17 Figure 4-11 Improvement factors listed in Table 4-2 and Table 4-4 per Equation 4-8

versus test duration..........................................................................................................4-19 Figure 4-12 Swedish RUB testing for crack initiation in Alloys 600 and 690 from [24] ............4-21 Figure 4-13 Worldwide causes of Alloy 600TT SG tube repair by degradation

mechanism from [31]........................................................................................................4-23 Figure 4-14 Worldwide causes of Alloy 690TT SG tube repair by degradation

mechanism from [31]........................................................................................................4-23 Figure 5-1 MHI test loop for uni-axial constant load studies of PWSCC initiation [4, 5].............5-5 Figure 5-2 “Active” loading mechanism for uni-axial constant load studies of PWSCC

initiation [4, 5] .....................................................................................................................5-6 Figure 5-3 Test specimens for uni-axial constant load studies of PWSCC initiation [4, 5].........5-6 Figure 5-4 Dependence of PWSCC initiation in Alloy 600MA on applied stress and lack

of cracking in 690TT BMI material after 58,000 hours of testing [4, 5]...............................5-7 Figure 5-5 Dependence of PWSCC initiation in Alloy 600MA on applied stress and lack

of cracking in 690TT CRDM nozzle material after 73,000 hours of testing [4, 5] ...............5-7 Figure 5-6 CGRs derived from a CERT study of various austenitic alloys in pure, de-

aerated SCW [6].................................................................................................................5-8 Figure 5-7 Cross-section of EPRI Alloy 600 after testing in 400°C/25.4MPa deaerated

pure SCW [7]......................................................................................................................5-9 Figure 5-8 Cross-section of EPRI Alloy 690 after testing in 400°C/25.4 MPa deaerated

pure SCW [7]....................................................................................................................5-10 Figure 5-9 Convention with regard to specimen orientation and the principle axis of cold-

working in Alloy 690 base material...................................................................................5-12 Figure 5-10 CGR response of cold-worked Alloy 690 plate (with low-temperature mill

anneal) tested at GE-GRC for >3000 h at constant stress intensity [9] ...........................5-13 Figure 5-11 CGR response of cold-worked Alloy 690 plate (with high-temperature mill

anneal) tested at GE-GRC for >3000 h at constant stress intensity [9] ...........................5-14 Figure 5-12 Predominantly transgranular morphology (but with some intergranular

facets) within band of PWSCC crack growth in cold-worked Alloy 690 plate initially tested at GE-GRC [9] .......................................................................................................5-15

Figure 5-13 Details of sample removal from cold-rolled Alloy 690TT plate tested at ANL [11 to 15] ..........................................................................................................................5-16

Figure 5-14 On-line data from ANL testing [11 to 15] of cold-rolled Alloy 690TT plate (S-L specimen orientation).......................................................................................................5-17

Figure 5-15 On-line data from ANL testing [11 to 15] of cold-rolled Alloy 690TT plate (S-T specimen orientation)....................................................................................................5-17

Figure 5-16 Macro- and microfractography for the ANL 690TT plate specimen with S-L orientation [11 to 15] ........................................................................................................5-18

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Figure 5-17 Macrofractography for the ANL 690TT plate specimen with S-T orientation [11 to 15] ..........................................................................................................................5-19

Figure 5-18 ANL test showing that the CGR behavior for cold-worked Alloy 690TT plate was unaffected by changing temperature within the range 320 to 300°C [13] .................5-20

Figure 5-19 Rapid crack growth in 1D cold-rolled (~26%) Alloy 690 plate supplied by ANL and tested at GE-GRC [17, 18] ................................................................................5-21

Figure 5-20 Second period of rapid PWSCC in 1D cold-rolled Alloy 690 plate from ANL and apparent lack of a CGR response to reducing temperature from 360 to 325, then to 290°C [17, 18] ......................................................................................................5-22

Figure 5-21 Change in macroscopic appearance of PWSCC region upon reducing temperature (c372) [17, 18]..............................................................................................5-22

Figure 5-22 Repeat test at GE-GRC on 1-D rolled ANL plate material showing a slight increase in CGR upon raising the dissolved hydrogen concentration [17, 18].................5-23

Figure 5-23 Moderately rapid PWSCC in 20% 1D cold-rolled Alloy 690 GE-GRC forged bar [17, 18] .......................................................................................................................5-24

Figure 5-24 Continued testing of 20% 1D cold-rolled Alloy 690 GE-GRC forged bar with reduction in both test temperature and applied stress intensity [17, 18] ..........................5-24

Figure 5-25 Lack of response to a major change in dissolved hydrogen during testing of a second specimen from a 20% 1-D cold-rolled Alloy 690 forged bar [17, 18] ................5-25

Figure 5-26 Results of testing the same material as in Figure 5-23, but this time in the S-T orientation [17, 18] ........................................................................................................5-25

Figure 5-27 Results from a further heat of Alloy 690 tested after 1-D cold-rolling at GE-GRC, this time showing expected response to a drop in temperature even for an S-L oriented specimen [17, 18]............................................................................................5-26

Figure 5-28 High-resolution fractography (c372) of the 26% 1D-cold rolled ANL 690 plate tested at GE-GRC [17, 18] ...............................................................................................5-27

Figure 5-29 High-resolution fractography of a 20% 1-D cold rolled Alloy 690 specimen from forged bar [17, 18]....................................................................................................5-27

Figure 5-30 Example of crack growth observed at INSS in 20% cold worked 690TT at 360°C (from [8])................................................................................................................5-30

Figure 5-31 Apparent effect of temperature and degree of cold work (CW) on measured CGR for 690TT (from [8]) .................................................................................................5-31

Figure 5-32 Increase in CGRs with increasing level of uni-directional cold-rolling in Bettis studies [22, 23].................................................................................................................5-35

Figure 5-33 Cold-rolling is more detrimental than tensile pre-straining (despite lower yield strength) [22, 23]......................................................................................................5-35

Figure 5-34 Effect of test temperature and degree of cold work on CGRs for the VIM/ESR plate [22, 23].....................................................................................................5-36

Figure 5-35 Apparent increase in CGRs at 50 cc/kg hydrogen (blue symbols) vs. 23 cc/kg (pink symbols) for Alloy 690 (left) and comparison (right) with opposite behavior for Alloy 600 [22, 23] .........................................................................................5-37

Figure 5-36 Predominantly intergranular crack advance in an S-L-oriented specimen of VIM/ESR TT plate subjected to only 12% cold-rolling [22, 23].........................................5-37

xviii

Figure 5-37 Detail from fractography of an S-T-oriented specimen of VIM/ESR plate subjected to 24% cold-rolling [22, 23] ..............................................................................5-38

Figure 5-38 Comparison of Alloy 600 and 690 CGRs measured by Bettis [22, 23] .................5-38 Figure 5-39 Attempted back-extrapolation by Bettis of cold-worked Alloy 690 CGRs on

the basis of assumed yield strength dependencies [22, 23].............................................5-39 Figure 5-40 Example of transition from transgranular cracking during cyclic loading in

SCW to intergranular SCC at ~ constant load in a 10% cold-worked specimen of Alloy 690TT [25] ...............................................................................................................5-43

Figure 5-41 Average SCC growth rate vs. stress intensity factor for Alloy 600 and 690 materials [25]....................................................................................................................5-43

Figure 5-42 Comparison of CGR measured for Alloy 600 MA and Alloy 690 TT in supercritical and primary water [25] .................................................................................5-44

Figure 5-43 Specimen orientation in ANL testing of Alloy 690 CRDM material [11 to 15] .......5-45 Figure 5-44 Reported CGRs during brief test periods at constant load in ANL work on

CRDM material [11 to 15].................................................................................................5-46 Figure 5-45 Absence of clear intergranular cracking on fracture surface of the first

CRDM specimen tested at ANL [11 to 15] .......................................................................5-47 Figure 5-46 Reported CGRs during constant-load test periods at two test temperatures

in further ANL work on CRDM material without deliberate cold working [11 to 15]..........5-48 Figure 5-47 Orientations of CT specimens used by PNNL for CGR testing on two Alloy

690 forgings: left is heat RE243; right is heat WP140 [26]...............................................5-49 Figure 5-48 PNNL data showing the patience used in transitioning to constant K loading

[26] ...................................................................................................................................5-50 Figure 5-49 PNNL data showing Alloy 690 CGR response for as-received TT versus a

carbide-modified SA condition [26] ..................................................................................5-50 Figure 5-50 PNNL data showing Alloy 690 CGR response for two further heats of

material in the as-received TT condition [26] ...................................................................5-51 Figure 5-51 PNNL fractography showing IG cracking limited to isolated grains in

specimens without cold-work [26] ....................................................................................5-51 Figure 5-52 Specimen location in GE testing of Alloy 690 CRDM material [17, 18] ................5-54 Figure 5-53 Data from one of two specimens of CRDM material with 20% homogeneous

cold work tested at GE-GRC [29] .....................................................................................5-55 Figure 5-54 Intergranular crack morphology in 20% cold-worked Alloy 690 CRDM

specimen [29]...................................................................................................................5-55 Figure 5-55 Data from a further specimen of CRDM material with 41% homogeneous

cold work showing tendency to crack arrest before reaching constant K conditions (top) and very low CGRs even under periodic partial unloading with a 24h hold time (bottom) [29].....................................................................................................................5-56

Figure 5-56 Macro – and micrfractography from the specimen of CRDM material with 41% homogeneous cold work showing extensive out-of-plane secondary cracking [29] ...................................................................................................................................5-57

Figure 5-57 Data from PNNL testing of CRDM material with deliberate cold working (here 17% in the S-L orientation) [26] ..............................................................................5-58

xix

Figure 5-58 Data from PNNL testing of CRDM material with deliberate cold working (here 30% in the T-L orientation) [26] ..............................................................................5-59

Figure 5-59 Details of an Alloy 690 HAZ specimen (CF690) under test at ANL [13] ...............5-61 Figure 5-60 Details of CT specimen being used at Studsvik to examine the CGR

behavior in Alloy 690 HAZ material [27] ...........................................................................5-62 Figure 6-1 Summary of laboratory SCC CGR data (as of November 2008) prepared by

PNNL [26] and showing the possibility of measuring moderate to high rates in Alloy 690 plate material subjected to non-uniform cold work ......................................................6-7

Figure 6-2 Evidence of microstructural banding in some areas of the 1-D cold-rolled ANL Alloy 690 plate material examined at GE-GRC [18] ...........................................................6-7

Figure 6-3 Microstructural banding perpendicular to the crack plane in Alloy 690 plate material from EPRI orginally tested for the MRP Program by GE-GRC [9]........................6-8

Figure 6-4 Very uniform, homogeneous microstructure in extruded Alloy 690 CRDM material [18] .......................................................................................................................6-8

xx

LIST OF TABLES

Table 1-1 First reported occurrence of Alloy 600 PWSCC for various PWR component items...................................................................................................................................1-2

Table 2-1 ASME specifications of Alloys 690 and 600 ..............................................................2-1 Table 2-2 ASME chemical composition requirement (wt%).......................................................2-2 Table 2-3 ASME room temperature properties ..........................................................................2-2 Table 2-4 ASME elevated temperature properties for Alloy 690 and Alloy 600.........................2-3 Table 2-5 Chemical composition (wt%) of Alloy 690 heats used by Sarver et al. [18] ...............2-6 Table 2-6 Effect of heat treatment on Alloy 690 carbide precipitation [18] ..............................2-11 Table 3-1 Estimated typical improvement factors vs. pHT considering all environments

[33, 34] .............................................................................................................................3-17 Table 3-2 Chemical composition of Alloys 600, 690, 82, & 52 tested by Brown and Mills

[36] ...................................................................................................................................3-18 Table 4-1 Examples of surface IG cracking in Alloy 690 and Alloy 52 specimens [2]................4-3 Table 4-2 Summary of Alloy 690 primary water stress corrosion test data to 2004...................4-7 Table 4-3 Weibull analysis for Alloy 600 tested with Alloy 690 [2] ...........................................4-15 Table 4-4 Summary of Alloy 690 hydrogenated and doped hydrogenated steam stress

corrosion test data to 2004...............................................................................................4-19 Table 4-5 Swedish RUB testing for crack initiation in Alloys 600 and 690 [23]........................4-20 Table 4-6 List of operating steam generators manufactured with Alloy 690 tubing (as of

December 2008) ..............................................................................................................4-24 Table 5-1 Examples of some replaced PWR RPV heads with CRDM penetrations in

Alloy 690 ............................................................................................................................5-1 Table 5-2 Examples of some relatively thick-walled Alloy 690 reactor coolant system

original equipment or replacement component items other than CRDM penetrations.......5-2 Table 5-3 Origins and heat treatments of Alloy 690 CRDM nozzles tested by EdF [3]..............5-3 Table 5-4 Chemical composition of materials used in MHI testing for PWSCC initiation

[4, 5] ...................................................................................................................................5-4 Table 5-5 Heat treatment and mechanical properties of materials used in MHI testing for

PWSCC initiation [4, 5].......................................................................................................5-4 Table 5-6 Environmental test conditions used in MHI testing for PWSCC initiation [4, 5] .........5-5 Table 5-7 Summary of maximum crack depths measured on cross-sectioned samples at

the University of Michigan [7] .............................................................................................5-9 Table 5-8 Chemical composition and mechanical properties of all the Alloy 690 materials

tested at GE-GRC in the MRP program...........................................................................5-11

xxi

Table 5-9 Chemical composition of Alloy 690TT plate material tested at ANL [11 to 15] ........5-15 Table 5-10 Material test matrix for Bettis Alloy 690 test program [22, 23] ...............................5-33 Table 5-11 Bettis summary of relative SCC susceptibility for Alloy 690 [22, 23]......................5-34 Table 5-12 Alloy 690 materials used in Westinghouse testing for CGRs in supercritical

water [25] .........................................................................................................................5-40 Table 5-13 Chemical composition of Alloy 690 materials used in Westinghouse testing

for CGRs in supercritical water [25] .................................................................................5-40 Table 5-14 Tensile properties (as reported by material vendors) of the Alloy 690

materials used in Westinghouse testing for CGRs in supercritical water [25] ..................5-41 Table 5-15 Tensile properties after trial forging of the Alloy 690 materials used in

Westinghouse testing for CGRs in supercritical water [25] ..............................................5-41 Table 5-16 Detailed Results for the Alloy 600 control samples and Alloy 690 plate

material used in Westinghouse testing for CGRs in supercritical water [25]....................5-42 Table 5-17 Detailed results for the Alloy 600 control samples and “as-received” Alloy 690

CRDM materials used in Westinghouse testing for CGRs in supercritical water [25] ......5-52 Table 5-18 Detailed results for the Alloy 600 control samples and Alloy 690 cold-worked

CRDM materials used in Westinghouse testing for CGRs in supercritical water [25] ......5-60 Table 6-1 Summary of results of CGR testing in simulated PWR primary water on Alloy

690 CRDM material without deliberate cold working (status: December 2008) .................6-3 Table 6-2 Summary of results of CGR testing in simulated PWR primary water on Alloy

690 CRDM material with deliberate cold working (status: December 2008) ......................6-4

xxii

LIST OF ACRONYMS

ALARA As low as reasonably achievable

APFIM Atom probe field ion microscope (microscopy)

ATEM Analytical transmission electron microscope (microscopy)

ASME American Society of Mechanical Engineers

ASTM American Society for Testing and Materials

AWS American Welding Society

AVT All volatile treatment

CEA Commisariat à l`Énergie Atomique (the French Atomic Energy Commission)

CGR Crack growth rate

CL Constant load

CLT Constant load test

CEDM Control element drive mechanism

CERT Constant extension rate test, also known as slow strain rate test (SSRT)

CM Carbide modified (heat treatment used in laboratory testing)

CRDM Control rod drive mechanism

CT Compact tension (specimen)

DCB Double cantilever beam specimen

DSA Dynamic strain aging

xxiii

DSC Differential scanning calorimetry

DTA Differentential thermal analysis

EdF Electricité de France (the French national electric utility)

EFPH Effective full power hours

EFPY Effective full power years

EPRI Electric Power Research Institute

FOI Factor of (relative) improvement (also known as IFR)

GTA Gas tungsten arc (welding)

HAZ Heat-affected zone

ID Inside diameter

IG Intergranular

IGA Intergranular attack

IGSCC Intergranular stress corrosion cracking

LTCP Low temperature crack propagation

MA Mill annealed

MRP Materials reliability program

NRC Nuclear Regulatory Commission

OD Outside diameter

OTSG Once-through steam generator

ppm Parts per million

ppb Parts per billion

PPU Periodic partial unloading (trapezoidal waveform used in CGR testing)

PWR Pressurized water reactor

xxiv

PWHT Post weld heat treatment

PWSCC Primary water stress corrosion cracking

RUB Reverse U-bend stress corrosion cracking specimens made from split half steam generator tubing, hence also known as split tube U-bend specimens

SCC Stress corrosion cracking

SCW Supercritical water

SEM Scanning electron microscope (or microscopy)

SG Steam generator

SSRT Slow strain rate test, also known as constant extension rate test (CERT)

STEM Scanning transmission electron microscope (or microscopy)

TEM Transmission electron microscope (or microscopy)

TGSCC Transgranular stress corrosion cracking

TT Thermal treatment

UNESA Spanish association of nuclear utilities

UNS Unified numbering system

WOL Wedge opening loading (specimen)

XPS X-ray photoelectron spectroscopy

1-D Unidirectional

xxv

1 INTRODUCTION

1.1 Background

Wrought Alloy 600 and its weld metals (Alloy 182 and Alloy 82) were originally used in pressurized water reactors (PWRs) due to the material’s inherent resistance to general corrosion in a number of aggressive environments and because of a coefficient of thermal expansion that is very close to that of low alloy and carbon steel. Over the last thirty-plus years, primary water stress corrosion cracking (PWSCC) has been observed in Alloy 600 component items and Alloy 82/182 welds such as steam generator tubes and plugs, pressurizer heater sleeves and welds, pressurizer instrument nozzles, reactor vessel closure head nozzles and welds, reactor vessel outlet nozzle welds, and more recently in a lower reactor vessel head instrumentation nozzle and weld. Table 1-1 provides a synopsis of the Alloy 600 PWSCC experience in commercial PWRs up to 2003. This table identifies the first commercially observed occurrence of PWSCC for each particular component item in a PWR and lists the approximate service life (in calendar years) at the time PWSCC was identified at that particular location. PWSCC was first observed at very highly stressed tube locations in the hot leg of steam generators in the 1970s. Pressurizer nozzles, which operate at the highest temperature in PWRs, were the next locations to have leakage and failures identified. Currently, PWSCC has been observed in nozzles and welds at nearly all locations where Alloy 600 is utilized throughout the reactor coolant system. The particularly complex service experience with cracking of Alloy 182 and 82 weld metals was recently reviewed [1].

The occurrence of PWSCC has been responsible for significant downtime and replacement power costs at PWRs. Notable examples of equipment failures include extended outages and repairs or replacements at Calvert Cliffs, V.C. Summer, Oconee Nuclear Station, Davis-Besse, and North Anna. Repairs and replacements since the late 1980s have generally utilized wrought Alloy 690 material and its weld metals (Alloy 152 and Alloy 52(M)), which have been shown to be considerably less susceptible to PWSCC in laboratory experiments and are thought to be resistant for all practical purposes under normal operating conditions. Nevertheless, U.S. NRC Bulletin 2002-02, “Reactor Vessel Head and Vessel Head Penetration Nozzle Inspection Programs,” dated August 9, 2002, originally indicated essentially that no inspection credit would be given for upgraded replacement reactor vessel heads, using Alloy 690 wrought and weld materials, over existing head penetrations using Alloy 600 wrought and weld materials. This situation has changed recently with the issuance of modifications to U.S. NRC Regulation 10 CFR Part 50 [2] that now reference ASME Code Case N729-1.

1-1

Introduction

Table 1-1 First reported occurrence of Alloy 600 PWSCC for various PWR component items

Component Item Date PWSCC Initially Observed

Service Life(a)

(Calendar Years)

Steam Generator Hot Leg Tubes 1971 2

Pressurizer Instrument Nozzles 1986 2

Steam Generator Cold Leg Tubes 1986 18

Pressurizer Heaters and Sleeves 1987 5

Steam Generator Channel Head Drain Pipes 1988 1

Pressurizer Heater Diaphragm Plate Weld 1989 16

Control Rod Drive Mechanism Nozzles 1991 12

Hot Leg Instrument Nozzles 1991 5

Power Operated Relief Valve Safe End 1993 22

Pressurizer Nozzle Welds (Repair) 1994 1

Steam Generator Tubesheet Plate Cladding 1995 13

Cold Leg Piping Instrument Nozzles(b) 1997 13

Hot Leg Nozzle Welds 2000 13

Reactor Vessel Hot Leg Nozzle Buttering/Piping Welds 2000 17

Pressurizer Instrumentation Nozzle Welds 2000 27

Control Rod Drive Mechanism Nozzle/RV Head Welds 2000 27

Surge Line Nozzle Welds 2002 21

Reactor Vessel Lower Head In-Core Instrumentation Nozzles/Welds

2003 14

(a) This listing identifies the first reported occurrence of identified cracking for each component item. Leakage has occurred in some component items in less than one year of service life and in other component items after nearly 30 years of service.

(b) One plant identified “suspect” visual evidence of boric acid leakage around two nozzles during a visual inspection; nozzles were preventively repaired without investigating whether leakage had in fact occurred.

1.2 Purpose and Scope of Revising MRP-111

Based on excellent test results and field performance, Alloy 690 has become the replacement material of choice for degraded Alloy 600 component items in PWRs. The objective of the original MRP-111 report [3], published in 2004, was to document everything that was then known about the resistance of Alloy 690 and its weld metals to PWSCC, both from laboratory testing and field experience, and to quantify factors of improvement in behavior over Alloys 600, 182 and 82. The summary conclusions reached were as follows:

1-2

Introduction

“Wrought Alloy 690 and its weld metals (Alloys 52 and 152) are acceptable and highly corrosion-resistant replacement materials for Alloy 600 and its weld metals in PWRs, although limited, further testing is needed to examine some specific knowledge gaps that have been identified. Wherever possible, the existing laboratory test data have been evaluated to estimate the improvement factor of Alloy 690 relative to Alloy 600. Average improvement factors of at least 26 relative to Alloy 600MA material and 13 relative to Alloy 600TT material can be derived, but these numbers are clearly conservative, due to an absence of PWSCC in most Alloy 690 specimens within the test duration.”

Of the various knowledge gaps mentioned in 2004, the most important clearly involved the Alloy 52 and 152 weld metals, where only one investigation each had been reported concerning PWSCC testing. The second most important gap concerned the possible effects of Alloy 690 product form (plate, tube, rolled bar, forged bar, and extruded bar), since nearly all testing up to that time had been performed with specimens obtained from thin-walled steam generator tubing. Mention was made in [3] of possible effects from subtle changes in chemical composition and thermo-mechanical processing. It was also noted in 2004 that

“Given the welcome difficulties in initiating cracking of Alloy 690 and its weld metals, no plausible estimates of crack growth rates for base metal, HAZ and weld metals are available at this time.”

Furthermore, a lack of data was identified concerning the combined effects of exposure to primary water and fatigue on Alloy 690 and its weld metals, although it was postulated that models developed for Alloy 600 may also be applicable here. Finally, it was mentioned that additional efforts might be needed to confirm the lack of relevance of low temperature crack propagation (LTCP) for Alloys 690/152/52 in PWR primary water, even though the available operating experience did not suggest that a practical problem exists.

Extensive studies, particularly on the crack growth behavior of thick-walled Alloy 690 materials, have been carried out since 2004 and were summarized in 2008 [4]. Many of these investigations are still ongoing and, together with the research branch of the U.S. NRC, EPRI is currently leading efforts to establish an international research collaboration on this topic. Nevertheless, it is judged that most of the abovementioned knowledge gaps have already been sufficiently closed for Alloy 690 base material to make revision of MRP-111 worthwhile. Note, however, that this revised report does not cover the weld metals (Alloys 152 and 52(M)), since too much additional data on these is expected shortly. These results will be incorporated in a separate report, currently scheduled for publication in 2010.

The present report includes some of the original material from MRP-111, but does not entirely replace it. The report structure has been re-arranged considerably to help focus on the key issue of demonstrating adequate, long-term, PWSCC resistance of thick-walled Alloy 690 components exposed to primary water environments. It is also intended to provide a technical basis for future development of the most appropriate inspection requirements for these materials.

1-3

Introduction

1.3 Causes of Alloy 600 PWSCC

Stress corrosion cracking of metals and alloys is caused by the synergistic effects of environment, material condition, and stress. In a PWR primary water environment, intergranular stress corrosion cracking of wrought Alloy 600 material and its weld metals (Alloy 182 and Alloy 82) is commonly referred to as PWSCC. The occurrence of stress corrosion cracking of Alloy 600 in high-purity water has been extensively studied since the first reported observation of cracking in laboratory tests by Coriou et al. [5] in 1959. The mechanism of this cracking phenomenon is still not completely understood, and prediction of crack initiation time has proven to be very difficult due to the uncertainty of numerous contributory variables. These include metallurgical condition, cold work, and residual stress. Although the crack initiation time can vary tremendously from heat to heat, wrought Alloy 600 material and its weld metals are generally susceptible to PWSCC when the total stress level is close to, or exceeds, the yield strength at operating temperatures.

PWSCC is a thermally-activated mechanism that can be correlated with an Arrhenius relationship (exponential) and is very temperature dependent. The vast majority of PWSCC at steam generator roll expansion transitions has occurred first on the hot-leg side of the tubesheet due to the 27-38°C (50-70°F) higher temperatures. However, failures of Alloy 600 material have also been reported in France [6] to have occurred in reactor vessel upper head nozzle material at a temperature of approximately 290°C (554°F). On at least one other occasion, PWSCC has been cited on a component item at a significantly lower water temperature of 217°C (423°F), but the details leading to this conclusion have not been independently verified [7].

The susceptibility of Alloy 600 depends on several factors including the chemical composition, metallurgical condition during manufacture of the material, heat treatment during fabrication of the component item, and its operating parameters [8]. The carbon and chromium contents appear to be the most important chemical composition variables. These, in turn, affect chromium carbide precipitation during thermo-mechanical processing. Microstructural conditions, such as grain size and location relative to carbide precipitation, are also important variables that determine the susceptibility of a particular material to PWSCC. Finally, fabrication parameters and heat treatment determine the overall yield strength and degree of cold work. Alloy 600 that has been low temperature mill-annealed, with grain boundaries poorly decorated with carbides, and has relatively high yield strength (due, e.g., to some degree of remaining cold work) is generally observed to be the most susceptible to PWSCC. Work is ongoing in the USA on an empirical/theoretical hybrid model [9] that shows promise in predicting the crack propagation rate, in particular of thick-walled Alloy 600 material, on the basis of the different engineering tensile properties of each heat (while taking into account changes in applied stress intensity factor, dissolved hydrogen, and water temperature). Efforts also continue in France to refine engineering predictions of the initiation of PWSCC in Alloy 600 on the basis of so-called “susceptibility index models” [10, 11].

Tensile stresses, resulting from both residual and operating stresses, can be significant for some Alloy 600 component items. A stress close to the high-temperature material yield strength is generally necessary for PWSCC to initiate. Operating stresses arise from mechanical (pressure) and thermal loading, while residual stresses are generated as a result of fabrication, installation, and welding processes. Residual stresses are more difficult to quantify than operating stresses. In many instances, they are of a higher magnitude and usually a major factor leading to premature failure. Note, however, that initiation of cracking may still occur only after very long periods of operation, particularly in heats of Alloy 600 exhibiting low susceptibility.

1-4

Introduction

Once PWSCC initiates in some thin wall components, such as pressurizer heater sleeves, the part concerned has essentially reached the end of its service lifeb. This is not true, however, for thick-walled components such as RPV head penetrations, where an appreciable amount of time may be taken up with crack growth before safety margins related to structural integrity are breached. One of the key factors governing crack growth rate (CGR) is the stress intensity developed at the crack tip and – in some cases – this may drop as the crack grows into a region of decreasing residual stress, thus leading to much slower growth, or even to crack arrest. Further consideration of Alloy 600 CGRs and their importance for determining the residual service life of components known to contain cracks from PWSCC are given in related MRP reports [12, 13].

In summary, PWSCC requires three key factors to be present simultaneously: an environment that promotes intergranular stress corrosion cracking, susceptible material, and significant, prolonged tensile stress. Eliminating any one of these three factors will mitigate cracking in principle, although in practice it is prudent to attack all of these factors at once, wherever feasible.

1.4 References

1. P.M. Scott and M.-C. Meunier, Materials Reliability Program: Review of Stress Corrosion Cracking of Alloys 182 and 82 in PWR Primary Water Service (MRP-220). EPRI, Palo Alto, CA: 2007. 1015427.

2. NRC, “Industry Codes and Standards; Amended Requirements; 10 CFR Part 50”, Federal Register: September 10, 2008 (Volume 73, Number 176).

3. H. Xu et al., Materials Reliability Program (MRP), Resistance to Primary Water Stress Corrosion Cracking of Alloys 690, 52, and 152 in Pressurized Water Reactors (MRP-111). EPRI, Palo Alto, CA: 2004. 1009801.

4. J. Hickling, EPRI Materials Reliability Program: Resistance of Alloys 690, 152 and 52 to Primary Water Stress Corrosion Cracking (MRP-237, Rev 1): Summary of findings from completed and ongoing test programs since 2004. EPRI, Palo Alto, CA: 2008. 1018130.

5. H. Coriou, et al., “High Temperature Stress Corrosion Cracking of Inconel in Water,” Third Metallurgical Symposium on Corrosion (Aqueous and Gaseous), 1959, North Holland Publishing Co., Amsterdam, published in 1960, pp. 161-169.

6. F. Champigny, F. Chapelier, et al., “Maintenance Strategy of Inconel Components in PWR Primary System in France,” paper presented at NRC/ANL Conference on Vessel Head Penetration Inspection, Cracking, and Repair in Gaithersburg, MD from September 29 – October 2, 2003.

7. B. Gronwall, L. Ljungberg, et al., “Intercrystalline Stress Corrosion Cracking of Inconel 600 Inspection Tubes in the Agesta Reactor,” Atomenergi, (Rapp.) AE, AE-245, 1966.

b An exception to this rule is that PWSCC in kiss-rolled transitions in SG tubes appears to grow so slowly, or even to arrest, such that many SG tubes remain in operation ten or more years following detection of PWSCC.

1-5

Introduction

8. C.A. Campbell and S. Fyfitch, “PWSCC Ranking Model for Alloy 600 Components,” Proceedings of 6th International Conference on Environmental Degradation of Materials in Nuclear Power Systems—Water Reactors, TMS, 1993, p. 863.

9. E. Eason, Program on Technology Innovation: A Preliminary Hybrid Model of Nickel Alloy Stress Corrosion Crack Propagation in PWR Primary Water Environments. EPRI, Palo Alto, CA: 2008. 1016546.

10. F. Vaillant et al., “Development of a Predictive Model for SCC Initiation of Alloy 600 in Primary Water”, Workshop on Detection, Avoidance, Mechanisms, Modeling, and Prediction of Stress Corrosion Cracking Initiation in Water-Cooled Nuclear Plants. EPRI, Palo Alto, CA; TEPCO R&D Center, Yokohama, Japan; AREVA NP, Technical Center, Le Creusot, France; Institut de Radioprotection et de Sûreté Nucléaire (IRSN), Fontenay-aux-Roses, France; Institute of Nuclear Safety System, Incorporated (INSS), Fukui, Japan; The Materials Aging Institute (MAI), Moret Sur Loing Cedex, France; and EDF R&D, Moret Sur Loing Cedex, France: 2009. 1018908.

11. I. de Curières and M.-C. Meunier, “Prediction of PWSCC initiation in Steam Generator Alloy 600 Tubes by the Index Model: Comparison between Model and Experience”, Workshop on Detection, Avoidance, Mechanisms, Modeling, and Prediction of Stress Corrosion Cracking Initiation in Water-Cooled Nuclear Plants. EPRI, Palo Alto, CA; TEPCO R&D Center, Yokohama, Japan; AREVA NP, Technical Center, Le Creusot, France; Institut de Radioprotection et de Sûreté Nucléaire (IRSN), Fontenay-aux-Roses, France; Institute of Nuclear Safety System, Incorporated (INSS), Fukui, Japan; The Materials Aging Institute (MAI), Moret Sur Loing Cedex, France; and EDF R&D, Moret Sur Loing Cedex, France: 2009. 1018908.

12. G. White et al., Materials Reliability Program Crack Growth Rates for Evaluating Primary Water Stress Corrosion Cracking (PWSCC) of Thick-Wall Alloy 600 Materials (MPR-55) Revision 1. EPRI, Palo Alto, CA: 2002. 1006695.

13. G. White et al., Materials Reliability Program: Reactor Vessel Closure Head Penetration Safety Assessment for U.S. Pressurized Water Reactor (PWR) Plants (MRP-110): Evaluations supporting the MRP Inspection Plan. EPRI, Palo Alto, CA: 2004. 1009807.

1-6

2 ALLOY 690 PROPERTIES AND METALLURGY

2.1 Material Specifications

Alloy 690 and its predecessor Alloy 600 were first developed by Inco Alloys International under the trade names Inconel Alloy 690 and Inconel Alloy 600 [1, 2]. The present owner of the trade name of Inconel is Special Metals Corporation. Currently, both alloys have been adopted by ASTM (American Society for Testing and Materials), ASME (American Society of Mechanical Engineers), and other international materials societies. Even though these alloys are listed under the UNS (Unified Numbering System) numbers in ASTM or ASME standards (see Table 2-1), these alloys are generically referred to as Alloy 690 or Alloy 600 in the nuclear power industry.

Table 2-1 provides a summary of the most commonly used product forms in the ASME Boiler and Pressure Vessel Code Section II material specifications for the wrought Alloy 690 materials. Because Alloy 690 was developed to replace its predecessor Alloy 600 for light water nuclear power reactors, both are listed in the same ASME material specifications.

Table 2-1 ASME specifications of Alloys 690 and 600 Reprinted from ASME 2001 BPVC, Section II-B and 2001 BPVC, Section II-D, by permission of The American Society of Mechanical Engineers. All rights reserved.

Specification Alloy Product Form

ASME SB-163c UNS N06690 (Alloy 690) UNS N06600 (Alloy 600)

Seamless Tubing

ASME SB-166d UNS N06690 (Alloy 690) UNS N06600 (Alloy 600)

Rod, Bar, Wire

ASME SB-167e UNS N06690 (Alloy 690) UNS N06600 (Alloy 600)

Seamless Pipe and Tube

ASME SB-168f UNS N06690 (Alloy 690) UNS N06600 (Alloy 600)

Plate, Sheet, Strip

c ASME SB-163, “Specification for Seamless Nickel and Nickel Alloy Condenser and Heat-Exchanger Tubes,” ASME Boiler and Pressure Vessel Code, Section II Part B, Nonferrous Material Specifications, 2001.

d ASME SB-166, “Specification for Nickel-Chromium-Iron Alloys (UNS N06600, N06601, N06603, N00690, N06025, and N06045) and Nickel-Chromium-Cobalt-Molybdenum Alloy (UNS N06617) Rod, Bar, and Wire,” ASME Boiler & Pressure Vessel Code, Section II Part B, Nonferrous Material Specifications, 2001.

e ASME SB-167, “Specification for Nickel-Chromium-Iron Alloys (UNS N06600, N06601, N00690, N06025, and N06045) Seamless Pipe and Tube,” ASME Boiler & Pressure Vessel Code, Section II Material Specifications, Part B, 2001.

f ASME SB-168, “Specification for Nickel-Chromium-Iron Alloys (UNS N06600, N06601, N06603, N00690, N06025, and N06045) and Nickel-Chromium-Cobalt-Molybdenum Alloy (UNS N06617) Plate, Sheet, and Strip,” ASME Boiler & Pressure Vessel Code, Section II Material Specifications, Part B, 2001.

2-1

Alloy 690 Properties and Metallurgy

2-2

Table 2-2 lists the chemical composition requirements for Alloys 690 and 600. The ASME chemical composition requirements for these alloys remain identical to the original specification developed by Inco Alloys International. However, stricter requirements on chemical composition, mechanical properties and heat treatment process are imposed on Alloy 600 or 690 by utilities and vendors for applications in PWRs (see below).

Table 2-2 ASME chemical composition requirement (wt%) Reprinted from ASME 2001 BPVC, Section II-B and 2001 BPVC, Section II-D, by permission of The American Society of Mechanical Engineers. All rights reserved.

Alloy Ni Cr Fe C Mn Si S Ti Nb + Ta

Cu P Al Mo Other

690 58.0 min

27.0-31.0

7.0-11.0

0.05 max

0.50 max

0.50 max

0.015 max

– – 0.50 max

– – – –

600 72.0 min

14.0-17.0

6.0-10.0

0.15 max

1.0 max

0.50 max

0.015 max

– – 0.50 max

– – – –

Table 2-3 lists the room temperature properties of Alloys 690 and 600 and Table 2-4 lists the ASME properties of these materials at elevated temperatures. Such properties are used in the design of Alloy 600 and 690 component items for repair or replacement and for calculating the operating stress.

International specifications for Alloy 690 may differ from the above in various respects. It should also be pointed out that some components in the USA may have been fabricated to older specifications that differ slightly from those given here.

Table 2-3 ASME room temperature properties Reprinted from ASME 2001 BPVC, Section II-B and 2001 BPVC, Section II-D, by permission of The American Society of Mechanical Engineers. All rights reserved.

Alloy 690 Alloy 600

lb/in3 0.293g 0.300g Density

kg/cm3 8.11g 8.30g

Poison’s Ratio – 0.29h 0.29h

ksi 85i 80i

Min. Tensile Strength MPa 586i 550i

ksi 30i 30i Min. Yield Strength

MPa 205i 205i

Min. Elongation % 35i 35i

g ASME Boiler & Pressure Vessel Code, Section II Part D, Materials Properties, 2001, Table NF-2. h ASME Boiler & Pressure Vessel Code, Section II Part D, Materials Properties, 2001, Table NF-1. i ASME SB-167, “Specification for Nickel-Chromium-Iron Alloys (UNS N06600, N06601, N00690, N06025, and N06045) Seamless Pipe and Tube,” ASME Boiler & Pressure Vessel Code, Section II Material Specifications, Part B, 2001.


2-3

Table 2-4 ASME elevated temperature properties for Alloy 690 and Alloy 600j Reprinted from ASME 2001 BPVC, Section II-B and 2001 BPVC, Section II-D, by permission of The American Society of Mechanical Engineers. All rights reserved.

Temp Young’s Modulus E(a), x106 ksi

Coefficient of Thermal

Expansion α(b), x10-6 in/in/°F

Design Stress Intensity(c) Sm, ksi

Yield Strength(d)

Sy, ksi

Tensile Strength(d) Su,

ksi

°F 690 600 690 600 690 600 690 600 690 600

70 30.3 31.0 7.7 6.8 23.3 23.3 35.0 35.0 85.0 80.0

100 30.1(a) 30.8 7.8 6.9 23.3 23.3 35.0 35.0 85.0 80.0

200 29.5 30.2 7.9 7.1 23.3 23.3 31.7 32.0 85.0 80.0

300 29.1 29.8 7.9 7.3 23.3 23.3 29.8 31.2 84.0 80.0

400 28.8 29.5 8.0 7.5 23.3 23.3 28.6 30.7 82.0 80.0

500 28.3 29.0 8.1 7.6 23.3 23.3 27.9 30.3 80.8 80.0

600 28.1 28.7 8.2 7.8 23.3 23.3 27.6 29.9 80.2 80.0

650 27.9(a) 28.5 8.2 7.8 23.3 23.3 27.5 29.7 80.0 80.0

700 27.6 28.2 8.3 7.9 23.3 23.3 27.5 29.4 79.8 80.0

(a) From Table TM-4 of j. Values of Young’s modulus for 100°F and 650°F are obtained by linear interpolation.

(b) Mean coefficient of thermal expansion going from 70°F to the indicated temperature, from Table TE-4 of j.

(c) From Table 2B of j. For annealed Alloy 690 and Alloy 600 of SB-163, SB-166, SB-167, and SB-168, some exceptions exist: see Table 2B of j for details.

(d) From Table Y-1 of j. For annealed Alloy 690 and Alloy 600 of SB-163, SB-166, SB-167, and SB-168, some exceptions exist: see Table Y-1 of j for details.

2.1.1 Typical PWR Specifications for Thin-Walled Alloy 690 SG Tubing

For Alloy 690 SG tubing, the EPRI Guidelines require the carbon content to be between 0.015% and 0.025% [3]. The lower limit ensures continuous or semi-continuous intergranular carbide precipitation while the upper limit keeps intragranular carbide precipitation to a minimum upon thermal heat treatment. Alloy 690 has higher yield and tensile strengths than Alloy 600 due to the increased Cr content. Therefore, the lowered carbon level does not affect the allowable stresses for Alloy 690.

The high cycle fatigue properties of Alloy 690 were recently investigated and found to be comparable to those of Alloy 600 [4].

j ASME Boiler & Pressure Vessel Code, Section II Part D, Materials Properties, 2001.


2.1.2 Specification and Manufacture of Thick-Walled Alloy 690 Components

Many replacement RPV head penetrations made of thick-walled Alloy 690 were manufactured prior to the 2008 publication of EPRI Guidelines for the manufacture of such components [5], which suggest limiting the allowable carbon content to between 0.015% and 0.035%. For Alloy 690 bars, plates, and heavy section tubing, however, the carbon content has routinely been specified to be in a range between 0.01% and 0.04% or narrower, instead of the max. 0.05% carbon specified in ASME Section II.

Because of the potential importance of both the exact chemical composition of Alloy 690 and the manufacturing process in ensuring good PWSCC resistance (as discussed in later sections), efforts were recently started to assess the actual characteristics of replacement CRDM nozzles that have been installed in U.S. PWRs within the last few years. To date, the picture remains somewhat incomplete [6], although further progress is expected shortly. The cited report provides generic information on the various practices that a melt shop could have used in the production of Alloy 690 pipe or rod material. Information from the three main suppliers (MHI, AREVA, and BWC), while not complete, identified the sources (melt shop) and some processing parameters of the Alloy 690 material, including the annealing and thermal heat treatments. Perhaps the most significant variable identified in this exercise was that straightening of the pipe or rod after final thermal treatment appears to be a common practice at the Alloy 690 suppliers. To avoid introducing cold work, such a procedure would not be allowed under the new EPRI guidelines [5] for procurement of such Alloy 690 material. Note, also, that the currently practised final thermal treatments (typically between 5 and 15 hours at around 700°C) to produce Alloy 690 TT, with extensive carbide decoration of the grain boundaries, were developed considering Alloy 600 PWSCC experience and were then adjusted a small amount to improve resistance to IGSCC for SG tubing in caustic environments. They have not yet been shown to be either necessary, or even beneficial, for obtaining PWSCC resistance in Alloy 690.

2.2 Phase Diagram of Alloy 690

The first comprehensive study of Alloy 690 was published by Sedriks et al. [7] of Inco Alloys International, Inc. in 1979. This was about the time that Alloy 690 began to be introduced for fabricating steam generator tubing in PWRs. The physical metallurgy and properties of Alloy 690 from this study and other sources are summarized below.

Figure 2-1 in [8] showed the Ni-Cr-Fe phase diagram with the location of the γ/γ+α’ solvus line indicated from 816 to 1260°C (1500 to 2300°F). The lowest temperature at which the γ/γ+α’ solvus has been determined is 816°C (1500°F). The locations of Alloy 690 and 600 compositions are both well within the austenite field (γ). The range of melting temperatures is 1343-1377°C (2450-2510°F) [2] for Alloy 690 and 1354-1413°C (2470-2575°F) [1] for Alloy 600. The α’ phase is a chromium-rich phase, which is very similar to the iron-chromium σ phase both in morphology and hardness. The α’ phase can cause embrittlement in alloys whose compositions lie in the two-phase γ+α’ field when precipitated during prolonged high temperature exposure. Precipitation of α’ phase is a mechanism of aging embrittlement of austenitic stainless steel welds or castings containing small amounts of δ ferrite and of martensitic, precipitation-hardenable stainless steel

2-4


when exposed to elevated operating temperatures in PWRs [9]. Since Alloys 600 and 690 are both stable austenitic solid-solution alloys from room temperature to the melting temperature, precipitation of α’ phase is not expected and has never been found in Alloys 600 and 690.

K. Smith et al. evaluated the possible occurrence of an “ordered” Ni2Cr phase in Alloy 690 [10]. This concern was raised due to the fact that the Ni/Cr (atomic) ratio for an alloy with 70% Ni-30%Cr by weight is about 2.1 to 1, which is very close to that of Ni2Cr [11]. The brittle and hard intermetallic Ni2Cr phase is stable below 580°C (1076°F) and could have a similar embrittling effect as the iron-chromium σ intermetallic phase in austenitic alloys. However, K. Smith’s research found that iron has an inhibiting effect on the formation of Ni2Cr phase and that the minimum specified iron content in Alloy 690 was high enough to inhibit any Ni2Cr formation. Later, Larsson et al. performed extensive testing for the same phenomenon including hardness tests, tensile tests, transmission electron microscopy (TEM), differential scanning calorimetry (DSC), and differential thermal analysis (DTA) on commercially produced Alloy 690TT that had been thermally aged for 3000 hours at 400 and 500°C (752 and 932°F), as well as on an Alloy 690 corrosion specimen exposed at 365°C (689°F) for 32,961 hours [12]. The study confirmed that no evidence of long range ordering of type Ni2Cr was found in the thermally aged Alloy 690. Nevertheless, reactor vendors usually specify higher minimum iron contents than the ASME Code allowed minimum in order to provide a significant margin against the formation of an ordered Ni2Cr phase in Alloy 690 and both the EPRI Guidelines for steam generator tubing [3] and the recent guidelines for pressure vessel nozzles [5] require a minimum content of 9 wt%. Some work is still in progress on defining the minimum iron content required more closely in the context of very long plant operating times [13, 14].

The microstructure of Alloy 690 is similar to Alloy 600, i.e., an austenitic matrix with the secondary phases being predominantly chromium carbides precipitated both intergranularly and intragranularly. The other minor secondary phases in Alloys 600 and 690 are titanium nitrides, titanium carbides, and carbonitrides. Because the austenitic matrix phase is stable up to the melting temperature, neither alloy is heat-treatable via phase changes (such as the austenite to martensite transformation in carbon and low-alloy steels) and cannot be hardened through secondary phase precipitation (such as through γ’ precipitation in Alloy X-750 or Alloy 718).

2.3 Carbon Solubility and Dynamic Strain Aging

The extent of intergranular and intragranular carbide precipitation depends on the thermal mechanical history and carbon content. Most research efforts on Alloys 690 and 600 have been focused on the grain boundary microstructure, especially due to observations of generally higher PWSCC resistance of Alloy 600 with a microstructure containing continuous intergranular carbides and few intragranular carbides [15]. The intergranular carbide precipitates are found to be both M7C3 and M23C6 types in Alloy 600 and mostly globular M23C6 type in Alloy 690 [16, 17].

It is generally recognized that the solubility of carbon in Alloy 690 is lower than in Alloy 600. This is due to the higher Cr content which lowers the solubility while increasing the propensity for carbide precipitation. The carbon solubility curve of Alloy 690 was investigated by Sarver et al. [18]. The Alloy 690 heats used are listed in Table 2-5 and the results are shown in Figure 2-1 with the Alloy 690 solubility line drawn to separate the specimens having no visible carbides from the specimens having visible carbides. Figure 2-1 also shows the comparison of carbon solubility curves for Alloy 690 and Alloy 600. The equations fitted to these curves are:

2-5


Alloy 690

°F = 2647.5 + 120 ln(%C) Equation 2-1

°C = 1453.1 + 66.9 ln(%C)

Alloy 600

°F = 2640 + 234.5 ln(%C) Equation 2-2

°C = 1449 + 130.3 ln(%C)

Table 2-5 Chemical composition (wt%) of Alloy 690 heats used by Sarver et al. [18]

Heat Ni Cr Fe C Mn Si S Ti Nb + Ta

Cu P Al Mo

Intergranular Carbide Precipitation Study

NX9217H Bal 27.27 10.22 0.01 0.15 0.10 0.003 0.25 0.15 0.17

NX9780H Bal 29.20 8.85 0.01 0.33 0.43 0.001 0.46 0.12 0.26

NX4459HG Bal 28.25 8.86 0.06 0.20 0.10 0.004 0.32 0.04 0.17

NX4401H Bal 29.25 10.22 0.01 0.21 0.25 0.001 0.28 0.26 0.15

Carbon Solubility Study

EXP 1 Bal 28.7 9.2 <0.001 0.02 0.001 0.001 0.03

EXP 2 Bal 28.8 9.8 0.01 0.06 0.06 0.003 0.02

NX4458H Bal 27.9 9.8 0.016 0.19 0.10 0.002 0.26

EXP 3 Bal 29.9 9.6 0.02 0.03 0.05 0.003 0.01

EXP 4 Bal 28.7 9.3 0.02 0.02 0.001 0.001 0.03

NX05E1H Bal 29.9 9.5 0.021 0.21 0.39 0.001 0.28

NX10C1H Bal 29.8 9.4 0.039 0.15 0.15 0.008 0.30

EXP 5 Bal 29.0 9.1 0.04 0.02 0.001 0.002 0.02

EXP 6 Bal 29.1 9.1 0.058 0.02 0.001 0.001 0.02

EXP 7 Bal 29.4 10.3 0.06 0.03 0.06 0.003 0.03

EXP 8 Bal 29.5 9.8 0.06 0.01 0.05 0.003 0.02

Heat Treatment Study

NX4401H Bal 29.25 10.22 0.01 0.21 0.25 0.001 0.28 0.26 0.15

NX4588H Bal 29.92 9.49 0.03 0.18 0.21 0.001 0.27 0.24 0.21

NX2184H Bal 28.82 8.98 0.01 0.18 0.24 0.001 0.26 0.24 0.30

NX4308 (Alloy 600)

Bal 15.11 7.60 0.03 0.35 0.21 0.007 0.26 0.29 0.50

2-6


It should be noted that thermo-mechanical processing of both of these materials may cause grain boundaries to move by recrystallization after all the available carbon has precipitated as carbides on the old grain boundaries, thus creating a “ghost” grain boundary carbide network. Coring, or non-equilibrium solidification of the original ingot, can also manifest itself in the form of carbide banding. Solution heat treatments at temperatures above the carbon solubility curves for prolonged periods followed by rapid quenches can minimize the observed effects, but cannot eliminate them. Prohibitively long solution heat treatment would be required to dissolve the carbides and allow the carbon to diffuse away from the original sites in order to prevent the carbon from re-precipitating at the old sites during carbide precipitation heat treatments.

1200

1400

1600

1800

2000

Tem

pera

ture

o F

2200

2400

0.001 0.010 0.100

Alloy 690

Alloy 600

No Carbides, Alloy 690

Carbides, Alloy 690

Carbon, wt%

f

Figure 2-1 Carbon solubility diagram for Alloy 690 and Alloy 600 from [18]

The phenomenon of dynamic strain aging (DSA) has been extensively studied in recent years (often using measurements of so-called “internal friction”) because of a suspected link between the resulting localization of deformation within a material and its SCC susceptibility. Because oits lower carbon solubility than Alloy 600, Alloy 690 might have been expected to show more susceptibility to DSA, but investigations by Hänninen et al. show that this is not the case [19]. The authors consider that the higher Cr content of Alloy 690 substantially increases its creep resistance, probably by reducing the stacking fault energy of the alloy, and that these positive effects of Cr are more pronounced than the possible negative effects of reduced carbon solubility. Additionally, higher Cr content will affect the oxidation reactions controlling vacancy injection and hydrogen uptake in the material (see Section 3.1.5), which also play key roles in strain localization and subsequent crack initiation.

2-7


2.4 Intergranular Carbide Precipitation and Sensitization

The time-temperature chromium carbide precipitation curves for Alloy 600 and Alloy 690k were determined by Yonezawa et al. [20]. The highest precipitation rate (the nose of the C-curves) isat approximately 850-950°C (1550-1750°F). The Alloy 690 and Alloy 600 heats were annealeat 1100°C (2012°F) and contained 0.03% carbon, but the exact chemical composition was not listed. From this work, chromium carbides precipitate first at the grain boundaries followed by precipitation inside the grains. In addition, carbide precipitation is faster in Alloy 690 than in Alloy 600 due to the higher chromium content of Alloy 690. For Alloy 600, the precipitation othe grain boundary carbides is complete after a short period of time at 704°C (1300°F). Howeveran extended 10-15 hours period is required to fully replenish the Cr-depleted grain boundary with Cr from the bulk of the grains to avoid a “sensitized” microstructure [21, 22]. It was reasoned that since Alloy 690 has a much higher Cr content, full replenishment of the grain boundaries is not necessary and a 4 to 5 hour thermal treatment should be adequate [23]. Recent thinking on this by one major reactor component vendor, as re

d

f ,

ported in 2005, also makes a distinction between the most suitable mill-annealing temperatures (as a function of carbon content) for either “cold”

.

istry Water Reactors).

or Alloy 600 steam generator tubing, a modified Huey Test (25% nitric acid with two 24-hour boiling periods) is often used, instead of the 65% nitric acid with five 48-hour boiling periods specified in ASTM A 262, Practice C. Due to the higher resistance of Alloy 690 to sensitization, Sarver et al. used 65% nitric acid for sensitization testing of several heats of Alloy 690 with varying carbon contents and heat treatments [18]. The time-temperature-sensitization diagrams plotted by Sarver et al. are shown in Figure 2-2 to Figure 2-4. Figure 2-2 shows very low corrosion rates for heat NX4459HG (0.06%C) with a 1038°C (1900°F) anneal before the sensitization treatment. The corrosion rate was highest (1.7 mil/month) after a sensitization treatment of 538°C (1000°F) for 100 hours. Figure 2-3 shows the corrosion rates for heat NX9217H (0.01%C), which was mill annealed at 1038°C (1900°F) for one hour. The corrosion rate was also low except when sensitized at 538°C (1000°F) for 20-100 hours. Figure 2-4 shows the results for heat NX9780H, also containing 0.01%C, but having a slightly higher Cr content than NX9217H (29.20% vs. 27.27%). Heat

or “hot” finished Alloy 690 TT [24].

Sedriks et al. [7] found that if the carbon content is kept to 0.02% or below, Alloy 690 cannot be “sensitized” as defined by the Huey test, i.e., the boiling nitric acid test per ASTM A 262 Practice C [25]. This means the Cr level near the grain boundary was sufficiently high in Alloy 690 containing ≤ 0.02% carbon to prevent intergranular attack by boiling nitric acid even after intergranular carbide precipitation. In the Huey test, specimens are placed in boiling 65% (by weight) nitric acid, typically for five 48-hour periods, with a fresh nitric acid solution being used for each period. It should be noted that the Huey test for sensitization is normally used for detecting susceptibility to intergranular attack in austenitic stainless steels, not for high-nickel alloys such as Alloy 690 or Alloy 600. “Sensitization”, as usually used, refers to grain boundary chromium depletion in austenitic stainless steels such as Type 304 and Type 316 exposed to a temperature range of 427-816°C (800-1500°F) for a prolonged period of time that varies with carbon contentPrecipitation of chromium carbides along the grain boundaries causes Cr depletion near the grainboundaries, which are preferentially attacked under certain oxidizing environmental conditions (for example in nitric acid, or in – in conjunction with stress – in the oxygenated water chemsometimes used in Boiling

F

k See Figure 2-3 in [8].

2-8


NX9780H also showed a very low corrosion rate, even when sensitized at 538°C (1000°F) for 10-20 hours. The highest degree of sensitization for Alloy 690, as defined by the Huey test, was produced by heat treatment at 538°C (1000°F) for 20 to 100 hours. The complete chemical compositions of heats NX4459HG, NX9217H, and NX9780H studied by Sarver et al. are summarized in Table 2-5.

Corrosion rate in 0.001inch/month

800

1000

S

0.5 0.7 0.6 1.70.6 0.9 0.4

1200

1400

1600

1800

2000

0.01 0.1 1 10 100 1000Sensitiz ion Time, hour

ensi

tizat

ion

Tem

p, o F

0.7

0.6

0.6

0.5

0.9

0.5

0.8

0.6

0.7

0.7

0.5

0.6 0.3

Figure 2-2 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690 Heat NX4459HG (0.06%C) from [18]

at

2-9


Corrosion rate in 0.001inch/month

800

1000

1200

1400

1600

1800

2000

0.01 0.1 1 10 100 1000Sensitization Time, hour

Sens

itiza

tion

Tem

p, o F

0.4

0.5

0.3

0.4

0.5

0.4

1.1

0.3

0.4

0.5

0.6

0.7

0.4

0.6

0.7

78.39.1 28.5

1.2

1.90.5 0.4

0.4

0.3

0.7

0.3

0.4

0.7

0.5

Figure 2-3 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690 Heat NX9217H (0.01%C) from [18]

Corrosion rate in 0.001inch/month2000

1400

1600

on T

emp,

800

1000

1200

0.01 0.1 1 10 100 1000

Sens

itiz

0.4

0.5

0.5

0.6

1800

Sensitization Time, hour

ati

o F

Figure 2-4 Time-Temperature-Sensitization diagram by modified Huey Test, Alloy 690 Heat NX9780H (0.01%C) from [18]

2-10


It should be noted that, for Alloy 600, a sensitized microstructure with continuous intergranular carbides (and associated grain boundary Cr depletion) is more resistant to PWSCC than the samheat without intergranular carbides, despite the grain boundary Cr depletion

e

Sarver et al. also investigated the effect of carbon content, annealing temperature, and one-hour prec eat treatments on the carbide morphology. The results are shown in Table 2-6. Both the at which heavy intergranular carbide precipitation occurred. Depending upon ed annealing tempcontent, a one-hour heat treatment produce heavy intergprecipitation. The high carb ore lcarbides than the lower ca .0 preci on oc ed equally in both materials h n ar p io ro ymany heat treatments, no significant intergranular sensitiz n occu and c osio n the H s were low (less than 0.4 mm/year) except in one case: the 0.03% carbon material anne 14 0° d hea eated and 7 °C (1200 and 1300°F) showed ratesof 1. m es tively, icating at interg ar ation had o

Table 2-6 e ent on Alloy 690 carbide precipitation [18]

0.03%C NX4588H

l.

ipitation hannealing temperature and carbon level affected the temperature

the desircould be chosen to

erature and carbon ranular carbide

on (0.03%C) material was much m ikely to show intragranular rbon (0

. Althoug1%C) ma heavy i

terial, although intergranular tergranul

pitatin was p

currduced b carbide recipitat

atio rred orr n rates iuey test

aled at 10 and 1.5

9°C (210 F) an t tr at 649 04 m/year, r pec ind th ranul sensitiz ccurred.

Eff ct of heat treatm

0.01%C NX4401H

Initial Anneal Temp, 20 Minutes Initial Anneal Temp, 20 Minutes

1800°F 1900°F 2000°F 2100°F 1800°F 1900°F 2000°F 2100°F

1200°F VL VL L VL VL H H Intra L

1300°F VL L H Intra L L H H Intra L

1400°F H H H H L Intra L H Intra H Intra Car

bid

e P

reci

pit

atio

n

Tem

p, 1

ho

ur

1500°F VL H H H L Intra L Intra H Intra H

(a) Heat treatment code:

VL – very light intergranular carbides; L – light intergranular carbides.

H – heavy intergranular carbides; Intra – intragranular carbides.

l For Alloy 600, the precipitation of intergranular carbides usually implies Cr-depletion and a sensitized microstructure, because of the limited Cr content of the alloy. This makes the alloy susceptible to intergranular attack in standard tests (such as the Huey test) and should also be avoided as the material condition in an oxidising HT-environment, such as a BWR on NWC. It is not relevant under low-potential conditions in PWR primary water, however, where the benefit of the carbides (e.g., in reducing grain boundary creep) dominates. This indicates that the mechanism of PWSCC is somewhat different from the slip dissolution/oxidation mechanism generally accepted to apply to SCC in oxidising BWR water.

2-11


2.5 Effect of Elevated Temperature Exposure

Figure 2-5 shows the results of short-term tensile tests performed on annealed Alloy 690 at temperatures ranging from room temperature up to 982°C (1800°F) [2]. The curves represent average values for both cold- and hot-worked products in the annealed condition and show that annealed Alloy 690 retains more than 90% of its room-temperature tensile properties (yield and tensile strengths, as well as elongation) up to 427°C (800°F). Only at temperatures over 540°C (1000°F) does the tensile strength start to decline substantially. This retention of room temperature tensile properties, together with the long-term, high-temperature stability discussed below, are reflected in the ASME design stress intensity values for Alloy 690 and Alloy 600 (see Table 2-4).

Figure 2-5 High-temperature tensile properties of annealed Alloy 690. Data shown are a composite of cold-and hot-worked products in the annealed condition and taken from [2]

m-temperature tensile results for annealed Alloy 690 after exposure

A(1050 to 1400°F). This range of temperatures is similar to that of post-weld, stress-relieving

This work showed that the room-temperature Charpy impact energy (189 J un-aged) was virtually unchanged after 12,000 hours or 500 days at elevated temperatures. This property is known to be a very sensitive indicator for the precipitation of α’ or σ phases in alloys with

Reference [2] also lists rooto elevated temperatures for various periods of time. The long-term, high-temperature stability of lloy 690 was demonstrated by Charpy impact testing after long periods at 566 to 760°C

heat treatment (PWHT) temperatures for carbon and low alloy steel vessels or piping in PWRs.

2-12


an austenitic matrix. For example, aging at 400°C (752°F) for 10,000 hours (417 days) reduced Charpy impact energy from ~220 J to 50 J for a CF-8 cast austenitic stainless steel equivalent of wrought Type 304 s

the (the tainless steel) containing 24% δ ferrite [26].

Pos[27pro

Hence, it can be concluded that thermal aging embrittlement is not a concern for Alloy 690

(inc

2.6

3. ubing Specification and Repair,” Vol. 2, Rev. 1: April 1999.

4. peratures up to 330oC in Very High Cycle Fatigue Regime”, Proceedings

ear n

5. ility Program: Guidelines for Thermally Treated

6. bility Program: Material Production and Component

7. al,” Corrosion Engineering (Boshoku Gijutsu), vol. 28, pp. 82-95, 1979,

8. s Reliability Program (MRP), Resistance to Primary Water Stress 11).

sible long-term aging of Alloy 690 is being investigated at EDF in France and a new paper ] shows that it is possible after 60000 h at 420°C for a material with 7.2% Fe, but that industrial ducts are expected to be free from aging.

and Alloy 600 components from exposure to high temperatures either during fabrication luding repair welding or PWHT heat treatment), or from long-term exposure at PWR

operating temperatures of up to 343°C (650°F).

References

1. INCONEL Alloy 600, Publication Number SMC-027, Special Metals Corporation, 2002 (Sept. 02).

2. INCONEL Alloy 690, Publication Number SMC-079, Special Metals Corporation, 2002(Sept. 02).

Guidelines for PWR Steam Generator TGuidelines for Procurement of Alloy 690 Steam Generator Tubing. Final Report, EPRI Report TR-016743-V2R1.

G. Chai, J. Frodigh and H. Törnblom, “Fatigue Behavior of Alloy 690 and Alloy 800 SGTubing at Temof the 13th International Conference on Environmental Degradation of Materials in NuclPower Systems, Whistler, British Columbia, August 19–23, 2007, Published by the CanadiaNuclear Society.

A. Mcllree and G. Ilevbare, Materials ReliabAlloy 690 Pressure Vessel Nozzles (MRP-241). EPRI, Palo Alto, CA: 2008. 1015007.

G. Theus et al., Materials ReliaFabrication and Installation Practices for Alloy 690 Replacement Components in Pressurized Water Reactor Plants (MRP-245). EPRI, Palo Alto, CA: 2008. 1016608.

A.J. Sedricks, J.W. Schultz, and M.A. Cordovi, “Inconel Alloy 690 – A New CorrosionResistant MateriJapan Society of Corrosion Engineering.

H. Xu et al., MaterialCorrosion Cracking of Alloys 690, 52, and 152 in Pressurized Water Reactors (MRP-1EPRI, Palo Alto, CA: 2004. 1009801.

9. H. Xu and S. Fyfitch, “Aging Embrittlement Modeling of Type 17-4PH at LWR Temperatures,” the 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE International, Houston, Texas (2001).

2-13


10. K. Smith, A. Klein, P. Saint-Paul, J. Blanchet, “Inconel 690, A Material with Improved Corrosion Resistance for PWR Steam Generator Tubes,” Proceedings of 2nd InternatioSymposium on Environmental Degradation of Materials in Nuclear Power Systems – WReactors, Monterey, C

nal ater

A, 1985 pp. 319-328.

12. d J. Frodigh, “On the Possibility of Forming Ordered Ni2Cr in

13. Matrix, Revision 1. EPRI, Palo Alto,

14. 1018400.

ey, CA, 1985

16. structure, Chemistry, and

A, 1991.

racking rgical and Materials Transactions, Vol.

18. ermal

s – Water

19. ing Of Ni-Base Alloys Inconel 600 and 690”, ls in

20.nd International Symposium on Environmental Degradation of Materials in

21.

2. J.R. Crum, “Effect of Composition and Heat Treatment on Stress Corrosion Cracking of Alloy 600 Steam Generator Tubes in Sodium Hydroxide,” Corrosion, vol. 36(1), 1982, pp. 40-45.

23. Proceedings: Workshop on Thermally Treated Alloy 690 Tubes for Nuclear Steam Generators. July 1986. EPRI report NP-4665M-SR.

11. A. Marucco, “Atomic ordering in the Ni-Cr-Fe system”, Materials Science & Engineering, Vol A189 (1994) pp 267-276.

T. Larsson, J. O. Nilsson, anAlloy 690,” Proceedings of 9th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Newport Beach, CA, 1999, pp. 143 to 147.

W. Lunceford et al., EPRI Materials DegradationCA: 2008. 1016486.

W. Lunceford et al., Materials Reliability Program: Pressurized Water Reactor Issue Management Tables – Revision 1 (MRP-205). EPRI, Palo Alto, CA: 2008.

15. S.M. Bruemmer and C.H. Henager, “Microstructure, Microchemistry, and Microdeformation of Alloy 600 Tubing,” Proceedings of 2nd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Monterpp. 293-300.

K. Norring, K. Stiller, and J. Nilsson, “Grain Boundary MicroIGSCC in Alloy 600 and Alloy 690,” Proceedings of 5th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Monterey, C

17. K. Stiller, J. Nilsson, and K. Norring, “Structure, Chemistry and Stress Corrosion Cof Grain Boundaries in Alloys 600 and 690,” Metallu27A-No.2, February, 1996.

J.M. Sarver, J.R. Crum, and W.L. Mankins, “Carbide Precipitation and the Effect of ThTreatments on the SCC Behavior of Inconel Alloy 690,” Proceedings of 3rd International Symposium on Environmental Degradation of Materials in Nuclear Power SystemReactors, Traverse City, MI, 1987, pp. 581-586.

H. Hänninen et al, “Dynamic Strain AgProceedings of 12th International Symposium on Environmental Degradation of MateriaNuclear Power Systems – Water Reactors, TMS, 2005, pp. 1423 to 1430.

T. Yonezawa et al., “Effect of Heat Treatment on Corrosion Resistance of Alloy 690,” Proceedings of 2Nuclear Power Systems – Water Reactors, Monterey, CA, 1985 pp. 593-600.

G.P. Airey, “Effect of Processing Variables on the Caustic Stress Corrosion Resistance of Inconel Alloy 600,” Corrosion, vol. 36(1), 1980, pp. 9-17.

2

2-14


24. T. Yonezawa et al., Materials Reliability Program: Proceedings of the 2005 International SCC of Alloy 600 Conference and Exhibit Show (MRP-154). EPRI, Palo Alto, CA,

VA, Lynchburg, VA, Westinghouse, Madison, PA, Structural Integrity Associates, Centennial, CO, and Welding Services, Inc., Norcross, GA: 2005. 1012089.

stenitic Stainless Steels”.

26. O.K. Chopra and H.M. Chung, “Aging Degradation of Cast Stainless Steels: Effects on Mechanical Properties,” Proceedings of the 3rd Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors,” pp. 737-748, September 1987, Traverse City, Michigan.

ontaining 30% Cr”, 14th Int Symp on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors,

PWAREInc.,

25. ASTM Standard A 262, “Standard Practices for Detecting Susceptibility to Intergranular Attack in Au

27. F. Delabrouille et al., “Long range ordering in Ni Alloys c

Virginia Beach (VA), August 2009 (to be published by ANS).

2-15

3CO ALLOY 690 APART FROM PWSCC

RROSION BEHAVIOR OF

3.1 General Corrosion Tests in Primary Water

ents on

ng primary water is nevertheless of particular interest for PWR steam generator tubing materials. Material lost to the water can become radioactive by coming in contact

e the metal release rate has a strong implication for a plant’s “ALARA” practice – to keep plant shutdown

s

2 4 inhibited with quinoline g

: the difference between the weight before descaling and the weight after descaling.

This section gives some examples of the available laboratory test data pertinent to general corrosion of Alloy 690 in PWR primary water. General corrosion can be defined as uniform deterioration of a metal surface by chemical or electrochemical reaction with the environment. Nickel-base Alloy 600and Alloy 690 are essentially immune to damage through general corrosion in PWR environmdue to the formation of an adherent Cr-rich oxide on the surface. The very low general corrosirate of Alloy 690 in flowi

with th reactor core and then redeposit on the surfaces of the primary loop. Hence,

dose rates to a level as low as reasonably achievable.

3.1.1 SG Tubing by Sedricks et al. 1979

Sedriks et al. [1] evaluated Alloy 690 general corrosion in two simulated PWR water environments using a high velocity test loopm. Borated water was added to simulate the primary side of the steam generator tubing and ammoniated water was used to adjust the pH value to simulate a secondary water environment. The Alloy 690 test specimens and Alloy 600 and Alloy 800 control specimenwere heat treated for ½ hour at 980°C (1800°F), followed by air cooling, and the surfaces were prepared by grinding on a wet 120 grit silicon carbide belt to a 0.75 µm finish. The specimens were weighed and then exposed to the borated water for 2,250 hours and to the ammoniated water for 1,000 hours at a flow velocity of 5.5 m/sec (18 ft/sec). After exposure, the specimens from the ammoniated water tests were descaled by cathodic charging in 5% H SOethiodide. The specimens from the borated water tests were descaled by the alkaline permanganateacid method. The terminology and reporting method followed NACE standard TM-02-74, coverinthe following terms:

1. “Descaled metal loss” (metal consumed): the difference between initial weight and weight after removal of adherent corrosion film.

2. “Corrosion film weight” (adherent corrosion film)

m See MRP-111 for further details.

3-1

Corrosion Behavior of Alloy 690 Apart from PWSCC

3-2

The difference between the descaled metal loss and the corrosion film weight represents the base metal lost to the flowing water. Alloys showing the least loss of material to the water would be expected to produce the lowest activity levels in the PWR primary coolant system. For the three alloys tested in the borated water, Alloy 600, Alloy 690, and Alloy 800, the general corrosion rate decreased with increasing Cr content. Alloy 690, having the highest Cr content of the three, lost the least amount of material to the high velocity simulated PWR water. The observed corrosion films formed on Alloy 690 after the borated and ammoniated exposures were of the thin tarnish type and appeared to be extremely adherent to the base metal. The study reported a standard experimental error of about 25% and suggested that the data should be used as a guide rather than a precise corrosion rate measurement in a high temperature PWR water environment.

3.1.2 SG Tubing by K. Smith et al. 1985

K. Smith et al. [2] investigated the corrosion rate and metal release rate of Alloy 690 and Alloy 600 SG tubing in flowing simulated PWR primary watern. The Alloy 690 specimens were fabricated from full length SG tubes (also called “pre-series” tubes) that were produced by an established industrial route and were intentionally chosen to span the limits of the specification requirements for C and Cr content. The Alloy 690 specimens for the general corrosion tests were in the mill annealed condition (MA) as no difference between the MA and the thermally treated (TT) condition was expected.

Weight changes before and after chemical de-filming were used to calculate the amount of total corrosion from the SG tube I.D. and the amount released to the water. The test results showed that Alloy 690’s general corrosion rate and metal release rate were reduced by a factor of 2 to 4 compared with Alloy 600 material.

3.1.3 SG Tubing by Yonezawa et al. 1985

Yonezawa et al. [3] investigated the weight loss of Alloy 690, Alloy 600, and Alloy 800L SG tubing in simulated PWR primary water chemistryn. The test coupons were made from split SG tubes and the test was conducted at 360°C (680°F) for up to 4,000 hours. Even though the data were scattered and the test duration was limited, the results showed that the general corrosion behavior of Alloy 690 and Alloy 600 in primary water is about the same.

3.1.4 Esposito et al. 1991

Esposito et al. [4] investigated the corrosion rate and metal release rate of Alloy 690 and Alloy 600 in a simulated PWR primary water environment with and without Zn addition (Zn was used in other tests to assess its effectiveness in mitigating Alloy 600 PWSCC initiation) at 330°C (626°F)n. Test coupons were fabricated from Type 304, Type 316, Alloy 600MA, Alloy 600TT, Alloy 690TT, Alloy 750, and Stellite, suspended in the autoclave and exposed for times up to 2,500 hours with or without adding zinc borate. Unlike in the general corrosion rate studies for SG tubing, the test coupons in this study were not subjected to any significant rate of coolant flow. At the conclusion of the test, some samples were descaled to a constant weight by cathodically stripping the oxide film. n See MRP-111 for further details.


The metal corrosion and corrosion release were calculated from the coupon weight change measurements.

Total metal corroded = W0 – Wd

Oxide film weight = Wa – Wd

Metal release = (W0 – Wd) – 70%*(Wa – Wd)

Where: W0 = original sample weight Wa = weight after coupon exposure Wd = weight after descaling 70% is the assumption that metals in the oxide corrosion film represent 70% of the total film weight.

The above values were then divided by the specimen area to obtain the corrosion release per unit area. The test results clearly show that the Alloy 690TT has the lowest general corrosion rate among all the materials tested in PWR primary water, both with and without zinc addition. The results also confirm that the general corrosion rate of Fe-Cr-Ni alloys in primary water environment decreases with increasing Cr content. In addition, the test results also indicated that zinc addition could significantly reduce the general corrosion rate of most commonly used materials in t e PWR primary coolant system.

3.1.5

F. Scenini et al. [5] have tested the intergranular oxidation model of PWSCC by exposing several Ni- to water and steam environments containing hydrogen, at temperatures up to 480°C, and without subsequent electropolishing, short-circuit outward diffusion of Cr occurred, and the differences between alloys or alloy conditions (600, 690, 600TT) were minimal. Only on deformation-free surfaces were the differences in oxidation response of the materials displayed clearly. For hydrogen partial pressures greater than that corresponding to the Ni/NiO equilibrium, Alloy 600 showed intergranular and internal oxidation with the growth of nodules of pure Ni within the grains. The oxidation morphology was strongly dependent on grain orientation. Nodule-free zones appeared near grain boundaries, which were also enriched in metallic Ni. Thermal treatment to precipitate grain boundary carbides eliminated or strongly modified the intergranular oxidation and the presence of nodules. No signs of internal oxidation were observed on Alloy 690, which underwent some thermal etching but formed a nearly continuous external oxide.

Extensive studies of the oxidation behavior of Alloys 600 and 690 in PWR primary water carried out in France have also identified important differences that support the expectation of better PWSCC resistance of Alloy 690 [6]. This includes ingenious work on the way in which applied stress acts to increase oxide thickness and decrease the chromium content of the oxide scale [7], as well as attempts to relate the oxides developed within SCC cracks to that found at the free surface of Ni-base alloys exposed to primary water [8].

h

Alloy Oxidation Studies related to the Mechanism of PWSCC

based alloys analyzing the resultant oxidized surfaces. If the surfaces were mechanically prepared

3-3


3.2 Corrosion Fatigue Tests in Primary Water

Psaila-Dombrowski et al. [9] performed S-N fatigue tests for Alloy 690 plate and Alloys 82, 182, and 152 weld metals in air and in simulated primary water at 315°C (600°F)o. A number of weld build-ups and a 1 m long composite plate were manufactured in a manner consistent with SG fabrication practices in which Alloy 52 is used to attach the Alloy 690 plate to the Alloy 82 build-up. The mock-up composite plate began as a 60 mm (2.36”) thick A508 carbon steel plate. Oneedge of the A508 plate was overlaid with a 5 mm thickness of Alloy 82 and then stress relieved.The test block was completed by making a modified double V-groove weld between the 5.5 mmthick Alloy 690 plate and the Alloy 82 overlay using Alloy 152. The plate was not stress relievfollowing the welding.

ed

odified tensile specimens was used to conduct the low cycle fatigue tests oriented such that the specimen axis was always perpendicular to the direction of welding. In all cases, the gauge sections of the specimens contained only the test material of interest. The tests were conducted in a 76 liter autoclave using a load frame with high lateral stiffness to maintain alignment of the specimen during testing. Most fatigue tests were completed using a strain range of 0.008 and a zero mean strain. However, several tests were performed at different strain levels to expand the data base. With one exception, all tests in the simulated primary water were performed at a strain rate of 0.001 sec-1. Higher strain rates were used in the air tests because no effect of frequency was expected in an air environment.

The fatigue test results showed that the Alloy 690 fatigue life in the primary water (one data point only) decreased by about 40% compared to the fatigue life in air. However, it was still well above the ASME design curve. The fatigue properties of Alloy 690 in the primary water and in air were in good agreement with predictions made using the following model for Alloy 600 at 150 to 350°C (302 to 662°F) [10].

In air, ln (N25) = 6.94 – 1.776 ln (εa – 0.12) + 0.498 Equation 3-1

In water, ln (N25) = 6.94 – 1.776 ln (εa – 0.12) + 0.498 – 0.401 Equation 3-2

Where, N25 = the fatigue life defined as the number of cycles for the peak tensile s o drop 25% from its initial value

Thi he primary water, all four materials exhibited roughly similar fatigue

ves, which were all well above the ASME design curve. Thus, it appeared from this early study that primary water has a measurable influence on Alloy 690 fatigue life at the strain range and rates used, but more data under other test conditions were clearly required to provide a better characterization.

Unf spe as not yet bee . In a 2006 paper [11], Higuchi et al. present a number of diagrams that include

for both Alloy 600 and 690 (as well as their weld alloys) in air (Figure 3-1) and simulated PWR rimary water (Figure 3-2).

M

tress t

εa = the applied strain amplitude in %

s work also showed that the Alloy 82 weld metal exhibited the highest fatigue life among tfour materials tested in air. Inli

ortunately, additional S-N studies of Alloy 690 corrosion fatigue using standard tensilecimens appear not to have been carried out, with the exception of work in Japan that hn fully published

datap

o See MRP-111 for further details.

3-4


Figure 3-1 Japanese fatigue data for Ni-base alloys in air at room temperature from [11]

Figure 3-2 Japanese fatigue data for Ni-base alloys in simulated PWR water at 325°C from [11]

3-5


The 52 wel eld metals. Considering that Ni-base alloys basically have a low sensitivity to environmental fatigue, it was decided in Japan

between nt (Fen)

effects of PWR primary water on fatigue life were slightly less for Alloy 690 (and Alloy 1d metal) than for Alloys 600/132 and essentially the same for base and w

to obtain the trend line for dependence upon strain rate (Figure 3-3) without any distinction the 600 and 690 alloys. Modeling to calculate the factor of environmental fatigue enhancemewas based upon the approach previously used for austenitic stainless steels (which show much larger effects of high-temperature PWR or BWR environments than Alloys 690 and 600 – see Figure 3-4). The comparison between predicted and experimentally determined environmental effects shown in Figure 3-5 for Alloy 690 exposed to PWR primary water appears to be very satisfactory over cyclic fatigue lives between ~ 103 and 105 cycles.

Figure 3-3 Japanese data from [11] on relationship between calculated factor of environmental fatigue enhancement (Fen) and strain rate for Ni alloys in simulated PWR water at 325°C

3-6


Figure 3-4 Japanese data from [11] on relationship between calculated factor of environmental fatigue enhancement (Fen) and temperature for various materials in simulated LWR water at 325°C

3-7


Figure 3-5 Japanese data from [11] showing results of model predictions for the corrosion fatigue behavior of Ni-base alloys in simulated PWR water at 325°C compared with experiments

Etien et al. [12] have also shown corrosion fatigue data in high-temperature, deaerated water for high-chromium weld alloys and Alloy 690 relating to the number of cycles at 0.2 Hz and ∆K = 31 MPa√m (28 ksi√in) required to initiate an approximately 0.5 mm (20 mil) deep crack in a notched 1TCT specimen (root radius = ~0.13 mm) exposed to hydrogenated water at 338°C. Figure 3-6 compares these results with baseline data for 308L stainless steel weld metal, showing improvement factors of 2 to 4 times.

It is also interesting, however, to compare the extent to which the presence of HT-water reduces the number of cycles to fatigue crack initiation under these loading conditions when compared with the values measured for the same alloy in HT-air. The three high-chromium weld metals appear to exhibit reduction factors of between 2.2 and 3.3, which are somewhat smaller than those measured for Alloy 690 (~ 3.9) or for 308L stainless steel (~ 4.5).

3-8


100

1000

10000

100000C

ycle

s to

20-

mil

Cra

ck E

xten

sion

MLTS-24Cr

MLTS-27Cr

MLTS-30Cr

308LStainless

A690Bar

Air: filled symbolsWater: open symbols

4 x2.5 x3 x2 x

Figure 3-6 Corrosion fatigue initiation data from [12] in high-temperature, deaerated water for high chromium weld alloys and Alloy 690 compared with stainless steels

A different aspect of corrosion fatigue relates to the degree of environmental enhancement measured during cyclic crack growth. This has been studied extensively over many years at Argonne National Laboratory [13] (see Figure 3-7) and results for Alloy 690 have recently been reported by Alexandreanu [14] in the context of ongoing studies of PWSCC behavior (described later). For two heats of Alloy 690 plate material in the “as-received” condition (with a final thermal treatment at ~ 720°C) there was no sign of any environmental enhancement during cyclic loading, even as the crack growth decreased (Figure 3-8)p. Instead, at growth rates < ~10-7 mm/s, signs of a ∆K threshold effect became apparent. In contrast, a different plate material, tested under loading conditi able to EACq, showed slight environmental enhancement even in the as-received con uced by u ant env yclic crack growth rates was observed starting at values > ~10-7 mm/s and continuing down to well below 10-8 mm/s. The nature of the changes that can be induced

thick-walled Alloy 690 by extensive, non-uniform cold working, and the extent to which they could be relevant to actual reactor component behavior, are discussed in more detail in later sections of this report dealing with PWSCC testing of such materials.

ons favordition and entirely different behavior after being subjected to ~26% cold work (introdnidirectional cold rolling in three passes). Figure 3-9 shows that, in this case, signific

ironmental acceleration of c

in

p Note, however, that these results were obtained using fast rise times and, in some cases, high R values

(i.e. under loading conditions where significant environmental enhancement of cyclic CGR would not be expected).

q Sawtooth waveform with slow rise time and less extreme load ratios.

3-9


Figure 3-7 Approach to the analysis of environmental effects on cyclic crack growth developed at ANL and now applied to PWSCC testing of Alloy 690 and its weld alloys (from [13])

Figure 3-8 No environmental enhancement of cyclic crack growth seen for thermally treated Alloy 690 material in either simulated PWR primary water or de-aerated pure water at 320°C [14], but under loading conditions not expected to favor EAC

3-10


Figure 3-9 Appearance of significant environmental enhancement at lower rates of cyclic crack

l (blue points) in simulated

y

ive as no

n easured in air either for Alloy 600 or for 690 (Figure 3-12). This work showed a threshold ∆K value of ~ 4

growth for unidirectionally cold-rolled Alloy 690 materiaPWR primary water at 320°C (from [14])

The ANL results on cyclic crack growth rate were obtained largely as a by-product of their experimental technique for SCC studies, but dedicated testing for such effects in simulated primarwater has now also been reported from both Sweden [15] and Japan [16] over a wider range of test conditions. The intent of the latter work was to propose draft fatigue crack growth rate curves for Ni-base alloys to the Japanese standards organization JSME. Figure 3-10 shows that a simulated PWR primary environment can enhance the fatigue crack growth rate of Ni-base alloys by some 5 to 10 times over a fairly wide range of test conditions, with the effect being slightly more pronounced for Alloy 600 than for Alloy 690. These experimental data breach the cyclic CGR curves that would be calculated from existing ASME rules (see Figure 3-11), but are described in a very conservative manner by the proposed Japanese disposition lines from the newly developed model. In contrast, the Swedish study reported in reference [15], which was carried out on archCRDM material from the replacement head for the Ringhals 2 plant, indicated that there wsig ificant influence of the environment on da/dN in comparison with the cyclic CGRs m

MPa√m for Alloy 690 in the primary water environment (Figure 3-13).

3-11


Figure 3-10 Japanese data from [16] indicating that a simulated PWR primary environment can enhance the fatigue crack growth rate of Alloys 600 and 690 by 5 to 10 times over a range of test conditions

3-12


Figure 3-11 Comparison of lines from a proposed (very conservative) Japanese model [16] for cyclic crack growth of Ni-base alloys in a PWR environment and experimental data that breach existing ASME curves

3-13


Figure 3-12 Cyclic CGR behavior of Alloys 600 and 690 in Swedish studies [15]

Figure 3-13Cyclic CGR threshold determined in simulated primary water for Alloy 690 in Swedish studies [15]

3-14


3.3 rro n a r in a

The l M 11 or lu p x si le u inf io on the c n behavior (in particular stress corrosion cracking) of Alloy 690 in both nominal and fault ond a em y. re h hi been exte , th y u e of Allo a m a G in t b t h l de ing nd predi WSCC behavior, which is the main goal of this report. The reason for this is the very diffe ter ha st he s da e, n th rd mi al compositi nce of lithium, boron and hydrogen additions, presence of volatile alkalizing agen t al th rd e h r ro ical corrosion potential that results at th y 69 o -s ur co re th if y w va s edthe primary side by deliberate addition of gaseous hydrogen. Nevertheless, the main conclusions of Appendix C 0 P r a v w h sake mp ss

1. In an AVT water environment, laboratory tests have demonstrated that Alloy 690 ( TT ve gh si t C

2. at Alloy 690TT material usually has significantly higher IGA/SCC resistance than Alloy 600MA and Alloy

ustic water (50% NaOH) at 315°C (600°F). The observed improvement for Alloy 690 is derived

MA

4. In caustic water doped with PbO , Alloy 690 TT material has sometimes shown a somewhat

dified chloride solutions, or when both oxygen and l has a higher SCC resistance than Alloy 600 material

Co sio Beh vio Second ry Water

origina RP- 1 rep t inc ded, in an ap endi , a con derab amo nt of ormat n orrosioed sec ary w ter ch istr The searc on t s has nsive given e earl sy 690 ascting P

stea gener tor (S ) tub g ma erial, ut is no very elpfu to un rstand a

rent wa chemistry t t exi s on t SG econ ry sid not o ly wi rega to che con (abse

ts), bu so wi rega to th much ighe elect cheme Allo 0 sec ndary ide s face mpa d wi the un orml lo lues e tablish on

from the 20 4 MR -111 eport re gi en belo for t e of co letene :

MA or ) is ry hi ly re stant o IGS C.

In deaerated caustic solutions, various laboratory tests consistently showed th

600TT materials. Alloy 690TT may have a slightly lower SCC resistance in highly ca

from longer crack initiation times and lower crack propagation rates compared to the Alloy600 materials. However, these tests also showed that Alloy 690TT material can undergoIGA/SCC in deaerated caustic environments.

3. In neutral or acidic water, or AVT water doped with lead, Alloy 690TT material is moreresistant to IGA/SCC than Alloy 600TT material, which is more resistant than Alloy 600material. However, all of these materials are considered by some investigators to be susceptible to cracking in neutral or acidic environments containing lead.

r

lower resistance than Alloy 600MA and Alloy 600TT, but all three materials are susceptible to IGA/SCC.

5. In chloride-containing solutions in the absence of oxygen, laboratory test data confirm that both Alloy 600 and Alloy 690 materials possess very high SCC resistance. However, they can be susceptible to SCC in highly acichloride are present. Alloy 690 materiaunder such extreme conditions.

6. In acidic solutions containing sulfate and chloride, laboratory tests show that Alloy 690TT material has a much higher IGSCC resistance than Alloy 600TT&MA materials.

7. In deaerated neutral, or slightly caustic, sulfate solutions, laboratory tests indicate that Alloy 690TT is highly resistant to IGA/IGSCC, even when copper or copper oxides arepresent (see, e.g., [17]).

r Such caustic environments are considered very unlikely actually to form in operating steam generators.

3-15


More recent work in this area has focused, in particular, on the role of lead contamination in impairing the SCC resistance of all Ni-base alloys [18, 19, 20, 21], whereby Alloy 690 is generally considered to be more resistant than Alloy 600. EPRI results [22, 23, 24] indicate that Alloy 690TT is only susceptible to cracking in environments with lead that are alkaline; in those environments where cracking was observed, Alloy 690TT was much less susceptible than Alloy 600MA or 600TT. Although there are numerous potential sources for low-level contamination of PWR secondary water by lead [25], the general opinion is that primary water will not be affected by such issues, although some slight uncertainty concerning possible effects on PWSCC remains (see Section C.3.5 of reference [26]).

Another area that continues to be studied extensively on the secondary side is the effect of reduced sulfur species on outer-diameter SCC of SG tubing (see Section 2.3.4.1 of [27] and reference [28]). In this case, however, there are also two confirmed plant incidents of sensitized Alloy 600 SCC having occurred as a result of such species being generated on the primary side following resin intrusions (TMI and Zorita). One of these (1994 incident in Spain) in fact led to PWSCC of 20 out of 37 CRDM penetrations. In part as a consequence of this, the current EPRI Primary Water Chemistry Guidelines limit permissible sulfate concentrations in the reactor water (see Sections 2.3.2.10 and 3.6.1 of [29]). Results from a recent destructive examination of a cracked Alloy 600 CRDM penetration from the Davis Besse reactor also implicated reduced sulfur species in the Alloy 600 PWSCC that was observed [30].

Rec ry water (both highly acidic and highly alkaline) includes a limited comparison with the behavior of

T material [31]. Even under very aggressive conditions that led to rapid rack propagation in Alloy 600, the 690TT tubing material failed to show SCC. A similar result was

reported by de Bouvier et al. [32] in so-called “complex” secondary side environments.

In terms of the implications of secondary side studies for the relative factors of improvement in SCC e 2003 EPR comes to the conclusions shown in Table 3-1. Overall, the following conclusions reached in the 004 MRP-111 report remain valid.

In summary, even though it cannot be considered immune to IGA/SCC, Alloy 690TT material has nevertheless demonstrated far superior IGA/SCC resistance compared to Alloy 600MA or Alloy 600TT materials under most conditions pertinent to faulted secondary water. Lead-doped caustic water is an exception here. Furthermore, Alloy 690 appears to be just as resistant to IGSCC in AVT water as it is to PWSCC in primary water.

ent Japanese work to determine the CGR of Alloy 600 SG tubing in severely faulted seconda

2% cold-worked Alloy 690Tc

behavior for Alloy 690 over Alloy 600 in various environments, reference is made to thI prediction methodology described in reference [33] and a more recent updates of this [34] that

2

s Note that a further update of references [33] and [34] is expected to be published by the end of 2009.

3-16


Table 3-1 Estimated typical improvement factors vs. pHT considering all environments [33, 34]

3.4 Low Temperature Crack Propagation (LTCP)

Anstructural materials, particularly nickel alloys, during (over)loading in water, rather than air.

here is growing evidence that the fracture resistance of some structural materials can be degraded significantly by stressing in water at low temperatures, especially after prior, long-term xposure to high-temperature water. Some data exist to characterize elements of the problem,

the broad nature and extent of potential concerns as a function of loading, temperature, material/microstructure, and environment are not yet well defined. At present, there is consensus that a hydrogen mechanism is involved, but insufficient information to quantitatively predict the effects [35].

3.4.1 Origins of the Phenomenon

J-R tests measure fracture resistance (both fracture toughness and tearing modulus), and - starting with classic work at the Bettis laboratories by Brown and Mills (see [36]) - various investigators have observed large reductions in these properties for precracked specimens of certain nickel alloys (see Table 3-2) when the tests were performed at specific strain rates in low-temperature (< 150°C) water. In these situations hydrogen was present either as an added constituent to the water, as a result of corrosion at the crack tip, or sometimes as a result of prior exposure to a high-temperature hydrogenated environment. These particular results are important in assessments of the integrity of nickel-base alloys operating in PWR primary circuits, since the extent of the decrease in fracture resistance depends on the alloy, and on environmental and dynamic loading conditions that are

emerging issue of concern with PWRs is a possible reduction in the fracture resistance of

T

ebut

3-17


rele r shut down. The phenomenon is usually referred to as “Low Temperature Crack Propro , subcritical crack growth due to the environment [37]. Both aspects may be involved. It is important to recognize that accelerated cracking at low

er

ot

• LTCP does not initiate at as-machined notches, but can initiate at sharp weld defectst.

• Decreasing the hydrogen content in LT water produces a large increase in fracture resistance sceptible materials.

Fe C Mn Si S Ti Nb + Cu P Al Mo Co

vant to reactopagation” (LTCP) and it is not yet clear whether it represents a genuine reduction in fracture perties of the material, or rather a form of rapid

temperature can also occur in some highly susceptible materials at constant load or K, not just undJ-R tearing conditions. Some key general aspects of LTCP behavior are:

• Fracture resistance increases with increasing water temperature. Above 149°C, LTCP is nan issue.

• Cracking resistance is recovered at stress intensity loading rates above 1000 MPa√m/h.(Insufficient time to embrittle grain boundaries ahead of crack.)

for many su

Table 3-2 Chemical composition of Alloys 600, 690, 82, & 52 tested by Brown and Mills [36]

Heat Ni Cr Ta

EN82, heat 1

73.2 18.7 1.5 0.007 2.4 <0.1 0.0007 0.4 0.006 <0.001 0.09 <0.1

EN82H, 76.3 19.93 0.68 0.037 2.73

heat 2 0.06 0.001 0.32 2.44 0.11 0.014 0.02

EN82H, heat 3

72.9 20.00 1.10 0.04 2.90 0.06 0.002 0.41 2.52 0.02 0.003

EN82H, heat 4

73.6 19.75 0.73 0.04 2.93 0.09 <0.001 0.30 2.51 0.07 0.004

EN82H, heat 5

73.7 19.54 0.77 0.04 2.92 0.18 0.001 0.30 2.49 0.02 0.002

EN82H, heat 6

72.8 19.8 1.31 0.041 2.87 0.07 0.004 0.28 2.5 0.07 0.003

EN52 60.4 28.97 8.98 0.03 0.23 0.17 <0.001 0.56 <0.01 0.004 0.63 0.01

Alloy 600, plate

75.4 15.54 7.76 0.07 0.25 0.29 <0.001 0.35 0.11 0.007 0.17 0.04

Alloy 690, bar

59.8 29.54 8.25 0.026 0.29 0.01 0.0005 0.32 0.01 0.001 0.31 <0.001

(a) EN82 heat 1 is the as-welded chemistry. All other EN82 and EN52 were the filler metal chemistry.

3.4.2 Recent Studies to Assess the Possible Relevance of LTCP to PWRs

I-sponsored research is ongoing in an attempt to assess the relevance of such environmental uence on J-R tearing res

EPRinfl istance to reactor components during certain phases of PWR plant

conoperation [38] using both testing [39] and analysis. While historical plant operation provides

firmation of degradation predictions for SCC and corrosion fatigue, there is little (if any)

Or at the tip of pre-existing cracks, e.g. from high-temperature SCC, as studied in tht e EPRI work [39].

3-18


plan The al integrity margins and

s

Mo ent materials) rather than Alloy 690 anWith regard to the relative susceptibility of various materials in low temperature (<150°C) water,

rk sug

•

his rogen

t experience to characterize the nature and magnitude of the environmental fracture issue. consequence of the phenomenon might be a reduction in structur

an increase in the risk of loss-of-coolant accidents to an extent that depends on specific materialof construction and component operating conditions [40].

st of the ongoing work is focused more on Ni-alloy weld metals (both original and replacemd it would be premature to draw final conclusions at this time.

Brown and Mills [36] have proposed the classification scheme shown in Figure 3-14. Their wogests the following for Alloys 600/690:

Alloy 600 is not susceptible to LTCP.

• Alloy 690 exhibits mostly moderate (category II) degradation in fracture resistance but timproves to borderline category II/III behavior at lower concentrations of dissolved hyd(see Figure 3-15 and Figure 3-16).

• Degradation in Alloy 690 fracture resistance is due to hydrogen-induced intergranular cracking, but this is not seen in Alloy 600u with its lower Cr content (see Figure 3-17).

Figure 3-14 Classification scheme for categories of fracture resistance to LTCP after Brown and Mills [36]

is may be partly due to lower solubility of hydrogen in Alloy 600 and thus reduced hydrogen pick-up. u Th

3-19


Figure 3-15 J-R curves determined by Brown and Mills for Alloy 690 in RT air and water at various

temperatures [36]

3-20


Figure 3-16 J and T values deIC termined by Brown and Mills [36] for Alloy 690 in air and water at

various temperatures (values within bars indicate the dissolved hydrogen concentration)

3-21


Figure 3-17 Comparison of fractography from J-R testing under various conditions by Brown and Mills [36] of Alloy 690 (top) and Alloy 600 (bottom)

Bruemmer and Toloczko are performing limited testing for LTCP at the end of some of their HT PWSCC tests (see Section 5.1.2.2.2) and have observed some intergranular crack propagation at 50°C in Alloy 690TT CRDM material, but during fatigue loading, not at constant K [41]. This suggests that homogeneous material representative of plant components may show only limited susceptibility. In contrast, Paraventi and Moshier [42] have reported severe degradation in the fracture resistance for Alloy 690 plate material subjected to large amounts of additional, non-uniform cold working (see Figure 3-18)v. This mirrors the high SCC CGRs for such material discussed in Section 5.1.2.1.7.

v It is noteworthy in Figure 3-18 that even the L-T specimen orientation (where fracture resistance values were less affected) showed extensive, out-of-plane cracking along the 1-D rolling direction.

3-22


Figure 3-18 Results by Paraventi and Moshier [42] from J-R testing of Alloy 690 plate with additional,non-uniform cold work in 50°C water with varying contents of dissolved hydrogen

3.5 References

n

n of Materials in Nuclear Power Systems – Water Reactors, Monterey, CA, 1985 pp. 593-600.

SCC of

5. F. Scenini et al., “Alloy Oxidation Studies Related to PWSCC”, Proceedings of the Twelfth ference on Environmental Degradation of Materials in Nuclear Power

005.

ase Alloys in PWR Water: Oxide Layers and Associated Damage to the Base Metal”, Proceedings of 12th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, TMS, 2005, pp. 883 to 890.

1. A.J. Sedricks, J.W. Schultz, and M.A. Cordovi, “Inconel Alloy 690 – A New CorrosioResistant Material,” Corrosion Engineering (Boshoku Gijutsu), vol. 28, pp. 82-95, 1979, Japan Society of Corrosion Engineering.

2. K. Smith et al., “Inconel 690, A Material with Improved Corrosion Resistance for PWR Steam Generator Tubes,” Proceedings of 2nd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Monterey, CA, 1985pp. 319-328.

3. T. Yonezawa et al., “Effect of Heat Treatment on Corrosion Resistance of Alloy 690,” Proceedings of 2nd International Symposium on Environmental Degradatio

4. J.N. Esposito et al., “The Addition of Zinc to Primary Reactor Coolant for Enhanced PWResistance,” Proceedings of 5th International Symposium on Environmental Degradation Materials in Nuclear Power Systems – Water Reactors, Monterey, CA, 1991. pp. 495-501.

International ConSystems-Water Reactors, Salt Lake City, UT, TMS, 2

6. P. Combrade et al., “Oxidation of Ni B

3-23


7. F. Delabrouille et al., “Effect of the Chromium Content and Strain on the Corrosion of Nickel Based Alloys in Primary Water of Pressurized Water Reactors”, Proceedings of

in LWR Environments,” NUREG/CR-6335 (ANL-95/15), Argonne National Laboratory, 1995.

ver,

690/52/152 PWSCC Research Test Materials Meeting, Industry/NRC RES, July 17-18, 2008,

12th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, TMS, 2005, pp. 903 to 911.

8. L. Legras et al., “ATEM & SEM Study of the Oxides Developed in SCC Cracks and at the Surface of Nickel Based Alloys Exposed in Primary Water”, Proceedings of the 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems, Whistler, British Columbia, August 19-23, 2007, Published by the Canadian Nuclear Society.

9. M.J. Psaila-Dombrowski et al., “Evaluation of Weld Metals 82, 152, 52, and Alloy 690 Stress Corrosion Cracking and Corrosion Fatigue Susceptibility,” Proceedings of 8th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Amelia Island, FL, 1997. pp. 412 to 421.

10. J. Keisler, O.K. Chopra, and W.J. Shack, “Fatigue Strain-Life Behavior of Carbon and Low-Alloy Steels, Austenitic Stainless Steels, and Alloy 600

11. M Higuchi et al., “Revised and new proposal of environmental fatigue life correction factor (Fen) for carbon and low-alloy steels and nickel base alloys in LWR water environments”, Proc. ASME Pressure Vessels and Piping Division Conference July 23-27, 2006, VancouCanada, PVP2006-ICPVT-11-93194.

12. R. Etien et al., EAC Behavior and Mechanical Properties of Improved Alloy 690 Filler Metals. EPRI MRP PWSCC Expert Panel Meeting, St. Petersburg, Florida (Nov. 2007).

13. W.J. Shack and T.F. Kassner: Review of Environmental Effects on Fatigue Crack Growth of Austenitic Stainless Steels, NUREG/CR-6176 ANL-94/1, (May 1994).

14. B. Alexandreanu, “SCC CGRs of Alloys 690 and 52/152 Welds in PWR Water”, Alloys

Rockville, MD http://www.nrc.gov/reading-rm/adams/web-based.html, ADAMS Accession Number: ML082140693.

15. A. Jenssen, K. Norring and P. Efsing, Swedish Activities on Alloy 690 and its Weld EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

16. K. Tsutsumi, Fatigue Crack Growth Rate Curve for Nickel Based Alloys in PWR Environment. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

17. F. Vaillant et al., “Comparative behaviour of alloys 600, 690 and 800 in caustic environments”, 7th Int Symp on Environmental Degradation of materials in Nuclear Power Systems – Water Reactors, Breckenridge, Colorado, Published by NACE.

18. J. Gorman and C. Marks, Proceedings: 2005 EPRI/ANL/NRC Workshop on Effects of Leadand Sulfur on the

Metals.

Performance of Secondary Side Tubing of Steam Generators in PWRs.

EPRI, Palo Alto, CA: 2005. 1012780.

ce 19. B.T. Lu et al., “Passivity of Nuclear Steam Generator Tube Alloy in Lead Contaminated

Crevice Chemistries with different pH”, Proceedings of the 13th International Conferenon Environmental Degradation of Materials in Nuclear Power Systems, Whistler, British Columbia, August 19-23, 2007, Published by the Canadian Nuclear Society.

3-24

http://www.nrc.gov/reading-rm/adams/web-based.html


20. H.P. Kim et al., “Stress Corrosion Cracking of Steam Generator Tubing Materials in LeadContaining Solution”, Proceedings of the 13th International Conference on EnvironmenDegradation of Materials in Nuclear Power Systems, Whistler, British Columbia, August 19-23, 2007, Published by the Canadian Nuclear Society.

21. J. Lumsden and A. McIlree, “Factors affecting PbSCC in Alloy 600 and Alloy 690 Steam Generator Tubing”, Proceedings of the 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems, Whistler, British Columbia, August 19-23, 2007, Published by the Canadian Nuclear Society.

22. B.P. Miglin, Investigation of Lead as a Cause of Stress Corrosion Cracking at Support Plate Intersections. EPRI, Palo Alto, CA 1991. NP7367-S.

23. J. Lumsden, Resistance of Alloy 600 and Alloy 690 Tubing to Stress Corrosion Cracking in Environments With and Without Lead. EPRI, Palo Alto, CA: 2004. 1009532.

24. J. Lumsden, Factors Affecting PbSCC in Alloy 600/Alloy 690 Steam Generator Tubing. EPRI, Palo Alto, CA: 2007. 1014990.

tal

25. C. Marks, Pressurized Water Reactor Lead Sourcebook: Identification and Mitigation of R Secondary Systems. EPRI, Palo Alto, CA: 2006. 1013385.

26. G. White et al., Materials Reliability Program: Reactor Vessel Closure Head Penetration Safety Assessment for U.S. P ts (MREvaluations supporting the MRP Inspection RI, Palo Alto, CA: 2004. 1009807.

27. P ized W r Palo Alto, CA

28. C. Mansour et al., “Behavi r Cotowards an update of the Se del bTests in Sulfate Environme ional Environmental Degradation s, WColumbia, August 19-23, 2 So

29. Pressurized Water Reactor : Vol 6. EPRI, Palo Alto, CA: 2007

30. L. Thomas et al., Materials CRod Drive Mechanism (CR from Davis-Besse Reactor MRP-193. EPRI, Palo Alto

31. Y. Yamamoto et al., “Evalu TT SG Tubing in Primary and Faulted Secob roceedings of 12th International Symposium on EnvironmentaReactors, TMS 1

32. O. de Bouvier Stres s in “ d and Vapor) Environments” mpos vironmental Degradation of Materials in actors, TMS, 2005, pp.1255 to 1266.

33. J. Harris, V. Maroney and J. Gorman, Pressurized Water Reactor Generic Tube Degradation Predictions: U.S. Recirculating Steam Generators with Alloy 600TT and Alloy 690TT Tubing. EPRI, Palo Alto, CA: 2003. 1003589.

Lead in PW

ressurized Water Reactor (PWR) Plan Plan. EP

P-110):

ressur ater Reacto: 2004. 1008

Secondary Water Chemistry Guidelines –224.

or of Sulfur Species in Steam Generatocondary Side Corrosion Cracking Monts”, Proceedings of the 13th Internat of Materials in Nuclear Power System007, Published by the Canadian Nuclear

Primary Water Chemistry Guidelines. 1014986.

Reliability Program: Characterization ofDM) Nozzle Base Metal and Weldment, CA: 2006. 1013419.

ation of Crack growth Rate for Alloy 600dary Water Environments”, P

Revision 6. EPRI,

nditions of PWRs – ased on Laboratory Conference on histler, British ciety.

ume 1, Revision

racks in a Control

l Degradation of Materials in Nuclear Po243 to 1253.

s Corrosion Cracking of Nickel Alloy, Proceedings of 12th International Sy Nuclear Power Systems – Water Re

wer Systems – Water

Complex” (Liquiium on En

, 2005, pp.

et al., “

3-25


34. Alloy 690 Improvement Factor Update: Application of Improvement Factor Data A and

Constellation Energy Group, Inc., Baltimore, MD: 2006. 1013640.

35. C alua f AlterWater S

36. C rown a ll ion Be Alloy 82H Welds, Alloy 52 Welds, Allo RP LTCP Expert Panel Meeting, Los Angeles, CA

37. T. Yonezawa, The Possibil P of Alloy 690 and its Weld Metal in Hydrogenated Wa nel Meeting, Los Angeles, CA (Nov. 2008).

38. A. Demma et al., EPRI Low Temperature Crack Propagation Projects. EPRI MRP LTCP Expert Panel Meeting, Los

39. J. Peng et al., “Effects of D de on the Fracture Resistance of Weld Metals Shutdown Environment”, 13th International Conferen ation of Materials in Nuclear S , W 007C

40. W ceford PRI . ECA: 2008. 1016486.

41. S. Bruemmer and M. Toloc loy during Cyclic Loading at 5 eti geles, CA (Nov. 2008).

42. D.J. Paraventi and W.C. M ng. EPRI Workshop on Cold Work in Iron and N

C. Marks, to the Analysis of a Secondary System Chemistry Upset at Ginna. EPRI, Palo Alto, C

. Marks, Ev tion o ing the Hydrogen Concentration for Mitigation of Primary tress Corrosion Cracking

.M. B

. EPRI, Palo Alto, CA: 2007. 1015017.

s, Low Temperature Crack Propagaty 600 and Alloy 690 in Water. EPRI M

(Nov. 2008).

ity of Metallurgical Effects on LTCter. EPRI MRP LTCP Expert Pa

Angeles, CA (Nov. 2008).

issolved Hydrogen and Hydrogen Peroxi 182, 52 and 152 in Simulated PWR ce on Environmental Degrad

nd W.J. Mi havior of

ystems-WateNS.

. Lun

r Reactors

et al., E

histler, B.C., Canada, August 19-23, 2

Materials Degradation Matrix, Revision

zko, Intergranular Crack Growth in Al0°C. EPRI MRP LTCP Expert Panel Me

oshier, Alloy 690 SCC Growth Rate Testiickel Base Alloys, Toronto (2007).

, Published by

PRI, Palo Alto,

152 and Alloy 690 ng, Los An

3-26

4 PWSCC OF THIN-WALLED SG TUBING

The move to introduce Alloy 690 as a replacement for Alloy 600, starting during the late Eighties, was driven primarily by the need to replace the tube bundles in PWR steam generators (SGs) that had undergone cracking, either from the primary or the secondary side, although the new material was also used extensively from about 1989 for all small diameter nozzle and pressurizer heater sleeve repairs [1]. Accordingly, virtually all of the experience with the new material’s resistance to PWSCC up to 2004 was obtained with thin-walled SG tubing.

4.1 Laboratory Testing

The vast majority of laboratory testing for PWSCC behavior was carried out either on pre-production tubing, at a time when the manufacturing process was being fine-tuned, or on commercial tubing material. In isolated cases, laboratory heats of Alloy 690 were also manufactured and tested and the results from these have been included in this section, since – although not strictly SG tubing – they were also relatively thin-walled.

App f ava environments relevant to primary water for thin-walled Alloy 690 material up to 2004. Certain aspects were updated in [3] and further assessment is

tigations °C

, f

e test

ting crevice conditions), reverse U-bend (RUB), constant load test (CLT), four-point bend, and steam generator tubing mock-up

s are summarized hree groups:

endix A to MRP-111 [2] contained the detailed results from a comprehensive review oilable laboratory test data pertinent to SCC in

currently in progress in the context of the EPRI Steam Generator Program. Numerous inveshave been performed under a variety of environmental conditions including temperatures to 360(680°F) in water with dissolved oxygen levels < 20 ppb, lithium concentrations up to 3.5 ppmboron concentrations up to 1800 ppm, hydrogen concentrations up to 100 cc/kg H2O, additions ochlorides and zinc, and tests in doped 400°C (752°F) steam. Even though the wording “simulatedprimary water conditions” was often used by the authors of published work, it is clear that somconditions employed were outside the normal range of primary water chemistry.

Accelerated testing was performed on double U-bend (simula

specimens. The literature review of Alloy 690 PWSCC behavior, whose resultbelow, comprised essentially studies of crack initiation and was divided into t

1. Deaerated water

2. Deaerated water with additions of boron and lithium

3. Steam with the addition of hydrogen

4-1

PWSCC of Thin-Walled SG Tubing

4.1.1 Early Studies

The first published stress corrosion cracking test results on Alloy 690 material in deaerated high temperature water containing 20 ppb or less dissolved O2 were reported by Sedricks et al. of Inco in 1979 [4]. No cracking was produced in any of the Alloy 690MA U-bend specimens fabricated

me, the

k’s

C pure water than in primary water. Hence, a pure hydrogenated water

environment can be considered slightly more aggressive than hydrogenated water containing

A in [2] reviewed 15 different studies of Alloy 690 PWSCC in high temperature deaerated water with or

U-bend, double U-bend, and constant load specimens from about 40 heats of Alloy 690 have been tested. The carbon contents ranged from

e

pth was 70 µm (0.0028”) in the Alloy 690TT material

an by

observed in bent Alloy 690 and Alloy 52 specimens in several other studies. These observations and

after a maximum exposure of 10,000 hours at 360°C (680°F). Impressed by this test outcoauthors were among the first to suggest that, in high temperature deaerated water, Alloy 690 waseffectively “immune” to the SCC that had been observed in this environment in other high nickel alloys, especially Alloy 600. It is noted that a majority of the Alloy 600MA specimens in Sedrichigh temperature deaerated water tests did not develop cracking by the end of the tests, which may be an indication that the testing condition was not aggressive enough, or that the duration of the tests was not long enough. Airey later investigated Alloy 600 PWSCC in both pure hydrogenated water and primary water (containing boric acid and lithium) [5]. The results showed that the SCinitiation time was shorter in

boric acid and lithium when testing for susceptibility to PWSCC.

Since that time, laboratories in several countries have employed progressively more severe forms of stress corrosion testing under more and more aggressive conditions and to longer test durations in order to assess Alloy 690’s resistance to PWSCC. Most of these studies have confirmed that Alloy 690 material has extremely high PWSCC resistance, as no cracking was produced in the Alloy 690 specimens at the termination of testing in the vast majority of cases. Appendix

without boric acid and lithium. Approximately 300

0.001% to 0.065% and the heat treatment included MA, TT, and thermally aged conditions. Thvery few cases where cracking of Alloy 690 material were reported - and the reasons why they are not representative of in-service PWR conditions - are further discussed immediately below.

4.1.2 Single U-Bend Test in Saturated Hydrogen Water with B/Li

In 1987, Nakayama et al. reported results of Alloy 690 single U-bend tests in “saturated” hydrogenated water containing 1000 ppm boron and 2 ppm lithium at 330°C (626°F) for 3000 hours [6]. Slight intergranular cracking was detected in two out of the four Alloy 690 conditionstested. The maximum intergranular crack dewith an aging treatment of 24 hours at 500°C (932°F) and 30 µm (0.0012”) in the Alloy 690MA material with an aging treatment of 100 hours at 500°C (932°F). The Alloy 690 grain size was not reported, but the 70 µm maximum intergranular crack depth would correspond to two grains deep for a typical grain size of ASTM no. 6.5. Hence, these cracks were shallow and more comparable to an intergranular attack (IGA) depth than to the much deeper PWSCC cracks routinely seen in Alloy 600 material.

In addition, SEM examinations revealed that the intergranular cracks in the Alloy 690TT specimen (70 µm max. depth) were at the tips of dislocation pile-ups. This indicates that the Alloy 690 intergranular cracking could have been caused by mechanical strain from U-bending, rather thPWSCC. It is noteworthy that similar mechanically induced surface intergranular cracks have been

4-2


attributed causes are summarized in Table 4-1 for a side-by-side comparison. It is evident from these reports that Alloy 690 and Alloy 52 can be prone to develop surface defects (including intergranular microfissuring) induced by plastic straining from bending. Shallow, surface-modlayers from SG tubing fabrication processes are seemingly responsible for this mechanical cracking propensity. The CEA and EdF studies (see Appendix A in [2]) showed that the plastic strain at thapex of a RUB is typically between 30 to 40%. This level of plastic straining approaches the rootemperature ductility li

ified

e m

mit of Alloy 690. In addition, Alloy 690 shows a slight drop in ductility as the temperature is increased from room temperature to 315°C (600°F; see Figure 2-5). Alloy 690 also er cold work hardening rate as shown in Figure 4-1, which plots the Vickers hardness nummore sensitive to work hardening during cold work fabrication steps. Any highly orsurface is likely to re-crystallize during thermal treatment o to rto developing surface def p o A ud e 4-demonstrated that these surface defect micr pr ga ug SCCthe end of the tests. Henc IG ted by Nakayama et al. was also due to strain d p en .

Table 4-1 Examples of surface ing in Alloy 690 and Alloy 52 specimens [2

Origin (MRP-111) Ma Des ript rf Cr s an

Defe Attributed Cause

has a highber as a function of cold reduction percentage. Hence, an Alloy 690 surface is expected to be

cold-w be mo

h PW

ked e prone

1 ) by

layer and is, the st

opa

bservedies listed in Tablte (thro

ects com ared ts or

lloy 600. Howeverofissures did not

e, it appears that the short during s

cracking repor applie ecim bending

IG crack

tl.

]

c ion of Su ace IG ack d cts

Miglin, 1986 0M

SG Tubing

on dina a lunt defects wefound on the I.D. surfa f the Alloy 690RUB apex. These defects were not typical

I C. ap examinations showed tha these defects appeared to be like m nical g oves while other red epa

o e g da ce of further IG C crack emanating fromthe base of these defe ere fo d. SE

aminations of archive Alloy 69pe s i ngitudinal

defects, but they re shorter in th an those detected after the AVT autoclave

xp chive llo 0 R ~1. m in Alloy 690

RUB after the AVT exposure). The less ned leg of the RUB specimens

er ee fr d

length obse was due to an opening process of clos r part seddefects. During the high

osure, xide growth

cts uld nh an opening

c y a wedging

appearance of ng

Alloy 69 A&TT

L gitu l shallow nd b re The growth ince o

of GSC Metallogrt some of

hic

echa ros appearain boun

to be sries. No evi

rationdenal ng th

the su(in the defe

SC ing ects w un M pro

exs

0 RUB cimen found sim

welar lo

lengtheA

osure (0.1y 69

to 0.35 mm long in aUB vs.

r0 m

straiw

portionsom thesee fr efects.

rved

ed o ially clo

temperature exprface o

) woance ess b

effect to give an

lo itudinal growth.

Sui, 1997 Alloy

in

Short IG cracks (depth not reported) on the RUB specimen I.D. surface produced

nin t erature. The RUB specim een exposed to hydrogenated steam at 380°C (716°F)

Embrittlement of grain boundaries near the

rf m texposure.

690TT SG Tub g by flatte g at room

en had bemp

for 13,824 hours.

su ace fro he

4-3


Table 4-1 Examples of surface IG cracking in Alloy 690 and Alloy 52 specimens [2] (continued)

Origin (MRP-111) Matl. Description of Surface IG Cracks and

Defects Attributed Cause

Vaillant, 1999

Alloy 6

Shallow cracks (10 µm or 0.0004” deep) Cracking was due90TT SG

Tubing appeared at the apex of the Alloy 690 RUB specimens after bending, but before

red in the Alloy 690 RUB specimen apex surface was not well-defined. However, the stress level measured after removing the 10 µm deep

y 600

ast as high as the apex surface stress lloy 600 RUB specimens. No crack bserved in the Alloy 690 RUB

sting.

to bending.

immersion inside the autoclave. So the stress measu

crack was the same as for the AlloRUB specimens. Hence, the stress at the tip of the short crack caused by bending is at leof the Agrowth ospecimens during the te

Framatome, France

Alloy 690 TT, SG

IG cracks on the innetransition zones of the ki

2003 Tubing and their width 2 µm, see Appendix A for

r tube surface in the ss rolling (1.5 mm

details). Examination of the flattened half the presence of many

microcracks, some of the opened corresponded to those which

were also observed before flattening but

of the tube which were

µm thick in this case). These microcracks

Strain induced by rolling and are limited to the hard perturbed surface layer already present on the inside surface of the tube in the as-received condition.

tube revealed

microcracks

numerous others were new ones opened or created during the mechanical flattening operation. This type of flaw was also observed on partsnot rolled, that is remote from the rolled zone. The crack depth was always limited to the thickness of the perturbed surface layer on the internal surface of the tube (about 10

were intergranular but did not propagate in the tube during exposure to hydrogenated water at 360°C during 60,000 hours exposure.

4-4


400

350

Alloy 690Alloy 600

250

300

100

150

200

0

Vick

ers

Har

dnes

s N

umbe

r

10 20 30 40 50 60 70

Cold Reduction, %

Figure 4-1 Vickers hardness number as a function of cold reduction % from [2]. Alloy 690 has a higher work-hardening rate than Alloy 600

In aspe um induction furnace) Alloy 690 SG tubing. Only the chemical compositions of the tubing were reported and it is nclear whether remelting might have caused changes to the Alloy 690 chemical composition.

Interestingly, these Alloy 690 heats contained a rather high aluminum content of 0.39 or 0.40%.

mal

,

ld s

ons are not encountered in nuclear power plants. Angeliu et al. produced intergranular cracking in a 0.002%

ddition to the factors discussed above, the Alloy 690 plate used to fabricate the U-bend cimens tested by Nakayama et al. was made from remelted (in a vacu

u

(Note, aluminum is not specified in ASME Section II and is often unreported). A. Smith et al. reported that aluminum increases the propensity to caustic IGA in Alloy 690 [7, 8]. Unfortunately, the Alloy 690 microstructure (cracked or not cracked) was not examined to confirm that the thertreatment had produced intergranular carbide precipitations.

4.1.3 CERT Tests in Hydrogenated Water with or without B/Li

Other than the studies by Nakayama et al., and some CEA and EDF tests using high deformation U-bend or RUB specimens in which only mechanical surface intergranular cracking was detectedreports of intergranular cracking in deaerated water originate from CERT studies at strain rates onthe order of 10-7 sec.-1 or less. These are very severe tests for evaluating the susceptibility of Alloy 600 or Alloy 690 to stress corrosion cracking. The specimens are loaded well past the original yiestrength to maintain a continuous plastic strain rate, and are often continued until the specimenfail by ductile overload, if no SCC occurs. Such continuous plastic deformation situati

4-5


carbon laboratory heat of Alloy 690 in deaerated water at 360°C (680°F) during CERT tests [9, 10]. Since the same small amount of intergranular cracking (2%) was also produced in argon, the intergranular cracking in the deaerated water should not be considered an indication of Alloy 690 susceptibility to PWSCC.

Other than the study by Angeliu et al, a study by CEA and EdF also reported Alloy 690 cracking during CERT tests (see [2] for more details). Boursier et al. considered the CERT test as measuring the stress corrosion crack growth rate, rather than propensity to crack initiation [11]. They concluded that Alloy 690 material did not necessarily have a lower CGR than Alloy 600 material and suggested that the CGR obtained in the CERT tests is mainly dependent on the grain boundary carbide precipitation and the creep rate, while the crack initiation time in the RUB tests was mainly a function of Cr contents. It is noted that the most susceptible Alloy 690 heat in this study was a pre-production heat involving a heat treatment that gave rise to a very low density of grain boundary carbides.

Compact tension (CT) specimens, wedge opening load (WOL) specim le cantilever beam (DCB) specimens, or other forms of precracked specimens per ASTM E 399 or E 813 are commonly used for PWSCC CGR testing of Alloy 600 or Alloy 182 materials [12, 13]. During the CGR test, the specimen is loaded under plane strain condition to a pre-determined stress intensity factor which does not necessarily increase during the test. On the other hand, the tensile load in CE reases with test time to keep the elongation rate constant. Hence, CERT tests are very

4.1.4 RUB Test in Hydrogenated Steam

In 1997, Sui et al. reported stress corrosion cracking of Alloy 690 specimens in hydrogenated steam at 380°C (716°F) [14, 15]. Two RUB specimens from the same heat of Alloy 690 developed cracks between 12,600 and 13,824 hours of exposure. Both heats of Alloy 690 had been thermally treated at ~715°C (1319°F) for 15 hours. However, the cracked Alloy 690 Heat B had an unusually (deliberately) low final mill anneal temperature compared to the un-cracked Alloy 690 Heat A (1769°F vs. 1958°F). The average grain size was 25 µm (ASTM no. 7.5) for Heat A and 15 µm (ASTM no. 9.0) or less for Heat B. The small grain size in Heat B is reflected in its higher room temperature yield strength. All the grain boundaries in Heat A had almost continuous carbide coverage. The carbides were determined to be the M23C6 type by TEM. In contrast, Heat B was essentially free of intergranular carbides.

The cracked Alloy 690 heat, Heat B, also had a higher aluminum content (0.14%), which tends to result in a finer grain size and therefore a lower density of intergranular carbides. Interestingly, the Alloy 690 heats used by Nakayama et al. also had a rather high aluminum content. The results reported by Sui et al. show that Alloy 690 could be susceptible to IGSCC in a hydrogenated steam environment and potentially susceptible to PWSCC if the final mill anneal temperature is too low to produce intergranular carbide precipitation upon subsequent thermal heat treatment, or in the event of re-crystallization after thermal treatment. Because such a low temperature solution anneal heat treatment is prohibited by material specifications for Alloy 690 (or Alloy 600) for use in PWRs, and the intergranular carbide precipitation must be verified by optical microscopy and/or SEM, cracking of the type seen by Sui et al. will be prevented. On the other hand, re-crystallization may be a detrimental factor to take into account for the heat affected zones of welds. In any event, the times

ens, doub

RT tests inc different from the typical CGR tests used for Alloy 600 and Alloy 182 materials.

4-6


to failure at 380°C (716°F) reported by Sui et al are such that when extrapolated to normal PWR ope rgy for PWSCC activation energy commonly used for Alloy 600, see Table 4-2) would be well in excess of current expectations for plant life extension. Nevertheless, this possible limited susceptibility induced by the absence of grain boundary carbides is an aspect of Alloy 690 behavior requiring further consideration.

Table 4-2 Summary of Alloy 690 primary water stress corrosion test data to 2004

Ref. # in MRP-111[2]

Test Test Environ.

Test Temp(°F)

Alloy 690 Heat

Number

Alloy 690 Heat

Cond.

Total Speci No.(a)

Test Time at

Test Temp. (Hour)

Equation Test

Time at 600°F(b)

(Year)

Time to First Alloy

600 Failure (Hour)

ImproveFactor

(IFR)

rating temperatures using an Arrhenius relationship and a standard value of the activation enePWSCC, the predicted lives (equivalent to 98 EFPYs at 315°C based on the 50 kcal/mole

Double U-Bend

Deaerated water

600 Y24A7L MA, TT, MA+CW, MA+Weld

52 8064 0.9 3024 2.7 7

Double U-Bend

Deaerated water

680 NX4458HNX4460H

MA, TT 8 23.2 No 600 control

N/A 8064

20 29.9 1500 6.5 13000 11 RUB

Deaerated water + B + Li

680 Three pre-series, Heats 1-3

MA, TT 20 16000 36.8 2000 6.4

MA 3 6600 15.2 19.4 RUB

TT 2 12000 27.6 340

35.3 17

CLT


680 One Alloy 690 heat

TT 5 7000 16.1 1144 6.1

32, 33 C-Ring Deaerated water

644 industrial heats A, E

TT 2 1500 0.9 No 600 control

Two N/A

F MA 6 33000 103.6 41.3

H, G, A, B, D, C, I

MA, TT 47 25000 78.5 31.3

PP TT 6 23000 72.2 28.8 35 – 37 RUB

Deaerated water

689 800

25.6 Y, Z MA 8 20500 64.3

RUB 40 10000 23.0 3000 3.3 43

CLT


680 One Alloy 690 heat

TT 20 10000 23.0 1142(c) 8.8

44 RUB

Deaerated water + B + Li wor wit

ith hout

Zn

662 752246 TT 20 7500 9.1 5500 1.4

46 RUB Deaerated water + B + Li

680

9.092Exp 9.592Exp 9.799Ind 9G4

TT, MA 4 54000 124.1 500 108

Framatome, SG Mocku

Deaerated 680

WE094, TT 16 100000 229.8 800 125

France (new data in [2])

p water Pre-series

(a) Total number of specimens of similar heat treatment and test condition and duration.

(b) The equivalent test time at 600°F for Alloy 690 is calculated based a PWSCC crack initiation time of 50 kcal/mole (Equation 4-7).

(c) 1142 hours is the equivalent of the 4179 hours failure time at 644°F for the Alloy 600 control CLT specimens based on Equation 4-7. Only two of four Alloy 600 MA series results were reported. Hence, the time to the first Alloy 600 failure could be less than 4179 hours.

4-7


4.1.5 Weibull and Weibayes Analyses of the Test Results

The results of the laboratory tests reviewed in Appendix A of [2] were examined to obtain a quantitative or semi-quantitative assessment of Alloy 690’s resistance to PWSCC relative to Alloy 600. This is desirable for developing a statistically-based estimation of a hypothetical PWSCC event in Alloy 690 component items that can be used in selecting and justifying aninspection method.

The two-parameter Weibull distribution of Equation 4-1 is by far the most widely used distribution for stress corrosion cracking analysis [16].

β

θ⎟⎠⎞

⎜⎝⎛−

−=t

etF 1)( Equation 4-1

where: F(t) = the cumulative fraction of cracking by time (t).

θ = the characteristic time (scale parameter) equal to a 63.2% cumulative fraction of cracking.

β = slope or shape parameter

e = 2.71828, the base for natural logarithms.

The significance of β is briefly summarized below:

β < 1 Implies infant mortality, failure rate decrease with time.

β = 1 Implies random failure.

1< β < 4 Implies early wear out.

Implies old age and rapid wear out. A steeper β means less material variation

Weibull analysis has been used to analyze the failures of the Alloy 600 control specimens in the lysis

θ

the time to

β > 4 or tighter quality control. A steeper β means subsequent failures will occur quickly after the first failure.

studies reviewed in [2]. If sufficient data were reported or available, a best-fit regression anawas then performed to determine the Weibull characteristic time (θ) and slope (β).

In addition to Weibull analysis for the Alloy 600 control specimens, a Weibayes analysis (an estimated Weibull) was used in [2] for the Alloy 690 specimens that did not develop crackingby the end of the test period. In the Weibayes method, the slope is assumed to be the same as that from the Alloy 600 analysis. The Alloy 690 specimen characteristic time is estimated by Equation 4-2 [16].

ββ

θ1

⎥⎦

⎤⎢⎣

⎡= ∑

/1

=

N

i

i

rt

Equation 4-2

4-8


where: β = the slope or shape parameter from Alloy 600.

t = test time of each Alloy 690 specimen tested. i

N = total number of Alloy 690 specimen ed.

r = number of Alloy 690 specimens cracked.

As crac 716°F)), it isimm lies to t 3%. The failure number r can be set to produce any level of confidence desired. For example, Weibest uses 0.693 for r to produce a confidence level of 50%, which is less conservative than the 63% confidence level used by Weibayes.

Figure 4-2 is the Weibull plot from [2] of some early Alloy 600 RUB test results at 360°C (680°F) generated in France. The differences in the water chemistry used (beginning-of-fuel-cycle vs. end-of-fuel-cycle) have little impact on the Alloy 600 (one heat) RUB cracking compared to the heat treatment. Hence, the Alloy 600 RUB cracking data in the two primary water environments are combined in Figure 4-2. The Alloy 690 Weibayes characteristic life is calculated based on a total population of 40 RUB specimens from three different heats in either the MA or TT condition since there was no failure of any Alloy 690 RUB specimens. The Alloy 690 Weibayes line in Figure 4-2 assumes a slope of 5.0 (β = 5.0).

s test

no constant deformation (deflection) or constant load Alloy 690 specimen developed SCCks by the end of the test period (except in one study in hydrogenated steam at 380°C ( conservatively assumed that the first (i.e., r = 1) Alloy 690 specimen failure could occur ediately should the test continue. The confidence level that the true Alloy 690 Weibull line

he right of the Weibayes line is 6

90

7563.2

50

25

10600MA

690MA&TTWeibayes

600TT

Cum

uab

ility

, %

Figure 4-2 Weibull plot from [2] of early Alloy 600 RUB results in primary water at 360°C generated in France. The Alloy 690 (three heats) Weibayes line assumes β = 5.0

lativ

e Pr

ob

1,000 10,000 100,000

Time to Failure, Hour

4-9


Figure 4-3 and Figure 4-4 are the Weibull plots of the Alloy 600 RUB results at d by Norring et al. [17]. Figure 4-3 plots Alloy 600 RUB results of different heats of real SG tubing whi -4 plots the Alloy 600 RUB results of EPRI special production tubes. Figure 4-4 sho as a higher PWSCC resistance than the Alloy 600MA annealed at 927°C (1700°F). Figure 4-5 compares the influence of t bing diamet e tim racking. Even though the RU rom t am he ~22mm(7/8”) dia. RUB developed cracking much earlier than the ~19mm (3/4”) dia. RUB. This difference in tim i rib eren ss r e of rent diamet r, n s calculat or ments were reported.Weib c life lated ba a n 7 R pecim fromthree different heats of Alloy 690. These specimens ed in the MA or TT condition with additional variations in heat treatment temperature and duration ne 7 A 690RUB s d cr to 33,000 hourWeib in gure 4-3, Fi e 4- (β)

689°F reporte

le Figure 4ws that the Alloy 600MA annealed at 1024°C (1875°F) h

he SG tuer on th e to c Bs were f he s e heats, t

e to cracker tubing. Howeve

ng was att uted to diffo stres

t streions

levels after the measure

evers bending The Alloy 690

diffe

ayes chara teristic is calcu sed on total populatio of 6 UB s ens were heat treat

. No of the 6 lloy specimenayes line

develope Fi

acking after 20,500 gure 4-4, and Fi

s of5 assumes a slope

exposure. The All of 5.0.

oy 690 gur

90

52

50

25

10

8"

6/8

" d

ghal

4 (3

/4" d

Ri

al 3

(di

a.)

6T

763.

Rin

00

MA

(7ghals

2 (7

/

ia.)

dia.)

Allo

y

Rin

ia.)

Allo

y

ngh

3/4"

00T

690MA&TT eib

100 0 10,000 100,000


Cla

tivac

kin

rob

y,

FWeibull plot from [2] for lts in deaerated water at 365°C reported by Norring e h loy 690 e β 0

W ayes

1,00

umu

e C

rg

Pab

ilit

%

igure 4-3 RUB test resu

t al. [17]. T e Al (many heats) Weibay s line assumes = 5.

4-10


90

7563.2

50

5

0B

&P

17 Hng

t70

0Fnd

vik

0F

TP

Fik

182

1W

T00

F

unti

on 1

Sa 1

70

B&

W18

75Sa

ndv

75F

690 TT We s

100 1,000 100,000

Tim ailure

Cum

ui

nob

ay,

MA&ibaye

10,000

, Houre to F

latv

e C

rack

ig

Prbi

lit%

Figure 4-4 Weibull plot from [2] for RUB test results rrin n hof Alloy 600 in deaerated water at 365°C. The Alloy 690 Weibayes line assumes β = 5.0

by No g et al. [17] on special productio eats

7/8" = Ringhals 2, NX1991 and NX2650, 18 specimens.3/4" = R

90

10

60

ing

3/4"

dia

.

0 t

g

75

inghals 3 and 4, 28 specimens

63.250 0 t

ub

ubin

25

7/8" d

ia. A

lloy

Allo

y 60

690MA&TT

100 1,000 10,000 100,000


Cum

ula

g P

ly,

%

itab

iro

btiv

e C

rack

in

Weibayes

Figure 4-5 Weibull comparison from [2] for different tube diameters from RUB testing by Norring et al. [17] in deaerated water at 365°C. The Alloy 690 Weibayes line assumes β = 5.0

4-11


Extensive constant load tests (CLT) using specimens of Alloy 600M, Alloy 600TT and Alloy 690TT have been performed in Japan (see [2] for further details) and the overall results from theare summarized in Figure 4-6. None of the Alloy 690TT CLT specimens tested at 360°C (680°F) cracked and the Alloy 690 Weibayes characteristic life was calculated in MRP-111 based on apopulation of 20 CLT specimens tested in four variations of primary water chemistry, again assuming a slope (β) of 5.0.

se

total

90

7563.2

50

25

10

600MA at 644oF

690TT at 680oFWeibayese

600MA&TT at 608oFWeibayese

ur

Cum

ulat

ive

Prob

abili

ty, %

Figure 4-6 Weibull plots from [2] of Japanprimary water. No failure was obsetested at 320°C (608°F) and in the Alloy 690TT (one heat) CLT specimens tested at 360°C

failed specimens assume β = 5.0

In addition to the CLT tests, RUB tests with various degrees of prestrain were performed under the riation of the primary

water did not have a significant impact on the Alloy 600 RUB specimens cracking. Hence, the RUB

T

specimens tested at the higher temperature of 360°C (680°F) developed cracking after 10,000 hours of ex sure. The Alloy 690TT Weibayes characteristic life was calculated basepopulation of 40 RUB specimens assuming a slope (β) of 5.0.

1,000 10,000 100,000

Time to Failure, Ho

ese Alloy 600MA (one heat) CLT results at 340°C (644°F) in rved in any Alloy 600MA or Alloy 600TT CLT specimens

(680°F). The Weibayes lines for the un

same primary water conditions. Table 27 in Appendix A of [2] showed that va

test data from the four primary water chemistries were combined before applying the Weibull analyses reported in MRP-111. Figure 4-7 compares the results of 20% prestrained RUB specimensfrom the Alloy 690TT, the Alloy 600TT, and the Alloy 600MA materials. None of the Alloy 690T

po d on a total

4-12


10

25

5063.2

75

90

20% 600MA at 608oF 20% 600TT

at 608oF

20%690TTWeibayes

lativ

e Pr

obab

ility

,

at 680oF

u

Comparison from [2] of Japanese data for 20% pre-strained RUB specimens of Alloy 600MA or Allo sted in primary water at 320°C (608°F), with those for Alloy 690TT, tested under the

The RUB tests in primary water at 360°C (680°F) reported by Vaillant et al. [18] were also evaluated with Weibull analysis. RUB specimens from different heats of A 60 epinto t T ndit and esu ott . N th 6and Alloy 690MA specim el c r u ,0 ur po he A690 Weibayes characteristic life is calculated based on a combined TT an pop n oRUB specimens listed in the report. The A 0 We s line ssum lope f 5.0.

Further information on the extensi udie SCC Alloy 600 and 690 SG ng atin France is given in references [19, 20, 21].

The results of SG mockups tested in deaerated, hydrogenated wat at 6 by F ome P, Fran n Figure 4-9. None of the Alloy 690TT mockups developed any cracking after 100,000 hours of sur Al yes line assumes a slope (β) of 5.0.

The results of the Weibull and Weibayes analyses are used in the next section for estimating the imp primary water conditions.

%

1,000 10,000 100,000


Cum

Figure 4-7

y 600TT, te same conditions, but at 360°C (680°F). The Alloy 690TT Weibayes line assumes β = 5.0

lloyone of

s of exd MA

0 were se Alloy

sure. Tulatio

arated 90TT

lloy f 4

he MA and T co ions ens dev

the roped

lts are plracking afte

ed in Figup to 54

re 4-800 ho

lloy 69 ibaye a es a s (β) o

ve st s of PW in tubi EDF

er 80°F ramat ANce are plotted i

expo e. The loy 690 Weiba

rovement factor (IFR) of Alloy 690 relative to the Alloy 600 in PWR

4-13


10

25

5063.2

75

90

600MA 690MA&TTWeibayese

600TT

100 1,000 10,000 100,000


Cum

ulat

ive

Prob

abili

ty, %

Figure 4-8 Weibull plot from [2] of RUB results at 360°C in primary water reported by Vaillant et al.

y 600 RUB specimens were from four different heats in the MA and TT

TT conditions, experienced no failure after up to 54,000 hours of exposure. The Alloy 690 Weibayes line assumes β = 5.0

[18]. The Alloconditions. The Alloy 690 RUB specimens, also from four different heats in the MA and

10

25

5063.2

75

90

600MAWD281

690TT

mul

ativ

e Pr

obab

ility

, %

Weibayes

600TT WD281

600TT NX3335

100 1,000 10,000 100,000 1,000,000Time to Failure, Hour

Cu

Figure 4-9 Weibull plot from [2] of SG mockups tested in deaerated water at 680°F by Framatome

ce. Alloy 690TT SG mockups experienced no failure after 100,000 hours of

ANP, Franexposure. The Alloy 690 Weibayes line assumes β = 5.0

4-14


4.1.6 Improvement Factor by Weibull Analysis

The Weibull parameters of Alloy 600 materials tested with Alloy 690 are listed in Table 4-3. Using the Weibull analysis, the relative improvement factor, IFR, can be defined by Equation 4-3.

600,690,)1(

AlloyWeibullAlloyrWeibayesIFR θ

θ==

Equation 4-3

Table 4-3 Weibull analysis for Alloy 600 tested with Alloy 690 [2]

Figure Matl. Heat Test Environ. Temp (oF)

Weibull, θ (Hour)

Weibull Slope,

β (a)IFR

(b)

Figure 4-2 690MA&TT

Pre-series 3 Heats RUB

Deaerated water + B + Li 680 30,950 5.00 N/A

Figure 4-2 + B + Li 600MA No info RUB Deaerated water

680 8,112 1.96 3.8

Figure 4-2 600TT No info RUB + B + Li 680 14,539 5.08 2.1 Deaerated water

Figure 4-3 Figure 4-4 Figure 4-5

690MA&TT Many heats RUB Deaerated water 689 59,748 5.00 N/A

Figure 4-3 600 PWR Ringhals 2, 7/8” dia. RUB Deaerated water 689 910 1.74 65.7

Figure 4-3 600MA Alloy 600MA, 7/8” dia. RUB Deaerated water 689 1,440 3.91 41.5

Figure 4-3 600 PWR Ringhals 4, ¾” dia. RUB Deaerated water 689 4,590 2.47 13.0

Figure 4-3 600 PWR Ringhals 3, ¾” dia. RUB Deaerated water 689 5,620 6.05 10.6

Figure 4-3 600TT Alloy 600TT, ¾” dia RUB Deaerated water 689 8,060 5.52 7.4

Figure 4-4 600MA EPRI Provided B&WTP, 1700oF

RUB Deaerated water 689 740 4.53 80.7

Figure 4-4 600MA EPRI Provided Huntington, 1700oF

RUB Deaerated water 689 2,480 10.43 24.1

FigEPRI Provided

ndvik, 00oF

RUB Deaerated water 689 3,490 5.77 17.1 ure 4-4 600MA Sa17

Figure 4-4 600MA EPRI Provided B&WTP, 1875oF


Figure 4-4 600MA EPRI Provided Sandvik,1875o

F RUB Deaerated water 689 10,990 5.27 5.4

Figure 4-5 600 7/8” dia., Ringhals 2, NX 1991, NX 2650


Figure 4-5 600 3/4” dia. Ringhals 3 and 4


Figure 4-6 690MA&TT

One heat CLT Deaerated water + B + Li

680 18,206 5.00 N/A

4-15


Table 4-3 Weibull analysis for Alloy 600 tested with Alloy 690 [2] (continued)

(oF) θ (Hour) Slope,

(a)IFR

(b)Figure Matl. Heat Test Environ. Temp Weibull, Weibull

β644 5,358 5.91 –

Figure 4-6 600MA One heat CLT Deaerated water + B + Li 680 1464 (c) 5.91 12.4

Figure 4-7 690TT One heat RUB 20% prestrain

Deaerated water + B + Li 680 20,913 5.00 N/A

608 6,051 2.10 – Figure 4-7 600MA One heat 20%

prestrain + B + Li 680 414

RUB Deaerated water

1.2 (c) 2.10 5

608 9,530 3.64 – Figure 4-7 600TT

RUB Deaerated water One heat 20%

prestrain + B + Li 680 652 3.64 32.2

Figure 4-8 690TT Four heats RUB Deaerated water + B + Li 680 54,284 5.00 N/A

Figure 4-8 600MA Three heats RUB Deaerated water + B + Li 680 1,776 1.04 30.6

Figure 4-8 600TT Four heats RUB Deaerated water + B + Li 680 3,945 1.85 13.8

Figure 4-9 690TT One heat SG Mock-up Deaerated water 680 174,110 5.00 N/A

Figure 4-9 690MA WD281 SG Mock-up Deaerated water 680 8,954 0.95 19.4

Figure 4-9 690TT WD281 SG Mock-up Deaerated water 680 11,199 1.28 15.5

Figure 4-9 690TT NX3335 SG Mock-up Deaerated water 680 20,231 2.29 8.6

(a) θ is time to Weibull 63.2% cumulative failure for Alloy 600 and the equivalent Weibayes for Alloy 690.

(b) IFR per Equation 4-6.

(c) The equivalent time at 680oF per Equation 4-7 based a PWSCC crack initiation time of 50 kcal/mole.

Hence, IFR is the ratio of the time to a 63.2% cumulative fraction of failure between the Alloy 690 and the Alloy 600 specimens being tested. Transformation of Equation 4-1 yields Equation 4-4, which in turn produces Equation 4-5.

⎟⎟⎠

⎞⎜⎜⎝

⎛

⎥⎦

⎤⎢⎣

⎡−

=β

θ

1

)(11

tFt Equation 4-4

600

690

)(600

690

θθ

=tFgivenato

tt

Equation 4-5

Equation 4-5 shows that, if the Alloy 690 Weibayes line assumes the same Weibull distribution slope (β) as for the Alloy 600 (either TT or MA condition), the IFR to any cumulative fraction of cracking would be a constant. However, the slope of Alloy 690, although unknown due to a lack of cracking in the tests, may not vary as much as seen in Alloy 600 materials. Assuming a small

4-16


Weibull slope value for Alloy 690 can result in a higher Weibayes θ (hence a higher IF ), making the R reases with increasing cracking resistance, i.e., Alloy 600TT tend to have a higher β than Alloy 600MA in the same investigation. However, significant variation exists in the PWSCC susceptibility of different heats of Alloy 600. The variation in the Alloy 600 material can be seen in Figure 4-10, which plots the Weibull θ vs. β listed in Table 4-3, except the data in Figure 4-5 which were derived from the data in Figure 4-3 and Figure 4-4. The average β for Alloy 600MA is 4.12 while the average β for Alloy 600TT is 3.28. Hence, the slope for the Alloy 690 Weibayes line when arbitrarily set to 5.0 is higher than the average value β for Alloy 600. This introduces additional conservatism in the calculation of the Alloy 690 Weibayes θ. No differentiation is made in this case between Alloy 690 in the MA condition and in the TT condition because neither has cracked in the hydrogenated, deaerated water. Therefore, the improvement factor, IFR, is redefined as

R

IF less conservative. A review of the Alloy 600 Weibull plots indicates that the slope inc

600,690,)0.5,1(

AlloyWeibullAlloyrWeibayesIFR θ

θβ === Equation 4-6

0

2

4

Allo

y

6

10

12

0 20,000 25,000

Alloy 600 Weibull (θ) , hour

600

Wei

bu, β

8

ll Sl

ope

0 5,000 10,000 15,00

Alloy 600MA

Removed from PWR

Alloy 600TT

FigureWeibu d β fa

The resulting IFR value 600 test data shown in Figure 4-6 and Figure 4-7 ar li 90 test temperat n the . Overall, the average IFR of Alloy 690 listed in Table 4-3 is 26.5 relative to Alloy 600M

4-10 ll θ an ctors for Alloy 600 tests listed in Table 4-3

s are listed in Table 4-3. The Alloye normaure i

zed per the Arrhenius relationship of Equation 4-7 to 360°C, the Alloy 6 same studyA and 13.3 relative to Alloy 600TT.

4-17


⎥⎥⎦

⎤

⎢⎢⎣

⎡

⎟⎟⎠refTT

⎞⎜⎜⎝

⎛−∝ i

ref RQ

tt 11exp Equation 4-7

where Qi o

4.1.7 Im emen

Unfortuna ot all t er obtained or reported s nt dat ll or Weibayes analysis. In order to include these data, an alternative way to calculate 8 can be used:

= 50 kcal/m le.

prov t Factor with Minimum Alloy 600 Crack Time

tely, n he studies reviewed in 2004 (see Appendix A of [2]) had eithufficie a to allow a Weibu

the improvement factor using Equation 4-

600690Alloyofime

AlloyincrackingfirstTIFR = Equation 4-8

0 d

600 eal

as less than the accumulated test time at inspection intervals. Often, the first Alloy 600 specimen cracking was observed at the first

nterval. In addition, the actual cracking time of Alloy 690 specimens

TheIn a ) for each test is also determined per

AlloThe oy 600 in Table 4-2 and Table 4-4 per Equation 4-8

Tabthe

TesttoTime

Here, the time to first cracking in Alloy 600 refers to the time when the first failure in the Alloy 60control specimens was observed. The following factors are considered for the IFR values determinefrom Equation 4-8.

1. The PWSCC resistance of Alloy 600 has been observed to vary greatly and depend on several factors. Like the Alloy 690 specimens, some Alloy 600 TT specimens did not develop any cracking by the end of the test. Hence, the calculated IFR is relative to Alloy materials that are highly susceptible to PWSCC, such as Alloy 600MA with low mill anntemperature and a high degree of intragranular carbides. The IFR so calculated would be a more realistic improvement factor, somewhat less conservative than that by Equation 4-6.

2. The IFR calculated per Equation 4-8 is still considered conservative because the actual cracking time of the Alloy 600 specimens w

scheduled inspection icould be much longer than the test duration, if-in fact-they were ever to develop cracking at all.

resulting IFR values calculated by Equation 4-8 are listed in Table 4-2 and Table 4-4. ddition, the equivalent EFPYs at 315°C (600°F

Equation 4-8. This normalization has no affect on the IFR per Equation 4-8 except when the y 600 and Alloy 690 specimens in the same study were tested at different temperatures. average IFR of Alloy 690 relative to All

is 27.1, which is about the same as the 26.5 relative to Alloy 600MA from the Weibull analysis per Equation 4-6. Figure 4-11 plots the duration of the test for Alloy 690 vs. the IFR listed in

le 4-2 and Table 4-4. The IFR is seen to increase with increasing test time, indicating that IF is limited by the test duration. R

4-18


Table 4-4 Summary of Alloy 690 hydrogenated and doped hydrogenated steam stress corrosion test data to 2004

Environ. Heat Number Heat Speci Test Time at Alloy 600 Improve.

Factor (IF)

Ref. # in MRP-111 Test Test Test

Temp Alloy 690 Alloy 690 Total Test

Time at Equation

Test Time to

First

[2] (°F) Cond. No.(a) Temp. (Hour)

600°F(b)

(year) failure (hour)

690 (A) TT, 1 13824 107.4 25.0

TT 1 12600 (c) 97.9 22.8 52, 53 RUB H2 716 Steam 552

22.8 690 (B)

TT+aged 1 12600 (c) 97.9

RUB 3 9720 237.4 336 28.9 Framatome, Germany H (new data in [2])

RUB surface scored

2

Steam 752 754380 TT

3 9720 237.4 336 28.9

(a) Total number of specimens of similar heat treatment and test condition and duration.

(b) The equivalent test time at 600oF for Alloy 690 is calculated based a PWSCC crack initiation time of 50 kcal/mole (Equation 4).

(c) Two of the three Alloy 690 specimens in Sui’s investigation cracked after 13824 hours of testing. To be conserit is assumed here that these two Alloy 690 specimens developed crack soon after the previous examination made after accumulating 12600 hours. None of the other Alloy 690 specimens in this table failed during the test duration.

vative,

140

160

00 20,000 40,000 60,000 80,000 100,000 120,000

20

40

60

Impr

80

ovem

en

100

120

t Fac

t, IF

Primary WaterH2 SteamDoped H2 SteamH2 Steam, crackedLinear (Primary Water)

Stress Corrosion Test Duration, hour

Figure 4-11 Improvement factors listed in Table 4-2 and Table 4-4 per Equation 4-8 versus test duration

4-19


Sininit out by Westinghouse

hea erall exp SCC in these tests, nor did two heats of

acc600relatively short test durations used.

Mo[23]. The testing was performed in high-purity water with 50 ml/kg H2 at 365°C as shown in

in t R o Equ 9 vs. Alloy 600MA and ~ 5 vs. Alloy 600TT.

ce MRP-111 was published in 2004, there have been at least two major additional studies of iation behavior in Alloy 600 and 690 SG tubing. The first was carried

in order to ascertain the possible effects on PWSCC of operating at a slightly higher pH value in the primary circuit [22]. A single heat of Alloy 690TT was included together with multiple

ts of Alloy 600 and the testing was carried out under different chemical conditions with ovosures of up to ~10,000 h. None of the Alloy 690 RUB specimens showed any signs of PW

Alloy 600TT (out of a total of 7 heats in all). Depending upon the exact Alloy 600 heat considered as a reference (and showing cracking), the IFR factors calculated

ording to Equation 4-8 would range from 2.6 to 13.3 vs. Alloy 600MA and 1.6 to 4.0 vs. Alloy TT in the most susceptible (pH 7.2) test environment. These factors appear low because of the

re recently, the experiences from Swedish RUB testing of Alloy 690 have been summarized

Table 4-5. No cracking was observed in either the mill-annealed (MA) or thermally-treated (TT) conditions for testing times up to 33000 h, whereas Alloy 600 cracked under these conditions, even

he TT condition (see Figure 4-12 from [24]). In this case, the IF factors calculated according tation 4-8 would be ~ 4

Table 4-5 Swedish RUB testing for crack initiation in Alloys 600 and 690 [23]

4-20


Figure 4-12 Swedish RUB testing for crack initiation in Alloys 600 and 690 from [24]

In addition, some limited testing of Alloy 600 and Alloy 690 steam generator tubing material in supercritical water (SCW) was carried out by Westinghouse as a precursor to a major study of CGRs in these materials (see Section 5.1.2.1.8). This program, carried out for the EPRI MRP, was designed to accelerate the IGSCC process in SCC-resistant, Ni-base alloys by testing at elevated temperatures and pressures [25]. Specifically, this study involved testing at temperatures beyond the critical point of pressurized aqueous environments in an attempt to thermally accelerate corrosion and produce measurable effects from relatively minor SCC processes. High-pressure/temperature test conditions (385°C/33 MPa) were chosen to have “water-like” fluid properties with additions of Li/B and H2. An initial series of crack initiation tests showed that SCC initiation occurred in one Alloy 600 heat in this supercritical water (SCW) environment after a period of only 200 h. Including data from similar testing at 325°C suggested an activation energy for crack initiation of ~146 kJ/mole over the temperature range from 325°C to 385°C. This is less than the 209 kJ/mole often quoted for Alloy 600 in primary water, but consistent with the most recent KAPL data reported by Richey et al. [26]. No cracking was found in any of the 14 Alloy 690TT RUB specimens after 500 h.

4.2 Field Experience

Over the past thirty years, PWSCC of Alloy 600/182/82 material has become an increasingly significant problem in PWRs. As shown in Table 1-1, PWSCC was first detected in Alloy 600 steam generator tubing in 1971. The first report of PWSCC in a pressure boundary penetration was the leak from a pressurizer level nozzle in 1986. Both of these instances occurred after approximately 2 years of service. These initial reports have subsequently been followed by similar occurrences at other PWRs using Alloy 600/182/82 materials for the same component items. These problems have led to significantly increased inspection efforts, have contributed to the early replacement of many steam generators, have led to long forced outages to repair problems at some units, and have resulted in the expenditure of a significant amount of engineering effort to determine the root cause of the problems, develop inspection and repair methods, and develop strategic plans to manage the problems.

4-21


In addition to the initial occurrences of Alloy 600/182/82 PWSCC given in Table 1-1, the following steam generator tubing failures are highlighted:

1. Cracking of Obrigheim steam generator tubes (1971) after approximately 2 years of service [27]. This was the first instance of PWSCC of Alloy 600MA.

2. Cracking of Kori-2 steam generator tubes (1986) after nearly 3 years of service. This was the first occurrence of PWSCC in a steam generator fabricated with Alloy 600TT tubing [28].

3. Cracking of steam generator mechanical tube plug at North Anna 1 (1989) after just over 3 years of service [29]. This was the first occurrence of PWSCC of Alloy 600TT plugging material.

4. Cracking was identified at Oconee Nuclear Station-1 (1995) after about 18 years of operation [30]. This was the first instance of PWSCC in stress relieved OTSG (Once-Through Steam Generator) tubing.

It seems quite clear that a PWR steam generator is a very demanding material application, involving high temperature, high stress and strain at roll transitions and at tight U-bends, and a primary water environment that facilitates stress corrosion cracking (at a corrosion potential close to the potential of the Ni/NiO equilibrium). In addition, the number of SG tubes per PWR varies between 8,000 and 30,000. Hence, steam generator tubing is statistically favored to be the first component item in a PWR to show PWSCC. Operating experience indicates that a PWSCC-susceptible Alloy 600 material (low temperature mill annealed) can crack after approximately 2 years of operation in this application. Even improved Alloy 600 materials (high temperature mill annealed and thermally treated or high temperature mill annealed and stress relieved) may also be susceptible to cracking in as little as 3 to 8 years of operation at very highly stressed and/or cold-worked locations (such as tube bulges, kiss rolls, etc.). Nevertheless, Figure 4-13 (taken from reference [31]) shows that there has been a marked decrease in repairs to Alloy 600TT tubing in the last 2 years.

As shown in Section 4.1, Alloy 690 material appears to be highly resistant to PWSCC when laboratory tested in conditions where Alloy 600 material exposed under the same conditions routinely showed comparatively rapid cracking. Also, as shown in Table 4-6, a number of steam generators manufactured with Alloy 690 tubing material have now been in service for significant times without crack indications due to PWSCC. This unblemished record with regard to the SCC resistance of Alloy 690 tubing has continued through 2008, even though a small number of repairs have been necessary for other reasons (see Figure 4-14). The oldest steam generators with Alloy 690 tubing have been operating now for approximately 18 years with no reported signs of PWSCC degradation. As of December 2008, a total of 89 PWR plants have replaced steam generators originally tubed with non-thermally-treated Alloy 600 and the vast majority of new staem generators have been tubed with Alloy 690TT.

4-22


Figure 4-13Worldwide causes of Alloy 600TT SG tube repair by degradation mechanism from [31

]

Figure 4-14 Worldw uses

ide ca of Alloy 690TT SG tube repair by degradation mechanism from [31]

4-23


Table 4-6 List of operating steam generators manufactured with Alloy 690 tubing (as of December 2008)

Installation Date PWR

1989 D.C. Cook-2, Indian Point-3, Ringhals-2

1990 Dampierre-1

1991 Ohi-3

1992 Penly-2

1993 Mills North Antone-2, na-1, Beznau-1, Ohi-4

1994 Mihama-2, Takaham Gravelines-1, Daya Bay-1, Daya Bay-2, Golfech-2, Ge

a-2, Genkai-1, V.C. Summer,nkai-3, Ikata-3

1995 Ohi-1, Tihange-1, North An erre-3, St. Laurent-B1, Sizew

na-2, Ringhals-3, Dampiell-B

1996 Miha Takaham nna, Gravelines-2 ma-1, a-1, Doel-4, Catawba-1, Gi

1997 McG McGuire Ohi-2, Tricastin-2, Genkai-4 uire-1, -2, Point Beach-2, Mihama-3,

1998 Byron-1, St. Lucie-1, Tihange-3, Braidwood-1, Kori-1, Ikata-1, Tricastin-1

1999 Beznau-2

2000 Farle C. Cook- Texas Project-1, Arkansas Nuclear One-2, Choo Civaux-1, avelines-4

y-1, D. 1, Krsko, Southz-B1, Civaux-2, Gr

2001 Genkai-2, Ikata-2, Kewaunee, Shearon Harris, Farley-2, Tihange-2

2002 Tricastin essenheim z-B2, Calvert Cliffs-1, South Texas Project-2 -3, F -1, Choo

2003 Calv s-2, Palo onee-1, t.-Laurent Des Eaux B ert Cliff Verde-2, Sequoyah-1, Oc S

2004 Tricastin 4, Oconee 2, Prai nd 1, Oconee 3, Dampierre 2, rie Isla

2005 Calla NO 1 way, A

2006 Beaver Valley 1, Fort Calhoun, Watts Bar 1

2007 Com e Peak 1manch , St. Lucie 2

2008 Diablo Canyon 2

Some other replacement and original equipment component items utilize relatively thin-walled

ction practices, optimizing material melting and manufacturing practices, and reducing residual stresses. However, the excellent operating experience to date provides assurance that Alloy 690TT is highly resistant to PWSCC. The overall factor of improvement relative to Alloy 600 MA was estimated in 2006 from field data as >20x [3]. However, it is clear that this number is actually very conservative and efforts are ongoing to justify higher values [32].

Alloy 690 material. E.g., the replacement pressurizer heater sleeves at Calvert Cliffs-2 have been in operation for about 12 years and visual inspections have not identified any problems to date.

Obviously, a number of additional improvements to design and fabrication practice have also been made to reduce the likelihood of initiating PWSCC with thin-walled Alloy 690 material. These include optimizing component item constru

4-24


4.3 References

1. WRs: Review of Cracking Events and Repair Service Experience”, Proceedings of 12th International Symposium on Environmental Degradation of Materi ea s – Water Reactors, TMS, 2005, pp. 959 to 966.

2. H. Xu et al., Materials Reliability Program (MRP), Resistance to Primary Water Stress Corrosion Cracking of Alloys 690, 52, and 152 in Pressurized Water Reactors (MRP-111). EPRI, Palo Alto, CA, 2004. 1009801.

3. C. Marks, Alloy 690 Improvement Factor Update: Application of Improvement Factor Data Second yst m p na

Constellation Energy Group, Inc., Baltimore, MD: 2006. 1013640.

4. A.J. Sedricks, J.W. Schultz, and M.A. Cordovi, “Inconel Alloy 690 – A New Corrosion Resistant Material,” Corrosion Engineering (Boshoku Gijutsu), vol. 28, pp. 82-95, 1979, Japan Society of Corrosion Engineering.

5. G.P. Airey, “The Stress Corrosion C in Pure and Primary Water Environments”, Procee st) Inon Environmental Degradation of Materials in Nuclear Power Systems – Water Rpp. 462-478, NACE, 1983.

6. Nakayama, H. Tomari, K. Fujiwara, K. Shimogori, H. Hamada, and K. Takaishi,

sion

International Symposium on Environmental Degradation of Materials in Nuclear Power l Island, GA, 1989. pp.5-33 to 5-46.

8. A. Smith and R. Stratton, “Thermal Treatment, Grain Boundary Composition and r m

on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors,

” s

Power Systems – Water Reactors, Monterey, CA, 1991. pp. 475-481.

lar

ional Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, American Nuclear Society, 2003.

W. Bamford and J. Hall, “Cracking of Alloy 600 Nozzles and Welds in P

als in Nucl r Power System

to the Analysis of a ary S em Che istry U set at Gin . EPRI, Palo Alto, CA and

racking (SCC) Performance of Incondings of the (1

el Alloy 600ternational Symposium

eactors,

“GC/IGSCC and General Corrosion Behavior of Alloy 800 As a PWR S/G Tube Material,” Corrosion 87, Paper No. 82, NACE, March 1987.

7. A. Smith and R. Stratton, “Relationship between Composition, Microstructure and CorroBehavior of Alloy 690 Steam Generator Tubing for PWR Systems,” Proceedings of 4th

Systems – Water Reactors, Jekyl

Interg anular Attack Resistance of Alloy 690,” Proceedings of 5th International Symposiu

Monterey, CA, 1991. pp. 855-860.

9. T.M. Angeliu, J.K. Sung, and G.S. Was, “The Role of Carbon and Chromium on the Mechanical and Oxidation Behavior of Nickel-Based Alloys in High Temperature Water,Proceedings of 5th International Symposium on Environmental Degradation of Materialin Nuclear

10. J.K. Sung and G.S. Was, “The role of Grain Boundary Chemistry in Pure Water IntergranuStress Corrosion Cracking of Ni-16Cr-Fe Alloys,” Proceedings of 4th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Jekyll Island, GA, 1989. pp.6-25 to 6-37.

11. J.M. Boursier, F. Vaillant, P. Saulay, Y. Brechet, and G. Zacharie, “Effect of the Strain Rate on the Stress Corrosion Cracking in High Temperature Primary Water: Comparison Between the Alloys 690 and 600,” Proceedings of the 11th Internat

4-25


12. Materials Reliability Program Crack Growth Rates for Evaluating Primary Water Stress Corrosion Cracking (PWSCC) of Thick-Wall Alloy 600 Materials (MPR-55) Revision 1. Final Report, November 2002. EPRI Report 1006695.

13. Crack Growth of Alloy 182 Weld Metal in PWR Environments. Final Report, January 1999. EPRI Report TR-111993.

14. B. Heys, and J. Congleton, “Stress Corrosion Cracking of Alloy 600 and Alloy 690 nal

ymposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Amelia Island, FL, 1997. pp. 274 to 281.

15. G. Sui, J.M. Titchmarsh, G.B. Heys, and J. Congleton, “Stress Corrosion Cracking of Alloy lloy 690 in Hydrogen/Steam at 380oC,” Corrosion Science, Vol. 39, No.3, pp.

16. bull Handbook,” 2nd Ed., Dr. Robert .B. Abernethy, Author and Publisher, July 1996.

17. K. Norring, J. Engstrom, and P. Norberg, “Intergranular Stress Corrosion Cracking in Steam Generator Tubing, Testing of Alloy 690 and Alloy 600 Tubes,” Proceedings of 3rd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Traverse City, MI, 1987. pp. 587-593.

18. F. Vaillant, EDF Report HT-44/95/013/A, 1996, “Résistance a la corrosion sous contrainte en milieu primaire des alliages 690 et 800 – Point des résultats en Décembre 1995”, (English translation of the title, “Resistance of Alloys 690 and 800 to Stress Corrosion Cracking in PWR Primary Water – Status of Results Available to December 1995”).

19. F. Vaillant et al., “Influence of chromium content and microstructure on creep and PWSCC resistance of nickel base alloys”, 9th Int Symp on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Newport Beach (CA), August 1999, Published by

20. F. Vaillant et al., Assessment of PWSCC Resistance of Alloy 690: Overview of Laboratory esults and Field Experience. EPRI Workshop Alloy 600, Santa Ana Pueblo (NM), March

21. F. Vaillant et al., Assessment of PWSCC resistance of Alloy 690: an overview of laboratory results and field experience EPRI MRP PWSCC Expert Panel Meeting, Saint Petersburg, FL, November 2007.

22. R.J. Jacko and R.E. Gold, “Crack Initiation in Alloy 600 SG Tubing in Elevated pH PWR Primary Water”, Proceedings of 12th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, TMS, 2005, pp. 925-936.

23. K. Norring and P. Efsing, “Influence of Material Parameters on Initiation of PWSCC in Nickel Base Alloys in Primary PWR Environment”, Workshop on Detection, Avoidance, Mechanisms, Modeling, and Prediction of Stress Corrosion Cracking Initiation in Water-Cooled Nuclear Plants. EPRI, Palo Alto, CA; TEPCO R&D Center, Yokohama, Japan; AREVA NP, Technical Center, Le Creusot, France; Institut de Radioprotection et de Sûreté Nucléaire (IRSN), Fontenay-aux-Roses, France; Institute of Nuclear Safety System, Incorporated (INSS), Fukui, Japan; The Materials Aging Institute (MAI), Moret Sur Loing Cedex, France; and EDF R&D, Moret Sur Loing Cedex, France: 2009. 1018908.

G. Sui, G.in Hydrogen/Steam and Primary Water Side Water,” Proceedings of 8th InternatioS

600 and A565-587, 1997.

“The New Wei

TMS.

R2005.

4-26


24. P. Efsing, “Alloy 690 Issue from a Utility Perspective”, EPRI MRP PWSCC Expert Panel Meeting, St. Petersburg, Florida (Nov. 2007). See also K. Norring and J. Engstrom, “Initiation of SCC in nickel base alloys in primary PWR environment: studies at Studsvik since mid 1980s”, Energy Materials Vol. 3 (2008) No. 2, pp. 113-118.

25. R. Jacko, Materials Reliability Program: Testing the Resistance to Stress Corrosion Cracking of Alloy 690 and its Weld Metal in Supercritical Boron/Lithium/H2 Solutions (MRP-225). EPRI, Palo Alto, CA: 2007. 1015004.

26. E. Richey, D.S. Morton and R.A. Etien: SCC Initiation Testing of Nickel-Based Alloys in High Temperature Water, Proc. 13th Int. Conf. On Environmental Degradation of Materials in Nuclear Power System – Water Reactors, CNS (The Canadian Nuclear Society), 2007.

27. Steam Generator Reference Book. December 1994. EPRI TR-103824s-V1R1.

28. Experience of U.S. and International Steam Generators with Alloy 600TT and Alloy 690TT Tubes and Sleeves. 2002. EPRI Document No. 1003589.

29. “NRC Bulletin No. 89-01: Failure of Westinghouse Steam Generator Tube Mechanical Plugs,” U.S. Nuclear Regulatory Commission, Washington, D.C. May 15, 1989.

30. Experience of U.S. and International Steam Generators with Alloy 600TT and Alloy 690TT Tubes and Sleeves. 2002. EPRI Document No. 1003589.

31. T. Kaul, Steam Generator Management Program: Steam Generator Progress Report: Revision 16. EPRI, Palo Alto, CA: 2008. 1016561.

32. nce in Steam J. Benson, Research on Material Improvement Factors for PWSCC ResistaGenerator Tubes. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

4-27

5 PWSCC OF THICK-WALLED ALLOY 690 MATERIAL

The predominant use of thick-walled Alloy 690 material as a replacement for Alloy 600 has beenRPV head penetrations, the vast majority of which are CRDM or CEDM nozzles. The first insof this followed PWSCC cracking and coolant leakage at the Bugey 3 PWR in France in 1992led to the replacement of vessel heads at 33 further plants there up to 2000 (see [1]). In contrast, the first U.S. plant to undergo RPV head replacement following PWSCC problems was North Anna 2 in 2002, but t

in tance and

he pace picked up rapidly after that and nearly 30 plants had been equipped with new heads containing Alloy 690 penetrations by the end of 2008 (see Table 5-1). As mentioned in Section 2.1.2 and discussed in more detail in refer [2], a variety of melting and manufacturing steps have been used to produce extruded pipe (with a typical wall thickness of ~25 mm) that has usually then been subjected to a subsequent thermal anneal (“Alloy 690TT”), although some cold stra

Table 5-1 Examples of some replaced PWR RPV heads with CRDM penetrations in Alloy 690

ence

ightening may also have been carried out as a final step.

In-Service Date PWR

1992 Bugey 3

1994 Bugey 5, Blaiyais ey 21, Gravelines 4, Bug

1995 Blayais 2, St. Alban 1, Fla ravelines 3, Blaiyai 3, Tricmanville 1, G s astin 1

1996 Tricastin 4, Paluel 4, St. Laurent B2, Blaiyais 4, Dam ierre 1, Fessenheim 1, St. A

plban 2

1997 B Dampierre 2, Da rre 4, Belleville 2, Cr nes 5 ugey 4, mpie uas 4, Graveli

1998 F lle 2, Dampierre luel 3, Cattenom 2, Fessenheim 2, Crualamanvi 3, Pa s 2

1999 C , Chooz B2 hooz B1

2000 C Civaux 2 ivaux 1,

2002 N na 2 orth An

2003 N na 1 and 2, Crystal River 3, Ginna 1, Ocon 1 and 3, Surry 1 and 2, Three Mile Island 1

orth An ee

2004 Farley 1, Kewaunee 1, Oco e 2, Turkey Point 3 ne

2005 ANO 1, Farley 2, Millstone 2, Point Beach 1 and 2, Prairie Island 2, Robinson 2, Salem 1 and 2, St. Lucie 1, Turkey Point 4

2006 Beaver Valley 1, Calvert Cliffs 1, D.C. Cook 1, Fort Calhoun 1, Prairie Island 1,

2007 Calvert Cliffs 2, Commanche Peak 1, D.C. Cook 2, St. Lucie 2,

5-1

PWSCC of Thick-Walled Alloy 690 Material

Apart from CRDM penetrations, relatively thick-walled Alloy 690 has been used in number of other Alloy 600 component replacements in U.S. plants, as indicated in Table 5-2.

Table 5-2 Examples of some relatively thick-walled Alloy 690 reactor coolant system original equipment or replacement component items other than CRDM penetrations

Component Item In-Service Date

PWR Material

1990 Calvert Cliffs 2

1992 Palo Verde 1,San Onofre 3, San Onofre 2, Palo Verde 1

1993 St. Lucie 2, Palo Verde 2

Alloy 690 Tubing

1994 Palo Verde 3, St. Lucie 2

1995 San Onofre 3, St. Lucie 1

1997 San Onofre 2, San Onofre 3

1998 Calvert Cliffs 2, San Onofre 3

1999 San Onofre 3

PressurizeComponents

r

(e.g., Instrument Nozzles, etc.)

Alloy 690 Tubing and Alloy 52 Weld

2001 Waterford 3

2005 St. Lucie 1

2006 Millstone 2, ANO 2, Fort Calhoun

RV Lower Head BMI Nozzles 2000 Civaux 1, Civaux 2 Alloy 690 Tubing and Alloy

152/52 Weld

2003 South Texas 1 Alloy 690 Tubing and AWeld

lloy 52

5.1 Laboratory Testing

Despite the widespread early use of thick-walled Alloy 690 material for CRDM penetrations, onlyone pre-production and one commercial heat of this product form were apparently tested early on in France (by EdF [3] – see Table 5-3). No indications were given that the PWSCC behavior of specimens taken from these differed in any way from that of the majority of thin-walled SG tubing materials examined previously (see Section 4.1).

5-2


Table 5-3 Origins and heat treatments of Alloy 690 CRDM nozzles tested by EdF [3]

CRDM Nozzles

Manufacturer Process Heat Number

Tube Maker

Extrusion Temperature

(°C)

Solution Anneal and Heat

Treatment Temperatures

Final Finish

Diameter/Thickness

(mm)

Experimental Tecphy ESR WJ 151 Valinox 1110-1230 1050 + 5 h x715°C

- Pre-straighten in a press 1 straightening pass withhyperbolic rolls

109/21,5

Industrial Tecphy ESR WJ 172 Valinox 1110-1210 1080+ 5 h x 715°C

- Pre-straighten in a press 2 straightening passes with VALTI rolls

110/23

With the exception of the Alloy 690 plate material used to simulate a SG divider plate in corrosiofatigue and SCC tests of various welds (see Section 3.2), no other tests on thick-walled Alloy 690 material appear to have been carried out prior to publication of the original MRP-111 report in 2004. That situation has changed rapidly, however, in the last 5 years, as discussed below.

n

5.1

5.1.1.1 Testing in Simulated Primary Water by Mitsubishi Heavy Industries (MHI)

Initiation testing at MHI using uniaxial tensile specimens under constant, active load has been carried out in simulated prima e 6 fo y s t o n in Table 5-4 and Table 5-5. The environmental test conditions are shown in Table 5-6 and detailsof th system, including specimen form and loading mechanism, are shown in Figure 5-1 to Fi 3.

Key s we a . i June 007 [ d e 20 5- shows the apparent absence of “crack initiation” after 58,000 h in Alloy 690TT taken from a bottom mou stru tion nozz re h i g e r 0 oAlloy 690TT taken from a CRDM nozzle. Note, however, that it has not been possible to confirm whet ot en l been xam e fr i cr I.e., the data shown in the above figures apparently refer to the absence of failure in any of the Alloy 690TT tensile specimens, many of which are still thought to be on test.

.1 Crack Initiation Studies

ry wat r at 3 0°C r man year using the ma erials sh w

e testgure 5-

result re presented by Asad et al n 2 4] an Nov mber 07 [5]. Figure 4

nted in menta le and Figu 5-5 s ows s milar ood b havio after 73, 00 h f r

her or n these specim s have actua ly e ined y t for eedom from inc pient acks.

5-3


Considering the highest stress level at which a direct comparison can appropriately be made, the

Table 5-4 Chemical composition of materials used in MHI testing for PWSCC initiation [4, 5]

minimum improvement factors calculated according to Equation 4-8 would be ~48 and ~30 vs. Alloy 600MA for the BMI and the CRDM materials, respectively. Note, however, that these numbers are undoubtedly conservative and expected to rise with increased testing time.

Table 5-5 Heat treatment and mechanical properties of materials used in MHI testing for PWSCC

, 5] initiation [4

5-4


Table 5-6 Environmental test conditions used in MHI testing for PWSCC initiation [4, 5]

Figure 5-1 MHI test loop for uni-axial constant load studies of PWSCC initiation [4, 5]

5-5


Figure 5-2 “Active” loading mechanism for uni-axial constant load studies of PWSCC initiation [4,

5]

Figure 5-3 Test specimens for uni-axial constant load studies of PWSCC initiation [4, 5]

5-6


Figure 5-4 Dependence of PWSCC initiation in Alloy 600MA on applied stress and lack of cracking i690TT BMI material after 58,000 hours of testing [4, 5]

n

Figure 5-5 Dependence of PWSCC initiation in Alloy 600MA on applied stress and lack of cracking in690TT CRDM nozzle material after 73,000 hours of testing [4, 5]

5-7


5.1 ing in Pure Supercritical Water at the University of Michigan

WaDep stenitic alloys for service in the Gen IV SCWR [6]. This included quasi CGR dataw on Alloy 690, derived from cross-sections of CERT specimens after testing at a slow strain rate in pure, de-aerated supercritical water (SCW) at temperatures

bove. Even then, it performed better than stainless steels and approximately as well as Alloy 625.

In conjunction with the Westinghouse test program in SCW described at the end of Section 4.1.7 and 1.2.1.8, a small, amount of focused testing at the University of Michigan was spo e two SCW studies [7]. CERT of tensile specimens machined from several Alloy 690 matin SCW at t ained (see Table 5-7) showed that all alloys were more susceptible to cracking in SCW at 400°C and a pressure of 25.4 MPa than in a lower temperature (385°C) environment at higher pressure (26.7 MPa). Alloy 600 was clearly SCC susceptible (Figure 5-7); the very marginal cracking susceptibility of Alloy 690 (Figure 5-8) was similar in the three material conditions tested. The decrease in cracking at 385°C appeared to be predominantly due to the drop in temperature; i.e., the water density had little effect on the extent of cracking.

.1.2 Test

s and Teysseyre reported in 2005 on an extensive test program, sponsored by the U.S. artment of Energy, to qualify various au

between 400 and 550°C (see Figure 5-6). Alloy 690 showed only very marginal SCC susceptibility under these severe test conditions until reaching temperatures of 500°C and a

in Section 5.nsored by EPRI to investigate the possible cause of differencesx between th The approach used involvederials (including one tested in the Westinghouse program) and an Alloy 600 control material

wo temperatures and pressures. The data obt

Figure 5-6 CGRs derived from a CERT study of various austenitic alloys in pure, de-aerated SCW [6]

w Note that attempts to derive CGRs from this type of testing are controversial and subject to different

interpretations. Given the small crack depths involved for Alloy 690, these were clearly related more to initiation than to crack growth.

x Higher cracking susceptibility in the Westinghouse CGR tests, which used additions of boron and lithium to SCW.

5-8


Table 5-7 Summary of maximum crack depths measured on cross-sectioned samples at the University of Michigan [7]

er of een

Alloy Test Condition Max Crack Depth Observed (µm)

Number of Cracks Measured

NumbCracks S

UM-690 400°C 4.4 3 >20

385°C 3.5 2 <5

UM-TT690 400°C 7.0 3 >20

385°C 3.5 1 <5

EPRI-TT690 400°C 5.3 2 >10

385°C 3.4 2 <5

EPRI-600 400°C 169 2 >40

385°C 17.6 2 >40

Figure 5-7 Cross-section of EPRI Alloy 600 after testing in 400°C/25.4MPa deaerated pure SCW [7]

5-9


Figure 5-8 Cross-section of EPRI Alloy 690 after testing in 400°C/25.4 MPa deaerated pure SCW [7]

5.1.2 Crack Growth Rate Studies

The approach taken in the 2004 MRP-111 report was to analyze the data from Alloy 690 laboratory studies according to the type of test environment that had been used and this scheme was retained in the 2008 description of ongoing research in this area prepared for the MRP by the present author [8]. It is increasingly apparent, however, that the product form and detailed thermo-mechanical processing of Alloy 690 material can have a profound influence on its PWSCC behavior. Accordingly, the CGR studies that have now been carried out (and are, in some cases, still ongoing) are reported here according to the type of thick-walled Alloy 690 material that was used. Since the majority of initial CGR tests were performed on plate materialy, this is dealt with first, even though it is of far less relevance in the field than the piping material used for CRDM penetrations. These are dealt with next, whereby a distinction is made between the base material that would actually be delivered to the manufacturer of a replacement RPV head (see Section 2.1.2) and similar material that has been deliberately subjected to additional cold work in the laboratory before testing. Finally, consideration is given to the expected behavior of the heat affected zones (HAZ) formed adjacent to the fusion line in thick-walled Alloy 690 components that have undergone welding.

5.1.2.1 Testing of Alloy 690 Material not directly related to PWR Components

The plate and other materials widely used to date in CGR studies of Alloy 690 are not generally thought to be representative of field components, but represent either a standard alloy from metal ven re seldom – an experimental heat of material.

dors, or – mo

y Later also on bar material.

5-10


5-11

5.1.2.1.1 Feasibility Studies in Simulated Primary Water by General Electric Global Research (GE-GRC)

GE-GRC has been working on the PWSCC behavior of Alloys 690/152/52 for the EPRI MRP program since 2003, initially by carrying out a feasibility study of the application of sophisticated CGR testing methods to study such resistant alloys [9].

Product form

The feasibility testing was carried out on 2 Alloy 690 plates supplied by EPRI (Heat NX8244 HK11 with the chemical composition and the mechanical properties shown in Table 5-8), one with a final mill anneal (MA) at 982°C (1800°F) and the other at 1093°C (2000°F). The low-temperature MA produced a rather heterogeneous grain size, together with a lower level of grain-boundary precipitates and significant compositional banding (mainly of carbides). Details of the extensive banding in this plate material have recently been reported by Morra et al. [10]. The high-temperature MA resulted in a more uniform microstructure and a higher density of grain-boundary carbides. Both plates were subsequently subjected to uniform cold work (by cross-forging at RT) to a level of approx. 25%. This was intended both to raise residual SCC susceptibility and to cover the maximum level of deformation expected in the HAZ immediately adjacent to the fusion line after weldingz.

Table 5-8 Chemical composition and mechanical properties of all the Alloy 690 materials tested at GE-GRC in the MRP program

Source Heat # Cr Fe Al Ti Mn Si C S Cu RT YS

MPa

RT UTS MPa

EPRI Plate 982C MA

NX8244 HK11

30.0 9.20 0.36 0.20 0.20 0.14 0.018 <0.001 242aa 210bb

656aa 595bb

Duke CRDM

WN415 29.1 10.1 0.31 0.29 0.018 <0.001 0.007 260 514

ANL Plate

NX3297HK12 29.5 9.9 0.20 0.07 0.03 <0.001 0.01

GRC bar

B25K 29.3 9.21 0.26 0.37 0.22 0.06 0.034 <0.001 <0.01 293 699

Other NX2865HK 29.3 9.99 0.16 0.03 0.03 <0.001 0.01 292 688

Specimen/crack-growth orientation

Compact tension type (0.5T CT) specimens, with 5% side grooves on each side, were machined from the cross-forged plates in the L-T orientation (see Figure 5-9). The crack plane was thus perpendicular to the banded microstructure.

z Note, however, that such deformation in the HAZ is likely to be more uni-directional than homogeneous. aa For material with MA at 982 C. bb For material with MA at 1093 C.


Figure 5-9 Convention with regard to specimen orientation and the principle axis of cold-working in Alloy 690 base materialcc

Tes

The en level chosen at 18 and 26 cc/kg, respectively, so as Ni/NiO phase stability transition. Particular attention is paid to transitioning from the transgranular air fatigue pre-crack to a uniform (and possibly intergranular) SCC crack front by gentle cycling with increasing R ratio and

Data evaluation

The test program relies on highly accurate DCPD measurements of on-line crack depth to assess CGRs in the various phases of testing, with a plausibility check being provided by a comparison of cumulative crack increments from on-line data and the overall length of the final crack on the fracture surface after breaking open the CT specimens, often after extended test periods of many thousands of hours.

The feasibility testing described in reference [9] showed stable, sustained crack growth – albeit at very low rates of < 10-8 mm/s – over long periods of time at constant K in the 25% cold-worked plates tested at 340°C (see Figure 5-10 and Figure 5-11). The highest CGR at ~ 0.25 mm/year was assessed to be almost two orders of magnitude lower than that expected for Alloy 600 under comparable conditions and thus of little or no engineering significance. The crack morphology was primarily transgranular over the very limited increments of crack advance obtained (see Figure 5-12) and no apparent difference was found between low and high temperature mill annealing of the starting material.

ting details

MRP program at GE-GRC uses simulated primary water at 340 or 360°C with the hydrog to be around the

hold time, before attempting to impose constant KI conditions (generally at ~27.5 MPa√m).

cc Rolling is done in the longitudinal (L) direction, with reduction occurring in the short transverse (S) direction

and some broadening in the width (T) direction. Forging reduces the S direction, but the T and L directions are indistinguishable.

5-12


SCC#2 - c248 - 690, 25%RA, NX8244HK111, 1800F Anneal

11.3

11.31

11.32

11.33

11.34

11.35

11.36

11.37

11.38

500

Test Time, hours1000 1500 2000 2500 3000 3500 4000 4500 5000

Cra

ck le

ngth

, mm

-1

-0.8

-0.6

-0.4

-0.2

0

0.2

0.4

Con

duct

ivity

, µS/

cm o

r Pot

entia

l, V

she

Outlet conductivity x 0.01

CT potentialPt potential

c248 - 0.5TCT of 690 + 25%RA, 340C25 ksi√in, 550 B / 1.1 Li, 18 cc/kg H2

To R

=0.7

, 0.0

01 H

z+

9000

s ho

ld @

506

h

To R

=0.7

, 0.0

01 H

z+

85,4

00s

hold

@ 6

65h

1.4 x 10-8

mm/s

NE

Pow

erO

uta g

e @

124

7h

2.4 x 10-9 mm/s

To C

onst

ant

K@

175

7h

5 x 10-9 mm/s

EN

D O

F TE

ST @

513

5h

At 340C, pH = 7.60. At 300C, pH = 5.93 and potential would be ~155 mV higher

Figure 5-10 CGR response of cold-worked Alloy 690 plate (with low-temperature mill anneal) testeat GE-GRC for >3000 h at constant stress intensity [9]

d

5-13


SCC#2 - c249 - 690, 20%RA, NX8244HK112, 2000F Anneal

11.31

11.33

11.35

11.37

11.39

11.41

11.43 0.4

500 1000 1500 2000 2500 3000 3500 4000 4500 5000

Test Time, hours

Cra

ck le

ngth

, mm

-1

-0.8

-0.6

-0.4

-0.2

0

0.2

Con

duct

ivity

, µS/

cm o

r Pot

entia

l, V

she



c249 - 0.5TCT of 690 + 20%RA, 340C25 ksi√in, 550 B / 1.1 Li, 18 cc/kg H21.4 x 10-8

mm/s

To R

=0.7

, 0.0

01 H

z+

9000

s ho

ld @

506

h

To R

=0.7

, 0.0

01 H

z+

85,4

00s

hold

@ 6

65h

NE

Pow

erO

utag

e @

124

7h8 x 10-9

mm/sTo

Con

stan

t K

@ 1

757h

3.7 x 10-9 mm/s

EN

D O

F TE

ST @

513

5h

At 340C, pH = 7.60. At 300C, pH = 6.93and potential would be ~155 mV higher

Figure 5-11 CGR response of cold-worked Alloy 690 plate (with high-temperature mill anneal) testeat GE-GRC for >3000 h at constant stress intensity [9]

d

5-14


Figure 5-12 Predominantly transgranular morphology (but with some intergranular facets) within band of PWSCC crack growth in cold-worked Alloy 690 plate initially tested at GE-GRC [9]

5.1 nvestigations in Simulated Primary Water at Argonne National Laboratory (ANL)

ata from the ANL test program, sponsored by NRC Research, has been reported at successive etings [11, 12, 13] as well as in the context of public NRC review meetings

A 5.4 cm (2.125 in) thick, VIM-ESR plate (Heat NX3297HK12, supplied by ATI Wah Chang, with the chemical composition shown in Table 5-9), annealed for 2h at 1038°C and apparently subj laboratory thermal stabilization treatment [15] at 720°C for 10h, was uni-overegi 1.5 in thick final plate being harder than the mid-plane by 12%.

.2.1.2 I

DEPRI Expert Panel me[14, 15].

Product form

ected to a subsequentdirectionally cold-rolled at Special Metals in 3 passes (10, 8 and 8%) to achieve ~26% rall reduction in thickness. The deformation was inhomogeneous, with the near-surface on of the

Table 5-9 Chemical composition of Alloy 690TT plate material tested at ANL [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Department of Energy under Contract No. DE-AC02-06CH11357.

5-15



1/2T CT specimens were removed as shown in Figure 5-13 and have been tested in both the S-L and S-T orientations, with similar crack growth rates being attained.

Testing details

The test environment is simulated primary water at 320°C, and the chosen H2 level (2 ppm = ~22 cc/kg) is in the region of Ni metal stability, i.e. above the expected level for peak CGR in Alloy 600. Pre-cracking was carried out by fatigue at temperature in the autoclavedd, with a gradual change (via increasing R-ratio and sawtooth loading with longer rise times) to approximately constant mechanical load. ANL analyze these cyclic data according to the superimposition model of Shack and Kassner [16] and consider clear acceleration of cyclic crack growth above the so-called true corrosion fatigue line (άCF) as frequency is lowered to be a reliable indicator of subsequent IGSCC behavior. In the case of cold-rolled Alloy 690 tested at 320°C, the degree of environmental enhancement observed was actually above the best-fit curve for previous ANL tests on Alloy 600 (see Figure 3-9).

During subsequent testing at constant load (at a KI value of 26.7 MPa√m), steady crack growth was pecimen with S-L orientation. After some 2000h, a brief period of fatigue loading was MP (see Figure 5-14). The crack advance during the two constant load test periods amounted to approx. 130 and 70 µm, respectively.

seen in the s used to increase the crack length. The final constant load test period at a KI value of 28.0a√m lasted ~600 h

3.5”

17”

rolling direction

Approx 1.5”1 2 3 4SL ST

Approx 1.5”

3.5”

17”

rolling direction

1 2 3 4Approx 1.5”

3.5”

17”

rolling direction

1 2 3 4SL ST

rtment

Figure 5-13 Details of sample removal from cold-rolled Alloy 690TT plate tested at ANL [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Depaof Energy under Contract No. DE-AC02-06CH11357.

dd n essential part of successful transitioning from transgranular cracking ANL regard this step in the procedure as a

to IGSCC.

5-16


Figure 5-14

nt

500 h under constant load at a KI value of 31.0 MPa√m (see Figure 5-15).

On-line data from ANL testing [11 to 15] of cold-rolled Alloy 690TT plate (S-L specimen orientation) Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Departmeof Energy under Contract No. DE-AC02-06CH11357.

For the S-T specimen orientation, steady crack growth of around 150 µm was obtained over nearly 1

Figure 5-15 On-line data from ANL testing [11 to 15] of cold-rolled Alloy 690TT plate (S-T specimen orientation) Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Department of Energy under Contract No. DE-AC02-06CH11357.

5-17


Data evaluation

The fracture surfaand predominantl

ce for the S-L specimen orientation showed fairly uniform crack advance overall y intergranular cracking during the constant load periods (see Figure 5-16),

whereas the macroscopic crack front appears somewhat more irregular for the S-T orientation(see Figure 5-17).

Figure 5-16 Macro- and microfractography for the ANL 690TT plate specimen with S-L orientation [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Department of Energy under Contract No. DE-AC02-06CH11357.

5-18


Figure 5-17 f tion [11 to 15]

Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Department

The following representative PWSCC CGRs were derived (apparently without any correction of the on-line DCPD data from investigation of the fracture surfaces):

1.5 E-8 mm/s at KI = 26.7 MPa√m and 2.8 E-8 mm/s at 28.0 MPa√m for the S-L specimen;

3.3 E-8 mm/s at KI = 31 MPa√m for the S-T specimen.

For comparison purposes, the MRP-55 disposition curve (75th percentile of the data on a heat by heat basis) would predict a CGR of ~ 1E -7 mm/s for Alloy 600 at KI = 31 MPa√m. Thus the ANL study could be interpreted to predict a factor of improvement of only ~3x for heavily cold-worked Alloy 690 material on this basis, or even no improvement at all if a comparison is made with the 50th percentile of the Alloy 600 database.

During testing of a second S-T specimen, the temperature was changed several times within the range 320 to 300°C without any apparent effect on the CGR as measured online (see Figure 5-18).

Macro ractography for the ANL 690TT plate specimen with S-T orienta

of Energy under Contract No. DE-AC02-06CH11357.

5-19


12.70

12.80

13.10

13.20

50

60

12.90

ack

Leng

13.00

h (m

103000 3500 4000 4500 5000

Time (h)

20

30

40

tm

)

Km

ax (M

Pa

m0.

5 )

max

Cr

K

Crack Length

Alloy 690Specimen # A690WC-ST-2320°C, Simulated PWR Water18

24 24 a2 h

24 b4 h

24c8 h

25C,

20°C

,

26

27

tem

p

erat re

0°C

,

erat

ure

310°

re-in

itial

ized

emp

ure

back

to 3

-initi

aliz

ed

pera

t 3

0re

-inedt

ure

toiti

aliz

tem

320°C 310°C 320°C 300°C

riodic Unloading

2.5 x 10–11 m/s26 MPa m0.5

CL w/ periodic unloading

nt

Tes fter uni-directional (1D) rolling, in one case of the same plate used at ANL (heat NX3297HK12 with approx. 26% cold work, see Section 5.1.2.1.2)

at GE-GRC of one-dimensionally cold rolled Alloy 690 plate supplied by ANL 2.

2.8 x 10–11 m/s26 MPa m0.5

Pe

Figure 5-18 ANL test showing that the CGR behavior for cold-worked Alloy 690TT plate was unaffected by changing temperature within the range 320 to 300°C [13] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Departmeof Energy under Contract No. DE-AC02-06CH11357.

5.1.2.1.3 Additional Studies in Simulated Primary Water by General Electric Global Research (GE-GRC) as part of the MRP Test Program

ting of Alloy 690 has been carried out a

and in another of Alloy 690 billet material (heat B25K) supplied by GE-GRC and cold rolled by 20% [17, 18]. A further heat (NX2865HK) of nuclear-grade Alloy 690 from another source has also been tested after undergoing ~19% 1D rolling. The composition and mechanical propertiesof these materials are shown in Table 5-8 and the basic test procedures used at GE-GRC have already been described in Section 5.1.2.1.1.

The first testproduced even higheree CGRs (~ 4E-7 mm/s) than those reported by ANL in Section 5.1.2.1.Crack advance was seen easily at a constant KI of 27.5 MPa√m after 700 h test time at 360°C and some 0.7 mm of growth then occurred over less than 1000 hours (see Figure 5-19).

ee Albeit at a higher temperature.

5-20


SCC#5 - c372 - Alloy 690, 26%RA 1D, NX3297HK12, ANL

12.9

13 0.4

12.2CT potentialPt potential

12.3 -0.8

12.4

12.5

h

-1

-0.6

-0.4

0

0.2

Con

duct

ivity

, µS/

cm o

rte

ntia

l, V s

he

c372 - 0.5TCT of 690 + 26%RA 1D, 360C25 ksi√in

12.6

ack

leng

t

12.7

, mm

-0.2 Po

12.8 Outlet conductivity x 0.01

650 750 850 950 1050 1150 1250 1350 1450

Test Time, hours

Cr

, 600 B / 1 Li, 26 cc/kg H2

Est. pH at 360C = 8.2 used for φcAt 340C, pH = 7. t 300C, pH = 6.86 53. A

tant

K

47h

at constant K

To C

onst

ant

Load

To C

ons

@ 6

~0.7 mm of growth

@

h

mm/s

690 plate supplied by ANL and tested

4 x 10-7

1339

3 x 10mm/s

-7

Figure 5-19 Rapid crack growth in 1D cold-rolled (~26%) Alloyat GE-GRC [17, 18]

After advancing the crack some 1.2 mm in fatigue, a second period of SCC crack growth at constant K gave CGRs that were initially somewhat lower, but gradually increased to ~ 2E-7 mm/s,at which point the test temperature was dropped, first to 325°C and then to 290°C. These changes produced no long-term reduction whatsoever in CGRff, but indications of “stepwise” growth (see Figure 5-20). The macroscopic crack front became less uniform at the lower temperatures (lighter areas in Figure 5-21). The observed high CGRs were duplicated in a separate, repeat test on this ANL plate material. This time, however, instead of changing temperature, the hydrogen concentration in the simulated primary water was increased from 26 to 80 cc/kg over a period of 1500 h (see Figure 5-22). This produced a slight increase in CGR, whereas such a change would have been expected to reduce CGR in Alloy 600 undergoing PWSCC.

ff For Alloy 600, the expected decrease in CGR upon this reduction in temperature would have been more than one

order of magnitude.

5-21


SCC#7a - c372 - Alloy 690, 26%RA 1D, S-L Orientation, NX3297HK12, ANL370

14.8

340

360

erat

Temperature15.8

14.42300 2500 2700 2900 3100 3300 3500 3700 3900 4100 4300

Test Time, hours

280

14.6290

300

15

Cra

310

320

Tem

p

15.2

ck le

ngth

, mm

330 ure,

C

15.4

15.6350

c372 - 0.5TCT of 690 + 26%RA 1D, 360C27.5 MPa√m, 600 B / 1 Li, 26 cc/kg H2

mm/s

After ~1.2 mm of growth

To C

on

by cyclic loading from prior constant K data2.2 x 10-7

2.2 x 10-7

s

mm/s

tant

K@

To R

=0.7

9000

s ho

0C @

347

3 163

5h

0.0

01 H

zld

@ 1

540

h

To

32

+ h 5C @

282

1h

,

To 2

9

NL and apparent Figure 5-20 Second period of rapid PWSCC in 1D cold-rolled Alloy 690 plate from Alack of a CGR response to reducing temperature from 360 to 325, then to 290°C [17, 18]

Figure 5-21 Change in macroscopic appearance of PWSCC region upon reducing temperature (c372) [17, 18]

5-22


SCC#2 - c400 - Alloy 690, 26%RA 1D, S-L Orientation, NX3297HK12, ANL19.5 0.2

12.5

13.5

14.5

15.5

16.5

17.5

18.5

1400 1900 2400 2900 3400 3900 4400

Test Time, hours

Cra

ck le

ngth

, mm

-1.1

-1

-0.9

-0.8

-0.7

-0.6

-0.5

-0.4

-0.3

-0.2

-0.1

0

0.1

Con

duct

ivity

, µS/

cm o

r Pot

entia

l, V s

he




Est. pH at 360C = 8.2 used for φcAt 340C, pH = 7.53. At 300C, pH = 6.86

To

Con

stan

t K

@ 8

02h

To 8

0 cc

/kg

H2

@ 1

858h

7 x 10-7

mm/s

2.5 x 10-7

mm/s

To 2

6 cc

/kg

H2

@ 3

391h

~5 x 10-7

mm/s

EN

D O

F TE

ST @

427

6h

sing the dissolved hydrogen concentration [17, 18]

neous before additional

C .5 to 17.3 MPa√m) in applied stress intensity also failed to slow down cracking at all in this specimen. A repeat test gave an

ogen concentration from 26 to 80 cc/kg (Figure 5-25).

en lower CGRs at constant K, indicating a clear effect of specimen orientation in the 1-D cold rolled material.

Figure 5-22 Repeat test at GE-GRC on 1-D rolled ANL plate material showing a slight increase in CGRupon rai

Figure 5-23 shows the moderate CGR (4.2E-8 mm/s) measured in the S-L orientation on a different heat of Alloy 690 material that had been 1-D cold-rolled by 20% at GE-GRC. The microstructure of this forged bar material was extremely homogecold-working. As for the ANL plate material, changes in test temperature produced no long-term

GR response (Figure 5-24). Interestingly, a significant drop (from 27gg

almost identical PWSCC CGR and no response whatsoever to an increase in hydr

The same material was tested in the S-T orientation (cf. Figure 5-9) and gave a somewhat slower CGR of ~8E-9 mm/s (Figure 5-26). Furthermore, in contrast to the S-L oriented specimens, this CGR dropped significantly upon reducing temperature to 325°C, more-or-less in line with the change that would be expected if the 130 kJ/mole activation energy for Alloy 600 were applicable here. Later on, a different area of this specimen was tested and produced ev

gg Occurring because the other (lead) test-specimen in the daisy chain within the same autoclave showed

significantly larger amounts of crack growth.

5-23


SCC#5 - c373 - Alloy 690, 20%RA 1D, S-L Orientation, Heat B25K

11.68

11.69 0.4

11.65

11.66

mm

Pot

e

11.59650 750 850 950 1050 1150 1250 1350 1450

Test Time, hours

-1

11.6

11.61

11.62

11.63

11.64

Cra

ck le

ngt

-0.8

-0.6

-0.4

Con

duct

ivity

, µS/

cm o

11.67

h, -0.2

0

0.2

rnt

ial,

V sheOutlet conductivity x 0.01




To C

onst

ant

K @

647

h

~0.09 mm of growthat constant K

To C

onst

ant

Load

4.2 x 10-8

mm/s

@ 1

33specimen's growth rate is slower

Figure 5-23 Moderately rapid PWSCC in 20% 1D cold-rolled Alloy 690 GE-GRC forged bar [17, 18]

9h

K has slowly decreased to 22.9 MPa√m because this

SCC#7a - c373 - Alloy 690, 20%RA 1D, S-L Orientation, Heat B25K

12.16

12.18

12.2

12.22

12.24

12.26

12.28

2300 2500 2700 2900 3100 3300 3500 3700 3900 4100

Test Time, hours

Cra

ck le

ngth

, mm

280

0

310

320

330

340

350

360

370

Tem

pera

ture

, C

Temperature, C

290

30

c373 - 0.5TCT of 690 + 20%RA 1D, 360Cby cyclic loading

27.5 MPa√m, 600 B / 1 Li, 26 cc/kg H2

After ~0.4 mm of growth

To

Con

stan

t K

@ 1

635h

To R

=0.7

, 0.0

01 H

z +

9000

s ho

ld @

154

0h

2 x 10-8

mm/s

To 3

25C

@ 2

821h

K has slowly decreased to 17.3 MPa√m because this

specimen's growth rate is slower

3 x 10-8

mm/s

To 2

90C

@ 3

473h

sity [17, 18]

Figure 5-24 Continued testing of 20% 1D cold-rolled Alloy 690 GE-GRC forged bar with reduction in both test temperature and applied stress inten

5-24


SCC#2a - c401 - Alloy 690, 20%RA 1D, S-L Orientation, Heat B25K

11.561400 1500

-1.1CT potentialPt potential

11.58

11.6

11.62

11.64

11.66

11.68

11.7

11.72

1600 1700 1800 1900 2000 2100 2200 2300 2400

Test Time, hours

Cra

ck le

ngth

, mm

-1

-0.9

-0.8

-0.7

-0.6

-0.5

-0.4

-0.3

-0.2

-0.1

0

0.1

0.2

0.3

Con

duct

ivity

, µS/

cm o

r Pot

entia

l, V s

he




To C

onst

ant

K @

802

h

4 x 10-8

mm/s

To 8

0 cc

/kg

H2

@ 1

858h

Figure 5-25 Lack of response to a major change in dissolved hydrogen during testing of a second specimen from a 20% 1-D cold-rolled Alloy 690 forged bar [17, 18]

SCC#3 - c393 - Alloy 690, 20%RA 1D, S-T Orientation, Heat B25K11.26

360

370

11.21

11.22

900 1400 1900 2400 2900

Test Time, hours

Cr

300

310

320

11.23

11.25

ack

leng

t

330

340

350

Tem

pera

ture

, C11.24

h, m

m



To C

onst

an

Ch

t K

@ 9

41h

mm/s

2 x 10mm/s

ange

tem

p 3

25C

-9

Close to ActivationEnergy for A600

&

8 x 10-9

to

10.

Figure 5-26

4 cc

/kg

H26

63h

Results of testing the same material as in Figure 5-23, but this time in the S-T orientation

2 @

[17, 18]

5-25


Generally similar CGRs overall were measured for a further heat of Alloy 690 that was also 1-D cold-rolled by ~20% at GE-GRC. This time, however, the S-L oriented specimen did respond to a drop in test temperature (see Figure 5-27).

SCC#3 - c394 - Alloy 690, 19%RA 1D, S-L Orientation, Heat NX2865HK

11.24900

300

11.26

11.3

11.32

11.34

11.36

11.38

11.4

11.42

1400 1900 2400 2900 3400 3900

Test Time, hours

Cra

ck le

ngth

, mm

310

320

330

340

350

360

370

Tem

pera

ture

, C

Est. pH at 360C = 8.12 used for φc

11.28

At 340C, pH = 7.53. At 300C, pH = 6.86


2.4 x 10-8

mm/s

To C

onst

ant

K @

941

h

4.6 x 10-9 mm/s

Close to ActivationEnergy for A600

tem

p to

325

C &

1 c

c/kg

H2

@ 2

663h

n

ture

Cha

nge

0.4

Figure 5-27 Results from a further heat of Alloy 690 tested after 1-D cold-rolling at GE-GRC, this time showing expected response to a drop in temperature even for an S-L oriented specimen [17, 18]

Figure 5-28 shows details of the fractography for the 26% 1-D cold rolled ANL plate that has produced the highest SCC growth rates measured to date for Alloy 690 in simulated primary waterhh. Comparing this picture with Figure 5-29 (from a 20% 1-D cold rolled Alloy 690 specimefrom forged bar that was originally very homogeneous and exhibited only moderate CGRs even after deliberate cold working), it would appear that there are clear signs for a difference in fracmorphology/crack path that may be an indication of a different fracture mechanism.

hh Even higher rates have, however, been measured at the Bettis laboratories in deaerated HT water – see

Section 5.1.2.1.7.

5-26


Figure 5-28 tion fractography (c372) of the 26% 1D-cold rolled ANL 690 plate tested at High-resolu

GE-GRC [17, 18]

Figure 5-29 High-resolution fractography of a 20% 1-D cold rolled Alloy 690 specimen from forged bar [17, 18]

5-27


5.1.2.1.4 Investigations in Simulated Primary Water at the University of Tohoku in Japan

These investigations are being carried out as part of a major research effort, co-sponsored by the EPRI Primary Systems Corrosion Research Program, and are oriented more to obtaining fundamental understanding of the SCC process than to developing engineering data [19]. Although some testing of Alloy 690 has been carried out in simulated primary water at 320 to 340C [8], much attention is also being paid to possible LTCP effects on this material at ~50°C [20] and many results

n.

lly

A 1 thick) 5% side-grooved specimen appears to have been used in the first CGR test (pla ess in the 90TT-B plate is high (~ 340 HV1) and relatively uniform. A 12.5mm thick circular double cantilevered beam (CDCB) specimen with 5% side groove on each sid sed with the notch direction parall th in ro ng directiwhen this ate ].

Testin

SCC b rm si d pri ry w in flo te, r hedstai eel autoclave. Gentle cycling has been used in order to facilitate transitioning from the air fati he SCC test periods appear to have been carried out with trapezoidal load an under purely constant loading.

Dat

Tes ly shown no or only exte intergranular cracks are 40% CW Allo r at 340°C for 4050 h [8]. m these findings and some of these cracks apparently propagated along a plane perpendicular to the main fracture surface.

have not yet been completely reported. The following represents relevant, preliminary informatio

Product form

Various heats of Alloy 690 are being studied, apparently including both mill-annealed and thermatreated (700°C for 15h) plate materials, one of which (plate B) was tested both for possibly enhanced CGRs in the HAZ after being welded to Alloy 52 material (see Section 5.1.2.4.5), as well as separately after undergoing 40% cold work (achieved by cross-rolling the plate eight times) [8].


T-CT (25mmte A), with the notch direction of the specimen parallel to the rolling direction. The hardnthrough-thickness direction of the 40% cold-worked Alloy 6

e was urial [8

el to e f al lli on testing latter m

g details

testing is eing perfo ed in mulate PWR ma ater a low w ra efres nless stgue pre-crack and some of ting (at an R-value of 0.7) involving 72h hold periods, rather th

a evaluation

ting of Alloy 690TT plate without additional cold work has apparentmely slowii crack growth through SCC. However, relatively long, partially thought to have been observed at separate locations on the fracture surface of they 690TT plate B specimen after testing in simulated PWR primary wate

No attempt appears to have been made to derive a CGR fro

ii A CGR of ~ 6E-10 mm/s at 320°C and a K-value of 28.5 MPa√m is indicated in [20].

5-28


5.1 tigations in Simulated Primary Water at the Institute of Nuclear Safety Systems (IN

These studies are being performed on behalf of the Kansai Electric Power Company (KEPCO) [8].

Product form

INSS are known to be testing thermally treated Alloy 690 plate with additional cold-work (20 and 50%) from rolling.


The program uses fatigue pre-cracked 1/2T CT specimens with T-L and T-S orientations. I.e., the crack advance plane is not coincident with the rolling direction, but oriented at 90° to this.

Testing details

The test environment is simulated primary water (with an unchanging H2 level of 30 cc/kg) over a range of temperatures from 360 down to at least 320°C. No attempt is apparently made to use mechanical load transitioning and testing is being carried out under constant load at a KI value of 30 MPa√m.

Data evaluation

Non ckets” of crack advance through intergranular PWSCC have been detected alon .

g CGRs is unclear, but reported values at 360°C are as high as 2 E-7 mm/s igure 5-31). Virtually no effect of temperature was seen for the 50% CW material, but CGRs for

.2.1.5 InvesSS) in Japan

-uniform “pog the crack front (Figure 5-30) and these apparently have significant dimensions at 360°C

The method of calculatin(Fthe 20% CW material apparently dropped to below 2 E-8 mm/s at ~330°C.

5-29


Figure 5-30 Example of crack growth observed at INSS in 20% cold worked 690TT at 360°C (from [8])

5.1.2.1.6 Further Investigations in Simulated Primary Water in Japan

CRIEPI is also known to be carrying out more fundamental studies of Alloy 690 [21], but further details (including the product form) are not yet known.

5-30


Figure 5-31 Apparent effect of temperature and degree of cold work (CW) on measured CGR for 690TT (from [8])

5.1.2.1.7 Investigations by Bechtel-Bettis in Deaerated High-Temperature Water

The nd Moshier [22] at thupdated this and included some key additional information. A main aim of the program is to identify potential areas of vulnerability for Alloy 690 by varying material heats and fabrication methods. Recently, Burke [24] gave a description of some of the microstructural characteristics of the materials that Bettis have used, paying particular attention to the development of so-called “banded” microstructures.

Product form

Bettis have tested Alloy 690 in plate form, representing three different heats from two distinctly different melting techniques:

• One heat (“VIM-1”) is vacuum induction melted with subsequent electroslag remelting (VIM/ESR), which gives the “cleaner” microstructure, and

• Two heats (“AOD-1” and “AOD-2”) are argon oxygen decarburization with subsequent electroslag remelting (AOD/ESR).

Final annealing temperature (1052 vs. 1093°C) is a test parameter for the VIM/ESR plate, as is the effect of omitting the final thermal stabilization heat treatment (TT for 10h at 718°C) in one case for the AOD/ESR plate.

Detailed material characterization [24], carried out on an entirely different heat of Alloy 690 material that was not used for CGR testing, showed the following general features:

large Bettis testing program on Alloy 690 was first presented by Paraventi ae Atlanta meeting of the EPRI PWSCC Expert Panel in Fall 2006. A 2007 presentation [23]

5-31


• Absence of Cr microsegregation in the as-cast structure implies that the development of carbide banding is related to subsequent thermomechanical processing treatments.

• Borides, carbides and boro-carbides are present in as-cast and as-homogenized A690.

• Thermomechanical processing should be designed to promote the dissolution of these coarse precipitates and optimize the controlled precipitation of carbides and nitrides to promote grain refinement in the final product.

A key aspect of the Bettis program is to investigate the effect of various levels of cold work. This is being carried out primarily by uni-directional rolling and was performed in one pass (for a target of 12% work) or two passes (for the target of 24% cold work) for heats VIM-1 and AOD-1. Heat AOD-2 was rolled in multiple passes. Table 5-10 shows the full test matrix and it can be seen that uni-directional tensile prestraining of the VIM material has also been used in 2 cases to obtain the highest levels of deformation (32 – 33% cold work). Material characterization [24] showed the following:

• Development of extensive slip band structure with a high proportion of dislocations.

• Fracture of intergranular and intragranular carbides (more fractured carbides as % deformation increases).

• “Bending” (i.e., curvature) of slip bands occurs to accommodate shape change.

• hich has a lower stacking fault energy (SFE), shows finer slip bands and more

Refregions near the plate surfaces (edges) where large grains were occasionally observed, but generally throughout the thickness the material was uniform in grain size and free of carbide banding. The AOD heats had different levels of banding, with AOD-2 being the most significantly banded both in terms of grain size and carbides. AOD-1 was much less banded than AOD-2 but still showed moderate grain size and carbide banding. In summary, the relative ranking of banding was characterized as follows:

• VIM-1: Very little to no banding.

• AOD-1: Intermediate banding (more grain size than carbides).

• AOD-2: Severe grain size and carbide banding, with a high concentration of inclusions.

Alloy 690, wplanar slip than Alloy 600.

erence [22] reported representative microstructures for each of the three heats. VIM-1 had some

5-32


Table 5-10 Material test matrix for Bettis Alloy 690 test program [22, 23]


Three orientations for crack growth in the precracked CT specimensjj were used (cf. Figure 5-9):

• L-T: this represents growth perpendicular to the rolling/straining direction on a plane also perpendicular to that direction and showed the lowest susceptibility to SCC.

• S-T: this represents growth in the plane of, but perpendicular to, the rolling/straining direction and produced intermediate susceptibility to SCC.

• S-L: this represents growth in the plane of, and parallel to, the rolling/straining direction

Tes

A to lacement) specimens were tested for times between 500 and nearly 15,000 h at 3 test temperatures (316, 338 and 360°C). Some load relaxation has to

d I or each specimen was calculated by dividing the average crack advance on the fracture

otal test time, i.e. ignoring any incubation time. It was then plotted with regard to an applied stress intensity (which also involves the assumption of uniform growth

ular)

and produced the highest susceptibility to SCC.

ting details

tal of 186 bolt-loaded (i.e. constant disp

be expecte here, especially during heat-up of the autoclave. Furthermore, crack advance in such specimens leads to a decrease in both applied load and resulting crack tip stress intensity value (K ).The CGR fsurface by the taverage value ofrate). As a result of the above issues, such test procedures usually tend to underestimate SCC susceptibility, both in terms of difficulties nucleating uniform crack fronts from the (transgranair fatigue pre-crack and in maintaining representative growth rates during crack advance. jj The specimens were without side grooves, so that some tendency for the crack growth plane to “wander” during

testing is to be expected.

5-33


Seven instrumented specimens were also tested under constant, active load, apparently to check the abovementioned influences [22]. None of these showed any evid

out ence of incubation times.The

e resd specimens

r he constant acement technique used in the majority of the Bettis tests might tend to underestimate

I ls without ve led to bias in some of the dependencies reported below.

Data evaluation

ng

• Only the more susceptible material (VIM/ESR plate) was tested in the fully annealed condition, and then showed no susceptibility to crack advance through SCC.

• Uni-directional cold rolling induced some SCC susceptibility in both materials, even at the 12% level, but much more strongly at the 24% level (see Figure 5-32) and in the VIM/ESR plate.

• Comparison of the highest cold worked materials seems to suggest that AOD heats may be somewhat more resistant; however, this may just reflect heat to heat variability.

• Uni-directional tensile straining (to around 33%) also resulted in moderate susceptibility of the VIM/ESR plate, but was not as detrimental as 24% cold-rolling (see Figure 5-33), even though it resulted in a higher yield strength (cf. Table 5-10).

• There was no apparent effect of a lower or higher final annealing temperature (before additional TT treatment) on behavior of the VIM/ESR plate.

• Results on omission of the additional TT treatment with one lot of the AOD/ESR plate do not appear to have been reported yet.

on ult shown in reference [23] indicates that cracking started quickly (at least for the medium value of stress intensity used) and it is stated that the CGR data from the actively loadewe e “consistent with rates from sister bolt-load tests”. Nevertheless, it is possible that tdisplrepresentative CGRs when these are low (e.g. at low K values, or for annealed materiadeliberate cold work) and may thus ha

Table 5-11 shows the qualitative evaluation of SCC susceptibility prepared by Bettis. The followikey points emerge with regard to material fabrication issues:

Table 5-11 Bettis summary of relative SCC susceptibility for Alloy 690 [22, 23]

5-34


Figure 5-32 Increase in CGRs with increasing level of uni-directional cold-rolling in Bettis studies [22, 23]

Figure 5-33 Cold-rolling is more detrimental than tensile pre-straining (despite lower yield strength) [22, 23]

The above findings clearly imply that SCC susceptibility is associated with the creation of heterogeneous microstructures in the test materials, in which case strong effects of specimen orientation on CGRs would be expected. These were indeed found:

• The S-L orientation, with crack advance in the direction and plane of deformation, was much more susceptible even than the S-T direction (see Figure 5-33).

• The L-T orientation was very resistant to cracking (cf. Table 5-11).

5-35


For the most susceptible materials and worst CGR orientation, only a low dependence on stress inte t-hand dia report KI exponents here in the range of 0.7 to 1.9, which is similar to the known dependency for Alloy 600. Higher values (up to an exponent of 6.2) were quoted for lower susceptibility conditions, but it is likely that these have been biased strongly by the difficulty in initiating uniform crack advance at low KI values, as discussed earlier.

The Bettis test program has shown a surprisingly low dependence on temperature for crack advance in cold-worked Alloy 690 material, with apparent activation energies as low as 5 kcal/mole being derived in one case of high SCC susceptibility (see Figure 5-34, right-hand side). This is very different from the values of around 35 kcal/mole usually reported for Alloy 600 and implies that, if cracking were to occur in service with cold-worked Alloy 690, it would not necessarily be limited to components with high operating temperatures.

Another apparent difference relates to the effect of increasing dissolved hydrogen in the test medium, where the trend, in fact, appears opposite to that observed for Alloy 600 cracking (see Figure 5-35).

nsity was found over the tested range of approx. 13 to 42 MPa√m (see top curve in the lefgram in Figure 5-33). Bettis

Figure 5-34 Effect of test temperature and degree of cold work on CGRs for the VIM/ESR plate [22, 23]

The biggest concern arising from the Bettis data is that very high CGRs (~ 6E-7 mm/s) have beregistered for the most susceptible material at quite moderate stress intensities (~ 30 MPa√m), as shown by the blue curve on

en

the left-hand side of Figure 5-33. This is all the more surprising given the frequent tendency of constant displacement test techniques to under- rather than overestimate CGRs, as discussed earlier. The fractography shown in reference [23] confirms that these rates have

lower levels of cold work) and the morphologies shown have been predominantly intergranular (see Figure 5-36).

vide any direct explanation of the high SCC CGRs that were measured.

been derived from appreciable increments of crack advance (even for the

In some cases, colonies of cracked carbides have apparently been detected on grain facets of the IGSCC fracture surface (see Figure 5-37), but this finding – although somewhat unusual – does not in itself pro

5-36


Figure 5-35Appa t increas t 50 cc/kg c/kg (pink symbols) fo 69 ) an mp n ( t) p e b vior llo [22, 23]

e in CGren Rs a hydrogen (blue symbols) vs. 23 c

r Alloy 0 (left d co ariso righ with o posit eha for A y 600

Figure 5-36 Predominantly intergranular crack advance in an S-L-oriented specimen of VIM/ESR TT plate subjected to only 12% cold-rolling [22, 23]

5-37


Figure 5-37 Detail from fractography of an S-T-oriented specimen of VIM/ESR plate subjected to 24% cold-rolling [22, 23]

Bettis themselves have indicated two possible approaches to assessing the relevance of their SCCresults to date. The first makes use of a direct comparison with similar data they have obtained in a parallel test program using Alloy 600 material [22]. As indicated in Figure 5-38, this leads to

usion that the Alloy 690 CGRs are a factor of 5 to 10 lower than for Alloy 600 material f comparable yield strength. This margin is considerably smaller than industry expectations.

However, given the relatively low yield strength (c k) dependence of Alloy 600 (σys

3) compared to the apparently very high cold work d nce in Alloy 690(σys

7-8), the margin betw oys would be expected to get larger as the extent of cold work decreases.

the conclo

old worepende

een the two all

Figure 5-38 Comparison of Alloy 600 and 690 CGRs measured by Bettis [22, 23]

5-38


A second approach assumes that the major role of cold-rolling is simply to increase the yield strength of the 690 material and attempts to extrapolate back (see Figure 5-39) to determine a representative CGR for the annealed VIM/ESR plate (yield strength = ~46 ksi), which did not show any crack advance at all in the actual testing. The resulting rate shown would be acceptably low, but any such extrapolation (on a logarithmic-linear plot) is fraught with uncertainty and the correctness of that chosen in Figure 5-39 can be challenged in the absence of additional data at other cold work levels. Furthermore, as pointed out above, a unique dependency on yield strength is not indicated in the complete test data, because of the differences found when cold work was introduced in a different way (rolling versus tensile pre-straining).

Ultimately, a proper assessment of the practical relevance of the Bettis data will require more understanding of the heterogeneity produced in Alloy 690 by uni-directional deformation and more confirmation of some of the key dependencies.

Figure 5-39 Attempted back-extrapolation by Bettis of cold-worked Alloy 690 CGRs on the basis of assumed yield strength dependencies [22, 23]

5.1.2.1.8 Investigations by Westinghouse in Supercritical Water with Additions of Boron, Lithium and Hydrogen

A Westinghouse program, carried out for the EPRI MRP, was designed to accelerate the IGSCC process in SCC-resistant, Ni-base alloys by testing at elevated temperatur s and pressures [25]. Specifically, this study involved testing at temperatures beyond the critical point of pressurized aqu ments in an attempt to thermally accelerate corrosion and produce measurable effe(385°C/33MPa) were chosen to have “water-like” fluid properties with additions of Li/B and H2.

est materials were selected to be representative of thick-walled Alloy 690 (see Table 5-12) and a limited number of Alloy 600 control samples were also tested. Table 5-13 shows the chemical composition of the materials and Table 5-14 shows their tensile properties as supplied. It can be

e

eous environcts from relatively minor SCC processes. High-pressure/temperature test conditions

T

5-39


seen that the yield strengths are rather low (between ~ 230 and 278 Mpa). Since this would have limited the stress intensity factors that can be achieved without violating ASTM testing requirements and in view of the anticipated difficulty in achieving and sustaining meaningful and measurable CGRs in these crack resistant materials, which concern implies the need for high stress intensities, the decision was taken to elevate the yield strengths by cold-working. The means chosen to effect the required cold work was high speed forging and Table 5-15 shows the properties achieved during trial efforts. It was concluded that engineering strain (in terms of a thickness reduction) of 10-11% would produce the desired strain hardening.

Table 5-12 Alloy 690 materials used in Westinghouse testing for CGRs in supercritical water [25]

Table 5-13 Chemical composition of Alloy 690 materials used in Westinghouse testing for CGRs in supercritical water [25]

5-40


Table 5-14 Tensile properties (as reported by material vendors) of the Alloy 690 materials used in Westinghouse testing for CGRs in supercritical water [25]

Table 5-15 Tensile properties after trial forging of the Alloy 690 materials used in Westinghouse testing for CGRs in supercritical water [25]

The microstructure of all the alloys was examined in some detail (see [25]) and all the 690 CRDM materials exhibited predominantly intergranular carbides. However, the microstructure of the one plate material used (which had been given a laboratory thermal treatment) was slightly different: although most of the M23C6 carbide precipitation seemed to be intergranular,”stringers” of what are most likely carbonitrides were aligned along what appears to be the rolling direction. This is not an uncommon feature for heavy section hot rolled plate and the role, or lack thereof, of these stringers in the crack propagation process was deemed to be of interest.

Two crack growth rate (CGR) test series were conducted in SCW containing 900 ppm boron and 3.15 ppm lithium using (overall) 2 compact tension (CT) specimens fabricated from Alloys 600, 1 CT specimen fabricated from 690 TT plate (with additional cold-work), 2 CT specimens fabricated from CRDM piping in the as-received condition and 4 CT specimens also from CRDM material, but with deliberate cold working. The hydrogen partial pressure in the test was chosen to be reducing compared to the Ni/NiO stability line and the boron and lithium conditions were close to median concentrations for a U.S. PWR. Stress corrosion crack propagation occurred in every specimen tested in the supercritical environment. The results obtained for the mill annealed Alloy 600 specimens (see Table 5-16) indicated higher CGRs in SCW than are observed in pressurized

5-41


primary water; however, the acceleration factor was lower than originally anticipated, possibly dto the chosen hydrogen level. The general features of crack morphology and oxide chemistry wereessentially the same as those observed at lower temperatures and pressures.

Clear intergranular crack growth was observed in b

ue

oth the as-received (690 TT) and the ~ 10% cold-worked Alloy 690 specimens tested at stress intensities of approximately 32 MPa√m

hed from the transgranular fatigue that occurred in the pre-cracking phase, or during gentle cycling in the autoclave environment. No meaningful differences were observed between

been

.5.

Table 5-16 Detailed Results for the Alloy 600 control samples and Alloy 690 plate material used in Westinghouse testing for CGRs in supercritical water [25]

Calc.1 mm

Meas.2

mm Meas.3

mm Mode4 Time5

Hour Time6

Hour

phase 2

MPa√m average

MPa√m

final

9 MPa√

m

AverageCGR10 m/s

(Figure 5-40). The slow, but steady, IGSCC measured for both material conditions could be clearly distinguis

the cold-worked plate material and the three different heats of thick-walled Alloy 690TT (thought to represent a range of thermo-mechanical processing typical for current CRDM nozzles inreplacement RPV heads) with subsequent deliberate cold working, so although the results have tabulated separately, they are shown together in Figure 5-41. Note from this figure, however, that lower CGRs were seen in the as-received Alloy 690TT, as discussed further in Section 5.1.2.2

Material Spec ID

da Fatigue

da Fatigue

da SCC Crack Total SCC K 7 K 8 K

600 MA 510-04a 0.037 n/m 1.690 IG 1920 1546 16.9 19.2 20.7 3.0E-10

600 MA 510-04b 0.037 n/m 1.540 IG 1920 1546 16.7 19.3 20.7 2.8E-10

690TT cw F8244-cw3 0.313 0.612 0.359 IG 4806 4323 31.9 31.3 32.0 2.3E-11

Notes:

1. Calculated fatigue crack growth

2. Average transgranular fatigue crack increment averaged across the crack front

3. Average SCC increment averaged across the crack front based on optical and SEM images

4. IG = intergranular; ID = interdendritic; TG = transgranular

5. Total time exposed to 385C supercritical water under applied load and pressure

6. SCC time is time under conditions with steady loading (e.g. unloading period ≥ 3600 s)

7. K phase 2 is the stress intensity after initial rapid cycling complete, (e.g. unloading period ≥ 3600 s)

8. Kaverage based on mid-crack a/w ratio and load conditions in the middle of the test (e.g. during Phase 3)

9. Kfinal based on final a/w ratio and load conditions in the end of the test (e.g. end of Phase 4)

10. Average CGR = da SCC based on destructive exam divided by SCC time

5-42


Figure 5-40 Example of transition from transgranular cracking during cyclic loading in SCW to intergranular SCC at ~ constant load in a 10% cold-worked specimen of Alloy 690TT [25]

Alloy 600/690 Tested in 385°C Supercritical B/Li/H2

1.0E-09

m/s

)

1.0E-12

1.0E-11SCC

G

1.0E-10

owth

Ra

1.0E-08

15 20 25 30 35 40 45Stress Intensity Factor (MPa√m)

rte

(

Alloy

+ coldwork

Alloy 600MA

690TT

K Phase 2

Average K

Final K

Alloy 690TT

Figure 5-41 Average SCC growth rate vs. stress intensity factor for Alloy 600 and 690 materials [25]

5-43


The d Alloy 690 TT CGRs are approximately 20 times lower than those measured for stress inte tor here at 385°C in SCW is ~20. As mentioned earlier, the thermal acceleration factor for Alloy 600 was lower than originally anticipated. Th pr no establi d b rap la e meas e CGdown ri p s r ke le n h e 600 in primary water and the lowe ted in Figure 5-42 should be regarded only as one possib s of l be ior

cold-workeAlloy 600 MA in the supercritical environment, after adjusting the latterkk to a commonnsity factor of 32 MPa√m (see Figure 5-42). I.e., the material improvement fac

ere is erature

esently in order dot

she detai

asis for extd com

o ting s with t

the MRP

ur ddatabas

Rs for to subc tical tem to ma

line pariso

le repre entation actua hav .

Alloy 600/6 m on up riti s s d W r

1.0E-12

1.0E-09

1.0E-

m

90 Co paris s - S erc cal v . Pres urize ate

08

1.0E-11

1.0E-10

CG

R (m

/s) f

or K

= 3

2 M

1.40 1.45 1.50 1.55 1.60 1.65 1.70 1.75 1.80

(1000/T)

Pa√

93510 - MRP55

93511 - MRP55 & W data

93510 - Supercritical B/Li/H2

93510 - Average Primary H2O

93510 - Anticipated SCrit

A690 10%CW in Supercritical B/Li

A690 0%CW in Supercritical B/Li

Q = 130 KJ/moleQ = 80 KJ/mole

Supercritical

Kavg = 37 MPa√m

3 tests

2 Testing of Alloy 690 CRDM Material without Deliberate Cold Working

ection on,

e on the part of both the investigator and the program sponsor.

Figure 5-42 Comparison of CGR measured for Alloy 600 MA and Alloy 690 TT in supercritical and primary water [25]

5.1.2.

The majority of the thick-walled Alloy 690 material currently in service in PWRs worldwide is in the form of CRDM penetrations that are not thought to have significant cold work during fabrication, although some uncertainty remains regarding possible pipe straightening (see S2.1.2). Nevertheless, relatively few laboratory tests to date have used material in this conditisince it is expected to show extremely high resistance to PWSCC. This makes reliable testing a challenge requiring extreme patienc

kk Stress intensity correction factor for Alloy 600 data = (32/K)2.2

5-44


5.1.2.2.1 Investigations in Simulated Primary Water at Argonne National Laboratory (ANL)

The testing approach used at ANL has already been described in Section 5.1.2.1.2. Two 0.5T CT specimens with different orientations of through-thickness crack growth have been prepared from Alloy 690TT material (heat WP142) as shown in Figure 5-43. In the first test at 320°C, a CGR in the radial direction of ~ 5E-9 mm/s was reported at constant load early in the test (see Figure 5-44, top diagram). Subsequently, even slower crack growth (~ 2E-9 mm/s) was reported to have occurred during a relatively short period (~ 800 h) under constant load at the end of the experiment, but no clear signs of intergranular cracking were subsequently seen on the fracture surface of this specimen (Figure 5-45).

Figure 5-43 ]

prov C, for the U.S. Department Specimen orientation in ANL testing of Alloy 690 CRDM material [11 to 15Graphic ided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLof Energy under Contract No. DE-AC02-06CH11357.

5-45


Figure 5-44 Reported CGRs during brief test periods at constant load in ANL work on CRDM material [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Depaof Energy under Contract No. DE-AC02-06CH11357.

rtment

5-46


Figure 5-45 Absence of clear intergranular cracking on fracture surface of the first CRDM specimen

.S. Department

tested at ANL [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the Uof Energy under Contract No. DE-AC02-06CH11357.

Testing of a second specimen (with different orientation) gave ~ 3E-9 mm/s under constant load at320°C, but slightly more (~ 8E-9 mm/s) for ~ 500 h at a temperature of 350°C (see Figure 5-46).

5-47


12.70

12.71

12.72

12.73

12.74

12.75

2.76

2.77

12.78

40

45

50

5 )

1

1

10

15

20

25

30

35

1200 1400 1600 1800 2000 2200 2400

Cra

ck L

engt

h (m

m)

Km

ax (M

Pa m

0.

Time (h)

Kmax

Crack Length

Alloy 690Specimen # C690-LR-2Simulated PWR Water

3.2 x 10–12 m/s21.3 MPa m0.5

Constant load

8.0 x 10–12 m/s°C

21.3 MPa m0.5

Constant load

Incr

ease

d te

mpe

ratu

re to

350

Figure 5-46 Reported CGRs during constant-load test periods at two test temperatures in further ANL work on CRDM material without deliberate cold working [11 to 15] Graphic provided by Argonne National Laboratory, managed and operated by UChicago Argonne, LLC, for the U.S. Department of Energy under Contract No. DE-AC02-06CH11357.

5.1.2.2.2 Investigations in Simulated Primary Water at Pacific Northwest National Laboratory (PNNL)

Product form

Six heats of Alloy 690 CRDM thick-walled piping material were obtained from Valinox, all received in the thermally treated (TT) condition. A variety of thermomechanical treatments are being tested to explore what processing and microstructural conditions may influence SCC resistance in PWR primary water [26]. The first Alloy 690 tests compared the response of as-received TT material to a solution anneal (SA) heat treatment at 1100°C (chosen to alter the carbide microstructure) using the following three heats of material:

• Heat # RE243: 10.4Fe, 28.9Cr, 0.02C, 0.31Mn, 0.35Si, 0.14Al, 0.23Ti, Bal Ni

• Heat # WP142: 10.5Fe, 29.0Cr, 0.02C, 0.31Mn, 0.35Si, 0.18Al, 0.27Ti, Bal Ni

• Heat # WP140: 10.4Fe, 29.0Cr, 0.03C, 0.31Mn, 0.33Si, 0.18Al, 0.30Ti, Bal Ni


0.5T CT specimens with 5% side grooves are being used for all evaluations. Alloy 690TT and SA specimens were cut from the 690 piping material in the C-L orientation as shown in Figure 5-47. Since the grain size near the OD was only about half that observed at the midwall or ID, all specimens have been taken from one or other of the latter locations.

5-48


Figure 5-47 Orientations of CT specimens used by PNNL for CGR testing on two Alloy 690 forgings:

Tes

H2 l

CR inal

Data n

Onl h occ f intergranular morphology on the fracture surface (see Figure 5-51), have been

wh(i.eof g of no engineering significance.

left is heat RE243; right is heat WP140 [26]

ting details

The test environment is simulated primary water, predominately at 325 or 350°C. The chosen evel of 29°cc/kg is in the region of Ni metal stability, well above the expected level for peak

CGR in Alloy 600 at 325°C. Evaluations have also been performed on two of the Alloy 690 DM heats at 20 cc/kg and 350°C, which is near the expected peak CGR for Alloy 600. F

pre-cracking is carried out by fatigue at temperature in the autoclave followed by a very gradual transitionll in periodic partial unloading (see Figure 5-48) to a constant KI value of 30 or 40 MPa√m.

evaluatio

y very minimal amounts of crack growth at constant K (see Figure 5-49 and Figure 5-50), witasional signs o

seen in any of the specimens tested to date at PNNL without deliberate cold working. It is unclear ether genuine CGRs can, in fact, be derived here, but all would lie at or below 1E-9 mm/s . below the typical diameter of a single grain in the material being tested even after one year rowth) and are thus clearly

ll Modeled after the GE-GRC approach described in Section 5.1.2.1.1.

5-49


CT014 & CT015 CGR, 0.5TCT Alloy 690 Valinox, Heat RE243, pipe 2360, sample 1 & 2325°C, 30 MPa¦m, 1000 ppm B, 2.0 ppm Li, 29 cc/kg H 2

0.050

0.055

0.060

0.065

0.070

0.075

0.080

0.085

0.090

0.095

0.100

0.105

1000 1500 2000 2500 3000 3500 4000 4500 5000 5500 6000

time (hrs)

-10

0

10

20

30

40

50

60

ou

tlet

con

du

ctiv

ity (

µS

/cm

) o

rEC

P (

V1

0)

Pt ECP CT ECP

Alloy 690CM

Alloy 690TT

0.0

01

Hz, R

= 0

.7

2.1e-09 mm/s

2.6e-09 mm/s

0.001 Hz + 24 h,R = 0.7

~7e-10 mm/s

~6e-10 mm/s

corrected for actual crack length

0.001 Hz + 2.5 h,R = 0.7

constant K

5 month transitioning using cycle + hold steps

Figure 5-48 PNNL data showing the patience used in transitioning to constant K loading [26]

CT014 & CT015 CGR, 0.5TCT Alloy 690 Valinox, Heat RE243, pipe 2360, sample 1 & 2325°C, 30 MPa¦m, 1000 ppm B, 2.0 ppm Li, 29 cc/kg H 2

0.084

0.086

0.088

0.090

0.092

0.094

0.096

0.098

0.100

0.102

0.104

5500 5700 5900 6100 6300 6500

time (hrs)

-10

0

10

20

30

40

ou

tlet

con

du

ctiv

ity (

µS

/cm

) or

EC

P (

V1

0)

Pt ECP CT ECP

Alloy 690SA

Alloy 690TT

constant K

~3e-10 mm/s

~2e-10 mm/s

corrected for actual crack length

2 µm

Extremely slow (~1 µm/month) of "stable" propagation

Figure 5-49 PNNL data showing Alloy 690 CGR response for as-received TT versus a carbide-modified SA condition [26]

5-50


CT026 & CT027 CGR, 0.5T CT Alloy 690 Valinox, WP140, Pipe 2502 & WP142, Pipe2541

350°C, 30 MPa¦m, 1000 ppm B, 2.0 ppm Li, 29 cc/kg H 2

0.447

0.449

3800 3900 4000 4100 4200 4300 4400 4500 4600

time (h

-10

-6

0.451

0.455

0.459

0.461

0.465

0.467

rs)

-2

6

10

14

18

22

26

30

ou

tlet

con

du

ctiv

ity (

µS

/cm

) o

r

0.453 2

0.457

0.463

EC

P (

V1

0)

Alloy 690TT WP140

Pt ECPCT ECP

Alloy 690TT WP142

constant K

~5e-10 mm/s

0.0

01

Hz +

9 k

s,R

=0

.5

1.9e-08mm/s

~7e-10 mm/s~0 mm/s

~1e-9 mm/s

1.9e-08mm/s

2 µm

Figure 5-50 PNNL data showing Alloy 690 CGR response for two further heats of material in the as-received TT condition [26]

Figure 5-51 PNNL fractography showing IG cracking limited to isolated grains in specimens without

cold-work [26]

5-51


5.1.2.2.3 Investigations in Simulated Primary Water at Studsvik in Sweden

Two 0.5CT specimens of Alloy 690 taken from archive material for the CRDM penetrations Ringhals 2 replacment RPV head were tested at 320°C [27]. They were subjected to gentle cyclicloading with increased hold times (K

in the

es)

actual plant components and will include material with additional deformation simulating actual thick ection plant components.

Investigations in Simulated Primary water at MHI in Japan

CGR testing of Alloy 690 CRDM material is also ongoing at the MHI laboratories in Japan in collaboration with the Japan Nuclear Energy Safety Organization (JNES) and, in September 2006, had failed to produce any identifiable SCC after 4000 h in Alloy 690 [28]. Further details are not yet available.


The key features of this program were described already in Section 5.1.2.1.8 and the numerical results for the CT specimens of Alloy 690 in the as-received condition are tabulated in Table 5-17. Less cracking, and about 3 times lower growth rates, were observed in comparison with the results for the cold-worked Alloy 690 TT specimens (see Figure 5-41).

Table 5-17 Detailed results for the Alloy 600 control samples and “as-received” Alloy 690 CRDM materials used in Westinghouse testing for CGRs in supercritical water [25]

Material Spec ID

da Fatigue Calc.1 mm

da FatigueMeas.2

mm

da SCC

Meas.3

mm

CrackMode4

TotalTime5

Hour

SCCTime6

Hour

Kphase 2

7

MPa√m Kaverage

8 MPa√m

Kfinal

9 MPa√m

AverageCGR10 m/s

max = 25 MPa√m, R = 0.6, 0.05 Hz, 0 and 1000 s hold timfor about 400 and 800 hours, respectively, without any indications of intergranular cracking beingobserved. However, transgranular cracking, interpreted as corrosion fatigue, was observed on the fracture surfaces. Studies are continuing to determine FOI values relevant to

s

5.1.2.2.4

600 MA 510-04a 0.037 n/m 1.690 IG 1920 1546 16.9 19.2 20.7 3.0E-10

600 MA 510-04b 0.037 n/m 1.540 IG 1920 1546 16.7 19.3 20.7 2.8E-10

690TT 415-a 0.310 0.757 0.234 IG 4806 4323 32.1 31.3 31.8 1.5E-11

690TT 7a12-a 0.370 0.675 0.098 TG/IG 4806 4323 32.2 31.1 31.3 6.3E-12

Notes: see Table 5-16

The as-received Alloy 690 TT CGRs are approximately 60 times lower than those measured for Alloy 600 MA in the supercritical environment, after adjusting the lattermm to a common stress intensity factor of 32 MPa√m (see Figure 5-42). I.e., the material improvement factor here at 385°C in SCW is ~60. As mentioned earlier, the thermal acceleration factor for Alloy 600 was lower than

mm Stress intensity correction factor for Alloy 600 data = (32/K)2.2

5-52


originally an p h is p CGRs dow o sub ic p s o ase for 600 prima w e pos e represen ed Alloy 690 TT in SCW are sufficiently slow that significant SCC propagation in thick-walled Alloy 690 components exposed to pr ary water uld no expect , even i CC iniwer occur.

5.1.2.3 ing of Alloy 690 CRDM Material after Deliberate Cold Working

Fro e b ing, GE-GRC chose to use the introduction of deliberate cold work for the EPRI MRP Test P am in order both to r he chances of measuring any crack growth at all in a very SCC-resistant material and to simu he al s ins that might be ected i he HAZ of wel Alloy 0 immediately adjacent to the fusion line (see Section 5.1.2.4). Following the publication in 2006 of the Be es on te m terial (see Section 5.1.2.1.7), many other labo ories have follo d su d h a o ected towards the way in which subsequent cold working of various types might affect the CGR behavior of otherwise very resistant CRDM material.

5.1.2.3.1 Investigatio in S la ric Global Research (GE C)

The sic te ng app h at G ibed in Sect the feasibility tests on EPRI plate material and e or studies on additional heats of material that are re y rele nt to actual components. Some furth ing oy 690, however, used an hive DM en tio upplied Duke Power (Heat WN415) as the starting material (see Figure 5-52). Deliberate levels of cold work (20 or 41%) were then introduced by u orm forging of tangular blo s p to nufact ens 9]. Nothow r, that e material was not thick enough to align the crack growth plane with the plane of forging deformation.

For two specimens (tested together) with 20% homogeneous cold work, relatively little time (onl 800h as spen t con nt K see and growth r s were then very low (≤ 5 mm/ Nevert ess, c r enc as und of i rgranul crackin Figure 5-54), much of which must have taken place during earlier periodic partial unloading conditions (PP .e. w element of trapezoidal, cyclic l ding). T was al tru furth ecimensubjected to 41% uniform cold work and tested u er vario loading nd or ,000 in t s cases, n n to a st a y when movin a e of 24h durinPPU, or growth rates a dy remained bel the oveme d 5E-9 mm sheng g ificance (see Figure 5-55). Interestingly, the fracture surface of is speci n revealed “fi s” of c k gro h with mu ou f-plane ondary cracking (see Figure 5-56).

ticicritry

ated. Tal temater andtation o

ere is perature the lowf actual b

resentl in order dotehavi

y no eer to mted linor. Ne

stabliake d

e in Fverthe

shed basetailed cigure 5-42less, the C

for extramparisons should beGRs mea

olating th with the M regarded sured in th

e measuredRP datab

only as one as-receiv

n t in sibl

e to

m th

ded

rat

-GR

ba

not necessarily di

nifeve

y ~E-9

U, i

otal. In

ineerin

im wo t be ed f S tiation

Test

eginnrogr aise t

late t loc tra exp n t 69

ttis rit an

ults muc

plattenti

an is currently dirwe

ns imu ted Primary Water by General Elect

sti roac E-GRC was already descr ion 5.1.2.1.1 forin S

va CR

ction 5.1.2.

p

ck

1.3 f

n s

ma

ctlarc

rec

er test of extruded All

[2

etra

rior

by

ure of CT specim e, th

) ws).

t ahel

stalea

( evid

Figure w

e 5-53) fo

atear nte g (

ith an

ome

oandlread ab

hisus

ntione

so co

g to

e of aitionhold

/s t

er spover

hold of no

h g

s f timhre

11 crackilrea

g te ded rreow

signnger

th merac wt ch t-o sec

5-53

PWSCC of Thick-Wa

5-54

lled Alloy 690 Material

FiSp

gure -52 eci ation in GE sting of All 690 CRDM material [1 8]

5men loc te oy 7, 1


FigDattes

ure 5-53 a fro ne of tw ecimens of CRDM material with 20% homogeneous cold work ted a RC [2

m ot GE-G

o sp9]

FigureInterg

5ra crack morphology in 20% cold-worked Alloy 690 CRDM specimen [29]

-54 nular

5-55


Figure 5-55 Data from a further specimen of CRDM material with 41% homogeneous cold work showing tendency to crack arrest before reaching constant K conditions (top) and very

old time (bottom) [29]

low CGRs even under periodic partial unloading with a 24h h

5-56


Figure 5-56 Macro – and micrfractography from the specimen of CRDM material with 41% homogeneous cold work showing extensive out-of-plane secondary cracking [29]

5-57


Thethis r of improvement of at least 70x in comparison to as-received Alloy 600 and probably more like 400x if the Alloy 600 material had been subjected to similar levels of cold work.

5.1(PN

ribed

very low CGRs measured even after introducing heavy, but homogeneous, cold work into actual CRDM material can be considered to indicate a facto

.2.3.2 Investigations in Simulated Primary Water at Pacific Northwest National Laboratory NL)

The approach to CGR testing at PNNL is similar to that at GE-GRC and has already been descin Section 5.1.2.2.2. They have also tested four specimens (2 each in the TT and CMnn conditions)of Alloy 690 CRDM material with deliberate, non-uniform cold working [26]. In one case, bar material was 17% unidirectionally cold rolled and tested in the S-L orientation. These specimens exhibited no enhanced SCC susceptibility. CGRs were already low under PPU and virtually arrested at constant K, even after raising this value from 30 to 40 MPa√m (Figure 5-57).

CT019 & CT020 CGR, 0.5T CT Alloy 690 Valinox, Heat RE243, Pipe 2360, 17% CW S-L325°C, 30 MPa¦m, 1000 ppm B, 2.0 ppm Li, 29 cc/kg H 2

0.006

0.007 25

0.000

0.001

0.002

0.003

0.004

0.005

1775 1875 1975 2075 2175 2275 2375 2475-10

-5

0

5

10

15

ou

tlet

con

du

ctiv

ity (

µS

/cm

)

time (hrs)

20

Alloy 690SA+17% CW S-L

Alloy 690TT+17%CW S-L

5.1e-09mm/s

0.001 Hz + 9 ksR = 0.7

constant K1 µm

3.4e-09mm/s

~4e-10 mm/s

~2e-10 mm/s

Figure 5-57 Data from PNNL testing of CRDM material with deliberate cold working (here 17% in the S-L orientation) [26]

nn CM = carbide modified (by solution annealing).

5-58


In the second case, these same materials were 30% unidirectional cold-rolled and tested in the T-L orientation. This did have had some effect on SCC susceptibility, particularly for the TT specimenat K values of 30 and 40 MPa√m (Figure 5-58). However, even here the measured CGRs were below an engineering significance threshold of 5E-9 mm/s. Only isolated grains of cracking were

seen on the fracture surfaces of all but the TT+30%CW specimen, where more intergranular cracking was reported, consistent with the slightly higher observed CGR.

CT022 & CT023 CGR, 0.5T CT Alloy 690 Valinox, Heat RE243, Pipe 2216, 30% CW T-L350°C, 40 MPa¦m, 1000 ppm B, 2.0 ppm Li, 29 cc/kg H 2

0.010

0.012

0.014

0.016

0.018

0.020

20

30

tivit

y (

µS

/cm

) o

r

0.0

01

Hz +

9 k

s,R

= 0

.5

constant K

2.9e-09mm/s

0.000

0.002

0.004

0.006

0.008

4200 4300 4400 4500 4600 4700 4800 4900 5000 5100 5200-10

0

ou

tlet

c

time (hrs)

10

on

du

c

EC

P (

V1

0)

Aloy 690TT+30%CW T-L3

.8e

mm

/s -0

8

1.4e-09 mm/s2 µm

9e-10 mm/s

Alloy 690SA+30%CW T-L

Pt ECPCT ECP

2.4e-08mm/s

e 5-18.

esults ost

Figure 5-58 Data from PNNL testing of CRDM material with deliberate cold working (here 30% in the T-L orientation) [26]


The key features of this program were described already in Section 5.1.2.1.8 and the numericalresults for the CT specimens of Alloy 690 in the cold-worked condition are tabulated in TablMore cracking, and about 3 times higher growth rates, were observed in comparison with the rfor the cold-worked Alloy 690 TT specimens (see Figure 5-41). Note, however, that although mof the cold-worked specimens displayed CGRs of around 3 E-8 mm/s, one specimen cracked appreciably faster (albeit at a somewhat higher stress intensity).

5-59

PWS

5-60

CC of Thick-Walled Alloy 690 Material

mm mm mm

eCGR 0 m/s

Table 5-18 Detailed results for the Alloy 600 control samples and Alloy 690 cold-worked CRDM materials used in Westinghouse testing for CGRs in supercritical water [25]

Material Spec ID

da Fatigue Calc.1

da FatigueMeas.2

da SCC

Meas.3CrackMode4

TotalTime5

Hour

SCCTime6

Hour

Kphase 2

7

MPa√m Kaverage

8 MPa√m

Kfinal

9 MPa√m

Averag1

600 MA 510-04a 0.037 n/m 1.690 IG 1920 1546 16.9 19.2 20.7 3.0E-10

600 MA 510-04b 0.037 n/m 1.540 IG 1920 1546 16.7 19.3 20.7 2.8E-10

690TT cw F415-c 0.308 0.689 0.309 IG 4806 4323 31.5 31.1 31.9 2.0E-11

690TT cw F7a21-c 0.312 0.561 0.313 IG 4806 4323 31.7 31.1 31.8 2.0E-11

690TT cw F415-b 0.319 0.204 0.290 IG 1920 1546 28.8 32.4 33.6 5.2E-11

690TT cw F7a21-a 0.454 0.400 2.144 IG 1920 1546 30.6 37.0 41.4 3.9E-10

Notes: see Table 5-16

5.1.2.4 Testing of Heat Affected Zones (HAZ) from Welding of Alloy 690 Material

ch of the justification for testing Alloy 69Mu 0 base material with deliberate cold working relates

fusion line afHAZ, areHAZ specimens [30]. Thus testing of actual Alloy 690 HAZ material was recommended to fill in a knowledge gap in the 2004 MRP-111 report. Details of known efforts in this direction and preliminary results are reported below, but it is clear that this is work-in-progress and final conclusions cannot yet be drawn.

5.1.2.4.1 Investigations in High-Temperature Water at KAPL

As mentioned above, KAPL drew attention some years ago to the possibility that the heat-affected zone (HAZ) in Alloy 600 might be significantly more susceptible to SCC than unaffected base material and have confirmed this finding more recently [31]. Detailed investigations revealed, however, that Alloy 690 HAZ microstructures are different from Alloy 600 HAZ microstructures due both to the higher chromium content and higher solvus temperature of Cr23C6. Neither unconstrained, nor constrained welds (with up to 14% plastic strain in the weld metal) exhibited SCC susceptibility of the Alloy 690 HAZ under conditions where Alloy 600 HAZs cracked readily [31].

5.1.2.4.2 Investigations in Simulated Primary Water at Argonne National Laboratory (ANL)

The general ANL testing approach was described Section 5.1.2.1.2. Figure 5-59 shows the geometry of a specimen that was under test in November 2008 [13]. No further details are yet available.

to the possibility of additional tensile strains being present at relatively high levels adjacent to the ter welding (see [17]). These, together with possible changes in microstucture in the

thought to explain the more rapid cracking seen in KAPL testing of Alloy 600


[13] Argonne, LLC, for the U.S. Department

Investigations in Simulated Primary Water at Studsvik in Sweden

Figthe no e ent of CGR, but a detailed comparison with both base metal from this particular heat

Figure 5-59 Details of an Alloy 690 HAZ specimen (CF690) under test at ANLGraphic provided by Argonne National Laboratory, managed and operated by UChicago of Energy under Contract No. DE-AC02-06CH11357.

5.1.2.4.3

ure 5-60 shows details of a 1T CT specimen on test at Studsvik in November 2008 [27] with crack growing parallel to, but ~ 1mm away from, the fusion line. Preliminary results suggestnhancem

and with an Alloy 600 HAZ specimen has yet to be carried out.

5-61


Figure 5-60 Details of CT specimen being used at Studsvik to examine the CGR behavior in Alloy 690

lity b

out in simulated primaryth s

[33]thes

5.1

As parallel to, the fusion line has apparently been details and current status are not known, but it is believed that SCC was not detected after 4360 h at 340°C for this specimen, indicating that this particular HAZ probably retained high PWSCC resistance despite significant hardening from the welding process.

HAZ material [27]

5.1.2.4.4 Investigations in Simulated Primary Water at CIEMAT in Spain

CIEMAT is carrying out a MRP- and UNESA-sponsored test program on the CGR susceptibiof oth Alloy 600 and Alloy 690 HAZs [32]. Two CGR tests on Alloy 600 HAZs have been carried

water at 360°C. Specimen evaluation proved complicated, however, due bo to fabrication defects in some of the test welds and incomplete linearity of the weld fusion line

. Subsequent testing of Alloy 690 HAZ specimens has been delayed pending clarification of e issues, but further results are expcted shortly.

.2.4.5 Investigations in Simulated Primary Water at Tohoku University in Japan

Information on the main testing program at Tohoku University was reported in Section 5.1.2.1.4.part of this, a 12.5mm thick CDCB specimen with the notch located 2 mm away from, and

used for an Alloy 690TT HAZ test [8]. Further

5-62


5.2 Field Experience

In contrast to the situation with thin-walled SG tubing described in Section 4.2, new or replacemecomponents made from thick-walled Alloy 690 material have only been in service for much shorteperiods. Although no problems whatsoever with cracking have been reported to date, there is also

nt r

not yet the same body of data available from repeat inspections to demonstrate conclusively the

eplaced in September 1994) was first inspected in 2002, i.e. after around 8 years of operation.

– Gravelines 4 (head replaced in March 1994) and Blayais 2 (head replaced in March 1995) were first inspected in 2003, i.e. after 9 and 8 years of operation, respectively.

– Next inspections for all 3 lead units will be in 2013.

• The primary inspection uses eddy-current techniques and indications are noted if their lengths are ≥ 2 mm.

• Any such indications would then be sized and characterized by further ultrasonic inspection.

• No flaws have been detected so far.

5.3 References

1. F. Cattant, Inspection of Reactor Pressure Vessel Heads equipped with A690 penetrations at EDF. Presentation to the EPRI Primary Systems Corrosion Research Committee, Long Beach, CA: January 2009.

2. G. Theus et al., Materials Reliability Program: Material Production and Component Fabrication and Installation Practices for Alloy 690 Replacement Components in Pressurized Water Reactor Plants (MRP-245). EPRI, Palo Alto, CA: 2008. 1016608.

3. F. Vaillant, EDF Report HT-44/95/013/A, 1996, “Résistance a la corrosion sous contrainte en milieu primaire des alliages 690 et 800 – Point des résultats en Décembre 1995”, (English translation of the title, “Resistance of Alloys 690 and 800 to Stress Corrosion Cracking in PWR Primary Water – Status of Results Available to December 1995”).

4. S. Asada et al., PWSCC Life Time Evaluation on Alloy 690,52 and 152 for PWR Materials. EPRI PWSCC of Alloy 600 2007 International Conference & Exhibition, June 11-14, 2007, Atlanta, GA.

5. S. Asada et al., PWSCC Life Time Evaluation on Alloy 690,52 and 152 for PWR Materials. EPRI MRP PWSCC Expert Panel Meeting, St. Petersburg, Florida (Nov. 2007).

6. G.S. Was and S. Teysseyre, “Challenges and Recent Progress in Corrosion and Stress Corrosion Cracking of Alloys for Supercritical Water ReactorCore Components”, Proc. 12th Int. Conf. On Environmental Degradation of Materials in Nuclear Power System – Water Reactors, TMS (The Minerals, Metals & Materials Society), 2005.

absence of incipient PWSCC. However, the efforts made in France to assess the field performance of CRDM penetrations are of considerable value here and some key points from these were recently described by EDF [1]:

• Repeated inspection of RPV head penetrations is taking place during every 10-year unit outage for 3 lead units as follows:

– Bugey 3 (head r

5-63


7. G.S. Was and S. Teysseyre, Materials Reliability Program: Constant Extension Rate SCC Testing of Alloys 600 and 690 in Supercritical Water. EPRI, Palo Alto, CA (2007). Technical Report 1016154 (MRP-233).

8. J. Hickling, EPRI Materials Reliability Program: Resistance of Alloys 690, 152 and 52 to Primary Water Stress Corrosion Cracking (MRP-237, Rev 1): Summary of findings from completed and ongoing test programs since 2004. EPRI, Palo Alto, CA: 2008. 1018130.

9. P.L. Andresen, Development of Advanced Testing Techniques to Quantify the Improved PWSCC Resistance of Alloy 690 and its Weld Metals. EPRI, Palo Alto, CA: (2004). Technical Report 1010269 (MRP-123).

10. M. Morra et al., Quantification of grain size and banding in differently Thermo-Mechanical-Processed (TMP) Heats of Alloy 690 using Image Analysis. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

11. B. Alexandreanu et al., Crack Growth Rates of Alloys 690 and 152 in PWR Environment at 320°C. EPRI MRP PWSCC Expert Panel Meeting, Atlanta, GA (Oct. 2006).

12. B. Alexandreanu et al., SCC CGRs of Alloys 690 and 152 Weld in PWR Water. EPRI MRP PWSCC Expert Panel Meeting, St. Petersburg, Florida (Nov. 2007).

13. B. Alexandreanu et al., SCC CGRs of Alloys 690 and 52/152 Welds in PWR Water. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

14. B. Alexandreanu et al., “Task 3: Cracking of Nickel Alloys and Welds – CGRs of Alloys 600 and 690 in PWR Water”, ANL presentation at NRC review meeting, Argonne, IL, September 25-26, 2007.

15. B. Alexandreanu, “SCC CGRs of Alloys 690 and 52/152 Welds in PWR Water”, Alloys 690/52/152 PWSCC Research Test Materials Meeting, Industry/NRC RES, July 17-18, 2008, Rockville, MD http://www.nrc.gov/reading-rm/adams/web-based.html, ADAMS Accession Number: ML082140693.

16. W.J. Shack and T.F. Kassner, “Review of Environmental Effects on Fatigue Crack Growth of Austenitic Stainless Steels”, NUREG/CR-6176 ANL-94/1, (May 1994).

17. P. Andresen and M. Morra, SCC Growth Rate Data in Alloy 690 Base and Weld Metals. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

18. P. Andresen and M. Morra, Materials Reliability Program: Laboratory Testing to Determine Resistance of Alloys 690/52/152 to Stress Corrosion Crack Growth in Simulated Primary Water—An Update (MRP-253). EPRI, Palo Alto, CA: 2008. 1016604.

19. T. Shoji, Program on Technology Innovation: Prediction and Evaluation of Environmentally Assisted Cracking in LWR Structural Materials: Interim Report on PEACE-E, November 2008. EPRI, Palo Alto, CA, Tohoku University, EDF-SEPTEN, Hitachi Ltd., The Japan Atomic Power Company, The Kansai Electric Power Co., Inc., Mitsubishi Heavy Industries, Ltd., Swedish Radiation Safety Authority, Tohoku Electric Power Co., Inc., Tokyo Electric Power Company, Toshiba Corporation, and Ishikawajima- Harima Heavy Industries Co., Ltd.: 2008. 1016549.

5-64

http://www.nrc.gov/reading-rm/adams/web-based.html


20. T. Shoji, Program on Technology Innovation: Prediction and Evaluation of Environmentally sted Cracking in LWR Structural Materials: Interim Report on Phase III. EPRI, Palo

CA; Fracture and Reliability Research Institute, Tohoku University, Miyago, Japan; SEPTEN, Villeurbanne Cedex, France; Hitachi Ltd.; The Japan Atomic Power

vy Industries, Ltd.; Swedish

Electric Power Company; Toshiba Corporation; and Ishikawajima – Harima Heavy Industries Co.: 2006. 1013380.

rt Panel Meeting,

CC

. EPRI Workshop on

(MRP-225). EPRI, Palo Alto, CA: 2007. 1015004.

s. ting, Los Angeles, CA (Nov. 2008).

d

29. P.L. Andresen, Resistance of Alloy 690/52/152 to Stress Corrosion Crack Growth in lto, CA: (2006). Technical Report 1013516

30. G.A. Young, N. Lewis & D.S. Morton: “The Stress Corrosion Crack Growth Rate of Alloy 600 Heat Affected Zones Exposed to High Purity Water”, NUREG/CP-0191, The Vessel Penetration Inspection, Crack Growth and Repair Conference, Oct. 2, 2003, Gaithersburg, MD, Vol. 1, pp. 371-386.

31. G.A. Young et al.: Stress Corrosion Cracking of Alloy 600 and Alloy 690 Heat Affected Zone Material in High Purity Deaerated Water. EPRI MRP PWSCC Expert Panel Meeting, Atlanta, GA (2006).

32. D. Gomez-Briceno, Crack Growth Rate Studies of Weld HAZ of Alloys 600 and 690 Materials. EPRI MRP PWSCC Expert Panel Meeting, Atlanta, GA (Oct. 2006).

33. D. Gomez-Briceno and J. Lapena, Materials Reliability Program: Effect of Defects in an Alloy 600/82 Weld on Stress Corrosion Cracking in Testing of Heat Affected Zone Specimens (MRP-254). EPRI, Palo Alto, CA: 2008. 1016607.

AssiAlto, EDF-Company; The Kansai Electric Power Co., Inc.; Mitsubishi HeaNuclear Power Inspectorate, Stockholm, Sweden; Tohoku Electric Power Co., Inc.; Tokyo

21. T. Arai: CRIEPI Research Program on PWSCC, EPRI MRP PWSCC ExpeSt. Petersburg, Florida (Nov. 2007).

22. D.J. Paraventi and W.C. Moshier, Alloy 690 SCC Growth Rate Testing. EPRI MRP PWSExpert Panel Meeting, Atlanta, GA (2006).

23. D.J. Paraventi and W.C. Moshier, Alloy 690 SCC Growth Rate TestingCold Work in Iron and Nickel Base Alloys, Toronto (2007).

24. M.G. Burke, Microstructual Characterization – Alloys 690 and 600. EPRI MRP PWSCCExpert Panel Meeting, Los Angeles, CA (Nov. 2008).

25. R. Jacko, Materials Reliability Program: Testing the Resistance to Stress Corrosion Cracking of Alloy 690 and its Weld Metal in Supercritical Boron/Lithium/H2 Solutions

26. M. Toloczko and S. Bruemmer, Crack Growth Response of Alloy 690/152/52 in SimulatedPWR Water. EPRI MRP PWSCC Expert Panel Meeting, Los Angeles, CA (Nov. 2008).

27. A. Jenssen, K. Norring and P. Efsing, Swedish Activities on Alloy 690 and its Weld MetalEPRI MRP PWSCC Expert Panel Mee

28. Y. Yamamoto et al., “Development of the Crack Growth Rate Curves for Stress CorrosionCracking of Nickel Based Alloys in a Simulated Primary Water Environment”, Fontevrau6 (2006).

Simulated PWR Primary Water. EPRI, Palo A(MRP-196).

5-65

6 DISCUSSION

6.1 Resistance of Alloy 690 to PWSCC

The PWR industry selection of Alloy 690 as the material of choice to replace Alloy 600 components susceptible to PWSCC appears to have been a sound one for three main reasons:

Firstly, the field experience to date whatsoever of PWSCC. This is truenow having been exposed for up to nearly 20 years) and for thick-walled components suchhead penetrations (where service experience without detection of cracks has already been demonstrated in 3 lead units after nearly 10 years).

with Alloy 690 components has been exemplary, with no sign both for thin-walled SG tubing (with many thousands of tubes

as RPV

Secondly, Alloy 690 has proved to be almost entirely resistant to PWSCC initiation in a wide ing some

y, fatigue pre-cracked, thick-walled Alloy 690 CRDM material thought to be generally ative of real components has shown no, or only insignificantoo, amounts of crack growth

sted under virtually static load, i.e. at (nearly) constant, stress intensity in simulated PWR

lloy 690TT has shown SCC CGRs > 5E-9 mm/s r containing additions of Li

ual PWR operation remains -8 mm/s, i.e. a factor

of 2x higher than the threshold rate suggested as representing complete lack of engineering d some 60x lower than the SCC CGRs for Alloy 600 tested

er environment. Somewhat higher rates were measured, however,

variety of laboratory tests (see Sections 4.1 and 5.1.1), the only marginal exceptions beminor cracking in pre-production heats of SG tubing tested under very severe conditions (see Sections 4.1.3 and 4.1.4).

Thirdlrepresentwhen teprimary water. This has been found in multiple laboratories, even when the experimental conditions (temperature, crack transitioning, periodic partial unloading, test duration, etc.) have been chosen so as to maximise even the smallest degree of inherent SCC susceptibility (see Section 5.1.2.2 and Table 6-1).

The only real case to date where as-received Awithout deliberate cold working involves testing in supercritical wateand B (see Section 5.1.2.2.5) and the relevance of these results to actunclear. Even in this case, however, the CGRs observed were only ~ 1E

significance in primary water anin the same supercritical watin specimens from materials that had been subjected to deliberate cold working – see Section 5.1.2.3.3.

Alloy 690 materials.

oo It is suggested that complete lack of engineering significance might best be defined as a CGR ≤ 5E-9 mm/s (i.e. a maximum of ~0.15 mm/year); this often corresponds to around the average grain diameter in thick-walled

6-1

Discussion

The absence of any significant SCC growth from a pre-existing crack in the thick-walled Alloy 690 CRDM material specimens tested to date in simulated primary water is obviously very satisfactory. This continues to be the case even after the deliberate introduction of significant amountsadditional cold work (both uniform and non-uniform) to as-received material, as summarized in Section 5.1.2.3 and Table 6-2. Thus, it is somewhat difficult to report a meaningful factor of improvement (FOI) fodate for Alloy 690 CRDM penetration materials with the 75

of

r CGR with regard to Alloy 600, but comparison of all the data reported to th percentile curve for Alloy 600 CGRs

from MRP-55 [12] would suggest a FOI of at least 100x and probably higher.

6-2

Discussion

6-3

Table 6-1 Summary of results of CGR testing in s ter ate(status: December 2008)

Testing Lab

Material Supplier

Material Form

Heat # Final dition

Micro-Structure

pec

(

imulated PWR primar

Con

y wa

S

on Alloy 690 CRDM material without del

imen #

Pre-Crack Orientation

Test Temp.

(C)

H2 Level (cc/kg)

Eff. K (MPa

iber

Level√m)

cold working

Rep. CGR mm/s)

Remarks

ANL Valinox CRDM tubing WP142 TTpp homogeneous C6901 24 2

of

tant -CR- long/radial? 320 23 E-09

No clear sign IGSCC after 1000h at consload

ANL Valinox CRDM tubing WP142 TT homogeneous C690 21 3

d to ing o

-LR-2 circ/radial? 320; 350 23 E-09

CGR increase~8E-9 on raistemperature t350C

MHI ? CRDM tubing?

Not known TT homogeneous Not k known 1

d GR

ro nown Not known Not known

Not known Not E-10

No SCC founafter 4000h; Ceffectively ze

PNNL Valinox CRDM tubing RE243 TT homogeneous CT14 30; 40

ated

zero long/radial 325 29 3E-10

Only very isolsigns of SCC;CGR almost

PNNL Valinox CRDM tubing RE243 CMqq homogeneous CT15 30; 40

ated

zero long/radial 325 29 2E-10

Only very isolsigns of SCC;CGR almost

PNNL Valinox CRDM tubing WP140 TT homogeneous CT26 30 5E-10

ress, nly. sed = 20

long/radial 350 29; 20

Tests in progDCPD data oCGRs increaby ~2X at H2cc/kg.

PNNL Valinox CRDM tubing WP142 TT homogeneous CT27 l 30 7E-10

ress, nly. time t K

long/radia 350 29; 20

Tests in progDCPD data oCGRs ↓ with under constanconditions

Studsvik Sumitomo

archive CRDM penetration material D520906 TT homogeneous 182-1 25 1E-10

d s of ~ 0 long/circ 320 25-29

Tested under PPU with increasing holtimes; no signIGSCC; CGR

Studsvik Sumitomo

archive CRDM penetration material D520906 TT homogeneous 182-2 long 25 1E-10

d s of ~ 0 /circ 320 30-31

Tested under PPU with increasing holtimes; no signIGSCC; CGR

pp TT = Thermally Treated qq CM = Carbide Modified (TT plus an additional anneal)

sion

Table 6-2 Summary of u of CGR testing in simula P w A 690 CRD h i orking

t

tp

er t C i

c y E

Cevel (

c c T l

( m)

Rep. CGR

(mm/s) Remarks

resemb

erial plier

lts er 20

MatFo

ted

MiStru

WR

ro-cture

prim

T

ary

pe ofxtra CW

ater

M

on

ethod

lloy

L

M ma

. PrOri

teria

e-Craentati

l wit

k on T

del

est emp. (C)

bera

H2Lev

(cc/k

te co

el g)

ld w

Eff. KLeve

MPa√

(s atus

Testing Lab

: Dec

MaSu

08)

rial m Hea #

Final ondit on

W

%)

Spe#

GG

E-RC

DukEne

e rgy

extrCRDpenon

u

ehomogeneou g 6 1E-09

Long-term test under PPU with increasing hold times; IGSCC CGR effectively zero

ded M trati

WN415 TTrr s uniform crofor

ss-ing 41 c280 T-L 3 0 18 38.5

GG

E-RC

DukEne

e rgy

extruCRDpeneon

ded M trati

WN415 TT homogeneous uni 6 5E-09

Tested under PPU with <1000h at constant K; IGSCC CGR very low form

croforg

ss-ing 20 c285 T-L 3 0 18 27.5

GG

E-RC

DukEne

e rgy

extruCRDpeneon

ded M trati

WN415 TT homogeneous uni 6 3E-09

Tested under PPU with <1000h at constant K; IGSCC CGR very low form

croforg

ss-ing 20 c286 T-L 3 0 18 27.5

PN 3 i g 9 2 4E-10

Only isolated signs of SCC; CGR very low indeed NL Valinox

CRDtubin

M g RE24 TT homogeneous

noun

n-form 1-D rollin 17 CT1 S-L 3 5 29 30; 40

PN 3 C i g 0 2 4E-10

Only isolated signs of SCC; CGR very low indeed 30; 40 29 5 3 S-L CT2 17 1-D rollins

noun

n-form homogeneouNL Valinox

CRDtubin

M g RE24 Mss

Discus

6-4

rr TT = Thermally Treated ss CM = Carbide Modified (TT plus an additional anneal)

Discussion

6-5

Table 6-2 Summary of results of CGR testing in simulated PWR primary water A M material with deliberate cold w(status: December 2008) (continued)

Testing Lab

r Material Form

Finaonditi

y E

(

c e ci i

Test H2 Eff. K

(

( s) Remarks

on

ethod

lloy

L

690

CW evel %)

CRD

Spe#

orkin

l m)

g

Rep.CGRmm/

MateSuppli

ial er Heat # C

l on

MiStru

cro-cture

T pe ofxtra CW

M. Pr

Or-Cra

entatk

on Temp. (C)

Lev(cc/k

el g)

LeveMPa√

PNN L ValinoCRDM

E243 T homogeneou i D g 2 3

mewhat re IGSCC; R very low x tubing R T s

noun

n-form 1- rollin 30 CT2 T-L 350 29 30; 40 E-9

SomoCG

PNNL Valino E243 homogeneou i g 3 9

ly isolated ns of SCC; R very low eed x t

CRDM ubing R CM s

noun

n-form 1-D rollin 30 CT2 T-l 350 29 30; 40 E-10

OnsigCGind

Discussion

It should be emicrostructures of Alloy 690 material used for RPV head penetrations in the field is thought to be larger than is repr d by the approx. 10 heats tested worldwide to date in various laboratories. Similarly, the complete Alloy 690 components t clarified. Of primary concern issignifica lts of CGR testi

In an earlier the present author described a “bimodal” aspect of Alloy 690 laboratory PWSCC behavior whereto relatively rapid intergrepresentative of actual this has alwaSection 5.1.2 008 by PNNL [2 -55 reference disposition curve for cracking in thick-walled Alloy 600 material [12]. It can be seen that most of the Bettis data, in particular, actually lie above this curve, but that such results have also been reproduced inare thought to represent been superimposed on a with regard to e. If such b and crack propag ns of the ANL plate material (see Sections 5 ibility may result. In contrast, e(see Figure 6-3) did not induce high CGRs in the original feasibility studies on EPRI plate material (see Section 5.1.2.1.1). F s microstructure of an ext when subjected to delib

A further group of data points (from ANL, Bettis, GE-GRC and INSS) lie below the MRP-55 curve, but above 5E-9mm/s, thAs discussed above, the even for the PNNL tests with m. Points pl below 1E-9 mm/s can reallyload is so slo the pre-fatigue c to the restrainalso very ne

noted, how ver, that the range of material chemistries, mechanical properties and

esentespectrum of material variability for long-term exposure of all thick-walled o primary water in both existing and new plants has not yet been fully

the possible presence of inhomogeneous microstructures and nt amounts of unidirectional tensile strain in Alloy 690 material, given the further resu

ng discussed below.

report [4], it has been possible under certain circumstances to demonstrate moderate ranular crack growthtt in plate and bar materials, thought to be less plant components. With regard to testing in high-temperature water,

ys involved additional forms of non-uniform cold work, as reported in detail in .1. Figure 6-1 shows a summary diagram of these results prepared in November 2

6] and compares the CGRs measured on 1-D cold-rolled specimens with the MRP

other laboratories (GE-GRC, INSS). Some (but not alluu) of these points the situation where the additional, non-uniform deformation at RT has n inhomogeneous microstructure in the original material (particularly

banding originating both from segregation (e.g. of carbides) and grains of varying sizanding (see, e.g., Figure 6-2) is aligned with the direction of subsequent cold working

ation, as may have been the case with the S-L and S-T orientatio.1.2.1.2 and 5.1.2.1.3), it is intuitively plausible that high SCC susceptxtensive banding perpendicular to the direction of crack growth

or comparison purposes, Figure 6-4 shows the very uniform, homogeneouruded CRDM penetration that was extremely resistant to SCC growth, evenerate cold working (see Section 5.1.2.3.1).

e suggested threshold for CGRs that are of no engineering significance. CGR results from CRDM material all lie below that threshold value, 30% 1-D cold rolling and a relatively high stress intensity of 40 MPa√

otted at 1E-10 mm/s in Figure 6-1 indicate virtually no SCC growth at all, but all data be regarded as representing a situation where the CGR under constant

w that it usually corresponds to only very localized pockets of minor SCC along rack front. Over long periods of time, such growth is unlikely to be sustained, owing ing influence of adjacent uncracked areas. It thus represents an overall CGR that is

arly zero.

tt Uuu Bettis report an apparent absence of significant banding in the one heat of their 3 test materials that showed the

highest CGRs.

p to 1000x higher than in CRDM material without deliberate cold work.

6-6

Discussion

Figure 6-1 Summary of laboratory SCC CGR data (as of November 2008) prepared by PNNL [26] and showing the possibility of measuring moderate to high rates in Alloy 690 p

non-uniform cold work late material

subjected to

Figudence of microstructural banding in some areas of the 1-D cold-rolled ANL Alloy 690

al examined at GE-GRC [18]

re 6-2 Eviplate materi

6-7

Discussion

Figlar to the crack plane in Alloy 690 plate material from

rogram by GE-GRC [9]

ure 6-3 Microstructural banding perpendicuEPRI orginally tested for the MRP P

Figure 6-4 y uniform, homogeneous microstructure in extruded Alloy 690 CRDM material [18] Ver

ior microstructural inhomo Alloy 6 at

CGR onto disap

for such(Figure 2

thor’s opinion, all of these

from ve

The mechanism by which unidirectional cold work (with or without prgeneity) can induce SCC susceptibility in an otherwise very resistant material such as 90 is currently unclear and is the subject of much ongoing research. It is noteworthy th

several laboratories (ANL, Bettis, GE-GRC, INSS) have found that the expected dependency of test temperature (based upon Alloy 600 cracking experience – cf. Figure 5-27) appears pear almost entirely for the most susceptible materials when tested in the worst crack

orientation (see Figure 5-18, Figure 5-20, Figure 5-31 and Figure 5-34). Furthermore, the CGRs specimens appear to exhibit either little effect of dissolved hydrogen concentration

5-25), or the opposite behavior to that known from Alloy 600 cracking (see Figure 5-2and Figure 5-35), and little effect of varying stress intensity. In the auobservations suggest bimodal behavior, rather than a continuum of Alloy 690 cracking response

ry slow to high CGRs.

6-8

Discussion

There has been lively debate concerning the relevance of these moderate to hi Alloy 690 plate and bar materials to the behavior of actual

components. On the one hand, such materials do not appear to be representative of wh used in plants, as discussed in Section 2.1.2, although some further clarification (e.g. w

gh laboratory CGR data from unidirectionally deformedreactor at is being ith

ual

strains 690 after welding. Although the resulting local microstructure is likely to be very different from

high CGRs were also found during the Bet .1.2.1.7). This is t adjacen ted

on 5.1.2.4, but this presents real experimental challenges and the data available to date are ing, but inconclusive.

For the g placed both on protypical CRDM materials (whose extreme resistance to PWSCC requires very

cted to deliberate cold wounidirectthe bou quickly and reliably than co way in which nlaborato

In addit conclusive results l water containi to demo

Althoug exhibit significant growth at constant load in n of Alloy 690 PWSCC resistance, it should ctually involve the initiation data (lab and field) from SG tubi nitiation (describ iety of test method the metallurgical conditions examined to date. These are thought to be generally representative of most

s, although further clarification is neededstrains in weld heat-affected zoncarried suscept

alled

e

regard to possible cold-straightening of extruded CRDM piping) of the complete range of actmaterial conditions is definitely required. On the other hand, significant, highly localized plastic

can be expected immediately adjacent to the fusion line in the HAZ of thick-walled Alloy

that produced by 1-D cold rolling of plate, it is noteworthy that tis studies on material that had been subjected to RT tensile straining (see Section 5hought to be far more representative of what might be present in Alloy 690 base materialt to real welds [18]. Attempts are also being made to test actual HAZ material, as repor

in Sectipromis

abovementioned reasons, widespread research [1] is continuing with the emphasis bein

long testing times to obtain valid data) and on such materials that have been subjerking and may be more susceptible. If bimodal behavior is involved here, testing of ionally deformed material – even if unrepresentative of field conditions – may permit

ndaries of Alloy 690 PWSCC susceptibility to be established more ncentrating just on “good” material. Furthermore, a better understanding of theon-uniform deformation can induce significant susceptibility to crack growth during ry testing is clearly required, as is clarification of whether or not such materials would

undergo crack initiation from a smooth surface in simulated primary water.

ion, efforts are being pursued in at least one laboratory [2] to understand the inthat have been obtained to date during testing of Alloy 690 materials in supercriticang Li and B additions, since this was originally expected to be a reliable accelerated testnstrate satisfactory PWSCC resistance in advance of long-term field exposure.

h evidence that a pre-existing (fatigue) crack does notprimary water is clearly the best possible demonstratiobe noted that the overwhelming majority of cracking incidents with Alloy 600 have ad incipient crack formation at a smooth surface. Thus ng (described in Section 4) and the lab testing of thick-walled Alloy 690 for crack ied in Section 5.1.1) are also of considerable importance. Despite the large vars used, these show essentially no crack formation through PWSCC for materials in all

field components, including thick-walled CRDM penetration of some factors (such as cold work from final straightening operations and high residual

es). It should also be noted that no initiation testing has yet been out on the 1-D cold-rolled Alloy 690 plate material that has shown moderate to high ibility to crack growth through PWSCC.

In the 2004 MRP-111 report [3], the average FOI from lab-testing for crack initiation in thin-wAlloy 690TT was derived as 26x relative to Alloy 600MA and 13x relative to Alloy 600TT (see Section 4.1.7), although it was emphasized that these numbers were clearly very conservativ

6-9

Discussion

and likely to increase, owing to the absence of cracking in the Alloy 690 specimens. An assessment of the relevant lab data for both thin and thick-walled material in 2008 [4] concluded that minimum

G

ata for Allo d.

r ,

or somewhat better, than that sometim

same 1-

ement

52 weld show adequate fracture resistance even under extreme loading conditions in low-temperature water,

good. As disc or non-represe

d.

1. A. A690/152/52 in Nuclear EnvironmentsAng

Wor geles, CA

ram (MRP), Resistance to Primary Water Stress Co 11). EPRI, Palo Alto, CA: 2004. 1009801.

Pri m completed and ongoing test programs since 2004. EPRI, Palo Alto, CA: 2008. 1018130.

factors of 40 to 100x can now be justified. As described in Section 4.2, field experience with Stubing already justifies a FOI >20x, and the realistic value is expected to lie much higher.

6.2 Other Aspects of Alloy 690 Corrosion Behavior

One of the gaps identified in the 2004 MRP-111 report was the absence of corrosion fatigue dy 690 in primary water. As reported in Section 3.2, this deficit has largely been rectifie

Thick-walled Alloy 690 does show a reduction in fatigue life when tested under exposure to a simulated primary water environment, and the extent of the decrease depends upon a numbeof factors (especially testing frequency/strain rate). However, this behavior is about the same

of other Ni-base alloys. Similarly, cyclic crack growth rates are es affected (in this case, increasedvv) by simultaneous exposure to a high-temperature

water environment, but generally slightly less than for Alloy 600. The only exception to this involves the more accelerated cyclic crack growth in simulated primary water seen for the

D cold rolled plate of Alloy 690 from ANL that showed relatively high susceptibility to PWSCC (see Section 5.1.2.1.2), but - as discussed above - this material is not thought to be of direct relevance to plant components. In contrast to this finding, prototypical Alloy 690 CRDM material (that was highly resistant to PWSCC crack growth) showed no environmental enhancof fatigue crack growth at all in recent Swedish studies [15].

With regard to low temperature crack propagation (LTCP - see Section 3.4), there is a consensus among experts that Alloy 82 and 182 give rise to more potential concerns than the Alloy 152 and

metals, and that Alloy 690 base metal is of even less concern. The latter is expected to

particularly if dissolved hydrogen levels are kept low, as long as its metallurgical condition isussed above for PWSCC resistance, however, this assessment may no longer be valid f

ntative materials that have been subjected to severe unidirectional cold work. Additional clarification of the boundaries to enhanced LTCP susceptibility for Alloy 690 is recommende

6.3 References

hluwalia and R. Tregoning, Collaborative Research Activities: Performance of Alloy . EPRI MRP PWSCC Expert Panel Meeting, Los

eles, CA (Nov. 2008).

2. R. Jacko and J.K. McKinley, PWSCC Resistance of Alloy 690 and its Weld Metals: Cold k, Temperature and H2 Effects. EPRI MRP PWSCC Expert Panel Meeting, Los An(Nov. 2008).

3. H. Xu et al., Materials Reliability Progrrosion Cracking of Alloys 690, 52, and 152 in Pressurized Water Reactors (MRP-1

4. J. Hickling, EPRI Materials Reliability Program: Resistance of Alloys 690, 152 and 52 to mary Water Stress Corrosion Cracking (MRP-237, Rev 1): Summary of findings fro

vv E.g., by some 5 to 10 times in Japanese studies at slow loading frequencies [16].

6-10

7 CONCLUSIONS

Significthick-se

• Cra not hese are thought

lled CRDM pen of some factors (such as cold work fromaffe

0TT

• Suc o actual PWSCC originating at a smooth surface has yet been seen (except for marginal levels

• The racks g in-service inspections after nearly 20 years of plant operation with Alloy 690 steam

nd nearly 10 years with CRDM nozzles in replacement RPV hea

• Forions), susceptibility to crack growth through PWSCC

posspec /s

uniform coldby introducing non-uniform deformation (e.g. by uniCR

ial with respect to CGRs in All

ant progress towards demonstrating satisfactory, long-term, PWSCC resistance of ction Alloy 690 by laboratory testing has been made since the publication in 2004 of

MRP-111, but complete immunity to cracking should not be assumed:

ck initiation through PWSCC from a smooth surface appears to be very difficult, if impossible, for Alloy 690TT in the metallurgical conditions tested to date. Tto be generally representative of most field components, including thick-wa

etrations, although further clarification is needed final tube straightening operations and possible high residual strains in weld heat-

cted zones).

• Minimum relative factors of improvement for initiation of 40 to 100 times for Alloy 69can now be justified with regard to PWSCC of Alloy 600MA.

h estimates are conservative and likely to rise further with continued testing, since n

of crack initiation in some specimens subjected to severe deformation in CERT).

se findings from laboratory testing are in accord with the total absence of detected cduringenerator tubing materials a

ds.

thick-walled Alloy 690TT material in the normal metallurgical condition (e.g. as extruded piping for CRDM penetratappears to be very marginal at the most. Using sophisticated testing methods, it has been

sible to produce small amounts of intergranular cracking in some, but not all, of the imens tested to date. However, the measured CGRs are extremely low (< 5E-9 mm

or < 0.15 mm/yr) and thus of no engineering significance.

• Despite the small increase in measured CGRs sometimes seen, this statement remains valid even after the deliberate introduction of significant amounts (i.e., well over 10%) of

work into CRDM material. To date, even attempts to induce higher SCC susceptibility directional cold rolling) into actual

DM material have not produced CGRs that are of any concern.

• The relative factor of improvement for Alloy 690 CRDM materoy 600MA is thought to be over 100 times (and may actually be much higher).

7-1

Conclusions

Perhaps inevitably, however, certain conditions have been found that appear to lead to “bimodal” r of some Alloy 690 materials in the sense that a much higher susceptibilty to crack grbehavio owth

through SCC (up to 1000 times faster) is then observed during laboratory testing of pre-cracked

• Rap oy s

any cases, this also appears to lead to major changes in the dependence of measured CGa di

• The

suctho d

mstances.

• Sigwork) during testing in a

re is presently no e

of the further research needed to reach the original goal of demonstrating long-term absenceis ongo nal collaboration. The industry and regulatory prioritization of these is seen as follows:

sus

• Esta to that

mat

l y for

plan

specimens in high-temperature environments:

id, intergranular crack growth (sometimes even at rates comparable to those for All600) in both pure water and simulated primary water can be triggered in certain directionwithin some Alloy 690 plate and bar materials by the introduction of inhomogeneous cold work (particularly by uni-directional rolling, but also by tensile straining).

• In mRs on stress intensity, test temperature and dissolved hydrogen levels, possibly indicating fferent cracking mechanism.

limits of such behavior have not yet been satisfactorily established and there is no real mechanistic understanding of the phenomenon. To date, the Alloy 690 materials exhibiting

h high susceptibility to intergranular SCC after unidirectional cold working are not ught to be directly relevant to PWR plant components, but further efforts are require

to confirm that this interpretation will remain valid under all circu

nificant intergranular crack growth has been observed in Alloy 690TT (including in CRDM material after the introduction of 10 % uniform cold supercritical environment containing lithium, boron and hydrogen. The

stablished basis for extrapolating the measured CGRs at 385 oC down to subcriticaltemperatures in order to make detailed comparisons with the existing databases for primary water.

In terms of PWSCC in all thick-walled plant components made from Alloy 690 base material, work

ing in a number of areas and these efforts are being coordinated as part of an internatio

• Complete planned testing programs to investigate the possibility of enhanced PWSCC ceptibility in the HAZ adjacent to welds in thick-walled Alloy 690.

blish the limits to the apparently bimodal behavior of Alloy 690 (high susceptibilitySCC crack growth in some material that has been subjected to non-uniform cold work) has been observed to date in laboratory testing and understand why this occurs.

• Carry out testing for crack initiation on some of the uni-directionally rolled Alloy 690 erials that have shown appreciable susceptibility to crack growth through PWSCC.

• Investigate rigorously whether such material conditions could have any relevance to reaplant components, taking into account that the complete spectrum of material variabilitlong-term exposure of thick-walled Alloy 690 to primary water in both existing and new

ts has not yet been fully clarified.

7-2

Conclusions

Considerarding corrosion fatigue appears to have been nearly closed. Alloy 690 TT material ntative of thick-walled plant components is expected to exhibit a reduction in fatigue lin acceleration in cyclic crack growth rates as a result of exposure to primary water tha

ing other knowledge gaps for Alloy 690 identified in the 2004 MRP-111 report, that regreprese fe and/or a t is

n behavior for other Ni-base alloys (such as Alloy 600).

With reAlloy 8 and

ure olved

ome additional ative Alloy 690

materia

comparable with, or better than, the know

gard to low temperature crack propagation (LTCP), there is a consensus among experts that 2 and 182 give rise to more potential concerns than the Alloy 152 and 52 weld metals,

that Alloy 690 base metal is of even less concern. The latter is expected to show adequate fractresistance even under extreme loading conditions in low-temperature water, particularly if disshydrogen levels are kept low, as long as its metallurgical condition is good. Sclarification of the boundaries to enhanced LTCP susceptibility for less-represent

ls is recommended.

7-3

A TRANSLATED TABLE OF CONTENTS

DISCLAIMER OF WARRANTIES AND LIMITATION OF LIABILITIES

THIS DOCUMENT WAS PREPARED BY THE ORGANIZATION(S) NAMED BELOW AS AN ACCOUNT OF WORK SPONSORED OR COSPONSORED BY THE ELECTRIC POWER RESEARCH INSTITUTE, INC. (EPRI). NEITHER EPRI, ANY MEMBER OF EPRI, ANY COSPONSOR, THE ORGANIZATION(S) BELOW,

Y PERSON ACTING ON BEHALF OF ANY OF THEM: NOR AN

OWNED DOCUME

(B) ASSUMES RESPONSIBILITY FOR ANY DAM

ACCURA TERMS A

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Power Research Institute (EPRI)

A-1

Translated Table of Contents

材料信頼性プログラム:加圧水型原子炉におけるAlloy 690の一次冷却水応力腐食亀裂への耐性 (MRP-258)

1019086

終報告 2009年８月

製品説明

鍛錬用合金600およびその溶接金属(Alloy 182およびAlloy 82)は、多数の浸食性環境における一般的腐食に対する材料の特有の耐性の理由から、また低合金および炭素鋼に極めて近い熱膨張率のために、本来加圧水型原子炉(PWR)で使用されていました。近30年以上、PWR 一次冷却水内応力腐食亀裂(PWSCC)は、多数のAlloy 600構成項目および関連する溶接部で、場合によっては比較的長い潜伏期間の後で、観

A-2


察されてきました。PWSCCの発生は、深刻なダウンタイムと交換の発電コストの原因となります

部品の修理、交換は、一般に鍛錬用合金690材料とその適合する溶接金属（Alloy 152およびAlloy 52または52M）が使用されます。研究所の実験ではPWSCCに対する高い耐性を示し、すでに20年近くにもわたって原子炉の稼働中の亀裂はありません。信頼できる技術的基礎を修理または交換部品項目のさらなる検査要件の向上に役立てるため、腐食による経年劣化に関してこれらの材料の耐用年数の数値化を試みることに挑戦しています。この文書は、2004年に公開されたEPRI報告書1009801 (MRP-111)を改訂し、近得たAlloy 690ベース材料のPWSCC挙動に関する情報を考慮に入れた増補版です。Alloy 152および52溶接金属の性能はここでは考慮されていませんが、後日単独で報告される予定です。

2004年の報告書(MRP-111)で知識のギャップが確認されたため、厚肉Alloy 690材料の腐食疲労および低い温度の亀裂の伝播についても検討しています。

方法 MRP-111報告書を基に、PWSCCに主な焦点を当てて、PWR環境における腐食耐性に関連して近20年間にわたって行われたさまざまなテスト条件の下で鍛錬用合金690材料を使

されました。可能な限り、既存の実験室試験データが評価されAlloy 600と比較してAlloy 690の改善度が予測されました。また、PWRにおけるAlloy

90使用実験が報告されたため、実験結果を増補しました。

結論づけられました。これらには、不均一な冷間加工(特に

。

って行われる多数の実験室試験が検証

6

結果と研究成果 PWSCCに関して明らかになった特定の知識のギャップについて調査するためテストをさらに続ける必要がありますが、鍛錬用合金690は条件を満たしており高腐食耐性のPWRにおけるAlloy 600代替材料であることが

A-3


一方向冷間圧延/引張歪)の亀裂の拡大に対する耐性に有害な影響と、溶接後の加熱影響ゾーンで亀裂発生度が高まる可能性が含まれます。亀裂の始まりの場合Alloy 600に対比して40～100倍の相対的な改善要因が判明しましたが、試験期間内にはほとんどすべてのAlloy 690試料中にPWSCCがなかったため、これらの数は明らかに控えめです。典型的な厚肉Alloy 690材料（たとえば、原子炉圧力容器の上蓋貫通部の押し出しパイプなど）は、発生度が大になるよう設定された実験室試験の条件下でさえも、一次冷却水模擬水中でのSCC成長に対する耐性が極めて高いです。これまでに計測された亀裂拡大率は工学的重要性がない程きわめて低く(5E-9 mm/sまたは0.15 mm/yearより小さい)、Alloy

y トおよび通常代表的なプラントの構成部品と考えられていないバーの材料に

られる材料の使用実験は、検討する構成部品の種類に応じて約10～20年に及びしました。

600MAに比較してAlloy 690 CRDM材料の改善の相対的要因は100倍を大幅に上回ると考えられます。ただし、粒間亀裂は不均一な冷間加工をAllo690プレー導入した後の試験の一定の負荷の下で極めて高い割合で観察されました。調査結果が長期のPWRに関連する可能性が低いことを確認するために、この発生度境界の調査は続けられています。

Alloy 690の応力腐食劣化はこれまでの交換応用では観察されていません。検査のためにPWR一次冷却水にさ

応用、値、および使用この報告書はPWR一次冷却水におけるAlloy 600亀裂のすべての局面に関わる電力会社のエンジニアおよび科学者と、特に検査体制の発展、構成部品交換の決定、およびプラントの老朽化問題の取り扱いに関わる技術者が関心を持つべき内容です。これは、既存のプラントにおける厚肉なAlloy 600構成部品に影響を与えるPWSCCの増え続ける事故に取り組むために産業によって採用されたソリューションを規制の承認を得ることの直接の価値となります。

A-4


電力会社、原子炉ベンダー、およびエンジニア/研究機関に関連する継続中の、包括的なプログラムの一部、この報告書はニッケルベース合金の腐食劣化が耐用年数を制限せず、交換構成部品および新しい原子炉の両方の設計を改善することから十分な利益が得られることを確認する手助けにもなります。

EPRIの観点この報告書は、Alloy 690のPWSCC耐性に関する現在のナレッジベースに記載されています。データはAlloy

間加工などの要因があります。進行中および計画済みの研究はこれらの要素により明確な定義PWSCCへのAlloy

る脆弱性を特定することに焦点を絞っています。このさらなる研究

PWSCC

690がその先行材料よりもPWSCCに対して顕著により大きな耐性を持っていることを示しています。これらの結果は有望であり付随調査が継続されると見込まれますが、発生度を若干損なう可能性のある冷

690耐性と、あらゆが完了するまで、この報告書に発表された結果は情報に過ぎず、Alloy 690から構成される部品の部品寿命の予測には使用しないでください。

キーワード Alloy 600 Alloy 690

材料劣化 RPV貫通部

A-5


目次

1 序言.......................................................................................................................... 1-1

.................................................................................................................... 1-1

改訂の目的および範囲 ........................................................................... 1-2

y 600 PWSCCの原因 .................................................................................... 1-4

書............................................................................................................. 1-5

0の特性と冶金学........................................................................................ 2-1

2.1 材料の仕様.......................................................................................................... 2-1

2.1.1 薄肉Alloy 690 SG配管の典型的なPWR仕様 ................................................... 2-3

2.1.2 厚肉Alloy 690構成部品の仕様と製造............................................................... 2-4

2.2 Alloy 690の状態図 .............................................................................................. 2-4

2.3 炭素溶解度と動的ひずみ時効 ............................................................................ 2-5

2.4 粒界カーバイド析出と鋭敏化 .............................................................................. 2-8

2.5 高温暴露の影響 .................................................................................................. 2-12

2.6 参考文書........................................................................................................... 2-13

3 PWSCCを除くALLOY 690の腐食挙動..................................................................... 3-1

3.1 一次冷却水内での一般的腐食試験 .......................................................................... 3-1

3.1.1 SedricksほかによるSG配管 1979 ............................................................... 3-1

3.1.2 K. SmithほかによるSG配管 1985 .............................................................. 3-2

3.1.3 YonezawaほかによるSG配管 1985 ........................................................... 3-2

3.1.4 Espositoほか 1991 ..................................................................................... 3-2

3.1.5 PWSCCメカニズムに関する合金酸化の研究 ............................................ 3-3

3.2 一次冷却水内での腐食疲労試験 ............................................................................. 3-4

3.3 二次冷却水内での腐食挙動.................................................................................. 3-15

3.4 低温における亀裂伝播(LTCP) ........................................................................... 3-17

3.4.1 現象の発生源............................................................................................ 3-17

3.4.2 PWRに対するLTCPの関連度評価の近の研究 ...................................... 3-18

1.1 背景

1.2 MRP-111

1.3 Allo

1.4 参考文

2 ALLOY 69

A-6


3.5 参考文書........................................................................................................... 3-23

4 薄肉SG配管のPWSCC ............................................................................................. 4-1

4.1 実験室試験 ......................................................................................................... 4-1

初期の研究 ................................................................................................. 4-2

4.1.2 B/Liを含む飽和水素水中でのU字屈曲試験.................................................. 4-2

4.1.3 B/Liを含むまたは含まない水素添加水中でのCERT試験 ........................... 4-5

5

5.

バ

の調

5.1.2.1.6 日本における一時冷却水模擬水中でのその他の調査 ............. 5-30

4.1.1

4.1.4 水素を添加した水蒸気中でのRUB試験 .......................................................... 4-6

4.1.5 試験結果のワイブル分析およびWeibayes分析............................................. 4-8

4.1.6 ワイブル分析による改善要因....................................................................... 4-15

4.1.7 Alloy 600 小亀裂時間による改善要因.................................................... 4-18

4.2 現場経験 ............................................................................................................ 4-21

4.3 参考文書........................................................................................................... 4-25

厚肉ALLOY 690材料のPWSCC................................................................................ 5-1

1 実験室試験 ......................................................................................................... 5-2

5.1.1 亀裂開始研究................................................................................................ 5-3

5.1.1.1 三菱重工(MHI)によるシミュレートした一次冷却水模擬水中での試験5-3

5.1.1.2 ミシガン大学における純超臨界水中での試験 .......................................... 5-8

5.1.2 亀裂成長率の研究..................................................................................... 5-10

5.1.2.1 PWR構成部品に直接関連しないAlloy 690材料の試験 ..................... 5-10

5.1.2.1.1 ゼネラルエレクトリックグローバルリサーチ(GE-GRC)による一時冷却水模擬水中での実現可能性の研究 ....... 5-11

5.1.2.1.2 アルゴンヌ国立研究所(ANL)による一時冷却水模擬水中での調査5-15

5.1.2.1.3MRP試験プログラムの一部としてのゼネラルエレクトリックグロールリサーチ(GE-GRC)による一時冷却水模擬水中での付随研究5-20

5.1.2.1.4 日本の東北大学による一時冷却水模擬水中での調査 ............. 5-28

5.1.2.1.5日本の原子力安全システム研究所(INSS)による一時冷却水模擬水中で査 ............................................................................................ 5-29

A-7


5.1.2.1.7 高温脱気水中でのBechtel Bettis社による調査 ....................... 5-31

5.1.2.1.8ホウ素、リチウムおよび水素を加えた超臨界水中でのウエスチングハウス社に

45

パシフィック・ノースウェスト国立研究所(PNNL)による一時冷却水模擬水

ハウ

水模

ハウ

熱の影響を受けるゾーン

5-61

5.1.2.4.4

5.2 現場

よる調査................................................................................... 5-39

5.1.2.2 冷間加工を行わないAlloy 690 CRDM材料の試験............................. 5-44

5.1.2.2.1 アルゴンヌ国立研究所(ANL)による一時冷却水模擬水中での調査5-

5.1.2.2.2

中での調査.............................................................................. 5-48

5.1.2.2.3 スウェーデンのStudsvikによる一時冷却水模擬水中での調査 5-52

5.1.2.2.4 日本のMHIによる一時冷却水模擬水中での調査..................... 5-52

5.1.2.2.5ホウ素、リチウムおよび水素を加えた超臨界水中でのウエスチングス社による調査....................................................................... 5-52

5.1.2.3 冷間加工を行った後のAlloy 690 CRDM材料の試験 ......................... 5-53

5.1.2.3.1 ゼネラルエレクトリックグローバルリサーチ(GE-GRC)による一時冷却水模擬水中での調査 ............................. 5-53

5.1.2.3.2パシフィック・ノースウェスト国立研究所(PNNL)による一時冷却擬水中での調査....................................................................... 5-58

5.1.2.3.3ホウ素、リチウムおよび水素を加えた超臨界水中でのウエスチングス社による調査....................................................................... 5-59

5.1.2.4 Alloy 690材料の溶接からの (HAZ)の試験...... 5-60

5.1.2.4.1 KAPLによる高温水中での調査............................................... 5-60

5.1.2.4.2 アルゴンヌ国立研究所(ANL)による一時冷却水模擬水中での調査5-60

5.1.2.4.3 スウェーデンのStudsvikによる一時冷却水模擬水中での調査

スペインのCIEMATによる一時冷却水模擬水中での調査....... 5-62

5.1.2.4.5 日本の東北大学による一時冷却水模擬水中での調査 ............. 5-62

経験........................................................................................................... 5-63

A-8


5.3 参考

6 議論 ........ .

6.1 PWS る

6.2 A

6.3 参考

7 結論 ........ .

図リス

文書........................................................................................................... 5-63

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

CCに対す Alloy 690の耐性...................................................................... 6-1

lloy 690腐食挙動のその他の特徴................................................................... 6-10

文書........................................................................................................... 6-10

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

ト

図2 1 [18]か

図2 2 修正されたHue温度

図2 3 修温度NX921

図2 4 修正したHu温度NX978

図2 5 アニーリ

690高下の冷

図 3 1

図 3 2

関係

図に

ついての[11]からの日本のデータ ...................................................................... 3-7

らのAlloy 690およびAlloy 600炭素溶解度図 .............................................. 2-7

y Testによる時間-温度-析出図、[18]からのAlloy 690 NX4459HG(0.06%C) .................................................................................. 2-9

正したHuey Testによる時間-温度-析出図、[18]からのAlloy 690 7H (0.01%C) .................................................................................. 2-10

ey Testによる時間-温度-析出図、[18]からのAlloy 690 0H (0.01%C) .................................................................................. 2-10

ングされたAlloy 温引張特性。示されているデータはアニーリング条件および[2]から取得した条件

間加工および高温加工の合成物です。........................................................... 2-12

[11]からの日本のNiベース合金の室温空気中での疲労データ............................ 3-5

[11]からの日本のNiベース合金の325°CのPWR模擬水中での疲労データ ......... 3-5

図 3 3 Ni合金の325°CのPWR模擬水中での疲労増大(Fen)の計算した要因と歪み率の間のについての[11]からの日本のデータ................................................................... 3-6

3 4 325°CのLWR模擬水中での疲労増大(Fen)の計算した要因と各材料の温度の間の関係

A-9


図 3 5 3実験と比較モデル予測

して25°CのPWR模擬水中でのNiベース合金腐食疲労挙動のの結果を示す[11]からの日本のデータ ............................................. 3-8

図 3 6 ステンレス鋼に比べて、高温、高クロム溶接合金の脱気水、およびAlloy

はEA

図 3 9

図 3 10

があることを示す[16]からの日本のデータ ...................................................... 3-12

図 3 11環境におけるNiベースの周期的な亀裂成長のために提示された(非常に控えめな)日本の

図 3 12 スウェーデンの研究におけるAlloy 600および690の周期的CGR挙動[15]....... 3-14

図 3 13 スウ690の一次冷却水模擬水で定義される周期的CGRしきい値 ............................ 3-14

図 3 14

図 3 18

図 4 1

690の[12]からの腐食疲労開始データ ................................................................ 3-9

図 3 7 Alloy 690とその溶接合金のPWSCC試験に適用されるようになったANLで進展する周期的な亀裂成長への環境の影響の分析方法（[13]から）............................. 3-10

図 3 8 320°CのPWR一次冷却水模擬水か脱気純水のいずれかで、熱処理されたAlloy 690 材料に見られる周期的な亀裂成長への環境の影響がない[14]。しかし負荷条件下でCを支持することは期待できない。................................................................. 3-10

320°CのPWR一次冷却水模擬水中の、一方向に冷延されたAlloy 690材料のより低い率の周期的な亀裂成長重要な環境の増大の外観.............. 3-11

PWR模擬水一次環境がAlloys 600および690の疲労亀裂成長率を試験条件の範囲を超えて5～10倍に増大させる可能性

PWR

モデル[16]からの線と既存のASME曲線を破壊する実験データの比較............ 3-13

ェーデンの研究におけるAlloy

BrownとMills[36]後のLTCPに対する破壊抵抗のカテゴリのための分類体系.. 3-19

図 3 15 各温度のRT空気および水中のBrownとMillsによって Alloy 690に対して定義されたJ-R曲線[36] ......................................................................................................... 3-20

図 3 16 各温度のRT空気および水中のBrownとMillsによって Alloy 690に対して定義されたJICおよびT値(バー内の値は溶解水素濃度を示す)..... 3-21

図 3 17 BrownとMillsによる[36]各条件下でのJ-R試験からの破面解析のAlloy 690(上部)とAlloy 600 (底部)比較 .................................................................................................. 3-22

さまざまな溶解水素内容を持つ50°Cでの付随する不均一の冷間加工を使ったAlloy690プレートのJ-R試験からのParaventiとMoshier[42]による結果 .................. 3-23

[2]からの冷間圧延 %の機能としてのVickersの硬度数。Alloy 600よりも高い加工硬化率を持つAlloy 690。 .................................................... 4-5

A-10


図 4 2 フランスで作成された360°Cの一次冷却水中での初期のAlloy 600 RUB結果の[2]からのワイブルプロット。Alloy 690(3つの温度)のWeibayes線はβ = 5.0を前提とします。 ......................................................................................... 4-9

Norringほかによって報告された365°Cの脱気水中でのRUB 結果の[2]からのワイブル図 4 2

0

°Cの

図 4 6

では障害は観察されませんでした。失敗しない資料の場合Weibayes線はβ =

図 4 7 320°C (608°F)の一次冷却水で試験される、20%の歪み前のAlloy 600MAまたはAlloy

図 4 8

図 4 9 ット

図 4 10

図 4 11 試験期間に対する式4～8ごとに表4 2および表4 4にリストされている改善要因4-19

プロット。[17]. Alloy 690(多数の温度)のWeibayes線はβ = 5.0を前提とします。4-1

図 4 4 NorringほかによるRUB試験結果の[2]からのワイブル

プロット。365°Cの脱気水中でのAlloy 600の特別な製造温度の[17]。Alloy 690 Weibayes線はβ = 5.0を前提とします。 .......................................................... 4-11

図 4 5 NorringほかによるRUB試験からの異なる管径の場合の[2]からのワイブル比較。365脱気水中での[17]。Alloy 690 Weibayes線はβ = 5.0を前提とします。 ........... 4-11

日本の一次冷却水中での340°C(644°F)のAlloy 600MA (1つの温度) CLT結果の[2]からのワイブルプロット。320°C (608°F)のAlloy 600MAまたはAlloy 600TT CLT試料、および360°C (680°F)のAlloy 690TT (1つの温度) CLT

5.0を前提とします。 ....................................................................................... 4-12

600TTのRUB試料と、360°C(680°F)以外は同一条件下で試験されるAlloy 690TTのものを、日本のデータの[2]から比較。Alloy 690TT Weibayes線はβ = 5.0を前提とします。 ....................................................................................... 4-13

Vaillantほかによって報告された360℃の一次冷却水中でのRUB 結果の[2]からのワイブルプロット。[18]. Alloy 600 RUB試料はMAおよびTT状態での4つの異なるの温度からでした。大54,000時間の曝露後障害がないことを試験されるAlloy 600 RUB試料もMAおよびTT状態での4つの異なる温度からでした。Alloy 690 Weibayes線はβ = 5.0を前提とします。 .......................................................... 4-14

フランスのFramatome ANPによる360°C(680°F)の脱気水中で試験されるSG模型の[2]からのワイブルプロ。Alloy 690TT SG模型は100,000時間の曝露後に障害を経験しませんでした。Alloy 690 Weibayes線はβ = 5.0を前提とします。 .......................................................... 4-14

Alloy 600試験用に表4 3にリストされているワイブルθおよびβ要因.............. 4-17

A-11


図 4 12 スウェーデンの[23]からのAlloys 600および690での亀裂開始の場合のRUB試験4-21

[31]からの劣化メカニズムによるAlloy 600TT SG配管修理の世界的な原因... 4-23

[31]からの劣化メカニズムによるAlloy 690TT SG配管修理の世界的な原因... 4-23

PWSCC開始[4、5]の一軸性の一定負荷研究のためのMHI試験ループ ................... 5-5

図 4 13

図 4 14

図 5 1

図 5 3 PWSCC開始[4、5]の一軸性の一定負荷研究のためのテスト試料 ..................... 5-6

図 5 4

図 5 5

図 5 7

図 5 8

軸

図 5 11

図 5 12ある)

細......................................................................................................................... 5-16

図 5 14L試料オリエンテーション)のANL試験[11～15]からのオンラインデータ ....... 5-17

図 5 15

図 5 2 PWSCC開始[4、5]の一軸性の一定負荷研究[4、5]の「アクティブ」負荷メカニズム........................................................................................................................... 5-6

適用される応力のAlloy 600MAにおけるPWSCC開始の依存および58,000時間の試験[4、5]後の690TT BMI材料における亀裂の欠如.......................................... 5-7

適用される応力のAlloy 600MAにおけるPWSCC開始の依存および73,000時間の試験[4、5]後の690TT CRDMノズル材料における亀裂の欠如 ............................ 5-7

図 5 6 脱気純粋SCW [6]中でのオーステナイト合金のCERT研究からもたらされたCGR5-8

400°C/25.4Mpa脱気純粋SCW [7]中でのテスト後のEPRI Alloy 600断面図........ 5-9

400°C/25.4 Mpa脱気純粋SCW [7]中でのテスト後のEPRI Alloy 690断面図.... 5-10

図 5 9 試料オリエンテーションに関する規約およびAlloy 690ベース材料での冷間加工の主......................................................................................................................... 5-12

図 5 10 一定応力度[9]での3000時間のGE-GRCにより試験された、冷間加工されたAlloy 690プレート(低温圧延機アニール) のOGR 応答 ............................................. 5-13

一定応力度[9]での3000時間のGE-GRCにより試験された、冷間加工されたAlloy 690プレート(高温圧延機アニール) のOGR 応答 ............................................. 5-14

GE-GRC [9]で初めに試験された冷間加工されたAlloy 690プレートのPWSCC亀裂成長の集団内の主に粒内形態(ただし粒界ファセットも......................................................................................................................... 5-15

図 5 13 ANL [11～15]で試験された冷延されたAlloy 690TTプレートからのサンプル削除の詳

冷延されたAlloy 690TTプレート(S-

冷延されたAlloy 690TTプレート(S-T試料オリエンテーション)のANL試験[11～15]からのオンラインデータ ...... 5-17

A-12


図 5 16 S-Lオリエンテーション[11～15]を持つANL 690TTプレート試料用のマクロおよびマイクロ

フラクトグラフィー............................................................................................ 5-18

図 5 17 S-Tオリエンテーション[11～15]を持つANL 690TTプレート試料用のマイクロフラクトグラフィー......................................................................................... 5-19

図 5 18 冷間加工されたAlloy 690TTプレートのCGR挙動が320～300°C [13]の範囲内で変化する温度により影響を受けなかったことを示すANL試験 5-20

図 5 19 ANLにより提供されGE-GRC [17、18]で試験された1D冷延(~26%) Alloy 690プレートにおける迅速な亀裂成長 ............................................................. 5-21

図 5 20 1D冷延Alloy 690プレートにおけるANLからの2番目に迅速なPWSCC および温度を360°Cから325°C、その後290°C [17、18]に下げるためのCGR応答の明らかな欠如 .......................................... 5-22

図 5 21 温度(c372) [17、18]を低下させるためPWSCC内の巨眼的外観の変更 .......... 5-22

溶解水素濃度[17、18]を上げるとすぐにCGRのわずかな増加を示す1-D圧延図 5 22

......................................................................................................................... 5-24

応力度 0%の1

図 5 25 20%の1D冷延Alloy 690鍛造棒如

図 5 26 図 5 23内と同一材料で、S-T

ANLプレート材料に関するGE-GRCの繰り返し試験....................................... 5-23

図 5 23 20%の1D冷延Alloy 690のGE-GRC鍛造棒[17、18]において適度に迅速なPWSCC

図 5 24 試験温度および適用される [17、18]の両方を下げながら、継続する試験の2D冷延Alloy 690のGE-GRC鍛造棒.................................................................... 5-24

[17、8]からの2番目の試料の試験中に溶解水素における主要な変化への応答の欠......................................................................................................................... 5-25

向き[17、18]の場合のテスト結果、.................. 5-25

図 5 27 GE-GRCでの1D冷延後に試験されたAlloy 690のさらなる高温からの結果。S-L向きの試料[17、18]の場合でも温度の低下に対して期待される応答を示した場合です。...................................................................................................................... 5-26

図 5 28 GE-GRC [17、18]で試験された26%の1D冷延されたANL 690プレート高分解破面解析(c372) ................................................................. 5-27

図 5 29 20%の1D冷延されたAlloy 690の鍛造棒 [17、18]からの試料の高分解破面解析5-27

A-13


図 5 30 360°C ([8]から)で20%冷間加工された690TTにおけるINSSで観察される亀裂成長の例......................................................................................................................... 5-30

図5 31 690TT ([8]からの)に対するCGR測定時の温度と冷間加工(CW)の程度への明らかな効果......................................................................................................................... 5-31

図 5 32 Bettisの研究[22、23]における一方向冷延のレベルの高まりに伴うCGRの増加5-35

冷延は引張歪(耐力強度はより低いが) [22、23]より不利 .................................. 5-35 図 5 33

0

図 5 36

図 5 39All

図 5 41

図 5 43

7

図 5 34 VIM/ESRプレート[22、23]の場合の、CGRについての試験温度と冷間加工の程度の効果......................................................................................................................... 5-36

図 5 35 50 cc/kg水素(青色の記号)対23 cc/kg(桃色の記号)でCGRの明らかな増加。Alloy 69(左)とAlloy 600 [22、23]の反対の挙動との対比(右)です。.............................. 5-37

12%の冷延[22、23]しか受けないVIM/ESR TTプレートのS-L向きの試料における主に粒間亀裂の進展 ...................................................... 5-37

図 5 37 24%の冷延[22、23]を受けるVIM/ESR TTプレートのS-T向きの試料における破面解析からの詳細 ...................................................... 5-38

図 5 38 Bettisによって測定されたAlloy 600と690 CGRの比較[22、23] ..................... 5-38

Bettisにより試行された、推定される耐力強度依存[22、23]に基づいて冷間加工されたoy 690 CGRの逆外挿法..................................................................................... 5-39

図 5 40 SCWにおける周期的負荷時の粒内亀裂からAlloy 690TT [25]の10%の冷間加工された試料における一定の負荷での粒間SCCまでの移行の例......................................................................................................................... 5-43

Alloy 600および690材料[25]の平均SCC成長率に対する応力強度要因........... 5-43

図 5 42 超臨界水および一次冷却水[25]中におけるAlloy 600 MAおよびAlloy 690 TTで測定されたCGRの比較 ............................................................................ 5-44

Alloy 690 CRDM材料[11～15]のANL試験における試料オリエンテーション . 5-45

図 5 44 CRDM材料[11～15]でのANL作業の一定負荷の短期テスト期間に報告されたCGR......................................................................................................................... 5-46

図 5 45 ANL [11～15]で試験される初のCRDM試料の破断面に明確な粒間亀裂がない5-4

A-14


図 5 46 冷間加工を行わない[11～15]CRDM材料での追加的なANL作業における2つの試験温度での一定負荷のテスト期間に報告されたCGR................................................. 5-48

図 5 47 PNNLによって2つのAlloy 690 鍛造物についてのCGR試験に使用されるCT試料のオリエンテーション。左側は温度RE

図 5 48

図 5 51 解析

図 5 52

から

図 5 55 部) 41%

図 5 56試料

図 5 57 26]

図 5 58 59

図 5 60

......................................................................................................................... 5-62

243で、右側は温度WP140 [26]です。 ............................................................ 5-49

定数Kの負荷[26]に移行時に使用する耐用性を示すPNNLデータ.................... 5-50

図 5 49 受け取ったままのTT対カーバイド修飾したSA状態[26]のAlloy 690 CGR応答を示すPNNLデータ .......................................................................... 5-50

図 5 50 受け取ったままのTT状態[26]でその他2つの温度へのAlloy 690 CGR応答を示すPNNLデータ........................................................................... 5-51

冷間加工[26]を行わない試料内の結晶の遊離が限定されるIG亀裂を示すPNNL破面

......................................................................................................................... 5-51

Alloy 690 CRDM材料[17、18]のGE試験における試料の場所......................... 5-54

図 5 53 GE-GRC [29]で試験される20%の均質な冷間加工による2つのCRDM材料の試料のうちの1つのデータ........................................................................................................... 5-55

図 5 54 20%の冷間加工されたAlloy 690 CRDM 試料[29]中での粒間亀裂形態 ........... 5-55

定数Kの状態(上部)が到着する前に亀裂を捕捉し、また24時間のホールドタイム(底[29]を持つ周期的な部分的除荷の下でもきわめて低いCGR傾向を示す傾向を示す、の均質な冷間加工によるCRDM材料の追加的試料からのデータ..................... 5-56

広範囲に及ぶ面外の二次亀裂[29]を示す41%の均質な冷間加工によるCRDM材料のからのマクロ、およびマイクロフラクトグラフィー........................................... 5-57

冷間加工を行わない(ここではS-L向きで17%)CRDM材料のPNNL試験からのデータ[......................................................................................................................... 5-58

冷間加工を行う(ここでT-L向きで30%)CRDM材料のPNNL試験からのデータ[26]5-

図 5 59 ANL [13]での試験の下のAlloy 690 HAZ試料(CF690)の詳細 ........................... 5-61

Alloy 690 HAZ材料[27]でのCGR挙動を検査するためにStudsvikで使用されているCT試料の詳細

A-15


図 6 1 PNNL [26]により作成され、不均一の冷間加工を受けるAlloy 690プレート材料において中程度から高い率を測定する可能性を示す実験室のSCC

図 6 2

表リスト

CGRデータ(2008年11月現在)の要約................................................................. 6-7

GE-GRC [18]で検査された1D冷延ANL Alloy 690プレート材料内の微細構造のバンドの形跡................................................. 6-7

図 6 3 GE-GRC [9]によるMRPプログラムでもともとは試験されたEPRIからのAlloy 690プレート材料で亀裂面に垂直な微細構造のバンド ...................................... 6-8

図 6 4 押し出しAlloy 690 CRDM材料[18]での極めて均一で、均質な微細構造............ 6-8

表 1 1 初に報告されたさまざまなPWR構成部品でのAlloy 600 PWSCCの発生....... 1-2

表 2 1 Alloys 690および600のASME仕様..................................................................... 2-1

要約 .................................................................................................................. 4-19

表 2 2 ASME化学組成要件(wt%) .................................................................................. 2-2

表 2 3 ASME室温特性................................................................................................... 2-2

表 2 4 Alloy 690およびAlloy 600のASME高温特性....................................................... 2-3

表 2 5 Sarverほかによって使用されたAlloy 690の熱の化学組成(wt%) [18] ................. 2-6

表 2 6 Alloy 690カーバイド析出[18]への加熱処理の影響........................................... 2-11

表 3 1 推定される典型的改善要因対すべての環境を考慮した[33、34] pHT.............. 3-17

表 3 2 BrownとMills [36]によって試験されたAlloys 600、690、82、および52の化学組成......................................................................................................................... 3-18

表 4 1 Alloy 690およびAlloy 52試料[2]中でのIG表面亀裂の例...................................... 4-3

表 4 2 2004年までのAlloy 690一次冷却水応力腐食試験データの要約......................... 4-7

表 4 3 Alloy 690 [2]とともに試験されたAlloy 600のワイブル分析 ............................. 4-15

表 4 4 2004年までのAlloy 690水素添加および水素添加した水蒸気に応力腐食試験データの

A-16


表 4 5 スウェーデンのAlloys 600および690[23]での亀裂開始のRUB試験................. 4-20

Alloy 690配管とともに製造される蒸気発生器の操作リスト (2008年12月現在)4-24表 4 6

表 5 5

表 5 9

表 5 1 Alloy 690でCRDM貫通部を交換されたPWR RPV上蓋の例............................... 5-1

表 5 2 相対的に厚肉のAlloy 690原子炉冷却システムのCRDM貫通部ではなく元の装置または交換部品の例5-2

表 5 3 EdF [3]によって試験されるAlloy 690 CRDMノズルの発生源と熱処理 ............. 5-3

表 5 4 MHIがPWSCCの開始[4、5]を試験する際に使用する材料の化学組成............... 5-4

MHIがPWSCCの開始[4、5]を試験する際に使用する材料の熱処理および機械的特性........................................................................................................................... 5-4

表 5 6 MHIがPWSCCの開始[4、5]を試験する際に使用する試験環境条件 .................. 5-5

表 5 7 ミシガン大学[7]による断面図のサンプルで計測された亀裂の深さの大値の要約........................................................................................................................... 5-9

表 5 8 GE-GRCによりMRPプログラムで試験された、すべてのAlloy 690材料の化学組成と機械的特性 ....................................................................................................... 5-11

ANL [11～15]で試験されたAlloy 690TTプレート材料の化学組成 ................... 5-15

表 5 10 BettisによるAlloy 690試験プログラム[22、23]の材料試験マトリックス ........ 5-33

表 5 11 Alloy 690 [22、23]に対する相対的SCC発生度のBettisによる要約.................. 5-34

表 5 12 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 690材料......................................................................................................................... 5-40

表 5 13 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 690 材料の化学組成 ................................................................................................ 5-40

表 5 14 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 690 材料の引張特性(材料ベンダーによる報告どおり)............................................. 5-41

表 5 15 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 690 材料の試験的鍛造後の引張特性 ........................................................................ 5-41

表 5 16 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 600 制御サンプルとAlloy 690プレート材料の詳細な結果 ...................................... 5-42

表 5 17 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 600 制御サンプルと「受け取ったまま」のAlloy 690 CRDM材料の詳細な結果..... 5-52

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表 5 18 ウエスチングハウス社が超臨界水[25]中でのCGRを試験する際に使用したAlloy 600 制御サンプルと冷間加工されたAlloy 690 CRDM材料の詳細な結果.................. 5-60

表 6 1 冷間加工を行わないAlloy 690 CRDM材料についてPWR一次冷却水模擬水中での CGRテストの結果の要約 (状態:2008年12月) .................................................... 6-3

表 6 2 冷間加工を行うAlloy 690 CRDM材料についてPWR一次冷却水模擬水中での CGRテストの結果の要約 (状態:2008年12月) .................................................... 6-4

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소재 신뢰도 프로그램: 가압수형 원자로에서 1차 냉각수 응력 부식 균열에 대한 Alloy 690의 저항력(MRP-258)

1019086

최종 보고서, 2009년 8월

연구 소산물에 대한 설명

단련된 Alloy 600과 용접 금속(Alloy 182 및 Alloy 82)은 원래 다양한 공격적 환경에서 전면 부식에 대한 소재 고유의 저항력 및 저합금 탄소강과 매우 비슷한 열 팽창 계수 때문에 가압수형 원자로(PWR)에 사용되었다. 지난 30년에 걸쳐 PWR 1차 냉각수에서의 입계 응력 부식 균열(PWSCC)은 때때로 상대적으로 긴 잠복기간 이후에도 수많은 Alloy 600 구성품 및 관련된 용접에서 관찰되었다. PWSCC의 발생 원인은 상당히 긴 정지시간과 대체전력 비용에 있었다.


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구성품의 수리 및 대체재로는 일반적으로 단련된 Alloy 690소재 및 호환 가능한 용접 금속(Alloy 152 및 Alloy 52 또는 52M)을 사용해 왔으며 이들 소재 및 금속은 실험실 실험에서 PWSCC에 대한 저항력이 큰 것으로 나타났으며 최대 20년의 기간에 걸친 원자로 가동 시 균열이 발생하지 않았다. 풀어야 할 과제는 수리되거나 교체된 구성품에 대한 향후의 검사요건 개발을 위한 확실한 기술적 근거를 제공하기 위해 부식에 의한 노화 열화와 관련된 이들 금속의 수명을 정량화하는 것이다. 본 문서는 2004년 발간된 EPRI 보고서 1009801 (MRP-111)를 새로이 수정한 것이며 최근에 입수된Alloy 690 기본재의 PWSCC 거동을 다룬 정보를 고려한 확대 개정판이다. 본 문서에서는 Alloy 152 및52 용접 금속의 성능에 대해서 고려하지 않지만, 이에 대해서는 나중에 따로 보고할 예정이다. 또한 이 보고서에서는 후벽 Alloy 690 소재의 부식 피로 및 저온 균열 전파에 대해 부분적으로 다룬다. 2004년 보고서(MRP-111)에서 이들 주제에 대한 내용이 부족한 것으로 확인되었기 때문이다.

접근방법 최초 MRP-111 보고서를 토대로 이뤄진 지난 2년 동안의 수많은 실험실 시험은 PWR 환경에서의 부식 저항력에 관계된 다양한 시험 조건에서 단련된 Alloy 690 소재를 이용해 수행된 것으로서 주로 PWSCC에 초점을 맞춰 검토하였다. 가능한 경우 기존 실험실 시험 자료를 평가하여 Alloy 600과 관련된 Alloy 690의 개선 계수를 추정하였다. 또한 PWR 에서 Alloy 690를 사용한 경험은 실험실 결과를 강하게 뒷받침하는 것으로 보고되고 있다.

결과 및 결론 단련된Alloy 690은 PWR에서 Alloy 600을 대체할 수 있는 적절하고 매우 부식 저항력이 높은 소재라는 결론에 도달하였다. 그러나 PWSCC와 관련하여 나타난 일부 구체적인 지식의 부족을 조사하기 위해 추가적인 시험이 필요하다. 여기에는 비균질 냉간 가공(특히 단방향성 냉간 압연/인장변형)이 균열 성장에 대한 저항력에 미치는 유해한 영향 및 용접 후 열의 영향을 받은 구역의 균열 감수성이 증가할 가능성이 포함된다. 균열의 발생과 관련하여 Alloy 600과 비교해 현재 40~100배의 상대적인 개선 가능성이 존재하지만, 시험 기간 동안 거의 모든 Alloy 690 시료에서 PWSCC가 없었기 때문에 이것은 명백히 보수적인 수치이다. (예컨대 원자로 압력용기 상단헤드 관통을 위한 압출


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파이프 같은) 원형적 후벽 Alloy 690 소재는 모의 환경에서 1차 냉각수에 이미 존재했던 피로 균열로 인한 SCC 성장에 대한 저항력이 매우 높다. 측정된 균열 성장률은 매우 낮아서 (< 5E-9 mm/s 또는 0.15 mm/년) 공학적 유의성이 없으며 Alloy 600MA와 관련해 Alloy 690 CRDM의 상대적인 개선 계수는 100배를 훌쩍 상회할 것으로 여겨진다. 그러나 일반적으로 발전소의 구성요소들을 대표하지 않는 것으로 여겨지는 Alloy 690 판재 및 각재에 비균질 냉간 가공을 도입한 후 일부 시험에서는 일정한 하중을 받은 상태에서 상당한 입계균열이 관찰되었다. 이러한 결과가 장기적인 PWR 가동에 관계될 가능성이 없음을 확인하기 위해 이러한 민감성의 한계를 계속 조사하고 있다. 지금까지 대체 적용에서 Alloy 690의 입계 부식 열화는 전혀 관찰된 바 없다. PWR 1차 냉각수에 노출된 검증된 소재의 사용 경험은 고려된 구성품의 종류에 따라 대략 10년에서 20년의 범위에 이른다.

적용, 가치 및 이용 본 보고서는 PWR 1차 냉각수에서 발생한 모든Alloy 600 균열의 측면에 관계된 전력회사 엔지니어 및 과학자, 그리고 검증체계를 개발하고 구성품 대체에 대한 결정을 내리며 발전소의 노화 문제를 다루는 일에 관여하는 사람들에게 특히 중요할 것이다. 본 보고서는 기존의 발전소에 있는 후벽 Alloy 600 구성품에 영향을 미치는 PWSCC의 발생이 증가하는 문제를 해결하기 위해 업계가 채택한 해결책이 규제 측면의 적합성을 획득하는데 있어서 분명한 가치가 있을 것이다.

전력회사, 원자로 업체 및 공학/연구 기관을 포함한 지속적인 종합 프로그램의 일환으로 본 보고서는 또한 니켈 기반 합금의 부식 열화가 유효 수명을 제한하지 않고 대체 구성품과 새로운 원자로 모두에 있어 개선된 설계로 얻어지는 장점을 극대화하는데 일조할 것이다.

EPRI의 관점 본 보고서에서는 Alloy 690의 PWSCC 저항력에 관한 현재의 지식기반을 설명한다. 데이터는 종전의 Alloy 600와 비교하여 PWSCC에 대한 Alloy 690의 저항력이 상당히 증가하는 것으로 보여주고 있다. 이러한 결과는 고무적이고 추가적인 연구에서도 계속 유지될 것으로 예상되긴 하지만, Alloy 690의 우수성을 약간 훼손할 수 있는 냉간 가공 같은 요인들이 여전히 존재한다. 진행되고 있는 연구와 계획된 연구는 PWSCC에 대한 Alloy 690의 저항력을 더욱 분명히 정의하고 취약성들을 규명하기 위해 이들 요인에


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초점을 두고 있다. 이러한 추가 연구가 완료될 때까지 본 보고서에서 제시하는 결과들은 오직 참고를 위한 것으로 Alloy 690으로 구성된 구성품들의 수명을 예측하는데 사용되어서는 안 된다.

키워드 Alloy 600 Alloy 690 PWSCC 소재 열화 RPV 관통


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목차

1 머리말...................................................................................................................... 1-1

1.1 배경 .................................................................................................................... 1-1

1.2 MRP-111 개정의 목적 및 범위 ........................................................................... 1-2

1.3 Alloy 600 PWSCC의 원인 ................................................................................. 1-4

1.4 참고문헌............................................................................................................. 1-5

2 ALLOY 690의 속성 및 야금 ...................................................................................... 2-1

2.1 소재 사양 ........................................................................................................... 2-1

2.1.1 박벽 Alloy 690 SG 튜브에 사용되는 일반적인 PWR 사양 .......................... 2-3

2.1.2 후벽 Alloy 690 구성요소의 사양 및 제조 ................................................... 2-4

2.2 Alloy 690의 상태도 ............................................................................................. 2-4

2.3 탄소 용해도 및 동적 변형 노화 ........................................................................... 2-5

2.4 입계 탄화물 석출 및 예민화................................................................................ 2-8

2.5 고온 노출의 영향 .............................................................................................. 2-12

2.6 참고문헌........................................................................................................... 2-13

3 PWSCC를 제외한 ALLOY 690 의 부식 거동 ............................................................ 3-1

3.1 1차 냉각수에서 실시한 전면 부식 시험 .............................................................. 3-1

3.1.1 Sedricks 등(1979)에 의한 SG 튜브 ........................................................... 3-1

3.1.2 Smith 등(1985)에 의한 SG 튜브 ................................................................ 3-2

3.1.3 Yonezawa 등(1985)에 의한 SG 튜브 ........................................................ 3-2

3.1.4 Esposito 등(1991) ...................................................................................... 3-2

3.1.5 PWSCC 기제와 관련된 합금 산화 연구...................................................... 3-3

3.2 1차 냉각수에서의 부식 피로 시험 ...................................................................... 3-4

3.3 2차 냉각수에서의 부식 거동............................................................................. 3-15


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3.4 저온 균열 전파(LTCP) ...................................................................................... 3-17

3.4.1 현상의 원인 .............................................................................................. 3-17

3.4.2 LTCP와 PWR의 관련 가능성을 평가한 최근 연구들 ................................ 3-18

3.5 참고문헌........................................................................................................... 3-23

4 박벽 SG 튜브의 PWSCC.......................................................................................... 4-1

4.1 실험실 시험 ........................................................................................................ 4-1

4.1.1 초기 연구 .................................................................................................... 4-2

4.1.2 B/Li가 있는 수소 포화 냉각수에서의 단일 U-Bend 시험............................. 4-2

4.1.3 B/Li가 있거나 없는 수소화 냉각수에서의 CERT 시험 ................................ 4-5

4.1.4 수소화 증기에서의 RUB 시험..................................................................... 4-6

4.1.5 시험 결과의 Weibull 및 Weibayes 분석 ...................................................... 4-8

4.1.6 Weibull 분석에 의한 개선 계수 ................................................................. 4-15

4.1.7 최소 Alloy 600 균열 시간을 이용한 개선 계수 .......................................... 4-18

4.2 현장 경험.......................................................................................................... 4-21

4.3 참고문헌........................................................................................................... 4-25

5 후벽 ALLOY 690 소재의 PWSCC............................................................................. 5-1

5.1 실험실 시험 ........................................................................................................ 5-2

5.1.1 균열 발생 연구 ........................................................................................... 5-3

5.1.1.1 미쓰비시 중공업(MHI)의 모의 1차 냉각수 시험 .................................. 5-3

5.1.1.2 미시건 대학교의 순수 초임계수 시험.................................................. 5-8

5.1.2 균열 성장률 연구 ...................................................................................... 5-10

5.1.2.1 PWR 구성요소와 직접 관련되지 않은 Alloy 690 소재의 시험........... 5-10

5.1.2.1.1 제너럴 일렉트릭 글로벌 리서치(GE-GRC)의 모의 1차 냉각수 타당성 연구 ........................................................................................ 5-11

5.1.2.1.2 아르곤 국립 연구소(ANL)의 모의 1차 냉각수 조사 ................ 5-15

5.1.2.1.3 MRP 테스트 프로그램의 일환인 제너럴 일렉트릭 글로벌 리서치(GE-GRC)의 모의 1차 냉각수 추가 연구 ........................................ 5-20

5.1.2.1.4 일본 토호쿠 대학교의 모의 1차 냉각수 조사 .......................... 5-28

5.1.2.1.5 일본 원자력 안전 시스템 기술원(INSS)의 모의 1차 냉각수 조사5-29


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5.1.2.1.6 일본의 모의 1차 냉각수 심층 조사 .......................................... 5-30

5.1.2.1.7 Bechtel-Bettis의 탈기 고온수 조사.......................................... 5-31

5.1.2.1.8 웨스팅하우스의 붕소, 리튬 및 수소를 첨가한 초임계수 조사 . 5-39

5.1.2.2 의도적 냉간 가공 없는 Alloy 690 CRDM 소재 시험 .......................... 5-44


5.1.2.2.2 퍼시픽 노스웨스트 국립 연구소(ANL)의 모의 1차 냉각수 조사 5-48

5.1.2.2.3 스웨덴 Studsvik의 모의 1차 냉각수 조사 ................................ 5-52

5.1.2.2.4 일본 MHI의 모의 1차 냉각수 조사........................................... 5-52

5.1.2.2.5 웨스팅하우스의 붕소, 리튬 및 수소를 첨가한 초임계수 조사 . 5-52

5.1.2.3 의도적 냉간 가공 후 Alloy 690 CRDM 소재 시험.............................. 5-53

5.1.2.3.1 제너럴 일렉트릭 글로벌 리서치 (GE-GRC)의 모의 1차 냉각수에서의 타당성 연구 ............................................................................ 5-53

5.1.2.3.2 태평양 노스웨스트 국립 연구소(ANL)에서의 모의 1차 냉각수 조사 ................................................................................................ 5-58

5.1.2.3.3 Westinghouse의 브론, 리튬 및 수소를 첨가한 초임계수 조사 5-59

5.1.2.4 Alloy 690 소재의 용접으로 인한 열영향부(HAZ) 시험 ...................... 5-60

5.1.2.4.1 KAPL의 고온 냉각수 조사....................................................... 5-60


5.1.2.4.3 스웨덴 Studsvik의 모의 1차 냉각수 조사 ................................ 5-61

5.1.2.4.4 스페인 CIEMAT의 모의 1차 냉각수 조사 ................................ 5-62

5.1.2.4.5 일본 토호쿠 대학교의 모의 1차 냉각수 조사........................... 5-62

5.2 현장 경험.......................................................................................................... 5-63

5.3 참고문헌........................................................................................................... 5-63

6 고찰.......................................................................................................................... 6-1

6.1 Alloy 690의 PWSCC 저항력............................................................................... 6-1

6.2 Alloy 690 부식 거동의 기타 측면 ...................................................................... 6-10

6.3 참고문헌........................................................................................................... 6-10

7 결론.......................................................................................................................... 7-1


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그림 목록

그림 2.1 [18]에서 발췌한 Alloy 690 및 Alloy 600의 탄소 용해도 도표............................ 2-7

그림 2.2 [18]에서 발췌한 수정된 질산용액 무게감량 부식시험, Alloy 690 Heat NX4459HG (0.06%C)의 시간 – 온도-예민화 도표 ................................................................ 2-9

그림 2.3 [18]에서 발췌한 수정된 질산용액 무게감량 부식시험, Alloy 690 Heat NX9217H (0.01%C)의 시간 – 온도-예민화 도표 .............................................................. 2-10

그림 2.4 [18]에서 발췌한 수정된 질산용액 무게감량 부식시험, Alloy 690 Heat NX9780H (0.01%C)의 시간 – 온도-예민화 도표 .............................................................. 2-10

그림 2.5 소둔된 Alloy 690의 고온 인장 특성. 제시된 데이터는 소둔 조건 및 [2]에서 발췌한 냉간 가공 및 열간 가공 산물의 복합물이다. ............................................................ 2-12

그림 3.1. [11]에서 발췌한 실온 공기에서의 Ni기 합금에 대한 일본의 피로 데이터........ 3-5

그림 3.2. [11]에서 발췌한 325°C의 모의 PWR 냉각수에서 실시한Ni기 합금에 대한 일본의 피로 데이터 ................................................................................................................ 3-5

그림 3.3. [11]에서 발췌한 계산된 환경적 피로 증가율(Fen)과 325°C의 모의 PWR 냉각수에서의 Ni 합금 변형률 간의 관계를 다룬 일본 데이터.................................................... 3-6

그림 3.4. [11]에서 발췌한 계산된 환경적 피로 증가율(Fen)과 325°C의 모의 LWR 냉각수에서의 다양한 금속 온도 간의 관계를 다룬 일본 데이터................................................ 3-7

그림 3.5 [11]에서 발췌한 325°C의 모의 PWR에서의 Ni기 합금의 부식 피로 거동에 대한 모델 예측 결과를 실험 결과와 비교한 일본 데이터 ................................................... 3-8

그림 3.6 [12]에서 발췌한 고크롬 용접 합금 및 Alloy 690과 스테인리스 강의 탈기 고온수 내 부식 피로 발생 데이터 비교........................................................................................ 3-9

그림 3.7 ANL에서 개발되어 현재 Alloy 690 및 용접 합금에 적용되는 주기적 균열 성장에 대한 환경적 영향 분석 방법([13]에서 발췌) .............................................................. 3-10

그림 3.8 열 처리된 Alloy 690의 경우 모의 PWR 1차 냉각수 혹은 320°C [14] 의 순수 탈기수에서 주기적 균열 성장이 환경적으로 개선되지 않았으나 하중 조건에서는 EAC에 유리할 것으로 예상되지 않는다 ................................................................................... 3-10

그림. 3.9 320°C 의 모의 PWR 1차 냉각수에서 단방향 냉간 압연된 Alloy 690 소재(청색 점)에서 상대적으로 낮은 속도로 진행되는 주기적 균열 성장이 환경적으로 상당히 진전된 모습([14]에서 발췌) .......................................................................................... 3-11


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그림. 3.10 모의 PWR 1차 냉각수 환경이 Alloys 600 및690의 피로 균열 성장률을 시험 조건의 범위에서 5배 내지 10배 증가시킬 수 있음을 나타내는 일본 데이터([16]에서 발췌)3-12

그림 3.11 PWR 환경에서 Ni기 합금의 주기적 균열 성장에 대한 제시된 (매우 보수적인) 일본측 모델의 선과 기존의 ASME 곡선을 벗어나는 실험적 데이터의 비교................. 3-13

그림 3.12 스웨덴 연구[15]에서 나타난 Alloys 600 및 690의 주기적 CGR 거동 ........... 3-14

그림 3.13 스웨덴 연구[15]에서 나타난 Alloy 690에 대한 모의 1차 냉각수에서 결정된 주기적 CGR 한계값 ..................................................................................................... 3-14

그림 3.14 Brown 및 Mills [36]에 따른 LTCP의 파괴 저항력 범주에 대한 분류 체계..... 3-19

그림 3.15 다양한 온도[36]의 RT 공기 및 냉각수에서의 Alloy 690에 대해 Brown 및 Mills가 결정한 J-R곡선. .......................................................................................................... 3-20

그림 3.16다양한 온도의 공기 및 냉각수에서의 Alloy 690에 대해Brown 및 Mills [36]가 결정한 JIC 및 T 값(막대 안의 값이 용존 수소 농도를 나타낸다). ...................................... 3-21

그림 3.17 Brown 및 Mills[36]가 실시한Alloy 690 (상단) 및 Alloy 600 (하단)을 대상으로 한 다양한 조건에서의 J-R 시험을 통한 파면 비교 ............................................................ 3-22

그림 3.18 다양한 용존 수소 함량의 50°C 냉각수에서 추가의 비균질 냉간 가공을 이용한 Alloy 690 판재의 J-R 시험에서 도출된 Paraventi 및Moshier [42]의 결과 ................. 3-23

그림 4.1 [2]에서 발췌한 냉간압연률에 비례한 비커스 경도 수치. Alloy 690의 가공 경화율은 Alloy 600보다 높다...................................................................................................... 4-5

그림 4.2 프랑스에서 얻은 360°C의 1차 냉각수에서 나타난 초기 Alloy 600 RUB 결과에 대한 Weibull 도표([2]에서 발췌). Alloy 690(3가지 가열 온도) Weibayes 선은 β = 5.0으로 추정된다............................................................................................................. 4-9

그림 4.3 Norring 등 [17]이 보고한365°C의 탈기수에서의 RUB 시험 결과에 대한 Weibull 도표([2]에서 발췌).Alloy 690(여러 가열 온도) Weibayes 선은 β = 5.0으로 추정된다.......................................................................................................................... 4-10

그림 4.4 365°C의 탈기수에서의 Alloy 600의 특수한 산물의 열에 대한 Norring 등[17]에 의한 RUB 시험 결과에 대한 Weibull 도표([2]에서 발췌). Alloy 690 Weibayes 선은 β = 5.0으로 추정된다........................................................................................................... 4-11

그림 4.5 365°C의 탈기수에서 Norring 등[17]이 실시한 RUB 시험을 통한 상이한 튜브 변수에 대한 Weibull 비교 ([2]에서 발췌). Alloy 690 Weibayes 선은 β = 5.0으로 추정된다.......................................................................................................................... 4-11


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그림 4.6. 1차 냉각수340°C (644°F)에서 일본 Alloy 600MA (열 하나) CLT 결과에 대한 Weibull 도표([2]). 320°C (608°F)에서 시험된 Alloy 600MA 혹은 Alloy 600TT CLT 시료와 360°C (680°F)에서 시험된 Alloy 690TT (하나의 가열 온도) CLT 시료 모두 파괴가 관찰되지 않았다. 파괴가 없는 시료에 대한 Weibayes 선은 β = 5.0으로 추정한다. ...... 4-12

그림 4.7 320°C (608°F)의 1차 냉각수에서 시험된 Alloy 600MA 또는 Alloy 600TT의 20% 사전 변형된 RUB 시료에 대한 일본 데이터와 동일한 조건의 360°C (680°F)에서 시험된 Alloy 690TT의 데이터 비교. Alloy 690TT Weibayes 선은 β = 5.0로 추정된다. ....... 4-13

그림 4.8 Vaillant 등[18]이 보고한 360°C의 1차 냉각수에서 얻은 RUB 결과의 Weibull 선도([2]에서 발췌). Alloy 600 RUB 시료는 MA 및 TT 조건의 네 가지 다른 온도에서 얻은 것이다. MA 및 TT 조건에서 네 가지 다른 온도로 얻은 Alloy 690 RUB 시료는 최대 54,000시간의 노출 이후 어떤 파괴도 일어나지 않았다. Alloy 690 Weibayes 선은 β = 5.0으로 추정된다. ............................................................................................ 4-14

그림 4 9 프랑스 Framatome ANP 가 실시한680°F 탈기수에서 시험된 SG 모형의 Weibull 도표([2]에서 발췌). Alloy 690TT SG 모형에서는 100,000 시간의 노출 이후에도 파괴가 일어나지 않았다. Alloy 690 Weibayes 선은 β = 5.0으로 추정한다. ................ 4-14

그림 4 10 표 4.3의 Alloy 600 시험에 대한 Weibull θ 및 β 계수 ................................ 4-17

그림 4 11 시험 지속시간에 대비한 등식4-8 에 따른 표 4.2 및 표4.4의 개선 계수........ 4-19

그림 4 12 Alloys 600 및 690의 균열 발생에 대한 스웨덴의 RUB 시험([23]에서 발췌). 4-21

그림 4.13 열화 메커니즘에 의한 Alloy 600TT SG 튜브 수리의 전세계적 원인([31]에서 발췌)......................................................................................................................... 4-23

그림 4.14 열화 메커니즘에 의한 Alloy 690TT SG 튜브 수리의 전세계적 원인([31]에서 발췌)......................................................................................................................... 4-23

그림 5.1 PWSCC 발생에 관한 단축 정하중 연구를 위한 MHI 시험 루프[4, 5]................ 5-5

그림 5.2 PWSCC 발생에 관한 단축 정하중 연구를 위한 “ 활성” 하중 메커니즘[4, 5]. 5-6

그림 5.3 PWSCC 발생에 관한 단축 정하중 연구를 위한 시료[4, 5] ............................... 5-6

그림 5.4 적용된 응력에 대한 Alloy 600MA의 PWSCC 발생 의존도 및 58,000시간 시험 후690TT BMI 소재에 균열 없음[4,5].................................................................................. 5-7

그림 5.5 적용된 응력에 대한 Alloy 600MA의 PWSCC 발생 의존도 및 73,000시간 시험 후690TT CRDM 노즐 소재에 균열 없음[4,5] ..................................................................... 5-7

그림 5.6 순수 탈기 SCW에서의 다양한 오스테나이트 합금에 대한 CERT 연구에서 도출된 CGR[6] ............................................................................................................... 5-8


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그림 5.7 400°C/25.4Mpa 탈기된 순수 SCW에서의 시험 후 EPRI Alloy 600의 횡단면[7].5-9

그림 5.8 400°C/25.4 MPa 탈기된 순수 SCW에서의 시험 후 EPRI Alloy 690 의 횡단면[7]5-10

그림 5.9 Alloy 690 기재에서 시료의 방향성 및 냉간 가공 주축에 관한 도식................ 5-12

그림 5.10 일정 응력도에서 시험된 3000시간 이상 동안 GE-GRC에서 (저온 mill anneal을 이용하여) 냉간 가공된 Alloy 690 판재의 CGR 반응[9] ..................................... 5-13

그림 5.11 일정 응력도에서3000시간 이상 동안 GE-GRC에서 시험된 (고온 mill anneal을 이용하여) 냉간 가공된 Alloy 690 판재의 CGR반응[9] ...................................... 5-14

그림 5.12 GE-GRC에서 초기에 시험된 냉간 가공 Alloy 690 판재에서 발생한 PWSCC 균열 성장 띠 내에서 우세한 입내 형태 [9]......................................................................... 5-15

그림 5.13 ANL에서 시험된 냉간 압연된 Alloy 690TT 판재에서 시료 추출 상세도[11 - 15]5-16

그림 5.14 냉간 압연된 Alloy 690TT 판재(S-L 시료 방향)의 ANL 시험 [11 - 15]에서 도출한 온라인 데이터 .............................................................................................................. 5-17

그림 5.15 냉간 압연된 Alloy 690TT 판재(S-T 시료 방향)의 ANL 시험 [11 - 15]에서 도출한 온라인 데이터 .............................................................................................................. 5-17

그림 5.16 S-L 방향성을 가진 ANL 690TT 판재 시료의 거시적 및 미시적 파면[11 - 15]5-18

그림 5.17 S-T 방향성을 가진 ANL 690TT 판재 시료의 거시적 파면[11 - 15] ............... 5-19

그림 5.18 냉간 가공 Alloy 690TT 판재의 CGR 거동이 320°C 내지 300°C의 온도 범위 내에서 변화하는 기온에 영향을 받지 않았음을 나타내는ANL 시험[13] ...................... 5-20

그림 5.19 ANL이 공급하고 GE-GRC에서 시험된 ID 냉간 압연된 (~26%) Alloy 690 판재의 급속한 균열 성장[17, 18] .................................................................................. 5-21

그림 5.20 ANL의 1D 냉간 압연된 Alloy 690 판재에서 발생한 두 번째 급속한 PWSCC 및 360°C에서 325°C로의 온도 감소에 대한 CGR 반응 및 그 후290°C로의 온도 감소에 대한 CGR 반응 없음 [17, 18].................................................................................... 5-22

그림 5.21 온도 감소 즉시 PWSCC 구역에 나타난 거시적 외관(c372) [17, 18] ............ 5-22

그림 5.22 용존 수소 농도 증가 시 CGR의 미미한 증가를 나타내는 1-D 압연된 ANL 판재에 대한 GE-GRC에서의 반복된 시험[17, 18] ................................................................ 5-23

그림 5.23 20% 1D 냉간 압연된 Alloy 690 GE-GRC 압연봉에서 나타나는 완만하게 빠른 PWSCC......................................................................................................................... 5-24

그림 5.24 온도 및 적용된 응력강도 모두의 감소에 따른 20% 1D 냉간 압연된 Alloy 690 GE-GRC 압연봉에 대한 계속된 시험............................................................................... 5-24


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그림 5.25 20% 1D 냉간 압연된 Alloy 690 압연봉에서 추출한 또 하나의 시료에 대한 시험 중 용해된 수소에서 나타난 상당한 변화에 대한 반응 없음 [17,18] ....................... 5-25

그림 5.26 그림 5.23의 경우와 유사한 소재를 시험한 결과, 그러나 이번에는 S-T 방향성에서의 경우[17,18] ....................................................................................................... 5-25

그림 5.27 GE-GRC에서 1-D 냉간 압연 후 시험된 Alloy 690의 추가 가열 온도에서 도출한 결과. 이번에는 S-L 방향 시료도 온도 감소에 대한 예상된 반응을 나타내고 있다[17, 18].5-26

그림 5.28 GE-GRC 에서 시험된 26% 1D 냉간 압연된 ANL 690 판재에 대한 고해상도 파면(c372) [17, 18]........................................................................................... 5-27

그림 5.29 압연봉의20% 1D 냉간 압연된 Alloy 690 시료에 대한 고해상도 파면 [17, 18]5-27

그림 5.30 INSS가 관찰한 360°C의 20% 냉간 가공된 690T에서의 균열 성장의 예([8])5-30

그림 5.31 690TT의 경우 측정된 CGR에 미치는 냉간 가공(CW)의 온도 및 정도의 명백한 영향([8])............................................................................................................ 5-31

그림 5.32 Bettis 연구에 나타난 단방향 냉간 압연 정도의 증가에 따른 CGR의 증가[22,23]5-35

그림 5.33냉간 압연은 (상대적으로 낮은 항복강도에도 불구하고) 미리 변형한 인장보다 더 해롭다[22, 23] .................................................................................................. 5-35

그림 5.34 VIM/ESR 판재의 경우 CGR에 영향을 미치는 냉간 가공의 시험 온도 및 정도의 영향[22, 23] ...................................................................................................... 5-36

그림 5.35 Alloy 690(좌측)의 경우50 cc/kg 수소(파란 기호) 및 23 cc/kg에서 나타나는 CGR의 명백한 증가 및 Alloy 600에서의 반대적인 거동과의 비교(우측) [22, 23].......... 5-37

그림 5.36 불과 12%의 냉간 압연을 받은 VIM/ESR TT 판재의 S-L-방향 시료에서 발생하는 우세하게 분포된 입계 균열 진전[22, 23] .......................................................... 5-37

그림 5.37 24%의 냉간 압연을 받은 VIM/ESR 판재의 S-T 방향 시료에 대한 상세한 파면[22, 23]......................................................................................................................... 5-38

그림 5.38 Bettis가 측정한 Alloy 600 및 690 CGR의 비교 [22, 23] ............................... 5-38

그림 5.39 Bettis가 수행한 추정된 항복 강도 의존도에 따른 냉간 가공된 690 CGR에 대한 후외삽법 시도................................................................................................... 5-39

그림 5.40 주기적 하중을 받는 동안에 발생한 입내 균열에서Alloy 690TT의 10% 냉간 가공된 시료의 정하중에서 나타난 입계SCC로 이동시킨 예[25] .................................. 5-43

그림 5.41 Alloy 600 및 690에 대한 평균SCC 성장률 대 응력강도[25] ......................... 5-43

그림 5.42 초임계수 및 1차 냉각수에서의 Alloy 600 MA 및 Alloy 690 TT에 대해 측정된 CGR 비교 .................................................................................................................. 5-44


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그림 5.43 Alloy 690 CRDM 소재의 ANL 시험에서 나타난 시료 방향성[11 - 15] .......... 5-45

그림 5.44 CRDM 소재에 대한 ANL 연구에서 나타난 정하중을 간단한 시험 기간 동안의 보고된 CGR[11 - 15] .................................................................................................... 5-46

그림 5.45 ANL에서 시험된 최초의 CRDM 시료의 파면 위에 분명한 입계 균열 없음[11 - 15]......................................................................................................................... 5-47

그림 5.46 의도적 냉간 가공 없는 CRDM 소재에 대한 ANL의 추가연구에서 두 개 시험 온도의 정하중 시험 기간 동안 보고된 CGR[11 - 15] .................................................... 5-48

그림 5.47 단조된 두 개의 Alloy 690에 대한 CGR 시험에서 PNNL이 사용한 CT 시료의 방향: 좌측은 heat RE243, 우측은 heat WP140 [26] .................................................. 5-49

그림 5.48 일정 K 하중으로 전환할 때 사용되는 인내력을 나타내는 PNNL 데이터...... 5-50

그림 5.49 카바이드 수정된 SA 조건에 대비하여 무가공 TT에 대한 Alloy 690 CGR 반응을 나타내는 PNNL 데이터[26] .............................................................................. 5-50

그림 5.50 받은 그대로의 무가공 TT 조건에서 나타난 두 번의 추가적인 열처리에 책임이 있는 Alloy 690 CGR를 나타내는 PNNL 데이터[26] .................................................. 5-51

그림 5.51 냉간 가공 없이 시료의 격리된 입자들로 제한된 IG 균열을 나타내는 PNNL 파면[26]......................................................................................................................... 5-51

그림 5.52 GE의Alloy 690 CRDM 소재 시험에서 시료 위치[17,18] .............................. 5-54

그림 5.53 GE-GRC에서 시험된 20% 균질 냉간 가공을 한 CRDM 소재의 시료 두 개 중 한 개에서 얻은 데이터[29] ................................................................................................ 5-55

그림 5.54 20% 냉간 가공된 Alloy 690 CRDM 시료의 입계 균열 파면[29].................... 5-55

그림 5.55 상수 K 조건에 도달하기 전(상단), 그리고 유지시간이 24시간인 주기적인 부분 무부하에서 매우 낮은 CGR에 도달하기 전(하단)에 균열 정지 경향성을 나타내는 41%의 균질 냉간 가공 처리된 CRDM 소재의 추가 시료에서 얻은 데이터[29]............. 5-56

그림 5.56 광범위한 면외 2차 균열을 나타내는 41% 균질 냉간 가공을 한 CRDM 소재 시료에서 도출된 거시적 및 미시적 파면[29] .................................................................... 5-57

그림 5.57 계획된 냉간 가공을 받은 CRDM 소재의 PNNL 시험에서 얻은 데이터(S-L 방향성의 17%)[26]........................................................................................................... 5-58

그림 5.58 계획된 냉간 가공을 받은 CRDM 소재의 PNNL 시험에서 얻은 데이터(T-L 방향성의 30%)[26]........................................................................................................... 5-59

그림 5.59 ANL에서 시험된 Alloy 690 HAZ 시료(CF690)에 대한 상세 정보 ................. 5-61


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그림 5.60 Alloy 690 HAZ 소재의 CGR 거동을 연구하기 위해 Studsvik에서 사용 중인 CT 시료의 상세 정보[27].................................................................................................... 5-62

그림 6.1비균질 냉간 가공을 거친 Alloy 690 판재에서 나타나는 적정 비율 또는 높은 비율을 측정할 가능성을 나타내는PNNL이 작성한 (2008년 11월 현재) 실험실 SCC CGR 데이터 요약 .................................................................................................................... 6-7

그림 6.2 GE-GRC에서 연구한1-D 냉간 압연 ANL Alloy 690 판재의 일부 영역에서 형성된 미세구조의 띠에 대한 증거[18]........................................................................... 6-7

그림 6.3 본래GE-GRC의 MRP 프로그램을 위해 시험된 EPRI의 Alloy 690 판재에서의 균열면에서 수직으로 나타나는 미세구조의 띠[9] ............................................... 6-8

그림 6.4 압출된Alloy 690 CRDM 소재에서 나타나는 매우 균일하고 균질한 미세구조[18]6-8

표 목록

표 1.1 다양한 PWR 구성요소 항목에 대해 최초로 보고된 Alloy 600 PWSCC 발생....... 1-2

표 2.1 Alloys 690 및 600의 ASME 사양 ......................................................................... 2-1

표 2.2 ASME 화학적 조성 요구사항(wt%)...................................................................... 2-2

표 2.3 ASME 실온 속성.................................................................................................. 2-2

표 2.4 Alloy 690 및 Alloy 600에 대한ASME 증가된 온도 속성 ....................................... 2-3

표 2.5 Sarver 등[18]이 사용한 Alloy 690 가열의 화학적 조성(wt%) [18] ........................ 2-6

표 2.6 Alloy 690 카바이드 석출에 미치는 열 처리의 영향[18] ...................................... 2-11

표 3.1 추정된 일반적인 개선 계수 대 모든 환경을 고려한 pHT [33, 34]....................... 3-17

표 3.2 Brown 및 Mills가 시험한 Alloys 600, 690, 82 및52의 화학적 조성 .................... 3-18

표 4.1 Alloy 690 및 Alloy 52 시료에서 발생하는 표면 IG 균열의 예[2] ........................... 4-3

표 4.2 2004년까지 Alloy 690 1차 냉각수 응력 부식 시험 데이터 요약........................... 4-7

표 4.3 Alloy 690을 이용해 시험된 Alloy 600 에 대한 Weibull 분석 [2] .......................... 4-15


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표 4.4 2004년까지 Alloy 690 수소 처리 및 도핑된 수소화 증기 응력 부식 시험 데이터 요약......................................................................................................................... 4-19

표 4.5 Alloys 600 및 690에서 나타난 균열 발생에 대한 스웨덴 RUB 시험 [23]............ 4-20

표 4.6 Alloy 690 튜브로 제조된 가동 중인 증기 발생기 목록

(2008년 12월 현재) ...................................................................................................... 4-24

표 5.1 Alloy 690의 CRDM 관통이 있는 교체된 일부 PWR RPV 헤드의 예 .................... 5-1

표 5.2 원래 장비 또는 CRDM 관통 이외의 대체 구성요소 항목이 있는 비교적 벽이 두꺼운 Alloy 690 원자로 냉각 시스템의 예 ............................................................................. 5-2

표 5.3 EdF가 시험한 Alloy 690 CRDM 노즐의 원형 및 열처리 [3].................................. 5-3

표 5.4 PWSCC 발생에 대한 MHI 시험에서 사용된 소재의 화학적 조성[4,5].................. 5-4

표 5.5 PWSCC 발생에 대한 MHI 시험에서 사용된 소재의 열 처리 및 기계적 속성[4,5] 5-4

표 5.6 PWSCC 발생에 대한 MHI 시험에서 사용된 환경 관련 시험 조건[4,5]................. 5-5

표 5.7 미시건 대학교의 횡단면 시료에서 측정된 최대 균열 깊이 요약[7] ...................... 5-9

표 5.8 MRP 프로그램의 GE-GRC에서 시험된 모든 Alloy 690 소재의 화학적 조성 및 기계적 속성......................................................................................................................... 5-11

표 5.9 ANL에서 시험된 Alloy 690TT 판재의 화학적 조성 [11 - 15] .............................. 5-15

표 5.10 Bettis Alloy 690 시험 프로그램의 소재 시험 매트릭스[22,23].......................... 5-33

표 5.11 Alloy 690에 대한 상대적인 SCC 민감성에 대한 Bettis의 요약 [22, 23]............ 5-34

표 5.12 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 소재[25]5-40

표 5.13 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 소재의 화학적 조성[25] ............................................................................................................ 5-40

표 5.14 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 소재의 (소재 업체의 보고와 같은) 인장 특성[25] ................................................................... 5-41

표 5.15 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 소재의 시범 단조 후의 인장 특성[25]............................................................................................ 5-41

표 5.16 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 소재 및 Alloy 600 대조 시료에 대한 상세한 결과 [25] ................................................................... 5-42

표 5.17 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 “ 무가공” Alloy 690 CRDM 소재 및 Alloy 600 대조 시료에 대한 상세한 결과[25] ....................................... 5-52


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표 5.18 초임계수에서의 CGR에 대한 웨스팅하우스 시험에 사용된 Alloy 690 냉간 가공된 CRDM 소재 및 Alloy 600 대조 시료에 대한 상세한 결과[25] ....................................... 5-60

표 6.1 의도적인 냉간 가공을 하지 않은 Alloy 690 CRDM 소재에 대한 모의 PWR 1차 냉각수에서의 CGR 시험 결과 (2008년 12월 현재)............................................. 6-3

표 6.2 의도적으로 냉간 가공 처리된 Alloy 690 CRDM 소재에 대한 모의 PWR 1차 냉각수에서의 CGR 시험 결과 요약 (2008년 12월 현재)........................................................... 6-4



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