ML20209F704

an Overview of Environmental Materials Degradation in LIGHT- Water Reactors

ML20209F704
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Issue date:	08/31/1986
From:	Shaaban H, Wu P NRC OFFICE OF INSPECTION & ENFORCEMENT (IE)
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NUREG-1196, NUDOCS 8609120235
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NUREG-1196 An Overview of Environmental Materials Degradation in Light-Water Reactors

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U.S. Nuclear Regulatory Commission Offico of Inspecticn and Enforcement H. l. Shaaban, P. Wu

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NUREG-1196 An Overview of Environmental Materials Degradation in Light-Water Reactors Manuscript Completed: A; gust 1986 Date Published: August 1986 H. I, Shaaban*, P. Wu

' Affiliated with international Atomic Energy Agency Division of Inspection Programs Office of Inspection and Enforcement U.S. Nuclear Regulatory Commission Washington, DC 20565

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ABSTRACT This report provides a brief overview of analyses and conclusions reported in published literature regarding environmentally induced degradation of materials in operating light-water reactors. It is intended to provide a synopsis of subjects of concern rather than to address a licensing basis for any newly dis-covered problems related to reactor materials.

The subjects discussed here include degradation of materials in the following systems:

- reactor internals including nuclear fuels and other reactor core materials

- reactor pressure boundary including reactor vessel, piping systems, and bolting materials steam generators in pressurized-water reactors steam turbines condensers In each of these systems the degradation mechanisms and the suggested reasons for each mechanism are discussed. Possible remedies or methods for avoiding such degradation also are given.

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TABLE OF CONTENTS P. age ABSTRACT ........................................................ iii ACKNOWLEDGEMENTS ................................................ ix ACRONYM LIST .................................................... xi 1 INTRODUCTION ............................................... 1-1 2 REACTOR CORE MATERIALS ..................................... 2-1 2.1 Introduction .......................................... 2-1 2.2 Fuel Performance ...................................... 2-1 2.2.1 Pellet-Cladding Interaction .................... 2-2 2.2.2 Water-Side Corrosion of Zi rcaloys . . . . . . . . . . . . . . . 2-3 2.3 Failures of Austenitic Stainless Steels ............... 2-4 2.4 Failures of Nickel-Based and Cobalt-Based Alloys....... 2-5 2.5 References ............................................ 2-6 3 PRESSURE BOUNDARY MATERIALS................................. 3-1 3.1 Scope ................................................. 3-1 3.2 Reactor Pressure Vessel ............................... 3-1 3.2.1 Introduction ................................... 3-1 3.2.2 Corrosion Fatigue .............................. 3-3 3.2.3 Stress-Corrosion Cracking ...................... 3-4 3.2.4 Irradiation Effects ............................ 3-5 3.2.5 Aging Effects (Temperature and Strain) ......... 3-5 3.3 Piping ................................................ 3-6 3.3.1 Introduction ................................... 3-6 3.3.2 Causes of IGSCC in Austenitic Piping ........... 3-7 3.3.3 Remedies ....................................... 3-8 3.3.4 Safety Impact .................................. 3-10 3.4 Bolting and Threaded Fasteners ........................ 3-12 3.4.1 Introduction ................................... 3-12 3.4.2 Degradation of Low-Alloy Steel Bolting Materials ...................................... 3-12 3.4.3 Degradation of Age-Hardenable Austenitic Bolting Materials .............................. 3-14 3.4.4 Summary and Recommendations .................... 3-14 3.5 References ............................................ 3-15 v

TABLE OF CONTENTS (continued)

Pag 4 STEAM GENERATOR MATERIALS .................................. 4-1 4.1 Introduction .......................................... 4-1 4.2 Recirculating Steam Generators ........................ 4-1 4.3 Materials Degradation in Operating RSGs ............... 4-2 4.3.1 Wastage ........................................ 4-2 4.3.2 Pitting ........................................ 4-2 4.3.3 Denting ........................................ 4-4 4.3.4 Intergranular Attack . . . . . . . . . . . . . . . . . . . . . . . . . . . 4-4 4.3.5 Primary Side-Initiated Cracking . . . . . . . . . . . . . . . . 4-4 4.4 Once-Through Steam Generators ......................... 4-5 4.5 Materials Degradation in OTSGs ........................ 4-5 4.5.1 Erosion / Corrosion .............................. 4-5 4.5.2 Secondary Side Tube Cracking . . . . . . . . . . . . . . . . . . . 4-5 4.5.3 Primary Side-Initiated Tube Cracking . . . . . . . . . . . 4-7 4.6 Summary ............................................... 4-7 4.7 References ............................................ 4-7 5 STEAM TURBINE MATERIALS .................................... 5-1 5.1 Introduction .......................................... 5-1 5.2 Stress-Corrosion Cracking of Steam Turbine Discs ...... 5-1 5.2.1 Keyway Cracks .................................. 5-2

5. 2. 2 Rim Cracks ..................................... 5-2 5.2.3 Web-Face Cracks ................................ 5-2 5.2.4 Bore Cracks .................................... 5-3 5.3 Factors Affecting SCC of Turbine Discs ................ 5-3 5.3.1 Effect of Yield Strength ....................... 5-3 5.3.2 Effect of Stress ............................... 5-3 5.3.3 Effect of Cyclic Loads ......................... 5-3 5.3.4 Effect of Temperature .......................... 5-4 5.3.5 Effect of Electrochemical Potential ............ 5-4 5.3.6 Effect of Steam Chemistry . . . . . . . . . . . . . . . . . . . . . . 5-4 5.4 Failure of Turbine Blades ............................. 5-4 5.5 Conclusion ............................................ 5-4 5.6 References ............................................ 5-5 6 CONDENSER MATERIALS ........................................ 6-1 6.1 Introduction .......................................... 6-1 6.2 Erosion / Corrosion ..................................... 6-1 vi

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TABLE OF CONTENTS (continued)

Page 6.3 Corrosion Caused by Water Pollution ................... 6-3 i

6.4 Dealloying ............................................ 6-4 6.5 Crevice Corrosion and Pitting ......................... 6-5 6.6 Galvanic Corrosion .................................... 6-6 6.7 Environmental Cracking ................................ 6-6 6.8 Condensate Corrosion .................................. 6-7 6.9 Summary and Recommendations . . . . . . . . . . . . . . . . . . . . . . . . . . . 6-7 6.10 References ............................................ 6-7 l

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LIST OF FIGURES Page 4.1 Typical problem areas in PWR steam generators ................ 4-3 4.2 Babcock and Wilcox once-through steam generator .............. 4-6 LIST OF TABLES 1.1 Materials of major components in LWRs, related problems and their causes .................................................. 1-2 2.1 Material problems affecting fuel performance . . . . . . . . . . . . . . . . . . 2-2 2.2 Reactor ccre components fabricated from age-hardenable alloys and degradation mechanisms experienced . . . . . . . . . . . . . . . . . . . . . . . . 2-5 3.1 ASME Code specifications for the composition and mechanical properties at 20*C for SA533 Grade B Class 1 and SA508 Class 3 steels ........................................................ 3-2 3.2 Summary of all inspection findings on large piping in all operating BWRs inspected according to IE Bulletins 82-03 and 83-02 ......................................................... 3-11 3.3 Bolting materials and their applications ...................... 3-13 viii h

ACKNOWLEDGEMENTS The authors thank the members of the Engineering and Generic Communications Branch and the Reactor Construction Programs Branch, Office of Inspection and Enforcement, U.S. Nuclear Regulatory Commission, especially P. Cortland and W. Collins, for providing valuable guidance and meaningful discussions during the development of this report. They also thank R. L. Baer, R. Heishman, A. Dromerick, and J. Konklin for their help and encouragement.

One of the authors (H.I.Shaaban) wishes to thank the International Atomic Energy Agency for the fellowship offered to him during which this report was developed.

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ACRONYM LIST AIME American Institute of Mining and Metallurgical Engineering ANS American Nuclear Society ASME American Society of Mechanical Engineers ASTM American Society for Testing and Materials AVT all-volatile treatment BWR boiling-water reactor CNA Canadian Nuclear Association CRC corrosion-resistant cladding DBTT ductile-brittle transition te.nperature EPRI Electric Power Research Institute GE General Electric HAZ heat-affected zone HEC hydrogen embrittlement cracking HSW heat sink welding IAEA International Atomic Energy Agency ICCGR International Cooperative Group on Cyclic Crack Growth Rate IGA intergranular attack IGSCC intergranular stress-corrosion cracking IHSI induction heating stress improvement I-SCC iodine-induced stress-corrosion cracking LBB leak-before-break LOCA loss-of-coolant accident LPHSW last pass heat sink welding LTS low-temperature sensitization LWR light-water reactor MSIP mechanical stress improvement process NACE National Association of Corrosion Engineers NDTT nil-ductility transition temperature NG nuclear grade OTSG once-through steam generator PCI pellet-cladding interaction PTS pressurized thermal shocks PWR pressurized-water reactor RPV reactor pressure vessel RSG recirculating steam generator SCC stress-corrosion cracking SG steam generator SHT solution heat treatment SS stainless steel WO weld overlay i

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L 1 INTRODUCTION Service degradation in nuclear reactor materials represents one of the major technological factors that can limit the efficiency and viability of nuclear power. Worldwide research and development activities are currently under way to increase the performance of nuclear materials. The results of such activi-ties that have been completed are available in numerous publications. Most often authors have chosen to specialize in a narrow subset of issues pertaining to nuclear materials rather than attempt to cover a more complete set of the issues.

Review of some of the available literature has indicated a specific concern re-garding materials performance in light-water reactors (LWRs). This emphasis results from the fact that 86% of commercial nuclear power is obtained from LWRs [1]. This report is limited to presenting a broad view on the subject with specific emphasis on some problems that are of current interest. This report gives a general review of the materials-related problems observed in various LWR components including the reactor core, pressure boundary, steam generators, condensers, and steam turbines.

Most components of the LWRs operate in aggressive environments (high-temperature coolant); consequently, most of the materials-related problems are associated directly or indirectly with aqueous corrosion. A recent survey [2] shows that most of the degradation-related failures in commercially operated nuclear plants in the United States were the result of four mechanisms: erosion, corrosion, vibration, and the presence of foreign materials (suspended or deposited) in the cooling fluids. The latter failure mechanism also could be considered as a corrosion-related mechanism. However, there are other mechanisms involved such as irradiation embrittlement, temperature embrittlement (thermal shock), creep, and fatigue.

Table 1.1 summarizes the main component problems and materials-related causes in LWRs. It is clear that corrosion and corrosion-related mechanisms, such as stress-corrosion cracking (SCC) and corrosion fatigue, are common primary causes of materials-related problems in LWRs.

In this brief overview of published analyses and conclusions regarding environ-mentally induced degradation of materials in operating LWRs, the following systems have been included:

reactor internals including nuclear fuels and other reactor core materials reactor pressure boundary including reactor vessel, piping systems, and bolting materials steam generators in pressurized-water reactors (PWRs) steam turbines condensers 1-1

Table 1.1 Materials of major components in LWRs, related problems and their causes*

Reactor component Material Problem Primary cause Core: -

Fuel rod Zircaloy/U02 Cladding Stress-corrosion perforation cracking (SCC)

Wastage Corrosion /hydriding Control rod 304 stainless Cladding SCC steel (SS)/ perforation / Creep / swelling boron carbide B4 C leachout Creep / corrosion (B4C)

Fuel assembly SS/Ni-based Bowing and Creep / swelling /

alloy dilation corrosion Pressure boundary:

Vessel Low-alloy Demonstration of Radiation embrittle-steel 30 yr integ- ment and corrosion rity in fatigue presence of small cracks Piping 304 SS Cracking in weld Intergranular heat-affected stress-corrosion zone (HAZ) cracking (IGSCC)

(BWR)

C steel Cracking (PWR) Corrosion fatigue Steam generator C steel / Tube and tube IGSCC, waste corro-Ni-based support dis- sion, pitting, alloy tortion/ crack- denting, fatigue ing Condenser Cu-Ni alloys Tube failure SCC / erosion / wastage Turbine Ni-Cu-Mo-V Rotor bursts, Environmental fa-bainitic disc cracking, tigue temper em-steels, blade cracking brittlement, 12 Cr SS IGSCC, corrosion fatigue or SCC

• Adapted from Reference 3.

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In each of these systems degradation mechanisms and the suggested reasons for each mechanism are discussed. Possible remedies or methods for avoiding such degradation also are given.

References

1. " Power Reactors 1984," Nuclear Engineering International, 29, 64, October 1984.

2. J. A. Rose, R. Steele, Jr., K. G. DeWall, and Bruce C. Cornwell, " Survey of Aged Power Plant Facilities," NUREG/CR-3819, U.S. Nuclear Regulatory Commission (June 1985).

3. J. T. A. Roberts, Structure Materials in Nuclear Power Systems, Plenum Press, New York, NY (1981).

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2 REACTOR CORE MATERIALS 5

2.1 Introduction The' conditions inside the core of a, light-water reactor (LWR) that affect mate-rials include neutron irradiation, mechanical stresses, thermal stresses, and the effect of fission products within the fuel rods. The materials normally used in the LWR core include uranium dioxide (U0 2 ) (fuel pellets), Zircaloy, austenitic stainless steel, and nickel- and cobalt-based alloys [1].

The key component in the reactor core is the fuel pellet. The other core mate-rials are used either for structural purposes or for nuclear control. Within the high flux region of the core, Zircaloy is used because it has a low neutron absorption cross section, it- is slightly activated by irradiation, and it re-leases only a small fraction of its corrosion products. Other core materials j include austenitic stainless steels (Types AISI 304, 316 and 347) and Inconel i 600, which is_used for screens in the core structure, thermocouples, and detec-l tors. Highly stressed components such as springs and bolts are made of precipitation-hardened nickel-based alloys (Inconels X-750 and 718), whereas cobalt-based alloys are often used for bearings.

Experience with LWR core materials is extensive, and the results of different studies have been the topics of several important meeting proceedings [2-4].

Most of the early problems associated with inadequate design or manutacturing have been eliminated. Such problems include hydriding, cladding collapsing, unanticipated growth, and fretting. Thus, the intent here is to, focus on those limitations that are under current study, although a brief summary of past development is given.

2.2 Fuel Performance In the.last decade a variety of fuel-related problems have been experienced.

For plants up to 8 years old, total capacity losses from fuel problems were found to be about 7.0% for boiling-water reactors (BWRs), 1.5% for PWRs, and 3.6% for all units [5]. Table 2.1 (adapted from Reference 6) summarizes the main factors that lead to limited fuel performance and remedies that have been

applied to eliminate or minimize the effect on plant output. ,

l Today's performance of Zircaloy-clad U02 nuclear fuels in LWRs has generally been adequate with.some exceptions such as early displacement of certain batches in some reactors [2-4]. Concerns are focused now on the random, low i frequency fuel defects that appear to be strongly dependent on the details of core operation. Analyses [7] showed the economical need to eliminate all known failure mechanism's through optimized fuel design and operations. Recently, the emphasis in worldwide research has been on the effective solution of the pellet-cladding interaction (PCI) problem and the water-side corrosion of Zircaloys at high burnups.

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Table 2.1 Materials problems affecting fuel performance

• Component Problem Solution l

Fuel pellets (1) Enrichment errors Gamma scanning of fuel rods before (UO2 ) shipment (2) Pellet densification Stable microstructure for pellets Fuel cladding (1) Zr hydriding Moisture elimination in fabrica-(Zircaloy) tion

(2) Scale (crud) Elimination of copper alloys in

deposit (BWRs) feedwater heaters (3) Cladding corrosion Quality control of cladding Fuel rod (1) Cladding Collapse Prepressurized cladding - stable

-(PWRs) pellets (2) Fuel rod growth Avoiding preferred orientation in and bowing texture Axial clearance j Spacer design (3) Pellet-cladding Slow power rise interaction Local power shape control Fuel preconditioning Fuel pellet design changes

• Adapted from Reference 6.

2.2.1 Pellet-Cladding Interaction PCI is recognized as the major mechanism for cladding rupture during normal operation [8] and as a possible cladding failure mechanism during normal power transients [9].

Defects in Zircaloy. cladding tubes, enclosing U02 fuel pellets, usually were found to accompany power transients. Studies revealed that all post-transient investigations confirmed the low deformation nature of the failures, which often appeared as pinholes in the tubing. It was found that these usually occurred either at pellet-to pellet interfaces or opposite to pellet cracks.

Laboratory tests with iodine (and several iodine compounds) have shown that cracks in the Zircaloy, having similar fractographic characteristics as in pile-generated defects, can be simulated. This explains the reason for studying the

-iodine-induced stress-corrosion cracking (I-SCC) behavior of Zircaloys [10-13].

2 Studies revealed that iodine concentrations as low as 10 8 g/cm on the inner surface of the cladding are enough to create I-SCC in.Zircaloy [10]. The time to cladding rupture decreases linearly with increasing iodine concentration up 2-2 I

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to 10 4 g/cm2 , and at this concentration saturation occurs [11]. The studies suggested that I-SCC occurs through two different mechanisms, crack initiation tnd crack propagation. Acoustic emission studies [12] showed that crack initi-ation occurs within a short time after fuel loading. Proposed possible mecha-nisms for crack initiation include mechanical / chemical attack of the grain boundaries, surface imperfections, and cracking of radial hydrides [1]. Crack propagation can occur through several mechanisms including interconnection of smaller cracks for smaller crack sizes and growth of the individual crack for larger crack sizes [11]. Garzarolli et al. [1] used the results of various in-vestigators to establish a crack propagation map showing various crack-propagation mechanisms proposed at various stress intensities.

There is disagreement on the effect of irradiation on the stress level required to produce SCC. Lunde and Videm [14] attributed inconsistencies to small dif-ferences in testing techniques. They also showed that small amounts of oxygen in the cooling water result in a drastic increase of the failure stress. How-cver, if no passivative condition exists, irradiation embrittlement obviously enhances both crack initiation and growth. It may well be that variations in the oxidation potential within a fuel rod influence behavior during a transient.

Some studies suggested conditions to improve cladding resistance to SCC [13].

Also, remedies were considered that either reduce the local straining at pellet interfaces or hinder the fission products from coming into contact with the Zircaloy cladding. Two different methods are being used to eliminate PCI risk.

One concept is called "CANLUB," in which a graphite coating of the interior of the cladding is used [15]. This method was found to improve PCI behavior up to 10 mwd /kg. The other method uses a barrier layer of pure zirconium to protect the Zircaloy cladding [16] and was found to be an absolute remedy up to 14 mwd /kg.

2.2.2 Water-Side Corrosion of Zircaloys The water-side corrosion of Zircaloy cladding in the reactor can take any of the following three major forms:

(1) a thin, uniform cohesive layer,1- to 10 pm thick, that will not exceed 10 pm even after 10 years of exposure (2) nodular corrosion, that is, a localized attack sometimes building up to

>200 pm thickness (3) increased uniform oxidation that starts to spall at around 70 pm thickness.

The uniform corrosion of Zircaloy cladding that is most prevalent follows clas-sical corrosion laws [17]. Initially, a thin and tightly adherent black corro-sion film is formed. The growth rate is proportional to time cubed during this process. Further exposure produces a transition after which the corrosion rate is linear. Zircaloy corrosion in the reactor is influenced by irradiation and by water chemistry. Garzarolli et al. [18] found that the corrosion rate in PWRs is comparable to the rate for similar conditions outside the reactor up to a thickness of 2 to 5 pm--with higher oxide layer thicknesses, corrosion is enhanced. The environmental conditions inside PWRs are characterized by radio-lytic oxygen suppression resulting from hydrogen addition. In BWRs, where oxygenated conditions predominate, the rate is enhanced, especially at the start of corrosion.

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Zircaloy proved to-have good corrosion resistance at normal operating tempera-tures. At higher temperatures, corrosion of Zircaloy increases markedly. The Zircaloy cladding may overheat because of the following reasons:

(1) The increase of the oxide layer itself increases the temperature of the metal / oxide interface: The acceleration of the corrosion rate in this case depends on the heat flux, thermal conductivity of the oxide layer, basic oxidation rate, and activation energy.

(2) Heat flow is blocked as a result of excessive crud deposits on the fuel rod surface: The rate of crud deposit depends on water chemistry, composition of the cooling circuit structural materials, heat flux of the fuel rod, and the radiation field. A porous crud and a dense crud were found [1].

The porous crud, rich in iron, consists of irregular shaped particles ranging in size from 0.1 to 3 pm. Garlick et al. [19] suggest that porous crud contains numerous boiling channels, 2- to 4 pm wide, that permit an internal circulation of water and steam and impede the heat flow. The dense crud is attributed to the porous crud drying out, leading to further deposit of copper, silicon, manganese, and so forth. The dense crud may lift and form gaps filled with stagnant steam. Because stagnant steam has a very low thermal conductivity, this situation leads to a drastic in-crease in temperature. Crud deposits have reportedly contributed to some fuel rod failures [20] and increased Zircaloy corrosion [21].

The effect of fast neutron flux was found to be small at fluxes of 1012 to 1014 n/cm2 /s. Corrosion rates were observed to be almost linearly dependent on cool-ant conditions in PWRs [22] and BWRs [23].

Nodular corrosion, another type of Zircaloy corrosion, is often observed in coolants containing oxygen. The nodules are local formations of white zirconium oxide lenses about 0.1 to 1 pm in diameter and up to 200 pm thick. Hydrogen appears to play an important role in nodular corrosion; it was found that in-

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creasing hydrogen overpressure enhances nodular corrosion [24]. Outside the reactor, nodular corrosion was observed in high pressure steam but not in oxygen [25]. Nodular corrosion was reported to cause fuel failures in combina-tion with copper-rich crud [26]. These failures led to the development of re-search programs to study ways, such as special treatment of the Zircaloy clad-ding, to decrease nodular corrosion.

2.3 Failures in Austenitic Stainless Steel l

The experience with in-core heavy parts made of stainless steel (SS) is quite l good. The degradation of austenitic stainless steel in heavier components will be discussed in other sections of this report. The discussion here will be devoted to the neutron-induced SCC in thin cladding tubes used in control rods and in some of the earlier LWR fuels.

Failures of Type AISI 304 austenitic SS thin cladding were observed in some BWR and PWR fuel cladding [27] and in the boron carbide (B4 C) control rod cladding l

[1]. The normal failure mode is intergranular cracks initiated at the water side, although the material is not sensitized. The interaction between fuel and cladding is the source of stress in fuel rods, and the swelling of the B 4C control material is the source of stress in control rods. It was suggested that this type of irradiation-induced stress-corrosion cracking may be due to a 2-4

solute segregation of phosphorus and silicon along grain boundaries, which hardens the grain boundaries and lowers the corrosion resistance locally [28].

Thus, it is recommended that silicon be kept below 0.1% and phosphorus below 0.01%. Other studies [1] suggested that irradiation can accelerate chromium depletion of the grain boundaries, thus increasing susceptibility to SCC.

Current research is concentrated on two areas:

(1) proving that SS with low silicon and phosphorous content is SCC resistant (2) determining the critical stress and radiation level at which SCC occurs in nonsensitized SS 2.4 Failures in Nickel-Based and Cobalt-Based Alloys The behavior of Inconel 600 thin tubes used inside the reactor core will be dis-cussed in detail in another section of this report. However, studies of steam generator tubing have confirmed the sensitivity of Inconel 600 to SCC in small tubes and screws.

The high-strength austenitic age-hardenable alloys (X-750, 718, and A-286),

which were initially developed for aircraft engine applications, have been adapted for use in nuclear power service as bolts, pins, and springs. Some failures of these components have been experienced inside the reactor core as illustrated in Table 2.2. This comparison indicates that parts made of X-750 experienced more degradation than those made of 718 or A-286 alloys. Inter-granular stress-corrosion cracking (IGSCC) appears to be the main initiation mode of degradation in most of the failure cases. Most of the studies [29, 30]

recommend specific heat treatments of these alloys to increase their IGSCC re-sistance. A survey study by McIfree [31] concluded that specific heat treatments are available to decrease susceptibility to IGSCC. Section 3.4.3 of this report addresses this subject in more detail.

Table 2.2 Reactor core components fabricated from age-hardenable alloys and degradation mechanisms experienced

• Component Reactor systems Alloy Degradation mode Bolts Core baffel X-750 IGSCC/ fatigue (less Fuel assembly or frequent)

Core barrel A-286 Thermal shield Springs Fuel assembly holddown X-750 Fatigue Fuel assembly fingers X-750 IGSCC Control component holddown 718 Fatigue Beams BWR jet pump and steam X-750 IGSCC separator or A-286

• Adapted from Reference 31.

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No realized selective corrosion was reported for the cobalt-based alloys [1].

However, it was found that the parts made of cobalt are a source of high radioactivity in certain nuclear power plant systems. Thus, several cobalt-free alloys were developed to substitute for cobalt-based alloys including age-hardenable nickel-based alloys and austenitic-martensitic steels.

2.5 References

1. F. Garzarolli, H. Rubel, and E. Steinberg, " Behavior of Water Reactor Core Materials With Respect to Corrosion Attack," Proceedings of the Interna-tional Symposium on Environmental Degradation of Materials in Nuclear Power Systems--Water Reactors, Myrtle Beach, South Carolina, August 1983 National Association of Corrosion Engineers (NACE) (1984).

2. Proceedings of the Joint Topical Meeting on Commercial Nuclear Fuel Tech-nology Today, April 28-30, 1975, Toronto, Canada, CNS ISSN 0068-8517, 75 CNA/ANS-100, Canadian Nuclear Association and American Nuclear Society.

3. Proceedings on Water Reactor Fuel Performance, May 3-11, 1977, St. Charles, Illinois, American Nuclear Society, Chicago, Illinois.

4. Proceedings on Water Reactor Fuel Performance, April 29-May 3, 1979, Portland, Oregon, American Nuclear Society, Chicago, Illinois.

5. R. H. Koppe and E. A. Olson, " Nuclear and Large Fossil Unit Operating Experience," EPRI NP-1191, Electric Power Research Institute (September 1979).

6. M. Levenson and E. Zebroski, "The Nuclear Fuel Cycle," Annual Review of Energy, 1, 645-674 (1976).

7. D. Pomerey and J. Waring, " Methods for Determining the Cost of Fuel Failures in Nuclear Power Plants," EPRI NP-854, Electric Power Research Institute, Palo Alto, California (August 1978).

8. F. Garzarolli and R. Von Jan-H. Stehal, "The Main Causes of Fuel Element Failures in Water-Cooled Power Reactors," Atomic Energy Review, 17,31-128 (1979).

9. C. L. Mohr, P. J. Pankaskie, J. C. Wood, and G. Hessler, "PCI Fuel Failure Analysis," NUREG/CR-1163, U.S. Nuclear Regulatory Commission, PNL-2755, Pacific Northwest Laboratory (December 1979).

10. M. Peehs, H. Stehle, and E. Steinberg, "Out-of-Pile Testing of I-SCC in Zry in Relation to the PCI Phenomenon," Proceedings of the 4th International Conference on Zr in the Nuclear Industry, Stratford-Upon-Avon, England June 1978, ASTM-STP 681, pp 244-260, American Society for Testing and Material.

11. A. K. Miller et al., "SCCIG: A Phenomenological Model for Iodine Stress Corrosion Cracking of Zircaloy," EPRI NP-1798, Electrical Power Research Institute (1981).

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12. M. Peehs, W. June, H. Stehle, and E. Steinberg, " Experiments To Settle an I-SCC Hypothesis for Zry Tubing," Proceedings of the Specialists' Meeting on PCI in Water Reactor, RISO, pp 170-181 (September 1980).

13. D. Cubicciotti, R. L. Jones, and B. C. Syrett, " Stress Corrosion Cracking of Zircaloys," EPRI NP-1329, Electric Power Research Institute (March 1980).

14. L. Lunde and K. Videm, "The Influence of Testing Conditions and Irradia-tion on the Stress Corrosion Cracking Susceptibility of Zircaloy," Journal of Nuclear Material, 95, 210-218 (1980).

15. D. C. Hardy et al., " Performance of CAN0U Development Fuel in the NRU Reactor Loop," Proceedings on Water Reactor Fuel Performance, May 1977, St. Charles, Illinois, pp 198-206, American Nuclear Society, Chicago, Illinois.

16. J. H. Davies et al. , " Irradiation Tests on Barrier Fuel in Support of a -

Large Scale Demonstration," Proceedings on LWR Extended Burnua Fuel Performance and Utilization, Williamsburg, Virginia, April 1932, pp 551-568, American Nuclear Society, Chicago, Illinois.

17. E. Hillner, " Corrosion of Zirconium-Base Alloys - An Overview," Proceedings of the 3rd International Conference on Zirconium in the Nuclear Industry, A. Lowe and G. Parry (eds.), ASTM-STP 633, American Society for Testing and. Materials (1977).

18. F.-Ganarolli et al. , "KWU-Results on Waterside Corrosion of Zircaloy in PWRs and BWRs," Meeting on Influence of Water Chemistry on Fuel Element Cladding Behavior in Water Cooled Power Reactors, Leningrad, USSR, June 1983, Intp national Atomic Energy Agency (IAEA).

19. A. Garlick et al., " Crud Induced Overheating Defects in Water Reactor Fuel Pins," Journal of British Nuclear Engineering Society, 1_6,77-80,(1977).

20. F. Garzarolli, R. von Jan and H. Stehle, "The Main Causes of Fuel Element Failure in Water Cooled Power Reactors," IAEA-Atomic Energy Review, 1_7,31-128, International Atomic Energy Agency (1979).

21. H. W. Wilson et al., " Fuel Performance Characteristics at Extended Burnup,"

Proceedings on LWR Extended Burnup Fuel Performance and Utilization, Williamsburg, Virginia, April 1982, pp.121-138, American Nuclear Society, Chicago, Illinois.

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! 22. E. Hillner, "Long Term In-Reactor Corrosion and Hydriding of Zircaloy-2

) Tubing," ASTM-STP 754, pp 450-478, American Society for Testing and l Materials (1982).

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23. L. Lunde and K. Videm, "The Influence of Neutron Irradiation on the i

Corrosion of Zircaloy-2 at 240 C," 6th Scandinavian Corrosion Congress, l Gdteborg, Sweden (1971).

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24. L. Lunde and K. Videm, "Effect of Material and Environmental Variables on Localized Corrosion of Zirconium Alloys," ASTM-STP 681, pp 40-59, American Society for Testing and Materials (1979).

25. F. W. Trowse, R. Summerling, and A. Garlick, " Zirconium in the Nuclear Industry," ASTM-STP 633, pp 236-257, American Society for Testing and Materials (1977).

26. G. Vesterlund et al., " Experience of Interaction Between Fuel Cladding and Coolant Water in Swedish BWRs," IAEA Specialists' Meeting, International Atomic Energy Agency, San Miniato, Italy, October 1981. I

27. A. Strasser et al. , "An Evaluation of Stainless Steel Cladding for Use in Current Design LWRs," EPRI NP-2642, Electric Power Research Institute (1982).

28. R. N. Duncan, " Stainless Steel Failure Investigation Program," Final Summary Report, GE Report GEAP-5530, General Electric (1968).

29. G. Boiside, H. Coriou, L. Grall, and R. Mahnt, " SCC of Inconel-X-750 in PWR at 300 and 350 C," Proceedings of British Nuclear Society (BNS)

Corrosion Conference, London, 1972.

30. S. Floreen and G. L. Nelson, "The Effect of Heat Treatment and Composition on the Stress Corrosion Cracking Resistance of Inconel-X-750,"

Metallurgical Transactions, 14a, 133-139 (1983).

31. A. R. McIfree, " Degradation of High Strength Austentic Alloys X-750, 718, and A-286 in Nuclear Power Systems," Proceedings of International Symposium on Environmental Degradation of Materials in Nuclear Power Systems--Water Reactors, Myrtle Beach, South Carolina, August 1983, National Association of Cor osion Engineers (1984).

l l

i t

l 2-8

3 PRESSURE BOUNDARY MATERIALS 3.1 Scope The nuclear pressure boundary consists of those components of the primary and secondary systems of the power plant that contain pressurized coolant, includ-ing the reactor vessel, steam generators, steam and primary coolant piping, valves, pumps, nozzles, and assorted small components. The steam generators are discussed in Section 4 of this paper. The integrity of the pressure bound-ary in nuclear plants is necessary for the proper functioning of the plant and for providing protection for the safety of the public. Accordingly, the compo-nent failure probabilities that are considered acceptable are very low. For example, the U.S. Advisory Committee on Reactor Safeguards has specified a maxi-mum acceptable failure rate as 10 8 failures per reactor pressure vessel year [1].

This section will address the possible degradation in pressure boundary mate-rials, the published data on the materials-related failures in these systems, and the actions taken to avoid such failures. The pressure boundary components discussed in this section include the pressure vessel, piping, and bolting ma-terials, valves, and pumps.

3.2 Reactor Pressure Vessel 3.2.1 Introduction The reactor pressure vessel (RPV) is the most critical component in the LWR pressure boundary. It is the only component in this boundary whose failure is considered to be of such low probability that it is considered essentially in-credible. Thus, safety systems to provide adequate core cooling in the event of RPV failure are not required by regulation. Therefore, if the RPV fails the results could be catastrophic. In PWRs, the structural integrity of the RPV is of particular concern because of the high coolant pressure and high neutron doses.

This concern with RPV safety has long been recognized and is exemplified in international efforts to study factors that influence failure properties of the RPV steels. Such effort is aimed at decreasing the possible anticipated or hypothetical failure. Available information indicates that the performance of RPVs is quite satisfactory. For example, preliminary investigations showed that the pressure vessel of the Three Mile Island Unit 2 reactor, which expe-rienced the worst accident to date in the United States of commercial nuclear power, had no detectable defects.

The ferritic steels used in the RPV (see Table 3.1 adapted from Reference 2) are characterized by a ductile-to-brittle fracture transition at a specific temperature which is called the nil-ductility transition temperature (NDTT).

Below this temperature the fracture mode is usually cleavage; above NDTT the fracture is a dimpled rupture. Raising the temperature above NDTT first in-reases the toughness, followed by a region of relatively constant toughness, called the upper shelf.

3-1

Table 3.1 ASME Code specifications for the l composition and mechanical properties )

at 20*C for SA533 Grade B Class 1 and '

SA508 Class 3 steels Composition / SA533-1, SA508-3, mechanical properties plates forgings Carbon (%) 0.25 (max) 0.15-0.25 Manganese (%) 1.10-1.66 1.20-1.50 Phosphorus (%) 0.035 (max) 0.025 (max)

Sulfur (%) 0.040 (max) 0.025 (max)

Silicon (%) 0.13-0.32 0.15-0.40 Molybdenum (%) 0.41-0.64 0.45-0.60 Nickel (%) 0.37-0.73 '0.40-1.00 Chromiura (%)

0.25 (max)

Vanadium (%)

0.05 (max)

Tensile strength 552-689 550-725 (MPa)

Yield strength -> 344 -> 345 l (0.2% offset) l (MPa)

Elongation in -> 18 -> 18 50 mm (%)

Reduction -

-> 38 area (%)

Source: Reference 2.

l Fracture toughness is measured by various methods. Qualitative methods include impact energy absorption in the Charpy V-notch test and NOTT in the drop weight test. The theory of fracture mechanics specifies fracture toughness in terms of other quantities. The stress-intensity factor (k) is used to predict the combination of loading and crack size that leads to the onset of the crack growth in structural components [3]. Once crack growth starts, it will continue unless the stress-intensity factor at the crack tip falls below the " crack arrest toughness," which is a material property.

When a crack in an RPV or nozzle is detected during in-service inspection, Sec-tion XI, Appendix A of the American Society of Mechanical Engineers (ASME) 3-2 i ___ _ . _ - _ _ - - . _ - - - _ _ _ _ _ - _ _ _ _ _ _ _ _ _ _ _ _ _ _ _ _ _ _ ___ _ _ _ _ _ _

Boiler and Pressure Vessel Code gives a procedure for estimating the remaining useful service life. The procedure combines a fatigue crack growth analysis with a failure margin analysis to determine whether or not repair of the crack is necessary to ensure an adequate margin against vessel failure for the remain-der of the vessel's planned service life. Because several variables are not accounted for in the Appendix A procedure, there are doubts about the accuracy of the calculated "end-of-life" flaw sizes.

Ideally, reliable fracture toughness data should be determined for all relative material forms (such as plates, forgings, weld region, and heat affected zone) and for all appropriate service conditions (such as temperature, strain rate, radiation, and coolant). Currently, many of the values incorporated into the ASME Code are based on very limited data. For example, fracture toughness and crack arrest toughness are based on only a few heats of steel, mostly plate and forging material and no weld material at all [4]. In addition, it is extremely important to prove that the data for the properties of materials developed from small laboratory specimens are truly representative of the behavior of the mate-rials in pressure vessels in both the unirradiated and irradiated conditions.

The environmental factors that may affect the RPV steel properties and produce possible environmental degradation include aqueous corrosion, neutron irradia-tion, and temperature aging. The environmental degradation could be controlled by controlling material selection and water chemistry. A brief summary of the present understanding of the environmental degradation of RPV steels is given in this section.

3.2.2 Corrosion Fatigue Generally, the initiation and growth of cracks under cyclic loading is caused by fatigue under dry conditions. The presence of a corrosive medium in contact with the surface can accelerate both crack initiation and growth. Corrosion pits or intergranular penetration can produce sharp stress concentration that may help crack initiation. In LWRs, water chemistry is controlled to avoid such effects. Moreover, the RPV is internally protected by a stainless steel cladding. In spite of these protective actions it has been realized that marked acceleration in fatigue crack growth may occur because of the aqueous environ-ment inside the RPV [5, 6].

The multiplicity of factors influencing fatigue crack growth in the RPV steel and the long test time required led to the formation of the International Co-operative Group on Cyclic Crack Growth Rate (ICCGR) to coordinate the inter-national efforts in this area. This group includes over 40 members from 11 countries. A summary of the activities of this group is given in References 5 and 7.

The effects of various variables on the enhancement of corrosion fatigue crack growth were studied. The most important metallurgical factor affecting fatigue crack growth was found to be the sulfur content and distribution in the steel.

Amzallag et al. [8] have shown that the crack growth rate acceleration factor (i.e., da/dN at specific environment divided by da/dN in air) rises systemati-cally with sulfur content. Other studies [9] showed that spheroidized sulfides lead to a lower acceleration factor than sulfide stringers at a given sulfur level. A report on PWR pressure vessel integrity published in the United 3-3

Kingdom [10] indicated that low sulfur steel forging should be more resistant to corrosion fatigue crack growth.

For materials having a low sulfur content (<0.01%), test variables only slightly affect the rate of cyclic crack growth [7]. However, at a high sulfur content, variables such as R-ratio (Kmin/Kmax), frequency waveform and water chemistry, can be substantial.

Recent studies [11, 12] have shown that sulfide inclusions play an important role in environment sensitive cracking of RPV steels. Manganese sulfide (MnS) inclusions may dissolve in aqueous solutions and provide initiation sites for cracking. Moreover, the dissolution produces sulfuric acid, which leads to a  ;

more acid condition inside the crack. Inclusions also can act as preferential '

trapping sites for hydrogen. The trapped hydrogen may cause brittle fractures nucleated around inclusions.

Cullen and Torronen [13] showed that under low frequency cyclic fatigue condi-tions a considerable enhancement of crack growth rate occurs in a relatively pure water environment. The maximum rate is observed at a frequency of 0.01 Hz.

Such enhancement is considered a unique corrosion fatigue phenomenon and one not associated with SCC. However, under the same testing conditions, with a con-stant rate ramp waveform rather than a sine function, the change in frequency was found to have no effect on the crack growth rate [14]. It is important to note that the tests on which the Appendix A (Section XI of the ASME Code) sur-face flaw curves were based used sine wave loading at a frequency of 0.0167 Hz.

These test conditions appear to be close to a worst combination of waveform /

frequency. Consequently, the Appendix A curves are expected to greatly over-predict crack growth at other frequencies and waveforms.

Another variable requiring more systematic study is fluid flow. Studies [15]

suggest that a highly turbulent water flow rate minimizes crack growth enhance-ment. Most of the earlier data were obtained from static or low flow tests.

As for the effect of oxygen in the cooling water, some investigators [16] find that it enhances the rate of crack growth. However, a considerable scatter in data was found in the information obtained from nine laboratories [5]. Irradia-tion levels simulating the end-of-life properties of PWR vessels have been found to have no effect on corrosion fatigue crack growth [17]. Irradiation effects on other properties are detailed in Section 3.2.4.

3.2.3 Stress-Corrosion Cracking It was found that cracks in RPV steels can grow under a combination of applied static stress in water environments. These are typical SCC conditions. Thus, several studies have been made of SCC behavior of RPV steels. The relationship between this mode of subcritical crack growth and fatigue is currently under study [18].

Slow strain rate tests of RPV steels have been ccnducted in oxygenated high purity water. It was found that SCC susceptibility depends on temperature and oxygen concentration [19]. Stress-corrosion cracks were observed at oxygen levels as low as 60 parts per billion (ppb) at 288 C [13]. Anodic polarization roughly corresponds to oxygen-induced SCC at oxygen levels of 60 ppb [20].

These results suggest that SCC initiation and growth are possible under BWR 3-4

coolant conditions only. However, it was found that very large plastic strains are required to initiate SCC [7]--these strains are most unlikely to be encoun-tered in an RPV in service.

SCC behavior in fracture mechanics specimens under constant applied loads has been analyzed [21]. It was found that the threshold stress intensity factor (k) is reduced by increasing the level of dissolved oxygen. However, more work is required to better understand possible SCC growth in the absence of both oxygen (PWR water) and a mechanical crack sharpening process (static load) and to better define the rates of crack growth above the threshold.

3.2.4 Irradiation Effects During reactor operation, the RPV steel and weldments are subjected to neutron irradiation, especially the region opposite the core midplane (beltline region).

The radiation damage causes changes in mechanical properties; both yield stress and hardness increase and the work hardening rate decreases. There are also changes in fracture properties that include an increase in the ductile-brittle transition temperature (OBTT) and a decrease in upper-shelf toughness.

The harmful effects of both copper and phosphorus on irradiation susceptibil-ity were recognized. Thus, specifications for nuclear beltline materials now restrict copper and phosphorus levels. Data on impact toughness changes using modern RPV steels [22] showed less susceptibility of such steels to ir-radiation damage compared with the older steels. The overall implication for modern pressure vessel steels is a 40C increase in DBTT for an end-of-life dose of 2.3 x 10" n/cm-2 (E >l MeV) at 29 C at the inner surface of the belt-line region. However, studies indicate the importance of compositional homo-geneity in determining the behavior of these modern steels.

Few data are available on the effects of irradiation on upper-shelf fracture toughness [6]. No changes have been observed in compact tension specimens of modern steels in either initiation toughness or subsequent resistance to crack growth. Similarly, no significant effect of irradiation on the upper-shelf behavior of modern steels has been detected [23].

The specification for materials resistant to irradiation embrittlement has re-sulted from emperical data rather than an understanding of the physical basis of the irradiation-induced embrittlement. The true sequence of events during irradiation has not been elucidated. However, theories have been postulated for the role of copper, and to a lesser extent those of phosphorus and nickel, in promoting irradiation embrittlement. It is postulated that a positive bind-ing between copper atoms and vacancies stabilizes microvoids or small vacancy dislocation loops against thermal annealing, thus preserving them as barriers to dislocation motion. More recently, ideas based on enhanced copper precipi-tation have been advanced [24], leading to enhanced modulus hardening, which is the main reason for the observed irradiation effects.

3.2.5 Aging Effects (Temperature and Strain)

The wall of the pressure vessel is held at about 290 C during operation of the reactor. Temper embrittlement may occur because of the precipitation of carbon or phosphorus along grain boundaries. Although RPV steels are less susceptible 3-5

to temper embrittlement, there is still evidence of some degree of embrittlement 4 [25].

, It has been confirmed that large grain size (about 200 pm) and impurity content (P As, Sn, and Sb) have deleterious effects during the thermal aging of RPV steels [26], particularly weldments and segregated regions in forgings. It has

been suggested [6] that an increase of 30C* in DBTT as a result of thermal aging 4 during the lifetime of the RPV should be considered. However, with adequate control of segregation and heat-affected zone (HAZ) microstructure, thermal "

4 aging effects could be decreased. Fracturetoughnesstestsshowthatthereis no significant effect of long-term aging at 300 C on upper-shelf fracture j toughness properties [25].

l Strain aging is a result of the interaction of the interstitial atoms, like carbon and nitrogen, with the dislocations. Thus, modern RPV steels are speci-

! fled so that the nitrogen content is very low (<10 ppm). However, possible j strain aging. effects can be found that include reduction in ductility and upper-and lower-shelf toughness as well as an increase in yield strength and DBTT.

Authors indicated a modest increase of about 10% in ductility for a 2% prestrain

! [6] and an increase in D8TT of 20C* [27] as a result of strain aging. Impact tests showed no significant change in upper-shelf energy for pre-strains up to 5%. However, tensile and fracture toughness tests (low strain rates) showed

evidence of dynamic strain aging [28]. Also, large dynamic strain-aging effects occur in certain stress relieved welds [29]. More work is required in this area to obtain a better understanding of the reasons for dynamic strain-aging effects.

l Recent concern has focused on the effect of pressurized thermal shock (PTS)

in some of the embrittled reactor vessels (neutron irradiated, thermal, and i

strain aged) that were fabricated with steels containing high contents of copper, phosphorus, and nickel. PTS may occur as a result of system transients that .

significantly cool the vessel while fluid pressure remains at or near normal f

! uperating values.

In such accidents, the aged vessels may possess a potential for extensive pro-i pagation of pre-existing inner surface flaws before the vessels' normal end of life. Extensive studies have been performed to evaluate this problem; results have shown that there may be some vessels that exhibit a potential for failure in a few years if subjected to the 1978 Rancho Seco-type transient [30]. How-

! ever, it is possible that the calculation model may be excessively conservative because of the relative uncertainties of data. This possibility and other issues related to the PTS problem are under investigation [4, 19].

3.3 Piping 3.3.1 Introduction

}.

i The main degradation problem in LWR piping is (IGSCC) in the sensitized parts

( of austenitic stainless steel--a problem that has been recognized since the j introduction of the first commercial BWR. In 1965, furnace-sensitized stainless i steel components cracked in the Dresden Unit 1 BWR [30]. The generic nature of j this problem in BWRs was recognized in 1974 when many cracks were found in the 1

HAZs in small-diameter piping [31] (4 inches and less) and later in piping with l

larger diameters [32] (10-inches). In 1978, the Gundermmingen BWR Unit A in i

3-6 1

Germany reported its first incident of IGSCC in a 24-inch Type 304 stainless steel pipe [33]. Further incidents of cracking in piping more than 20 inches in diameter did not occur until 1982 [34]. Subsequently, a number of other BWR incidents of cracking in large-diameter lines were reported [35].

In PWRs some minor cases of IGSCC have been reported in low pressure piping that is not part of the reactor coolant system [36]. These systems include the residual heat removal system, safety injection systems, and the containment spray system. The difference in behavior of main reactor coolant piping in BWRs and PWRs is attributed to maintaining the oxygen level in PWR water sub-stantially below that of BWRs through hydrazine additions during startup and a hydrogen overpressure during operation [37]. Also, most PWRs use centrifugally cast austenitic stainless steel or austenitic-clad ferritic steel piping, which is more resistant to IGSCC than the wrought stainless steel piping used in BWRs.

In the United States only, over $100 million has been expended in research pro-grams to develop remedies for the IGSCC problem in BWRs [37]. For one BWR (Nine Mile Point), the cost to replace the affected recirculating piping was estimated to be $65 million and an additional $100 million plus to purchase replacement power for the 15-month outage [38].

) In Japan most of the IGSCC in austenitic piping was discovered before 1978.

Since 1982, no cracks have been found [39], probably because of the extensive mitigating actions, including corrosion-resistant cladding (CRC), solution heat treatment (SHT), induction heating stress improvement (IHSI), or replacement with Type 304 NG or 316 NG stainless steel for operating plants and those under construction.

3.3.2 Causes of IGSCC in Austenitic Piping A combination of material, stress, and environmental factors causes IGSCC in BWR piping. Each of these factors can affect the initiation of a stress-corrosion crack and the rate of its subsequent propagation.

The main material factor in the occurrence of SCC in Type 304 or Type 316 SS is the formation of grain boundary networks of chromium carbides accompanied by j adjacent areas depleted in chromium. Electrochemical cells will set up between

the lower narrow areas of chromium adjacent to grain boundaries and the higher chromium bulk grain material. The formation of those depleted zones is a func-tion of steel composition and the time that the steel is heated in the tempera-ture range in which reaction kinetics are sufficiently rapid and thermodynamics i favor production of chromium carbides. Typically, this occurs in the temperature range of 500 -800 C. Such depletion and precipitation occur most commonly under the thermal conditions encountered during welding or furnace heat treatment [37].

l The material in this condition is referred to as " sensitized." Even when using welding processes with low heat input and fast cooling rates, some small carbide particles may be formed during welding. These particles can grow under the tem-i perature conditions of BWR operation, following what is called low-temperature

! sensitization (LTS) [40].

Sensitization can occur in steel with carbon levels as low as 0.02% [40]; how-ever, IGSCC has not been observed in stainless steels containing less than 0.04% in BWRs [39]. The probability of IGSCC occurrence increases rapidly at l

3-7 l

carbon concentrations between 0.04% and 0.05% and then remains constant up to the 0.08% limit.

Stresses required to initiate SCC may be either applied or residual. The pri-mary source of tensile stress in piping of an operating BWR is considered to be the residual stresses from welding. Other sources for stresses include operational pressure, vibrational, and thermal stresses. Generally, in stain-less steels there is no sharp elastic limit; that is, there is no sharp stress limit before which it could be said that the deformation is elastic. Some plas-tic deformation can occur at stresses considerably below the nominal design yield strength (0.2% offset). Tensile stress produces rupture in the protec-tive oxide film, followed by a rapid bare metal corrosion until the film re-heals. The lower chromium content along the grain boundary makes these areas more susceptible to a corrosive environment with the possibility of great metal loss locally before the film is rehealed. Repeated rupture of the film pro-duces a crack along the grain boundary [40].

The presence of oxygen in the cooling water of BWRs is supposed to be the main cause of SCC. Oxygen appears to establish a chemical potential for the stain-less steel so that cracks can initiate and propagate. Species responsible for conducting the electrochemical currents in IGSCC are assumed to come from the environment (leakage of impurities from the condenser for example) or from the corrosion of the metal itself producing salts of sulfur, phosphorus, silicon, or other anions. Such species need not be present in high concentrations; often small fractions of parts per million (ppm) may be sufficient [38],

especially for long-time applications.

3.3.3 Remedies A number of potential remedies for IGSCC in BWRs have been developed [41,42]

on the basis of reducing or eliminating one or more of the three factors related to IGSCC. Thus, it is convenient to categorize these solutions as being mate-rial, stress, or environment related.

(1) Material The simplest remedy is solution-heat treatment of the materials by increasing the material temperature (to about 1050 C for Type 304 SS) to drive the chromium carbide back into solution, then quenching in water to prevent reprecipitation.

This method is applicable to most shop welds but it is impractical for field welds. All stainless steels are required to be in the solution-annealed state before fabrication [43].

Another material remedy is the application of the corrosion-resistant cladding (CRC). Type 308 weld metal is weld deposited on the inside surfaces of the ends of the pipes to be welded for a few inches beyond the expected HAZ to pre-vent the HAZ from coming into contact with the coolant. The main drawback of CRC is that the cladding makes subsequent inspections of the weld more difficult.

The most important material remedy has been the development of nuclear grade (NG) stainless steels as substitutes for Type 304 SS [44]. Type 304 NG SS has a maximum concentration of 0.02% carbon (compared with 0.08% carbon in Type 304 SS and 0.03% carbon in Type 304L SS). To compensate for the low strength result-ing from low carbon content, nitrogen content is increased (0.06%-0.1%), but it 3-8 a _ _ _ _ _ _ _ _

is still within the maximum limit specified in the ASME Boiler and Pressure Vessel Code. Thus, the lower carbon content makes the material less susceptible to sensitization, and staying within the limits of the ASME Code-accepted mate-rials makes it possible to avoid costly and time-consuming qualification pro-grams. NG stainless steels have been used successfully in the Japanese nuclear reactors [39].

Steels containing alloying elements that form carbides more stable than chromium carbide (such as niobiua or titanium) are less prone to sensitization. How-ever, even with these steels, there is a potential area immediately adjacent to the weld fusion line in which chromium-depleted zones can be produced if the carbon levels are high (knife-line attack). The relatiye freedom of the newer German BWRs from IGSCC is thought to be the result of using a low-carbon, niobium-stabilized stainless steel that contains molybdenum (Type 347 NG)[39].

(2) Stress The heat sink welding (HSW) technique has been used to add a compressive stress to the inner surface of the piping, thus decreasing the tensile residual stress at this surface. In this process, flowing cold water is applied to the inner surface after the root pass of the weld is made. The water cools the inner surface, this stretches the inner-surface ligantents during the subsequent weld passes and leads to compressive residual stress at the inner wrface after weld cooldown. HSW also helps to reduce sensitiz9 tion in the HAZ. The last pass heat sink welding (LPHSW) could be used where the heat sink is applied only during the last welding pass on the outer surface. However, recent studies

[45] show that austenitic stainless steel (Type AISI 304) could possibly under-go SCC under residual compressive stresses. It was found that the incubation period for SCC under the compressive stresses is much longer and the crack propagation is much slower than under tensile stresses. Such results indicate that further consideration has to be given to the validity of the HSW technique.

An extension of the idea behind LPHSW is to employ a weld overlay (W0) on exist-ing welds. The WO is usually Type 308L SS. This was used as a temporary solu-tion for weldments with IGSCC indications to reduce tensile stresses in the weld below and to provide strength by adding to the net section thickness. It was thought until recently that the layer of weld metal on the outer surface of the pipe makes subsequent inspection very difficult, if not impossible. However, studies by Electric Power Research Institute (EPRI) [46] show that cracks reach-ing within the upper 25% of the original pipe wall can be effectively detected.

Thus, overlay flaws unacceptable within the limits of the ASME Code are now generally detectable.

Induction heating stress improvement (IHSI) is another technique used to introduce compressive stresses on the inner surface of the pipe weld. The technique is applied by heating the outer surface of the weld to about 550 C using induction heaters while keeping the inner surface at 150 C by means of flowing water. The Japanese used this technique in their newer BWRs, and it was also used in some U.S. BWR plants [47]. The main drawback is the substan-tial radiological exposure involved if applied at existing plants. Also, when operating loads are superimposed, the resultant stress at the inner surface of the pipe will be less compressive or even tensile.

3-9

The mechanical stress improvement process (MSIP) is the only nonthermal process used to introduce compressive stresses on the inner side of the pipe welds. In this process, the outside diameter of the pipe weld is squeezed. This estab-lishes compressive axial and hoop stresses on the inside diameter of the pipe welds. It is thought that the MSIP is simpler than thermal processes and pro-vides a more repeatable method of inducing uniform compressive residual stresses.

This process was applied in a U.S. nuclear power plant in 1986 [48] and will be evaluated during subsequent outages.

(3) Environment Oxygen levels as low as 0.2 ppm (the steady level during normal operation of BWRs at 289 C) were found to be high enough to produce IGSCC. Recent studies

[47, 49] propose the reduction of BWR oxygen levels to about 20 ppb along with more stringent controls on other minor water impurities to eliminate IGSCC.

Deaeration can effectively reduce the startup oxygen level (8%). However, other methods are required to reduce oxygen to the desired level. Hydro-gen addition to the feedwater at a concentration of 1.5 ppm was found feasible

[44,49]. However, it was found necessary to use oxygen instead of air in the offgas recombiner system. The cost of doing this was estimated as $1000 a day, and the initial equipment cost was about $1 million [49].

The main drawback of hydrogen addition is the greater carryover of radioactive nitrogen N18 (half-life of 7 seconds) to the turbine. Studies [50] have indi-cated that radioactivity resulting from the N16 increased by a factor of 5 when hydrogen was added to the feedwater. Also, hydrogen addition may have long-term effects, such as hydrogen embrittlement, on fuel performance.

Other species have to be controlled in the BWR feedwater so that the water con-ductivity is maintained at 0.2 ps/cm [38, 44] instead of the current NRC require-ment of 1 ps/cm [51]. Chlorides, sulfur species, and carbonates have been found to induce susceptibility to IGSCC in Type 304 SS when present in feedwater at levels of only a few tenths of a ppm. The benefits of adding hydrogen to the feedwater may be lost without tighter control of impurities [44].

3.3.4 Safety Impact l The NRC responded to the detection of IGSCC in recirculation system piping in BWR plants by publishing two Office of Inspection and Enforcement (IE) bulle-tins [35, 52]. These bulletins required that effective ultrasonic test (UT) procedures be used and that the effectiveness of these procedures and the com-l petence of the UT examiners be demonstrated on cracked piping samples. Inspec-tions according to the new NRC requirements showed that cracking in most BWRs was extensive (see Table 3.2). As a result of this inspection program, utili-ties are planning to take positive actions to preclude or minimize IGSCC.

Several utilities are either making positive plans or are in the process of l replacing susceptible piping with Type 316 NG stainless steel.

Along with increasing the requirements for inservice inspection of piping pres-sure boundaries and changing susceptible piping to less susceptible material, other actions are needed. These include:

i 3-10 l

Table 3.2 Summary of all inspection findings on large piping in all operating BWRs inspected according to IE Bulletins 82-03 and 83-02 Extent of inspection Inspection results

(% of welds ins)ected) (No. of cracked welds) No. of Residual Residual weld heat heat overlays Plants Recirc. removal Recirc. removal repaired Big Rock Point 20% (11/59) ---

0 ---

0 Browns Ferry 1 98% (103/105) 90% (36/40) 33 14 42 Browns Ferry 2 27% (25/91) 28% (9/32) 2 0 0 Browns Ferry 3 98% (103/105) 28% (9/32) 0 0 0 Brunswick 1 25% (29/115) 75% (3/4) 3 0 3 Brunswick 2 100% (102/102) 100% (5/5) 15 1 8 Cooper 100% (108/108) 100% (7/7) 20 0 13 Dresden 2 47% (47/101) 10% (4/40) 10 0 7 Dresden 3 100% (115/115) 90% (45/50) 53* 11* 61 Duane Arnold 42% (49/117) 40% (2/5) 0 0 0 FitzPatrick 47% (49/106) 45% (5/11) 1 0 0 Hatch 1 47% (47/100) 100% (11/11) 5 2 6 Hatch 2 94% (97/103) 100% (11/11) 36 3 27 Millstone 1 11% (*11/100) 0% (0/46) 0 0 0 Monticello 100% (106/106) 78% (18/23) 6 0 6 Nine Mile Pt. 1 82% (62/76) ---

53 0 0 Oyster Creek 39% (31/80) ---

0 0 0 Peach Bottom 2 100% (91/91) 91% (32/35) 19 7 21 Peach Bottom 3 91% (77/85) 92% (35/38) 10 5 15 Pilgrim 1**

Quad Cities 1 8% (9/110) 20% (9/44) 0 0 0 100% (106/106) 90% (45/50) 20 2 9 Quad Cities 2 Vermont Yankee 66% (58/88) 7% (2/30) 33 1 22

• lt should be noted that 18 welds originally reported to be cracked were later reevaluated and were determined not to be cracked; therefore, they are not included in these totals.

• • After inspecting about seven welds and finding cracks in four of them, the utility decided to replace the piping with Type 316NG; therefore, the exami-nation has not been completed.

(1) The current work on reducing oxygen through hydrogen additions should be closely followed, taking into consideration the possibility that this method may be used to further reduce the electrochemical potential of the stainless steel to a level at which SCC will not occur. It appears that hydrogen water chemistry is an effective IGSCC countermeasure, but consid-eration should be given to the adverse effects on other reactor components, such as hydrogen embrittlement and the greater carryover of the radioac-tive nitrogen isotope N16 to the turbines [37].

3-11

(2) Since operating experience and fracture mechanics evaluation indicate that the leak-before-break (LBB) concept is the most likely mode of piping failure [41], it is recommended that reasonably achievable leak-detection procedures be in effect in operating plants. However, studies [53] show that the LBB concept has not been adequately demonstrated to warrant any reduction in the present requirements for inservice inspection of piping.

3.4 Bolting and Threaded Fasteners 3.4.1 Introduction In the past several years the nuclear industry has experienced an increase in the number of reported bolt failures [54-56]. Failure or degradation has been reported in the following three generic areas:

(1) Bolting used for pressure boundary manway and flanges in primary components (steam generators, reactor coolant pumps, valves, etc.): Failure of bolts in these applications could result in a loss-of-coolant accident (LOCA) and thus affect the safe operation of the nuclear plant.

(2) Bolting used in component supports and embedments: These bolts are essen-tial for withstanding transient loads created during abnormal or acciden-tal conditions. Degradation of these bolts affects the reliability of the component support structure following a LOCA or earthquake.

(3) Bolting used in reactor pressure vessel internal applications. It was reported that the consequences of the failure of these bolts may not constitute a significant reduction in public health and safety [57].

However, such failure will affect the availability of the nuclear power plant.

It is important to notice that only two of the three above-mentioned categories are related to the primary pressure boundary. The third category (bolts used in reactor pressure vessel internals) is discussed here for the sake of com-pleteness and is also discussed briefly in Section 2.4 of this report.

Selection of the materials used for bolting applications depends on the specific use. The low-alloy, high strength steels are specified for the primary coolant pressure boundary [58]. Structural steel and low-alloy steel are used for com-ponent supports. High-strength, austenitic, age-hardenable alloys are used for reactor internal bolts [59]. Table 3.3 gives the alloys used in fabricating bolting materials and their specific applications in the nuclear industry.

On the basis of the results of inspections performed and reported by the utili-ties [60], it was concluded that the dominant mechanism for degradation of threaded fasteners in the primary coolant system of PWRs was general corrosion wastage and pitting rather than SCC. However, the study [59] noted the incom-pleteness of the data and stated that SCC should not be overlooked.

3.4.2 Degradation of Low-Alloy Steel Bolting Materials In this part of the review, the degradation mechanisms in bolting materials and the possible actions to be taken to avoid such degradation are discussed.

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I. <

b Table 3.3 Bolting materials and their applications

• Alloy. Component Low-alloy steels -Reactor vessel closure

_ Steam generator (SG) manway j

Reactor coolant pump closure 2 Main steam isolation valve SG support anchors Valve body to bonnet 7.

i X-750 Core baffle bolts.

(nickel-based alloy). Fuel assembly in BWRs Guide tube support pins i

4 A-286 Thermal shield

'l

. (austenitic stainless steel) Fuel assembly in BWRs Core barrel e

• Adapted from Reference 59.

There are two main mechanisms involved in the degradation of the bolting materials fabricated from low-alloy steel: general corrosion and SCC [61].

4 In a number of PWRs, gaskets around pumps have sometimes allowed seepage of small amounts of the primary coolant. The water flashes to steam leaving behind i a sludge or paste' consisting of dissolved boric acid (H3 80 3 ) and lithium hydrox-I ide-(LiOH) from the primary coolant. Studies [62] show that ferritic materials

'are susceptible to' corrosion attack by the H 3B03 + LiOH solution up to a temper-ature of at least 178*C. The rate of corrosion rapidly increases up to 100*C l

and then starts to decline as water is boiled off, but metal loss is still high.

~

Because this type of corrosion is associated with leaking gaskets, leak-tight seals at gasket interfaces should be maintained. Inspection intervals should-be adapted so that they'are consistent with the higher observed corrosion rates such as 130 mils / year [62]. In addition, licensees should continue to improve methods of detecting leakage at bolted connectia%0 to minimize the number of connections where leakage is allowed to persfit t 'r prolonged periods.

Another mechanism involved in the degradat ur, rf ow-alloy steel bolts is SCC.

In many cases [63, 64] cracks were observed in' oesraded bolts. Because inci-dents of boric acid leakage have not been associated with cracking, the observed cracking must be related to other environmental causes. Studies [65] show that cracking is a result of the SCC factor, which could be promoted by the presence of sulfur resulting from the partial decomposition of the'MoS 2 lubricant.

Research on turbine disc steels [66] and on bolting steel [67] has shown that MoS 2 lubricants can have a marked effect on lowering the ultimate tensile strength of high-strength, low-alloy steels when exposed to a steam environment. There-fore, the effect of the interactions between the leaking steam and the sulfur-containing lubricants on the bolts appears to be the reason why bolts fail as a result of SCC.

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,- ,.,m,.-. ,m,--y .,- - - . - __

3.4.3 Degradation of Age-Hardenable Austenitic Bolting Materials The degradation of austenitic high-strength alloys X-750 (nickel-based alloy) and A-286 (stainless steel) in LWR internal bolts has been initiated by either IGSCC or corrosion fatigue.

Alloy X-750 under different heat-treated conditions experienced IGSCC. Thus, an assessment of component heat treatment does not indicate that variations in heat treatment significantly affect resistance to IGSCC [58]. However, labora-tory testing results [68, 69] show that a high-temperature solution annealing (1060-1150*C) plus a single aging treatment (704-718*C) for alloy X-750 may increase the resistance of materials to IGSCC. As for A-286 alloy, most of the service experience has been with material given the 900*C annealing treatment, which has shown considerable susceptibility to IGSCC. It could be that A-286 alloy might respond favorably to higher annealing temperatures that might enhance resistance to SCC. More research is required to determine the effect of various annealing treatments on the susceptibility of A-286 alloy to SCC.

In qualitative terms the role of stress is easy to define. That is, the greater the stress, the greater the chance of degradation by either SCC or fatigue.

The components that have experienced the greatest degradation in service have always been the most highly stressed. Many failures also have contained threads or geometric discontinuities where stresses can be intensified. In exact quan-titative terms, it is difficult to determine the actual stresses on many com-ponents in service. Stresses used in design do not always account for the residual stresses or unanticipated overloads. Some studies [68] show that for alloy X-750, stresses of less than 0.5 times the yield stress might be required to totally avoid IGSCC. For A-286 alloy, cracks have occurred in specimens only loaded to 0.8 yield or higher in a BWR environment [70].

3.4.4 Summary and Recommendations The degradation mechanisms of bolting materials fabricated from low-alloy steels are mainly general corrosion and SCC. Degradation was observed in cases of leakage. To minimize such degradation the following recommenda-tions are offered:

(1) Leak-tight seals should be maintained at gasket interfaces.

(2) Inspection intervals should be consistent with the highest observed cor-rosion rate.

(3) Methods of detecting leakage at bolted corners should be improved.

(4) Users of MoS2 lubricant should give special consideration to reducing leak-age developed in lubricated bolted connections.

For the bolts fabricated from age-hardenable austenitic materials, degrada-tion mechanisms observed are either'the result of IGSCC or corrosion fatigue.

Recommendations for decreasing degradation in these types of bolts include:

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(1) Specific heat treatment should be performed as specified by various re-searchers [68, 69] (heating to 1060-1150 C then aging at 704-732 C) for bolts fabricated from alloy X-750.

(2) Research programs are required to determine the best heat treatment for bolts fabricated from A-286 alloy to provide good resistance to IGSCC.

(3) Design stresses of 0.5 times yield stress and less are required to avoid IGSCC.

3.5 References

1. U.S. Atomic Energy Commission, Advisory Committee on Reactor Safeguards, "The Integrity of Reactor Vessel for LWRs," WASH-1285 (1974).

2. American Society of Mechanical Engineers Boiler and Pressure Vessel Code,Section II, Part A, " Materials Specification - Ferrous Metals," American Society of Mechanical Engineers, New York (1983).

3. J. F. Knott, Fundamentals of Fracture Mechanics, Butterworths, London (1973).

4. M. Vagins and A. Taboada, "Research Program Plan - Reactor Vessels,"

NUREG-1155, U.S. Nuclear Regulatory Commission, Vol.1 (1985).

5. G. Slama and R. L. Jones, " Structural Integrity of Light Water Reactor Components," L. E. Steele, K. E. Stahlkopf, and L. H. Larsson (eds. ),

Applied Science, pp 261-273 (1982).

6. B. Tomkins and J. A. Hudson, " Environmental Factors Influencing the Failure Properties of Pressure Vessel Materials," Proceedings of the International Symposium on Environmental Degradation of Materials in Nuclear Power Systems -- Water Reactors, Myrtle Beach, South Carolina, August 1983, pp 25-52, National Association of Corrosion Engineers (NACE), Houston, Texas (1984).

7. R. L. Jones, "Recent Advances in Corrosion-Assisted Crack Growth Research in LWR Structural Integrity," K. Stahlkopf and L. Steele (eds.), Applied

, Science (1984).

8. C. A. Amzallag, J. L. Bernard, and G. Salama, "Effect of Loading and Metal-lurgical Parameters on the Fatigue Crack Growth Rates of Pressure Vessel Steels in PWR Environment," Proceedings of the International Symposium on Environmental Degradation of Materials in Nuclear Power Systems, Myrtle Beach, South Carolina, National Association of Corrosion Engineers, Houston, Texas (1984).

9. K. Torrenen, T. Saario, H. Hanninen, M. Kemppainen, and S. Salamon, " Fracture and the Role of Microstructure," Vol. II, p 539, Fatigue, K. L. Maures and F. E. Matzer (eds.), Engineering Materials Advisory Services (1982).

10. W. Marshall et al., "An Assessment of the Integrity of PWR Pressure Vessels,"

United Kingdom Atomic Energy Authority, London (June 1982).

3-15

(-

11. K. Klemetti et al. , "On the Role of Inclusions in Environment Sensitive Cracking of Reactor Pressure Vessel' Steels," Proceedings of the Inter-national Symposium on Environmental Degradation of Materials in Nuclear Power Systems, Myrtle Beach, South Carolina, August 1983, pp 368-375, g

National Association of Corrosion Engineers, Houston, Texas (1984).

12. G. M. Pressouyre and C. Zmundzinski, " Influence of Inclusions on Hydrogen Embrittlement," Proceedings of the 22nd Mechanical Working and Steel Pro-cessing Conference, pp 534-553, American Institute of Mining and Metallur-gical Engineering (AIME-ISS), Toronto, Canada (1980).

13. W. H. Cullen and K. Torrenon, "A Review of Fatigue Crack Growth of Pressure Vessels and Piping Steels in High Temperature Pressurized Reactor Grade Water," NUREG/CR-1576, U.S. Nuclear Regulatory Commission (October 1980).

14. W. H. Cullen, H. E. Watson, R. A. Taylor, F. J. Loss, and H. A. Spencer, Proceedings of the Meeting on Subcritical Crack Growth, pp 55-73, International Atomic Energy Agency, Frieburg, Germany (1981).

15. B. Tomkins, United Kingdom Atomic Energy Authority, Report ND-R-848(S),

London (1982).

16. T. A. Prater and L. E. Coffin, Proceedings on Subcritical Crack Growth, pp 640-654, International Atomic Energy Agency, Freiburg, Germany (1981).

17. P. M. Scott and A. E. Traswell, " Structural Integrity of LWR Components,"

-L. E. Steele et al. (eds.), Applied Science, pp 287-310 (1982).

18. U.S. Nuclear Regulatory Commission, " Heavy-Section Steel Technology Program, 5-Year Plan," NUREG/CR-4275, ORNL/TH-9654, Oak Ridge National Laboratory (July 1985).

19. D. D. MacDonald, "The General and Localized Corrosion of Carbon and Low Alloy Steels in Oxygenated High Temperature Water," EPRI NP-2853, Electric Power Research Institute, (February 1983).

20. W. A.. Van der Sluys, " Corrosion Fatigue Characterization of Reactor Pres-sure Vessel Steels," B&W Progress Report on EPRI 1325-1, Babcock and Wilcox (July 1983).

21. W. H. Bamford, D. H. Moon, and L. J. Ceschini, " Studies of Statistically and Dynamically Loaded Cracks in Pressurized Water Environmant," 1983 Annual Corrosion Conference, Paper No. 12, National Association of Corro-sion Engineers, San Francisco, California.

22. "Co-ordinated Research Programme on Analysis of the Behaviour of Advanced Reactor Pressure Vessel Steels Under Neutron Irradiation," IWG RRPC-78181, International Atomic Energy Agency (IAEA), Vienna (1977).

23. J. R. Hawthorne, " Evaluation of IAEA Coordinated Program Steels and Welds for 288*C Radiation Embrittlement Resistance," NUREG/CR-2487, NRL-Memo Rpt. 4655, U.S. Nuclear Regulatory Commission (1982).

3-16

l

24. S. B. Fisher, J. E. Harbottle, and N. B. Aldridge, Proceedings on Dimen tional Stability and Mechanical Behaviour of Irradiated Metals and Alloys, British Nuclear Engineering Society, Brighton, England, (May 1983).

25. R. A. Swift and J. A. Gulya, Welding Research Supplement, pp. 57-68 British Welding Society, London, England (February 1973).

26. A. W. Pense, R. D. Stout and E. H. Kottcamp, Welding Research Supplement, pp. 541-546 British Welding Society, London, England (December 1963).

27. B. Houssin, G. Slama and P. Moulin, Proceedings of the 1st International Seminar on Assessing Structural Integrity of Steel Reactor Pressure Vessel, Berlin, pp. 57-67 (1979).

28. B. 0stensson, " Reliability Problems of Reactor Pressure Components,"

IAEA-SM-218/17, p 303, International Atomic Energy Agency, Vienna (1977).

29. W. A. Van der Sluys, R. H. Emanuelson, and R. J. Futato, ASTM Symposium on Users Experience With Elastic Plastic Test Methods, Louisville, Kentucky, April 1983, American Society of Testing Materials, Meadows Park, Ohio.

30. C. F. Cheng, "Intergranular Stress-Assisted Corrosion Cracking of Austenitic Alloys in Water-Cooled Nuclear Reactors," Journal of Nuclear Material, 57, 11-33 (1975).

31. H. H. Klepfer et al., " Investigation of Cause of Cracking in Austenitic Stainless Steel Piping," NED0-2100, General Electric Company, (1975).

32. NRC Pipe Cracking Study Group, " Investigation and Evaluation of Cracking in Austenitic Stainless Steel Piping of BWR Plants," NUREG-75/067, U.S.

Nuclear Regulatory Commission (1975).

33. H. Flache and R. Ettemeyer, " Piping Repair Case Study in KRB Unit A,"

Transactions of the American Nuclear Society, 3_5, 487-488 (1980).

34. J. C. Danko, "Recent Observations of Cracks in large Diameter BWR Piping:

Analysis and Remedial Actions," Proceedings of the International Symposium on Environmental Degradation of Materials in Nuclear Power Systems - Water Reactors, Myrtle Beach, South Carolina, August 1983, National Association of Corrosion Engineers (1984).

35. U.S. Nuclear Regulatory Commissioner, Office of Inspection and Enforcement, IE, Bulletin 83-02, " Stress Corrosion Cracking in large-Diameter Stainless Steel Recirculation System Piping and BWR Plants" (March 4, 1983).

36. PWR Pipe Crack Study Group, " Investigation and Evaluation of Cracking Incidents in Piping in PWRs," NUREG-0691, U.S. Nuclear Regulatory Commis-sion (1980).

37. R. W. Weeks, " Stress Corrosion Cracking in BWR and PWR," Proceedings of the International Symposium on Environmental Degradation of Materials in Nuclear Power Systems - Water Reactors, Myrtle Beach, South Carolina, August 1983, National Association of Corrosion Engineers (1984).

1 3-17

38. W. Childs et al., " Plant Materials Program: Progress Report, June 1981 to May 1982," EPRI NP-2879-SR, Electric Power Research Institute (1983).

39. NRC Piping Review Committee, " Investigation and Evaluation of Stress Cor-rosion Cracking in Piping of BWR Plants," Vol 1, NUREG-1061, U.S. Nuclear Regulatory Comission (August 1984).

40. F. A. Nichols, " Mechanistic Aspects of Stress-Corrosion Cracking of Type 304 Stainless Steel in LWR Systems," NUREG/CR-3220, U.S. Nuclear Regulatory Commission, ANL-83-15, Argonne National Laboratory (April 1983).

41. NRC Piping Review Committee, " Summary: Piping Review Committee Conclusions and Recommendations," Vol. 5, NUREG-1061, U.S. Nuclear Regulatory Commission, (April 1985).

. 42. J. C. Danko and K. E. Stahlkopf, " Status of Research on Pipe Cracking in

BWRs," Nuclear Safety, 2_3, No. 6, (1982).

t

! 43. U.S. Nuclear Regulatory Commission, Regulatory Guide 1.44, " Control of the l Use of Sensitized Stainless Steel" (1973).

! 44. W. J. ' Shack et al. , " Environmentally Assisted Cracking in LWRs," NUREG/

I CR-4287, U.S. Nuclear Regulatory Commission; ANL-85, Argonne National

! Laboratory (August 1985).

45. W. Y. Chu, J. Yao, and C. M. Hsaio, " Stress Corrosion Cracking of Austentic l Stainless Steel Under Compressive Stresses," Corrosion 40 National AssociationofCorrosionEngineers, Houston, Texas (198U.,302-306, >

46. Electric Power Research Institute "EPRI NDE Center: 1979-1984 Review of Operations and 1984 Annual Report," EPRI-NP-4211 (December 1985).

47. Electric Power Research Institute, Proceeding of the Seminar on Counter measures for Pipe Cracking in BWRs, EPRI W5-79-174 (1980).

I

48. P. Cortland, U.S. Nuclear Regulatory Commission, private communication, l April 1986.

49. M. Fox "BWR Water Chemistry," R&D Status Report, EPRI Journal, Electric Power Research Institute, (January / February 1983).

50. General Electric Company, " Oxygen Suppression in BWRs Phase 2" (Final Report), DOE /ET 34203-47, NEDC-23856-7 (1982).

51. U.S. Nuclear Regulatory Commission, Regulatory Guide 1.56, " Maintenance of the Water Purity in BWRs," Rev. 1 (1978).

52. U.S. Nuclear Regulatory Commission, Office of Inspection and Enforcement, IE Bulletin 82-03, " Stress Corrosion Cracking in Thick Wall, Large Diameter, Stainless Steel Recirculation System Piping at BWR Plants" (October 1982).

53. E. C. Rodabaugh, " Comments on the Leak-Before-Break Concept for Nuclear Power Plant Piping Systems," NUREG/CR-4305, U.S. Nuclear Regulatory Commis-

! sion (September 1985).

3-18

54. U.S. Nuclear Regulatory Commission, Office of Inspection and Enforcement, IE Bulletin 82-02, " Degradation of the Threaded Fasteners in the Reactor Coolant Pressure Boundary of PWR Plants" (June 1982).

55. International Atomic Energy Agency, " Damage of Studs of Steam Generator Heads," IAEA Report No. CS-02, presented at TC/ Workshop on National System for Incident Reporting and Their Interconnections with the IABA-IRS, Budapest, Hungary (June 1984).

56. A. I. Belyaev, "An Abnormal Event in Unit 1 of the Ravno Nuclear Power Plant," Report No. L82-10351, presented at IAEA Technical Committee, Madrid, Spain (November 1982).

57. B&W Owner's Group Meeting on " Reactor Vessel Internals Bolting," Summary Report No. BAW-1784 (1983).

58. U.S. Atomic Energy Commission, Regulatory Guide 1.65, " Materials and Inspections for Reactor Vessel Closure Studs" (1973).

59. A. R. McIlree, " Degradation of High Strength Austenitic Alloys in Nuclear Power Systems," Proceedings of International Symposium on Environmental Degradation of Materials in Nuclear Power Systems Water Reactors, Myrtle Beach, South Carolina, August 1983, National Association of Corrosion Engineers (1984).

60. W. Anderson and P. Sterner, " Evaluation of the Responses to IE Bulletin 82-02: Degradation of Threaded Fasteners in Reactor Coolant Pressure Boundary of PWR Plants," NUREG-1095, U.S. Nuclear Regulatory Commission (May 1985).

61. C. J. Czajkowski, " Corrosion and Stress Corrosion Cracking of Bolting Materials in LWRs," Proceedings of the International Conference on Degra-dation of Nuclear Materials, Myrtle Beach, South Carolina, National Association of Corrosion Engineers (1984).

62. C. J. Czajkowski, " Boric Acid Solution of Ferritic Reactor Components,"

NUREG/CR-2827, U.S. Nuclear Regulatory Commission, (1982).

63. W. H. Koo, " Threaded Fastener Experience in Nuclear Power Plants,"

NUREG-0943, U.S. Nuclear Regulatory Commission (January 1983).

64. C. J. Czajkowski, " Examination of Failed Studs From No. 2 Steam Generator at the Maine Yankee Nuclear Power Station," NUREG/CR-2993, U.S. Nuclear Regulatory Commission (February 1983).

65. V. Pasupathi, D. R. Farmelo, and E. O. Fromm, " Examination of Steam Generator Manway Studs From Oconee Unit 3 Reactor," BCL-585-20, Battelle Columbus Laboratories (1981).

66. C. J. Czajkowski and J. R. Weeks, Materials Performance, 22, No. 3, National Association of Corrosion Engineers, Houston, Texas (1983).

3-19

67. C. J. Czajkowski, " Testing of Nuclear Grade Lubricants and Their Effect on A540 B24 and A193 B7 Bolting Materials," NUREG/CR-3766, U.S. Nuclear Regulatory Commission, BNL-51707, Brookhaven National Laboratory (September 1984).

68. 5. Hattori, Y. Mori, I. Masaoka, R..Sasaki, R. Watanabe, and H. Itoh, " SCC of Age-Hardenable Ni-Base Alloy in High Temperature Water," Presented at the 1981 American Nuclear Society Meeting, San Francisco, California, December 1981.

l 69. T. Yonezawa et al., "Effect of Heat Treatment on Stress Corrosion Cracking j

Resistance of High Nickel Alloys in High Temperature Water," Proceedings of the International Symposium on Degradation of Nuclear Materials, Myrtle Beach, South Carolina, National Association of Corrosion Engineers, p. 345 (1984).

70. B. Bengtssen and S. Korhonen, "Behaviour of A-286 Alloy in BWR Environment,"

Proceedings of the International Symposium on Degradation of Nuclear Mate-rials, Myrtle Beach, South Carolina, National Association of Corrosion Engineers-(1984).

s 3-20

4 STEAM GENERATOR MATERIALS 4.1 Introduction The design life of the steam generator in operating PWRs is 30 to 40 years.

However, operating experience has shown that in the 15 to 20 years that PWRs have operated commercially, there have been reliability problems. Steam gen-erator components including the Inconel-600 tubing have suffered premature degradation by mechanical as well as chemical means. The mechanisms that led to the degradation can be categorized as vibration, fretting, water hammer, wastage, pitting, denting, stress-corrosion cracking (SCC), intergranular attack (IGA), corrosion-fatigue, and erosion / corrosion. As of 1982, approxi-mately 29 units experienced tube denting, 27 units had wastage of tubes, 5 had tube cracking within the tube sheet crevice,13 had tube cracking in the U-bend, and at least 20 experienced water hammer [1]. In some units, some of these problems have been reduced or arrested by equipment modification and/or operat-ing changes. In general, steam generator degradation is caused by lack of attention to maintenance and operating procedures and by chemistry excursion because of condenser leakage.

Depending on the supplier of the nuclear steam supply system (NSSS) and the power plant size, commercial operating PWRs may have two, three, or four steam gener-ators. The steam generator is of shell and tube construction. Steam is produced from the secondary water which extracts heat from the outer surfaces of the steam generator tubes that are heated by the primary water flowing through these tubes.

Consequently, the steam generator serves as the bounds between the. radioactive primary coolant and the less radioactive steam cycle water.

The four systems in operating PWRs that affect steam generator performance and its component integrity are the primary coolant system, the secondary coolant system, the condenser cooling water system, and the makeup water system. In addition, there are two types of steam generators in operating PWRs, recirculat-ing steam generators (RSGs) and once-through steam generators (OTSGs). Mate-rial degradation associated with operation of these types of steam generators is discussed below.

4.2 Recirculating Steam Generators In the RSG feedwater is mixed with 2 to 4 volumes of recirculating water and flows through the downcomer to the bottom of the steam generator, across the tube sheet, and then upward through the tube bundle where steam is generated.

The feedwater comes into the steam generator at about 226 C while the pressure on the secondary side is 770 to 1050 psia under normal operating conditions.

The primary coolant enters the steam generator through the hot leg of U-bend tubes at a pressure about 2250 psia. The inlet temperature of the primary cool-ant is about 316 to 327 C and the outlet temperature about 288 C. At the primary inlet, the temperature gradient across the tube wall is about 36 to 39 C while at the cold-leg outlet side a 11.2 to 13.8 C temperature gradient exists across the tube wall.

4-1

i 4

The U-bend in an RSG is usually made of nickel-based alloy Inconel-600, which is an austenic Ni-Cr-Fe alloy. Mill-annealed Inconel-600 tubing was used to manu-

, facture RSGs before the late _1970's. Since then, thermally treated (15 hours1.736111e-4 days <br />0.00417 hours <br />2.480159e-5 weeks <br />5.7075e-6 months <br /> i at 700*C) Inconel-600 tubing material has been used to improve resistance to SCC.

The tube sheet material continues to be stainless-steel clad carbon steel. The tube-support plates for the early RSGs were made of carbon steel with round i

drilled holes. After these were found to be especially susceptible to denting,

[ 405 or 409 stainless steel has been used to manufacture the support plates. In addition, a nonround broached hole with lands, or an egg crate structure, has been used instead of the previous drilled-hole design. Early RSGs used coor-dinated phosphate water treatment, which is the conventional water treatment method used for fossil-fired boilers with similar steam pressure. After corro-sion problems were experienced as a result of the coordinated phosphate water treatment, a change was made to an all-volatile treatment (AVT) of water. The knowledge _of the materials of construction and operating environment of RSGs will help to understand the various materials degradation in operating RSGs.

4.3 Materials Degradation in Operating RSGs There have been extensive studies of various corrosion phenomena that have been

- experienced in operating RSGs [2-4]. The corrosion problems in RSGs are illus-trated in Figure 4.1, and a brief description of the field observations and the causes or mechanisms of corrosion is discussed.

4.3.1 Wastage Wastage or tube thinning, as it is sometimes called, occurs primarily in the sludge pile region above the tubesheet and in areas of high heat fluxes in operating RSGs under coordinated phosphate secondary water chemistry control.

In a few instances tube thinning has occurred within the first tube-support plate penetration of some RSGs under AVT chemistry control. In these regions, corrosive conditions developed because of the chemical concentration processes resulting from alternate wetting and drying. Several European RSGs that use phosphate chemistry control for the secondary coolant with Incoloy 800 tube material also have experienced tube-thinning problems [5].

A second type of tube-thinning problem, caused by resin breakthrough and its subsequent thermal decomposition, occurred on the periphery of the tube bundle limited to the cold-leg side in the first and second tube-support plate crevices

[6]. However, this seems to have been an isolated case'rather than a major damage degradation phenomenon in operating RSGs. Furthermore, since converting to AVT chemistry control from the previous phosphate chemistry control practice, wastage or tube thinning of any kind is no longer a serious concern in operating '

RSGs.

4.3.2 Pitting Although a pitting problem has been observed from time to time in laboratory tests under faulted AVT conditions as well as on support plate materials such as carbon steel and Types 405 and 409 ferritic stainless steels [2], it has not been considered a serious threat to the integrity of operating RSGs. How-ever, in the early 1980's severe pitting occurred in the tube sheet sludge pile regions of several operating RSGs. The acie thloride environment' caused by condenser leakage is the main cause of pitting attack. The pressure of copper 4-2

U-bend g Fe3 40 Wagnetite e 3 (O i  ; l'-

i al

• Il@(C D.y' '/ ,

U-bend ovality

-' s .

resulting from -

,' Denting ,

flow slot I '

( -

1 j{ i

} deformation ,,'

)

~~.-

'p . p 'r Uf Denting-induced Support plate , i c deformation of tube holes ,

y 7[ h j support plate

.2. :.1,, flow stots oo o X" , _

e n $Yi y samm o T Tube to Tubesheet Thinning T

) Crevice [

(( g h Studge P le 'p - *

/A<===I V' ' IGA Cracking Figure 4.1 Typical problem areas in PWR steam generators l

4-3

oxide _resulting from air inleakage further enhanced the pitting corrosion.

There are many ways to reduce or eliminate pitting corrosion in operating RSGs.

The most effective approach is to maintain condenser integrity thus eliminating ingress of contaminants into the secondary system heat exchangers. Performing sludge lancing and/or chemical cleaning to remove the sludge pile also are effective means to reduce pitting attack.

4.3.3 Denting After the industry started to use the AVT secondary water chemistry control, denting of steam generator tubes in the tube-to-tube support plate crevice regions was the most severe material degradation problem that threatened the safe operation of RSGs. Denting is the term used to describe the deformation of the steam generator tubes caused by the corrosion of the support plate mate-rial in the tube-to-tube support plate region. The tube-support plates in most operating RSGs are made of carbon steel plate. As pointed out by Potter and Mann [7], the volume of iron increases by about a factor of three when it corrodes and forms magnetite, a form of iron oxide. As a result, the pressure building in the tube-to-tube support crevices caused the distortion of tubes observed in operating RSGs. Because the nonsymmetrical growth of the oxide, deformation or denting of tubes is irregular, the resultant strains on tube wall as well as on support plates are not uniform.

Although the problem of denting has been experienced in seawater- and freshwater-cooled PWRs, statistics show that the problem was especially prevalent in sea-water or brackish water plants because of condenser leakage. In addition, it was found that denting increases linearly with respect to total chloride expo-sure. There have been reports about the chemical reactions involved in the dent-ing process [2]. Operating experience with RSGs also indicates that sulfuric acid generated because of resin breakthrough was responsible for denting problems observed in several operating plants. The role of copper and nickel in the denting process is not known exactly. However, it is generally believed that these species increase the acidity in the crevice region.

4.3.4 Intergranular Attack Intergranular attack (IGA) and secondary side-initiated SCC have been observed in tube sheet crevices and tube-support plate intersections. The tube-to-tube-sheet' crevice, created by the design of the partially expanded tube joint, is an ideal location for caustic and other impurities that have been introduced into the secondary water to concentrate to the level that leads to IGA and SCC.

The problems are minimized in the new design of the tube-to-tube-sheet joint by expanding the tubes for the full length through the tube sheet. This new design eliminates crevices. For those operating RSGs with the partially expanded tube-to-tube-sheet joints, it is recommended to carry out crevice-flushing operations during outages to reduce the caustic and other impurities that concentrate in the

! crevices. In addition, maintaining condenser integrity and reducing air inleakage l are effective measures to reduce secondary side-initiated IGA and SCC.

i 4.3.5 Primary Side-Initiated Cracking l Primary side-initiated steam generator tube cracking was observed mainly at the i apex of the U-bends and also at the roll transition regions in the tube sheet l area. Residual stresses resulting from fabrication as well as from denting that 1

4-4

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occur at the tube-support plate intersections are significantly higher than stresses that occur in the straight section of the U-bend. Therefore, these regions with higher stress and strain combined with the susceptible microstruc-ture of the Inconel-600 tubing, have caused the observed primary side cracking problem. At the present time, this problem in operating RSGs has been con-trolled or eliminated by minimizing the denting problem described above and by using thermally treated Inconel-600 tubing with a microstructure more resistant to cracking.

In addition to the above chemically induced degradation problems, RSGs in operat-ing PWRs also have experienced some mechanically induced degradation problems such as fatigue, wear, and clad separation. These problems are discussed in numerous review papers [1-4].

4.4 Once-Through Steam Generators Unlike the RSGs, in the OTSG the primary water flows through straight Inconel-600 tubes instead of the U-bend tubes. Water enters the steam generator through a feed annulus above the ninth tube-support plate. This water, preheated to satu-ration by steam coming from the tube bundle area, flows down the annulus, across the lower tube sheet, and upward into the tube bundle where it becomes steam.

It reaches 100% quality in the ninth and tenth tube-support plate regions. The superheated steam leaves the OTSG at about 925 psia. The primary coolant enters through the top of the steam generator at above 316 to 332 C and leaves at the bottom of the steam generator at about 291 to 302 C.

Structural materials and the tubing alloy for the OTSGs are similar to those used in the RSGs. Like the early RSGs, support plates in the OSTGs also are made of carbon steel. However, the design differs in that broached holes with three lands supporting the steam generator tubes were used in the OTSGs instead of round holes commonly seen in early RSGs. This design combined with the AVT secondary water chemistry control, which has always been the case for OTSGs, helps to prevent the denting problems from occurring in operating OTSGs.

4.5 Materials Degradation in OTSGs Degradation problems experienced in OTSGs are illustrated in Figure 4.2 and the main types of degradation and the associated causes are briefly discussed below.

4.5.1 Erosion / Corrosion Although it is not wide spread, erosion / corrosion damage has been observed in operating OTSGs in isolated instances. This type of degradation occurs mainly near the top of the steam generator around the periphery. The damage manifests itself in the form of a large number of deep cut grooves and pits. The exact cause of the damage is not known at the present time, but it is believed that unvaporized water with a high concentration of corrosive agents plays an impor-tant role in the observed degradation [1]. Because it is not a generic problem in operating OTSGs, the damaged tubes are simply plugged and taken out of service.

4.5.2 Secondary Side Tube Cracking Secondary side-initiated cracking problems were experienced in the upper tube sheet crevice regions at several operating OTSGs. The cause of the cracks were 4-5

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4-6

thought to be related to sulfur attack. The possible source of sulfur contami-nation was from resin break through or leakage from condensate polisher regen-eration chemicals. The corrective actions involve control of secondary water chemistry and prevention of contamination by resin breakthrough or other leakages.

4.5.3 Primary Side-Initiated Tube Cracking The sulfur-induced primary side tube cracking was an isolated incident and it occurred mainly at one plant. The intergranular stress-corrosion cracking (IGSCC) occurred in the upper tube sheet region near or at the roll-transition area. The corrosion attacks most rapidly at the air / water interface and curing layup. Sulfur in its reduced form has been identified as the active corrosive agent causing the tube cracking. The source of sulfur contamination has been traced to leakage of sodium thiosulfate which has been used in the past for iodine removal in the containment spray header during post-accident conditions.

The corrective actions involves removal of sodium thiosulfate or post-accident containment spray additive and avoidance of sulfur contamination from any other source.

4.6 Summary The various materials degradation resulting from corrosion attack in operating PWRs seems to be the result of a combination of inherent susceptibility of Inconel-600 tubing to stress-corrosion cracking, inadequate design and fabrica-tion, and poor operating and layup practices. Responding to these challanges, NSSS vendors and utilities have been working together to come up with corrective measures in operating practices and major modifications in equipment and system design to improve the steam generator integrity. However, it is important to note that no quick and simple solution will likely be found in the near future to eliminate all the problems that have confronted the industry. Therefore, continuing efforts in the areas of design improvements, materials optimiza-tion, and stringent operating and layup practices are essential in maintaining steam generator integrity.

4.7 References

1. S. J. Green et al., " Steam Generator Performance History," Chapter 2, Steam Generator Reference Book, Steam Generator Owners Group and Electric Power Research Institute, May 1985.

2. S. J. Green and J. Peter N. Paine, " Materials Preference in Nuclear Pres-surized Water Reactor Steam Generators," Nuclear Technology, 55, pp 10-29, l No. 1 (1981).

l

3. L. Frank et al., " Regulatory Bases for S/G Operation and Operating Experi-ence Update 1982-1983," Session 1.1 Proceedings on Specialist Meeting on j Steam Generators, Stockholm, Sweden (October 1984).

l

4. D. G. Eisenhut et al. , " Summary of Operating Experience with Recirculating Steam Generators," NUREG-0523, U.S. Nuclear Regulatory Commission (January 1979).

I l 4-7

5. G. Schucktanz, L. Stieding, and R. Ries, "KWU Steam Generator Experience,"

International Conference of Materials Performance for Nuclear Steam Gener-ators, St. Petersburg, Florida, October 5-10, 1980.

6. C. L. Williams and S. J. Green, " Thermal and Hydraulic Aspects of PWR Steam Generators," presented at ANS/ASME International Topical Meeting on Nuclear Reactor Thermal Hydraulics, Saratoga, New York (October 1980).

7. E. C. Potter and G. W. Mann, "The Fast Linear Growth of Magnetite on Mild Steel in High Temperature Aqueous Conditions," British Corrosion Journal (January 26, 1965).

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5 STEAM TURBINE MATERIALS d

5.1 -Introduction Corrosion fatigue and stress corrosion cracking of steam turbine components were

, . discovered in many foreign plants some 10 to 20 years before the first cracking

incident occurred in the United States. Although degradation of turbine blades, j rotor shafts, and numerous other components have been experienced in many domes-i tic plants since the early 1970's, stress corrosion cracking of steam turbine

,' discs was by far the most severe problem in operating nuclear power plants.

, Beginning in the late 1970's, the Electric Power Research. Institute (EPRI)

. launched an extensive program to address this issue. As a result, a series of j reports on operating history, material characterization, and simulated labora-tory tests were published in 1982 (1-7].

i j In general, the failure mode for turbine discs was intergranular stress corro-4 sion cracking (IGSCC). In the case of turbine blades, the failure mode has been

) predominantly corrosion-fatigue. Cracking occurred at locations such as the

] keyways, bores, web faces, and rim attachment areas. The cracks are character-

! istically filled with iron oxides. However, impurities such as hydroxides i chlorides, sulfates, copper, and copper oxide also were detected occasionally in I

or near the cracks. In some cases, caustic deposits also have been identified.

At the present time, no specific corrodent has been singled out as the sole

! contaminant uniquely responsible for the turbine disc cracking.

! Steam turbine discs are made of primarily 3.5 NiCrMoV steels while the turbine blades are predominantly made of AISI Type 410 ferritic stainless steel. The rotors of low pressure turbines in LWRs are designed to operate at 1,800 rpm.

The steam inlet temperature for low pressure turbines is about 149'C. The

! Westinghouse rotors employ discs with semicircular keyways and tightly-fitted I keys, while the General Electric rotors employ rectangular keyways with loosely-fitted keys. Various material degradation problems in nuclede low pressure j turbine components are briefly discussed in this section.

I 5.2 Stress-Corrosion Cracking of Steam Turbine Discs i Although cracking has been found primarily at four locations on low pressure l turbine discs, it has occurred most frequently in the key and rim attachments, j Cracking can occur at some intermediate stage where the steam becomes saturated

j. and wet, and it can occur in more than one location in any one disc.

i The discs have their highest stresses along the surface of the inner bore. At 1

the keyway where the discs are locked in place, these high stresses are increased even more, and stress corrosion cracking is often initiated at this point be-cause of the high stress concentration. In addition to this area, initiation of l stress corrosion cracking also has been observed from keyways, from the bore, j from the hub of the face of the discs, and in the blade attachment area. Details about location and orientation have been described and illustrated by McMinn et al. [8].

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5.2.1 Keyway Cracks Metallurgical characterization of keyway cracks in three discs at two PWRs h:s shown that cracking had occurred in the No.1 discs from the generator and g;vernor ends [7]. The keyway regions of hub samples showed well-defined crack indications. In one case, the crack extended almost the whole length of the k:ying and over the corner at the open end and outward on the outlet face of the hub for a distance of 18 mm.

Destructive examination of the disc samples indicated that the crack surfaces were covered by a tight, dark, oxide film. Furthermore, cracking morphology indicated that the initial stages of cracking involved initiation and growth of several independent colinear cracks and the cracks showed many short branches.

Both the main crack and the branch cracks were intergranular. Significant pitting of the keyway surfaces in the vicinity of the cracks also was evident.

5.2.2 Rim Cracks Rim cracks have been found in 30 discs from 13 rotors in 17 plants (1 BWR and 16 PWRs) [8]. All the rotors were of the same general design and were made of the same alloy with the same heat treatment and similar tensile properties. Two types of cracking morphology were observed. In the oxide entry fir-tree config-urations, cracks propagate across the fir-trees or " steeples." In the axial-radial configuration, cracks propagate into the disc. In all cases, pitting corrosion occurred within the steeple serrations. The cracking was predomi-nantly intergranular and developed by the initiation of multiple independent cracks along the steeple serrations. Pitting was observed and the cracking cxtended through the pitted regions, but there was no evidence to indicate that the pitting necessarily developed before the cracking or that the cracking was initiated by the pitting.

Analysis of the corrosion deposit showed that it consisted mainly of iron oxides with significant amounts of copper. This is consistent with the results reported in other investigations on rim cracks. Another characteristic of the rim cracks is that they tend to start within the blade grooves; thereby being difficult to detect during inservice inspection, especially at early stages.

5.2.3 Web-Face Cracks Web-face cracks have been observed in once-through steam generators (OSTGs) at one PWR plant. The cracking occurred on the outlet faces of the No. 1 discs cnd on the inlet faces of the No. 2 discs, both locations being immediately downstream of the saturation line. The cracks were predominantly intergranular cnd had essentially the same appearance as cracks found in rims and keyways of other discs. However, no pitting or other corrosive attack was found on the intergranular or disc faces. In addition, analyses indicate that the cracking was not related to material abnormality or chemical contamination; the orienta-tion of the face cracks indicates that the cracking occurred in response to a sustained tangential stress. However, no information concerning the as-installed hoop stress was available. On the bases of these observations, the exact cause of the cracking was not known.

l 5-2 l

5.2.4 Bore Cracks Bore cracking has been limited to four discs in four rotors in four PWRs with recirculating steam generators. Like the other types of cracking on turbine discs, bore cracking occurred downstream of the normal saturation line location.

Although there was no evidence of pitting or general corrosion of the bore sur-face, significant corrosive attacks are found on the intergranular crack surface.

It is believed that the failure resulted from long-term growth by IGSCC of Eultiple cracks at the bore surface. In one instance, catastrophic fracture occurred when one particular crack reached the critical crack size for fast fracture under the influence of the stress / temperature conditions encountered during startup. The critical size of 49 mm is consistent with the inherent low toughness of the material.

5.3 Factors Affecting the SCC of Turbine Discs The term " stress-corrosion cracking" usually relates to failure which requires conjointly a tensional stress and a specific environment. In the case of the intergranular characteristics of the turbine disc cracking, a unique grain boundary microstructure and alloy condition also are necessary for the observed failure mode. Consequently, any workable failure mechanism must include these three factors: stress, environment, and susceptible alloy. The large number of variables and specific conditions entering the phenomenon of SCC support the conclusion that the mechanism of SCC is very complicated and not easily derived.

The chemistry, metallurgy, and mechanics of metals and alloys undergoing fracture in an aggressive environment are all essential ingredients that lead to SCC.

5.3.1 Effect of the Yield Strength The yield strength of steam turbine disc materials is easily determined by se-lecting the appropriate tempering temperature. Low alloy steels used for low-pressure rotor discs in nuclear power plants vary in yield strength from over 90 ksi to over 190 ksi. In general, the stress corrosion crack growth rate increases with increasing yield strength as well as with increasing fracture toughness of the structural alloys.

5.3.2 Effect of Stress Laboratory tests on smooth tensile specimens have shown that reducing the tensile stress will increase the time to failure because the crack initiation time was increased as the result of reducing tensile stress. The beneficial effect of low tensile stresses extending the lifetime of a naterial in a stress corrosion situation is systematically investigated by most turbine vendors. Lowering stresses and stress concentrations through design modifications and operational changes have become a major effort in attempts to overcome the steam turbine disc SCC problem.

5.3.3 Effect of Cyclic Loads In a noncorrosive environment, small and less frequent cyclic loads have little or no effect on the turbine disc alloy. However, if such a cyclic load is superimposed with a high mean stress in an aggressive environment, it would 5-3

almost certainly contribute to turbine disc cracking. The combined effects are the main cause of corrosion fatigue.

5.3.4 Effect of Temperature Statistics on crack growth rate as a function of temperature indicate that higher temperatures result in higher crack growth rates of the turbine disc material.

Because the service temperatures for turbine discs vary significantly during startup and normal operation especially at or near the cracked area, it is difficult to predict accurately the crack growth rate for turbine disc cracking.

5.3.5 Effect of Electrochemical Potential Stress-corrosion crack growth rates depend strongly on the electrode potential of the cracking steel. For instance, 100 mV differences in electrochemical potential can increase the crack growth rates by a factor of 10 or more.

Because the electrochemical potential of a cracking disc while in service also is affected by temperature and purity of the steam, it is extremely difficult to know the exact potential of a turbine disc. Also, because of the changing temperature and varying chemistry at the localized cracked region, it is dif-ficult to quantify the electrochemical effect on the crack growth rates of a turbine disc under operational condition.

5.3.6 Effect of Steam Chemistry Sodium and chlorine were frequently identified in deposits within or near cracks on turbine discs. Because high sodium and/or chloride levels resulting from condenser leakage have been a frequent occurrence in many plants in which low-pressure turbine disc cracking has been found, it can be concluded that secondary coolant impurities such as sodium and chloride played an important role in tur-bine disc failure under operational conditions. Air in-leakage also is expected to have contributed to the observed turbine disc failures.

5.4 Failure of Turbine Blades Almost all steam turbine blades are made of AISI Type 410 stainless steel. The 12% chromium steel has been known to have a fairly good fatigue strength even at high mean stresses. However, because of steam chemistry excursion and cyclic loads caused by vibrations, occasional failure of steam turbine blade by corrosion fatigue has been reported. In May 1985 a generator fire caused by turbine blade failure forced an extended outage of a foreign PWR with recirculating steam generators. Laboratory tests have shown that in the presence of impurities such as chlorides or hydroxides, the fatigue life of the 12% chromium alloy decreased significantly. However, in pure steam and condensate the alloy retained its fatigue strength reasonably well. Consequently, good steam chemistry and mini-mizing vibrationally induced cyclic loads are essential steps toward preventing turbine blade failure.

5.5 Conclusion The failure mode of low pressure steam turbine components has been of an inter-granular nature and has occurred by a SCC mechanism. For turbine rotors and discs, the failure occurred predominantly at or near the steam saturation line where impurity concentration occurred. Experimental tests and microstructual 5-4

l characterization have shown that the structural alloys of turbine components are i inherently susceptible to SCC and/or corrosion fatigue under the steam turbine 4 operational conditions. Therefore, unless a new generation of structural alloys is discovered, minimizing component stresses through design modifications and fabrication improvements and reducing steam chemistry excursions through tighter

, coolant chemistry control and operational changes are essential steps toward reducing turbine failures.

5.6 References

1. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking

Experience," Vol.1

Literature and Field Survey, EPRI NP-2429-LD, Elec-tric Power Research Institute (June 1982).

2. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking l Experience," Vol. 2: Data Summary and Discussion, EPRI NP-2429-LD, Elec-tric Power Research Institute (June 1982).

3. F. F. Lyle, Jr. , and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking Experience," Vol. 3: Stress Corrosion Cracking of Low-Alloy Steels, EPRI l NP-2429-LD, Electric Power Research Institute (June 1982).

4. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking Experience," Vol. 4: Factors Determining Chemical Composition of Low-Pressure Turbine Environments, EPRI NP-2429-LD, Electric Power Research Institute (June 1982).

i 5. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking i Experience," Vol. 5: Characteristics and Operating Histories of U. S.

Power Plants, EPRI NP-2429-LD, Electric Power Research Institute (June 1982).

6. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking i Experience," Vol. 6

Description of Turbine Rotor Models, EPRI NP-2429-LD, Clectric Power Research Institute (June 1982).

l 7. F. F. Lyle, Jr. and H. C. Burghard, Jr. , " Steam Turbine Disc Cracking Experience," Vol. 7: Meta 11urigal Analysis of Cracked Discs from 10 U. S.

Power Plants, EPRI NP-2429-LD, Electric Power Research Institute (June I 1982).

l 8. A. McMinn et al., " Cracking of low-Pressure Steam Turbine Rotor Discs in Nuclear Power Plants," pp 243-272, Proceedings of International Conference

on Degradation of Nuclear Materials, Myrtle Beach, South Carolina, National i Association of Corrosion Engineers (1984).

9. H. C. Burdhard, Jr., " Metallurgical Analysis of Rim Cracking in an LP Steam Turbine Disc," EPRI NO-1532, Electric Power Research Institute (September 1980).

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6 CONDENSER MATERIALS 6.1 Introduction Any degradation in the condenser materials will affect the availability, and hence the economics, of the power plant. In a study made in the United States

[1], the Electric Power Research Institute (EPRI) has estimated that the loss of equivalent unit availability that is directly attributable to condenser problems is about 3.8%. This costs the U.S. utilities $600 million annually for power re-placement. Moreover, condenser performance significantly influences the perfor-mance and availability of many other components in the power plant, thus it af-fects the whole plant. This is particularly critical in nuclear power plants, where repairs are more expensive and plant shutdowns may be prolonged because of radiation protection requirements. Analysts have not yet quantified the costs of indirect effects of condenser problems, but these costs are probably as large or larger than the costs of direct losses [2].

The power plant components that may be affected as a result of condenser leakage include steam generators, turbines, pumps, feed lines, and feedwater heaters [3].

In a nuclear power plant, steam generators are probably the components most susceptible to damage by contaminants from condenser leakage. The denting phenomenon is believed to be influenced by the presence of chlorides. The intergranular corrosion of the steam generator tubes has been ascribed to the presence of caustic material. Caustic materials could originate from phosphate compounds used to condition steam generator water or could be the result of de-composition of bicarbonate and carbonate salts entering via condenser leakage.

Concerning turbines, contaminants of foreign substances in steam phase in concentrations, even in parts per billion level, can condense as concentrated solutions on turbine component surfaces. These concentrates, especially'chlo-rides, accelerate the degradation mechanisms occurring in turbines such as pitting attack, stress corrosion cracking, and corrosion fatigue. Also, the inleakage of carbon dioxide through the condenser may cause a marked reduction in feedwater pH. This, in combination with high flow velocities in downstream piping could accelerate erosive / corrosive mechanisms in feed lines, pumps, and feedwater heaters.

Because various types of cooling water (from sea, river, lake, or tower) and different condenser tube materials are used, a variety of degradation mechanisms could be encountered in power plant condensers [4]. Various environmentally related degradation mechanisms of condenser materials will be discussed herein.

Some of the methods of controlling such degradation also will be summarized.

6.2 Erosion / Corrosion Erosion / corrosion is recognized as a major cause of failure of copper-alloy steam condenser tubing [5]. This phenomenon is defined as an accelerated attack caused by the removal or breakdown of protective surface films resulting from the flow of cooling water. Thus it occurs only in areas where the turbulence in-tensity at the metal surface is high enough to cause mechanical or electrochemical 6-1

disruption of the protective oxide film. This problem is known to occur only in condensers with copper-alloy tubes.

A number of factors contribute directly or indirectly to erosion / corrosion.

These include [6] fluid flow, properties of passive film on the alloy surface, the presence of pollutants that may exacerbate local breakdown, metallurgy, and galvanic effects. The presence of entrained air, sand, fouling products, or other foreign bodies lodged in the tubes would increase the erosion / corrosion process, as could geometric factors, particularly in the inlet program.

The turbulence intensity at tube inlets is much higher than it is several feet down the tubes, resulting in the phenomenon of inlet and erosion / corrosion or what is called " inlet impingement attack." This is known to be the most prev-alent cause of cendenser-tube failures [6]. The sudden increase in corrosive attack at a critical flow velocity is attributed to shear stripping of the pro-tective film from the metal surface. The magnitude of the surface shear stress is a function of system geometry and the distance from the tube entrance [6].

The presence of pollutants, especially sulfides, has a major influence on the susceptibility of condenser tubes to erosion / corrosion. Most of the research to date [7,8] has assumed that the sulfides present are the result of point sources of industrial pollutants, although now it is well established that bio-films also are capable of producing sulfides [9]. Moreover, the biofilms pro-duce the sulfide right at the metal / water interface where it can do the most damage.

Abrasive particles (e.g., sand), which may penetrate the boundary layer at the wall water interface, can greatly increase the rate of erosion / corrosion by rupturing the protective film at flow velocities well below that which corre-sponds to the critical shear stress. It was found that the severity of the attack of the abrasive particles increases with increasing particle size [11].

This is due to the higher momentum possessed by big particle sizes that allows for higher penetrating power. Erosion / corrosion attack also was found to be highly enhanced by specific sand content in the penetrating water [11].

The following recommendations are some proposed to avoid or at least decrease the erosion / corrosion in condensers:

(1) The use of more protective alloys: Minor alloying elements capable of increasing protective film on the metal surface may be beneficial in re-ducing the severity.of erosion / corrosion. It was found that additions of iron [12] or Cr [13] substantially reduces the rate of erosion / corrosion of copper-nickel alloys.

(2) The addition of inhibitors: Some copper alloys benefit from periodically dosing the water with ferrous ions (usually added as ferrous sulfate) [14].

Ferrous ion is believed to provide protection by either (or both) of the following mechanisms: (a) precipitation of sulfide as fron sulfide and (b) modification of the structure of the passive film providing a more effective barrier between the metal and the environment. Because intermit-tent additions of iron ion are effective, it appears that the second mech-anism is operative. This mechanism also gives an explanation of the bene-ficial effect of iron in the alloy matrix. .

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6-2 i

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(3) Good design of the water box: A major reason for condenser failure is poor water box design which may lead to severe turbulence in the box itself and in the inlets to the tubes. For instance, a design that includes the installation of vanes to direct the flow perpendicular to the tube sheet can be effective in reducing turbulence at the tube inlet [6]; inasmuch, a design that avoids elbows or venturies in the inlet lines may be helpful.

Tube inserts also have been used to circumvent inlet end erosion / corrosion problems. A tube insert is a tightly fitting internal sleeve, typically 15-30 cm long, made from an erosion / corrosion resistant material that shields the susceptible tube ends. However, unless there is a smooth transition between the end of the insert and the tube, the insert can create turbulent conditions and promote erosion / corrosion further down the tube.

(4) Use of screens to prevent foreign material: Screens have been effectively used in the inlet end.

(5) Periodic cleaning: Periodically reversing flow (backwash), manually cleaning the copper alloy tubes with brushes, balls, and scrapers, and preventing bio-fouling by chlorination or thermal shock are some of the cleaning techniques that have been used.

(6) Cathodic protection: Recent experiments demonstrate that inlet end erosion /

corrosion can be prevented by cathodic protection (15]. This method pro-vides protection by imposing thermodynamic immunity on the base alloy. The greatly reduced corrosion rate for cathodically protected alloys [16,17]

indicates the effectiveness of this method for combatting inlet end erosion / corrosion.

6.3 Corrosion Caused by Water Pollution The adverse effects of water pollution on corrosion of copper alloys were rec-ognized long ago [17] although the attack mechanisms are still not completely understood [18]. However, it is generally agreed that sulfide is the main aggressive species in the polluted water [4] although polysulfides and elemental sulfur [19] may produce the same effect. Depending on the alloy composition and the precise details of the exposure history, sulfide polluted waters may accelerate or induce erosion / corrosion, pitting, or intergranular attack in copper alloys and may change the polarity of galvanic couples.

As little as 10 mg/m3 (nearly 10 ppb) sulfide in the cooling water can have a detrimental effect [4]. Concentrations in this range often have been measured in polluted harbors. In addition to industrial pollution, biofouling can be a source of sulfides on the surface of copper alloys [9]. There is ample evidence (19,20) that sulfide is not particularly aggressive in the absence of oxygen.

Recent experiments (18] have demonstrated that accelerated corrosion rates occur when the polluted water contains both oxygen and sulfide.

No copper alloy to date seems to be resistant to sulfide attack, and the re-lative performance of copper alloys in polluted waters seems to depend on precise environmental conditions. The penetration rates in polluted waters can be extraordinarily high. It was reported [4] that a 6-mm thick wall of 90/10 copper-nickel alloy pipe was penetrated af ter an estimated period of 4 months in service.

6-3

If the problem is identified as sulfide induced attack and there is no obvious method of eliminating the source, the following are some of the methods recom-mended to reduce or prevent future problems [4,18]:

(1) If the sulfide forms within the condenser by putrefaction of organic mat-ter during shutdown periods, it might be eliminated by turning the pumps on for a short period every day to flush the system with fresh water.

(2) The incoming water can be treated with ferrous sulfate solution or another form of ferrous ion. This will help precipitate sulfides as flocculent ferrous sulfide. The excess ferrous ions would be available for condi-tioning the copper alloy surface and improving its corrosion resistance, as explained in the last section.

(3) Precondition and periodically recondition the copper alloy surfaces to make them more resistant to sulfides.

(4) Cathodic protection could be effective in preventing sulfide corrosion of tube sheet, water box, and tube inlets and outlets.

(5) Periodic cleaning of the copper alloy tubes may reduce the risk of sulfide attack by avoiding the presence of sulfate-reducing bacteria which produces sulfides under debris and deposits where the oxygen content is low.

6.4 Dealloying Dealloying (or selective leaching) is the preferential removal of the more active element from a solid solution alloy by electrochemical processes, leaving behind a surface layer that is rich in the more noble alloying element. The most com-mon example is the selective removal of zinc in brass alloys (dezincification).

Similar processes occur in copper-nickel, copper-aluminum alloys and cast iron

[21]. Stainless steel and titanium are not susceptible to this form of attack.

Relatively few failures in condensers have been reported after dealloying

[22,23]. However, this kind of failure, if it occurs, can be catastrophic; thus the subject warrants some discussion.

Cast iron water boxes occasionally undergo a form of dealloying referred to as graphitic corrosion or graphitization. In this process, the iron is selectively removed from the cast iron leaving a graphite layer [21]. This attack can be deleterious following graphitization of the water box because severe galvanic corrosion of copper-base alloy tubes and tube sheets can occur [22].

In copper-base alloys, dealloying depends on the alloying elements present.

For example copper-nickel is more resistant than copper-zinc [24]. Moreover, in copper-zinc alloys, dealloying is inhibited by addition of arsenit., antimony, and phosphorous [25]. For single phase copper alloys, heat treatment seems to have little effect; whereas, proper heat treatment for multiphase alloys is important to avoid dealloying [26].

Dealloying is generally found beneath deposits and at hot spots and is promoted by stagnant conditions. This situation suggests that a differential oxygen cell promotes the attack [27]. Oxygen is depleted under the deposits, which reduces ,

the protective film on the metal while simultaneously supplying a reducing 6-4 l

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species on the badly exposed surface. Hydrolysis, which lowers the pH, and chloride concentration beneath the deposits tend to depassivate materials.

Studies [23] showed that dealloying increases with increasing water temperature.

This dealloying increase is very rapid for some alloys (e.g. , Muntz Metal and Naval Brass) and quite limited for some others (Red Brass). It was found too that dealloying increases with increasing chloride concentration.

Some techniques have been suggested to avoid dealloying attack. Such techniques include: using stainless steel or titanium, frequent tube cleaning, minimizing the occurrence of stagnant conditions, and cathodic protection.

6.5 Crevice Corrosion and pitting Crevice corrosion and pitting in condensers are usually discussed together [4,28]

because both are a localized type of corrosion and the environmental parameters that affect one invariably affects the other in a similar fashion. Pitting is the localized corrosion that occurs following the local breakdown of a protective surface film on a fully exposed metal surface. Crevice corrosion, although a localized type of corrosion, is the corrosive degradation that results from the shielding of some parts of the metal surface from full exposure to the environ-ment because of the close proximity of the metal to a foreign material such as a fastener, a gasket, mud, or bacteria. Such covered areas are called crevices.

Copper-based alloys are subjected to a kind of crevice corrosion mechanism called metal-ion concentration [29]. This mechanism occurs if two areas of the alloy are in contact with solution having a different Cu+1 ion concentration. This may occur in the presence of a covered area on the surface of the metal. The covered area becomes a cathode because of the high concentration of the released Cu& ions that do not move as fast as those ions formed on the uncovered area that becomes the anode. This difference in ion concentration develops a difference in potential which causes the flow of an electrical current. The corrosion in this case occurs in the uncovered area and it is more severe if the uncovered area is small compared to the covered one. Failures have been reported resulting from pitting and crevice corrosion [28,29], particularly in brass condenser tubes as well as other copper alloy tubes. The presence of hydrogen sulfide in water is known to increase the attack [30].

Types AISI 304, 316, and 439 stainless steels have been widely used as condenser tubing with a good record in fresh water [22]. However, pitting and crevice corrosion cases have been reported in seawater service [31]. The chloride con-tent of the cooling water is the controlling factor of the pitting and crevice corrosion of stainless steels [32]. High alloyed austenitic and ferritic stain-less steels with high chromium (20% and above) and molybdenum (6% and above) content have been successfully used in sea water [33].

The risk of localized corrosion of copper alloys and stainless steels could be greatly reduced if tubes are kept clean because attack is the result of deposit buildup in most cases. This is accomplished by maintaining optimum velocities, regular and automatic cleaning, and drying during extended outages [28].

6-5

6.6 Galvanic Corrosion When materials having different corrosion potentials are exposed to seawater or a medium-conductivity water, a galvanic cell is created that accelerates corrosion of the more active metal (less noble metal). This galvanic corrosion usually occurs in tube sheet and water box materials like Muntz metal (60% Cu/40% Zn) or aluminum bronze because they are less noble than tube materials like copper-nickel alloys, stainless steel, or titanium. 1 Cases of galvanic corrosion in condensers occurred when Muntz ana' tube sheets were fitted with stainless steel tubes [34]. The increased use of titanium tubes, which have a potential more noble than carbon steel, copper alloys, and most stainless steels, has increased the occurrences of severe galvanic corro-sion [35]. Other failures reported [5] include galvanic attack of Admiralty brass (72% Cu/1% Sn/27% Zn) tubes promoted by carbon deposits on the tubes, attack of silicon bronz tubesheets caused by installation of AL6X (20% Cr/

25% Ni/6% Mo/0.02% Max. C) stainless steel. inserts in Admiralty brass condenser tubes, and considerably less severe attack found in similar installations with aluminum bronze tube sheets and AL6X inserts. ,

The recommended method of reducing the problem is to install a cathodic pro-tection system in the water box. Coating of the water box and tabesheets with a nonconducting material (e.g. , epoxy) is useful method of reducing galvaric corrosion [35]. If economically feasible, use of one alloy or alloys with small corrosion potential differences for tubes, tube sheets, and water boxes would minimize galvanic corrosion [36].

6.7 Environmental Cracking Stress-corrosion cracking (SCC) of copper alloys and hydrogen emorittlement

^

cracking (HEC) are both forms of environmental cracking that can affect con-densers. Only two incidences of waterside HEC were identified i.i ferritic stainless steel tubes [4]. It is believed that hydrogen was a product of the cathodic protection current.

The main environmental cracking problem is the SCC of copper alloys. Most of the SCC failures initiated on the steam side of the tubes in the presence of a high concentration of ammonia in the condensate. It was estimated [22] that SCC is responsible for 25% of the failures of Admiralty brass tube sets in the air cooling section of condensers cooled by fresh water.

The controlling parameters of SCC in condenser tubes are [37]:

(1) Susceptible metal: Some copper alloys, like Admiralty brass and arsenical aluminum brass, are highly susceptible to SCC; while other copper alloys, like copper-nickel alloys (70/30 and 90/10) and copper-iron alloy 134, are resistant to SCC [3/]. ,

(2) Corrosion environment: Water ammonia and oxygen are important factors in promoting SCC. Ammonia is derived from the all volatile treatment for boiler feedwater chemistry control. Whereas oxygen originates fron air k that leaks into the system through an imperfectly maintained air-tight system [4].

6-6

(3) Tensile stress: Sources of tensile stresses in the condenser tubes in-clude operating stresses, thermal stresses, fit-up stresses, and residual stresses.

Waterside SCC of brasses has occurred less frequently than steamside attack

[4]. Ammonia and its derivatives (nitrates, nitrites) are suspected to promote SCC. This type of SCC frequently occurs beneath surface deposits, probably because deleterious species can concentrate beneath deposits.

Both types of SCC (steam and water-side) in condenser tubes could be reduced by reducing residual tensile stresses. Sources of residual stresses after the final stress relief anneal should be identified and avoided, i.e., no bending, straightening, over-rolling, or mechanical abuse. Ammonia concentration should be maintained as low as possible to reduce SCC. In this respect design consid-erations and water showering systems have been proposed [37] to provide dilution of ammonia in the air cooling section. Air leakage also should be minimized to reduce oxygen and carbon dioxide in the steam side of the condenser tubes.

6.8 Condensate Corrosion Under the same environmental conditions that promote SCC, copper alloy condenser tubes are susceptible to condensate corrosion [38]. Sometimes, this type of steamside attack is referred to as " ammonia attack." Condensate corrosion occurs in places where ammonia could be concentrated in the presence of oxygen. Thus control of condensate corrosion could be achieved by reducing ammonia and oxygen concentrations in the condensate.

6.9 Summary and Recommendations Titanium is a less susceptible material to degradation in power plant condensers.

The only restriction on using titanium is the high capital cost involved. Highly alloyed stainless steels, either ferritic [ Sea-Cure (27% Cr/2% Ni/3.5% Mo/0.4 Ti/0.02 C) and 29-4C (28% Cr/3.6% Mo/0/4 Ti/0.02 C)] or austenitic (AL6X), also have been used successfully. Copper alloys are the least corrosion resistant alloys to be used as condenser tubes; however, some copper alloy condensers have performed admirably for 20 years or more [4].

There are several preventive measures to control corrosion in operating con-densers. A summary of these methods are given by Gehring [38] as follows:

(1) prevention of biofouling (2) periodic cleaning (3) cathodic protection (4) protective coating (5) inhibitor treatment (6) proper layup procedures 6.10 References

1. R. Whitaker, " Stamina for Stop and Go," EPRI Journal, 5 Number 8, Electric Power Research Institute (1980).

6-7

2. R. L. Coit, " Overview Failure in Steam Surface Condenser," Proceedings of the Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

3. K. A. Lehner, "The Impact of Condenser Failures on Other Power Plant Components," Proceedings of the Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

4. B. C. Syrett and R. L. Coit, " Materials Degradation in Condensers and Feedwater Heaters," Proceedings of the International Symposium of Nuclear Materials Degradation, South Carolina, August 1983, National Association of Corrosion Engineers (NACE) (1984).

5. J. A. Beavers, A. K. Agrawal, and W. E. Berry, " Corrosion-Related Failures in Power Plant Condensers," EPRI-1468, Electric Power Research Instit'ute (1980).

6. D. D. MacDonald, " Erosion-Corrosion in Steam Condensers," Proceedings Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

7. B. C. Syrett, "Pretection of Cu Alloys from Lorrosion in Sulfide Polluted Sea Water," Materials Performance 20, 50 (1981).

8. F. L. LaQue, " Marine Corrosion," Wiley Interscience, pp. 264-274, New York, 1985.

9. S. C. Dexter, " Fouling and Corrosion," Proceedings Condenser Biofouling Control-State of the Art Symposium, EPRI-CS-4339 Electric Power Research Institute (1985).

10. Z. I. Tanabe, "On the Erosion of Condenser Tube Alloys by Sand in Water,"

Sumitomo Light Metal Technical Reports, 9, 160-175, No. 3, Sumitomo Heavy Industry, Tokyo (1968).

11. S. Sato and K. Nagata, " Factors Affecting Corrosion and Fouling of Con-denser Tubes of Copper Alloys and Titanium," Sumitomo Light Metal Technical Reports, ,1_9, 83-94, No. 3-4, Sumitomo Heavy Industry, Tokyo (1978).

12. W. C. Stewart, " Corrosion Resisting Characteristics of Iron Modified 90:10 Cupro Nickel Alloy," Corrosion, 8, 259-277, National Association of Cor-rosion Engineers, Houston, Texas (1952).

13. D. B. Anderson, F. A. Badia, " Chromium Modified Copper-Nickel Alloys for Improved Sea Water Impingement Resistance," Journal of Engineering for Power, Series A, 9, 160-175, No. 2, American Society of Mechanical Engi-neers (1968).

14. S. Sato, T. Nosetani, " Allowable Water Velocity and Cleanliness Factor of Aluminum Brass Condenser Tube with Ferrous Ion Additions into Seawater,"

Sumitomo Light Metal Technical Reports, 11, 271-270, No.4, Sumitomo Heavy Industry, Tokyo (1970).

6-8

15. B. C. Syrett, "Recent Developments in Controlling Corrosion in Electric Power Plants," Proceedings of the Third Asian-Pacific Materials and Corro-sion Association, Taipei, Taiwan, p.93 (1983).

16. S. Sato, T. Nosetani, Y. Yamaguchi, and K. Onda, " Factors Affecting the Sand Erosion of Aluminum Condenser Tubes," Sumitomo Light Metal Technical Reports, 1_6, 23-37, No. 1-2, Sumitomo Heavy Industry, Tokyo (1975).

17. P. T. Gilbert, "The Resistance to Failure of Condenser and Heat Exchanger Tubes in Marine Service," Transaction of Institute of Marine Engineering, 6_6, Institute of Marine Engineering, Miami (1954).

18. B. C. Syrett, " Corrosion Due to Water Pollution," Proceedings of the Seminar on Prevention of Condenser Failures, EPRI CS-4329-SR, p. 11.1, Electric Power Research Institute (1985).

19. B. C. Syrett, D. D. MacDonald, and S. S. Wing, " Corrosion of Copper-Nickel Alloys in Seawater Polluted with Sulfide and Sulfide Oxidation Products,"

Corrosion, 35, 409 (1979).

20. B. C. Syrett and S. S. Wing, "Effect of Flow on Corrosion of Copper-Nickel Alloys in Aerated Seawater and in Sulfide-Polluted Seawater," Corrosion, 3_6, 73 (1980).

21. L. Kenworthy, "Some Corrosion Problems in Naval Marine Engineering,"

Transaction of Institute of Marine Engineering, 77, 149-173, Institute of Marine Engineering, Miami, Florida (1965).

22. W. B. Lawrence, W. D. Ellis, F. J. Hekking, M. P. Lagache, A. M. Madagiri, and A. C. Madsen, " Steam Plant Surface Condenser Leakage Study," EPRI NP-481, Electric Power Research Institute (March 1977).

23. J. A. Beavers, A. K. Agrawal, and W. E. Berry, " Corrosion-Related Failures in Power Plant Condensers," EPRI NP-1468 (TPS79-730), Electric Power Research Institute (August 1980).

24. M. G. Fontana and N. D. Greene, " Corrosion Engineering," McGraw-Hill, 1978.

25. G. T. Colegate, "Dezincification," Matal Industry, ,7_3, 482 (1948).

26. B. Upton, " Corrosion Resistance in Sea Water of Medium Strength Aluminum Bronzes," Proceedings of the Second International Conference on Metallic Corrosion, pp. 806-811, National Association of Corrosion Engineers, South Carolina (March 11-15,1953).

27. J. A. Beavers, A. K. Agrawal, and J. H. Payer, "Dealloying and Stress Corrosion Cracking," Proceedings of the Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

28. G. E. Moller, " Pitting and Crevice Corrosion," Proceedings of the Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

6-9

29. T. P. May and B. A. Weldon, " Copper-Nickel Alloys for Service in Seawater,"

24th International Congress on Fouling and Marine Corrosion, Cannes, France (June 1964).

30. B. C. Syrett, "The Mechanism of Accelerated Corrosion of Copper-Nickel Alloys in Sulfide Polluted Seawater," Corrosion Science, 2 , 187, British Corrosion Society, London (1981).

31. S. W. Shor, E. E. Hanson, J. Rios, P. B. Lindsay, and N. F. Allard, " Steam Plant Surface Condenser Leakage Study Update," EPRI NP-2062, Electric Power Research Institute (March 1982).

32. R. C. Newman, H. S. Isaacs, and B. Alman, Corrosion, 38, 261 National Association of Corrosion Engineers, Houston, Texas (1952).

33. A. P. Band et al., " Stainless Steels for Sea Water Service," appeared in

" Stainless Steel '77," Climax Molybdenum Co. (1977).

34. G. A. Gehring and J. R. Maurer, " Galvanic Corrosion of Selected Tubesheet/

Tube Couples Under Simulated Seawater Condenser Conditions," NACE Annual l Conference, Toronto, Canada, paper 202, National Association of Corrosion Engineers (April 1981).

35. J A. Hanck, G. Nekoksa, and R. M. Chhatre, " Galvanic Corrosion in Con-densers," Proceedings Seminar on Prevention of Condenser Failures, EPRI-CS-4329-SR, Electric Power Research Institute (1985).

36. " Guidelines for Selection of Marine Materials," Report No. A-404, The International Nickel Co., New York (1971).

37. J. H. Payer, A. K. Agrawal, and J. A. Beavers, " Stress Corrosion Cracking of Copper Alloys in Condensers," Proceedings Seminar on Prevention of Con-denser Failures, EPRI CS-4329-SR, Electric Power Research Institute (1985).

38. G. A. Gehring, " Summary of Methods for Controlling Waterside Corrosion,"

l Proceedings Seminar on Prevention of Condenser Failures, EPRI CS-4329-SR, Electric Power Research Institute (1985).

l 6-10

NRCfORM338 U B, NUCLE AA E EIULATORY COMMISSION 1 P ErodY NuweEa rAmpaeaav rsOC. esa vos NJ, eranyJ f2 84)

• "o'i 2 3% BIBLIOGRAPHIC DATA SHEET NUREG-1196 SEE INSTRUCTIONS ON THE REVER$E 2 TITLE AND SUS tlTLE J LEAVE BLANE An Overview of En 'ronmental Materials Degradation /

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/ 4 DA TE REPORT COMPLETED

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[ 6 D AT E REPOR T 155UED H. I. Shaaban, P. Wu f O~T- vtAa fAugust l 1986 7 t'E AFORW4NG ORGAmilATION NAME AND M AeLIN ODR E55 fince w ar te Codes S PROJECTsT A$EfvWORE UNIT NUMSER Division of Inspection Pro ams Office of Inspection and En rcement . .i~ Oa Ga ANT NU .ER U. S. Nuclear Regulatory Co 'ssion Washington, D. C. 20555 10 5,0NSORING ORGAN 12ATION NAME ANO WAILING ADDRES$ ts ude le Coder lia TVPESSREPORT Technical Review Same as 7, above, b PERIOO COVERED fisses es,we s aprest 12 $UMLEMENT ARY NOTES

13. ASTR ACT (100 overes or 'essJ This report provides a brief overview o an ses and conclusions reported in published literature regarding environ ntal' induced degradation of materials in operating light-water reactors. I is inte ded to provide a synopsis of subjects of concern rather than to a ress a li nsing basis for any newly discovered problems related to react materials The subjects discussed include materials degradation in re tor internal reactor pressure boundary componcnts, steam generators, stearr turbines, and ndensers. In each of these systems, the degradation mechanism and the suggeste reasons for each mechanism are reviewed. Possible remedies methods for avoid g such degradation are also given.

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