The ageing of materials in the light water reactor fleet around the world is a major factor in ensuring not only the safe and economical operation of these power plants, but also preserving and extending their substantive contribution to carbon-free electricity production. ‘Ageing’ refers to the change in character or properties of components or systems with time. Ageing can occur in benign environments, such as those experienced by concrete under ambient conditions. It can also occur in harsh or extreme environments such as components in the core of a reactor where they are exposed to high temperature, high stresses, an aggressive chemical envi­ronment and a high level of radiation. These additional stressors can ini­tiate, accelerate and generally shape the ageing process over that in inert environments.

Classes of components for which an understanding of ageing is impor­tant are generally divided into four categories: core components including the reactor pressure vessel, other plant components, electric cables, concrete and piping. This book focuses on core components and electric cables and the processes by which ageing occurs. Couvant and Murty address two of the key ageing modes for core components — corrosion and creep deforma­tion — in chapters that illuminate the major processes and their consequences. Adamson and Rudling focus on the ageing of zirconium alloy fuel bundles that are critical components in the containment of the fuel and extraction of energy for electricity production. Next, Hashemian reviews the ageing of electric cables, hundreds of miles of which are installed in each nuclear plant.

But the last part of the book on assessment strategies for managing the ageing of materials in reactors is what makes it unique. While ageing of materials is unavoidable, and in fact ubiquitous, the key is to understand the ageing process and how the various components of the environment can affect or accelerate that process. Only by understanding the ageing process can mitigation or amelioration be considered. Jeong and Hwang, Katona, and Ray and Lahoda present perspectives on the evaluation of plant ageing including development of ageing management programs, proactive materi­als management, mitigation and repair methods, international cooperative activities and finally, integration of these programs into a system that ensures the safe, long-term operation of the power plant.

Gary S. Was

Walter J. Weber, Jr. Professor of Sustainable Energy, Environmental and Earth Systems Engineering University of Michigan

The concept of deformation mechanism maps was proposed by Ashby.7 6 Since different creep mechanisms operate or dominate in different stress, temperature and grain size regimes, Ashby envisioned that a deformation mechanism map would be an ideal representation of the materials consti­tutive behavior. Over the years, this concept has been extended to describe a variety of other physical phenomena such as sintering,77 wear78 and frac­ture.79 Figure 3.17 is a deformation mechanism map first reported by Ashby in 1972. The map was plotted as normalized stress (a/G) against homologous temperature (T/Tm) for a constant grain size. The map was then constructed by determining the stress or temperature boundaries where one mechanism would dominate others. To this end, the creep constitutive relations of dif­ferent mechanisms were compared and stress and temperature values where

2

to

.2!

ел

c

Л)

transitions from one mechanism to another would occur were determined. For example in Fig. 3.17, at low temperatures and low stresses, the material would resist plastic deformation and the material would behave elastically while in the later modifications this low stress regime was considered to be due to Coble creep.

However, as we continue to increase the temperature and approach higher homologous temperatures, diffusional processes become dominant. Also the applied stresses are sufficient to overcome the flow stress corresponding to that temperature and the material deforms plastically. Since diffusional creep can either be governed by Coble or N-H creep we find the map out­lining the regions where these mechanisms are dominant. As Coble creep is controlled by grain boundary diffusion, it is dominant at lower temperatures and the Coble creep field lies to the left of N-H creep on the map. Also, if we increase the stress at a given temperature, dislocation-based mechanisms come into play. Depending upon the homologous temperature, the defor­mation can be controlled by dislocation climb or glide. At low homologous temperatures dislocation climb is suppressed and hence dislocation glide becomes the dominant deformation mechanism; this is not to be confused
with the viscous glide creep discussed earlier which occurs along with climb creep in class-A alloys. For the sake of the reader, we present a small exam­ple of how the temperature and stress boundaries of different mechanisms can be determined in a given material. If we assume Coble creep and N-H creep as competing mechanisms for a given grain size, then Coble creep will be dominant when

^Coble > £N-H. [3.45]

From the relevant equations for Coble and N-H creep mechanisms, this would imply

Bc DB8B <°l dl oQ

> Bu.

П d3kT d2 kT

Cancelling the common terms we obtain

Db > Kl

Dl d

where K5 is a constant. At a constant grain size, and after expanding Db and Dl, the above equation will turn out to be

Dob (~Qb/rt)> ^

> Kc

Dol (-QlRT)

where K5 is a constant. Clearly the transition from Coble to N-H creep is temperature dependent and independent of stress. The transition is only dependent on the activation energies for grain boundary and lattice diffu- sivities. The temperature dependence of this cross-over is captured by the map where we can observe that a line parallel to the stress axis separates the Coble creep and N-H creep fields.

An alternate way of representing the deformation mechanism maps was proposed by Mohamed and Langdon.80 Since grain size is an important factor which governs the deformation behavior of materials, the mecha­nism map can also be plotted for normalized grain size (d/b) against nor­malized stress (a/G) for a given temperature (Fig. 3.18). As the plot shows, smaller grain sizes are favorable for Coble creep and as we increase the grain size N-H and H-D creep mechanisms become dominant. Since dis­location creep is independent of grain size, transitions from dislocation creep to other mechanisms are represented by lines parallel to the grain size axis. The climb-glide mechanisms are noted for larger grain sizes with

108

106

$ t5

105 104 103

climb occurring at lower stresses; this plot did not consider the climb region at the higher stress-end as described earlier. Also missing is the GBS that is expected between viscous and dislocation creep mechanisms.

One of the most unique aspects of material behaviour in a nuclear power plant is the effect of radiation (mainly neutrons) on the dimensional stabil­ity of the reactor components. In fast breeder reactors the Fe and Ni-based alloys creep and swell, that is, they change dimensions in response to a stress and change their volume in response to radiation damage. In LWRs, zirco­nium alloy structural components creep, do not swell, but do change their dimensions through the approximately constant volume process called irradiation growth. Radiation effects are not unexpected since during the lifetime of a typical component every atom is displaced from its normal lattice position at least 20 times (20 dpa). With the possible exception of elastic properties like Young’s Modulus, the properties needed for reliable fuel assembly performance are affected by irradiation. A summary of such effects is given by Adamson (2000).

Practical effects of dimensional instabilities are well known and it is rare that a technical conference in the reactor performance field does not include discussions on the topic. Because of the difference in pressure inside and outside the fuel rod, cladding creeps down on the fuel early in life, and then creeps out again later in life as the fuel begins to swell. A major issue is to have creep strength sufficient to resist outward movement of the cladding if fission gas pressure becomes high at high burnups. PWR guide tubes can creep downward or laterally due to forces imposed by fuel assembly hold down forces or cross flow hydraulic forces — both leading to assembly bow which can interfere with smooth control rod motion. BWR channels can creep out or budge in response to differential water pressures across the channel wall, again leading toward control blade interference. Fuel rods, water rods or boxes, guide tubes and tie rods can lengthen, possibly leading to bowing problems. (For reference, a recrystallized (RX or RXA) Zircaloy water rod or guide tube could lengthen due to irradiation growth more than 2 cm during service; a CWSR component could lengthen more than 6 cm.) Even RX spacer/grids could widen enough due to irradiation growth (if texture or heat treatment was not optimized) to cause uncomfortable interference with the channel. In addition, corrosion leading to hydrogen absorption in Zircaloy can contribute to component dimensional instability due, at least in part, to the fact that the volume of zirconium hydride is about 16% larger than zirconium.

The above discussion leads to the concept that understanding the mecha­nisms of dimensional instability in the aggressive environment of the nuclear core is important for more than just academic reasons. Reliability of materi­als and structure performance can depend on such understanding.

Comprehensive reviews of dimensional stability have been given in the ZIRAT Special Topical Reports (Adamson & Rudling, 2002; Adamson et al, 2009; Cox et al, 2005).The sources of dimensional changes of reactor com­ponents (in addition to changes caused by conventional thermal expansion and contraction) are: irradiation growth, irradiation creep, thermal creep, stress relaxation (which is a combination of thermal and irradiation creep), and hydrogen and hydride formation.

Irradiation effects are primarily related to the flow of irradiation-produced point defects to sinks such as grain boundaries, deformation-produced dis­locations, irradiation-produced dislocation loops, and alloying and impurity element complexes. In zirconium alloys, crystallographic and diffusional anisotropy are key elements in producing dimensional changes.

In the past, hydrogen effects have been considered to be additive to and independent of irradiation. Although this independency has yet to be defin­itively proven, it is certain that corrosion-produced hydrogen does cause significant dimensional changes simply due to the 16-17% difference in density between zirconium hydride and zirconium. A length change in the order of 0.20% can be induced by 1000 ppm hydrogen in an unirradiated material (Fig. 4.60) (King et al., 2002; Seibold et al., 2000).That the presence of hydrides contributes to the mechanisms of irradiation creep and growth is highly suspected but yet to be determined in detail.

Fuel rod diametral changes are caused by stress dependent creep pro­cesses. Fuel rod length changes are caused by several phenomena: [3]

4.60 Dimensional changes in unirradiated ZIRLO and Zircaloy-4 tubing and strip for different sample orientations as a function of hydrogen content. (Source: Reprinted, with permission, from King et al. (2002), copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, Pa 19428.)

heavily cold worked material, it has been reported that some shrinkage may occur. In a non-textured material such as SS, creep down of the cladding would only result in an increase in cladding thickness, with no change in length.

• Creep due to PCMI after hard contact between the cladding and fuel. This occurs in mid-life, depending on the cladding creep properties and the stability of the fuel.

• Hydriding of the cladding due to corrosion.

Bow of a component such as a BWR channel or PWR control rod assembly can occur if one side of the component changes length more than the other side. Such differential length changes occur due to differential stress and creep, to relaxation of differential residual stresses or to differential growth due to differences in flux-induced fluence, texture, material cold work and hydrogen content (and, although not usually present, differences in temper­ature or alloying content). This is described more in the ZIRAT10 Special Topics Report on Structural Behaviour of Fuel and Fuel Components (Cox et al, 2005).

The next section discusses the effect of irradiation on dimensional stability.

220 Materials’ ageing and degradation in light water reactors

The basic cable failure modes resulting from exposure to stressors include: short circuits between cable conductors, short circuits between one or more cable conductors or the shield and ground (ground fault), open circuits in the cable conductors, and breakdown of the cable insulation (AMS Corp., 2010). The most common failure mode is ground fault, in which the cable faults to ground from one or multiple conductors (U. S. NRC, 2001). Ninety-five per cent of cable problems occur at the cable connector where age factors are combined with mechanical damage and wear (AMS Corp., 2010).

For power cables, such failure modes can cause circuit protection devices to trip or partially discharge, resulting in excessive heating and degrada­tion of the cable insulation and ionization of the air around the discharge. This failure mode — degraded insulation resistance — can lead to a conduc­tor short-circuit to ground failure, conductor to conductor short circuit, or potentially both. Power cable failures have resulted in reactor trips, weak­ened engineered safety features, loss of redundancy, and reduced power operation (U. S. NRC, 2001).

For I&C cables, conductor short-circuit to ground failure and conductor — to-conductor-short-circuit failure interrupt the transmission of control sig­nals through the cable.

Degraded insulation resistance failure can impair the functioning of I&C cable and/or increase the rate of error (U. S. NRC, 2010a). The transmit­ted signal may become erratic, causing errors in measurement, spikes, noise, and other problems. When cables become bare, shunting and short circuits can occur, and if the cable insulator is degraded, the insulation material can become brittle and flammable (Hashemian, 2010). I&C cables are the most susceptible to ageing degradation (U. S. NRC, 2001).

In addition to signal anomalies and problems with plant control and safety systems, cable ageing has resulted in loss of critical functions and fire (AMS Corp., 2011). In light water reactors, the most severe cable failure scenario is loss of normal function during a LOCA when hot steam under pressure can cause cables to malfunction if insulation ageing, cracks, or other damage allow moisture to enter the cable. Hot steam combined with high pressure is the primary cause for cable malfunction in a LOCA, because steam pen­etrates smaller cracks more easily than water. Such consequences explain why the Hungarian Paks Nuclear Power Plant has described cable ageing as ‘the most significant I&C ageing issue’ in its plant (Hashemian, 2010).

A 2007 U. S. NRC report found that 93% of reported cable failures occurred in normally energized power cables: ‘More than 46% of the fail­ures were reported to have occurred while the cable was in service and more than 42% were identified as ‘testing failures’ in which cables failed to meet testing or inspection acceptance criteria’ (U. S. NRC, 2010a). The majority of these cable failures occurred between 11 and 30 years of service — less than the typical 40-year licensing period of a plant (U. S. NRC, 2010a).

While many cases of cable failure are identified through routine cable testing, some occur before a failure is identified (e. g. on cables that are not normally tested or powered). This fact underscores the importance of imple­menting a cable condition monitoring program (U. S. NRC, 2010b).

The previous section described the different regions of a creep curve and their corresponding characteristics. Mechanistically, the creep curve is a result of the changes occurring in a material at a microstructural level.

The creep curve is basically the outcome of the competition between the processes of strain hardening and recovery. Materials usually strain harden during plastic deformation due to dislocation multiplication. The strain hardening is a kind of ‘defense’ mechanism in response to an applied stress. Further plastic deformation can occur only if the applied stress exceeds the increase in flow stress of the material due to strain hardening. Alternatively, deformation can proceed at the initial applied stress if the material soft­ens. The mechanism of recovery acts to soften a deformed specimen thus allowing further plastic deformation. In the primary stage, the rate of strain hardening is greater than the rate of recovery. This is due to the formation of a more resistant creep substructure. The substructure could be the forma­tion of dislocation networks or the arrangements leading to the formation of subgrains. In the secondary stage of creep, the rate of strain hardening is balanced by the rate of recovery due to dislocation annihilation and defor­mation occurs at a constant strain rate. In the tertiary stage of creep, the increase in applied stress due to a reduction in specimen cross-sectional area surpasses the increase in flow stress due to strain hardening. The reduc­tion in specimen cross-sectional area can be due to necking or internal void formation. The tertiary creep is often associated with metallurgical changes such as the coarsening of precipitate particles, recrystallization or diffusional changes in phases present, void formation and so forth.

Figure 3.2 depicts four types of creep curves that have been generally observed.4 The shape of the creep curve is dependent on the initial condition

of the material prior to deformation. Curve A is a typical creep curve observed in several materials. The curve consists of a normal primary stage characterized by a decreasing strain rate, a secondary stage where defor­mation proceeds at a constant rate and a tertiary stage where the material deforms at increasing strain rates with time/strain leading to eventual fail­ure. Such creep curves are usually exhibited by annealed metals and cer­tain alloys (known as class-M or class-II type). In comparison to curve A, curve B depicts a very small primary creep stage. In fact, it appears as if the material enters the steady state immediately. Such a type of curve is obtained when the substructure pertaining to creep remains constant such as in some alloys (known as class-A or class-I type). Curves of type C are obtained from materials that have been previously crept at a higher stress. The increasing creep rate over the primary creep stage is due to the recov­ery of the substructure corresponding to the previous steady-state condi­tion. The sigmoidal type of creep curve (curve D) suggests the nucleation and spread of slip zones until a steady state is achieved. Such creep behavior has been exhibited by certain dispersed phase alloys.

In certain cases, it is possible that the total creep curve is in the primary stage, and the secondary and tertiary creep stages are not attained at all. Such a curve has been seen for materials tested at low temperatures (T < 0.3 TM) where effects due to diffusion are suppressed (no annealing or recov­ery) and the entire deformation is due to work hardening (dislocation con­trolled). The primary creep strain rate tends towards a value of zero at long periods. The strain hardening due to long range dislocation interactions pre­cludes a constant rate of creep deformation. The absence of recovery pro­cesses due to the low test temperatures allows the strain hardening process to be the creep-controlling mechanism. This eventually leads to an increase in strength of the material to a value greater than the applied stress value. Further deformation can only occur under the application of a higher stress or in the presence of a higher temperature. Such a behavior is described as an exhaustion creep behavior and the creep curve can be described by a logarithmic creep equation.4

An intermediate cooling rate from в phase, in a Zr-2.5wt.%Nb alloy, has resulted in a decrease in creep rate by 100 times over cold worked mate­rial and rendered higher anisotropy at 450°C. This increase in creep rate is attributed to segregation of Nb in grains that are favorably oriented for easy slip.133

The recent work on the thermal creep of Zr-2.5%Nb alloy by Kishore et at.134 indicates that a microstructure containing a stable phase creeps faster than one with a meta-stable phase and a phase redistribution is established (Fig. 3.31). During creep deformation the stable в phase (with 80 wt.%Nb) dissolves and re-precipitates as в phase (with ~35wt.%Nb), this resulting strain due to phase change adds to the creep strain. Similar phase transfor­mation is reported by Griffiths wherein the Zr-2.5wt.%Nb alloy after 2-14 years of in-reactor service shows that the в phase has a distribution of com­position, the Nb concentration varying from 37% to 75%.135 However, the effect of this phase change on the creep deformation is not well studied.

Stresses which induce both PCI (usually denotes combined mechanical and chemical pellet-cladding interaction) and PCMI (usually denotes pel­let cladding mechanical interaction) are caused by expansion of the fuel pellet against the cladding during power increases (Adamson et al, 2006/7; Strasser et al., 2010a). PCI failures are driven by a stress corrosion crack­ing (SCC) assisted component resulting from fission product release from the fuel, while PCMI failures are generally due to purely mechanical crack­ing, often enhanced by a reduction in cladding ductility due to formation of local hydrides at the clad outer surface. At the micro level, the PCI crack always starts at the cladding inner surface and propagates towards the outer cladding surface while the PCMI crack propagates from the outer to inner surface.

PCI is associated with local power ramps during reactor start-up or power manoeuvring (e. g. rod adjustments/swaps, load following) as shown sche­matically in Fig. 5.3, and is caused by the combination of cladding stress due to the power increase and the influence of iodine, caesium and cadmium released during the power increase in a susceptible material (Adamson et al., 2006/2007; Strasser etal., 2010a). This combination of stress, embrittling

fission products and susceptible material may result in SCC of the fuel clad­ding, as shown in Fig. 5.4.

PCI failures may occur in PWRs/VVERs and BWRs (Strasser et al, 2010a). The failure mechanism is much more prevalent in BWRs, since reactor power is controlled in part by control rod movements that subject the fuel to rapid power level changes. (The reactor power in both BWRs and PWRs is also reg­ulated by flow control.) In PWRs and VVERs, reactor power is not normally controlled by insertion and extraction of the control rods in the core; rather, reactor power is controlled by the boron concentration that is continuously decreased during operation to compensate for the decrease in reactivity. This type of reactor power control is much smoother than in the BWR case and, consequently, PCI failures are less common in PWRs. However, during reac­tor power increases, and specifically during a class II transient (anticipated operational occurrences, AOO), PCI failures may occur in a PWR.

To prevent PCI failures, it is necessary to remove at least one of the fun­damental conditions (tensile stress, sensitive material, aggressive environ­ment) which cause SCC. There are two principal types of remedy (Strasser et al., 2010a):

1. One is to develop reactor operation restrictions that will ensure cladding stresses are always below the PCI threshold stress during power increases. This is the main measure in avoiding PCI defects and the only measure used in PWRs. Operating rules (also called management recommenda­tions, or pellet-cladding interaction operating management restrictions (PCIOMRs)) to limit local power increases and ‘condition’ fuel for power ramping were implemented in BWRs during the late 1970s to mit­igate the PCI issue. The rules are usually a function of exposure and were developed by the different fuel vendors, so they differ between various fuel types. To establish and validate these rules, extensive power ramp tests were performed by the fuel vendors in experimental reactors.

2. The second remedy — design improvement — consists of two approaches: 2a. Cladding design

2a1. Development of radial cladding texture and small grain size that may increase cladding PCI resistance.

2a2. Development of the barrier/liner concept, initially with a ‘pure’ zirconium (Zr) metal barrier at the cladding inner diameter (ID). The barrier is soft and serves to reduce the local stress, hence giv­ing the cladding resistance to SCC. Later, fuel vendors realized that the Zr could be alloyed with Fe to improve the secondary degradation resistance in case of rod failure. The Fe in the Zr will dramatically improve the corrosion resistance of the liner/barrier but may reduce the PCI performance. Although this remedy has so far only been used in BWRs, it should be equally applicable to PWRs.

2b. Pellet design

2b1. Reducing the cladding local strains (and stresses) by shortening the pellet, chamfering the corners and eliminating the dishing.

2b2. Pellets with additives are being developed both for BWRs and PWRs that will increase the margins towards PCI failures (Adamson et al., 2006/7; Patterson, 2010). The additives of inter­est fall into two general categories, the first category involves materials that are essentially insoluble in the fluorite lattice and exists as a separate, grain boundary phase, for example, mix­tures of alumina and silica (aluminosilicates or Al-Si-O). The second category involves materials that are soluble in the cat­ion sub-lattice, such as chromia, or involve a mixture of soluble and insoluble materials, such as chromia and alumina. Although many other additives fall into both categories, attention is directed to the aluminosilicate additives and chromia-base dop­ants as they appear to be the closest to large-scale application.

2b3. Aluminosilicate additives consist of a mixture of SiO2 and Al2O3 and is offered by GNF. During the pellet sintering process, the additive forms a glassy phase that collects on the grain bound­aries. It appears that the Al-Si-O additive at the pellet grain boundaries will chemically react with I, Cs and Cd, thus prevent­ing these SCC-promoting elements from accessing the fuel clad inner surface (Matsunaga et al, 2009, 2010). Additive fuel has been irradiated in commercial and test reactors in the US and in Europe. Ramp tests under BWR conditions in the R-2 and Halden reactors of segmented additives rods from commercial reactors show excellent resistance to the PCI failure mechanism (Davies et al.,1999).

2b4. Chromia (Cr2O3) is the dopant of greatest commercial signif­icance in this class of additives. Two types of chromia-based additives are being offered. The first consists of Cr2O3 in UO2 as offered by AREVA (Delafoy et al, 2003). The sec­ond consists of Cr2O3 and Al2O3 in UO2 as offered by Westinghouse (Arborelius et al., 2005). Alumina is reported to be used in the second form to minimize the effects of chromium on the fission cross-section of doped pellets while enhancing grain growth. In both cases, chromia is expected to reside largely within grains as interstitial Cr3+ and as insoluble Cr2 O3 depending on the concentration and tem­perature. The alumina in the mixed Cr-Al-O dopant should exist as a grain-boundary phase as in the Al-Si-O additive. The cation dopants were developed to increase grain size to reduce fission gas release (FGR) at extended burnup(see for example Delafoy et al., 2007). In addition to improved FGR, chromia-based dopants are reported to improve PCI resistance. Information available in this area is less exten­sive for the chromia-based dopants than for the aluminosil­icate additives. However, ramp tests indicate that the resis­tance to PCI failures of fuel with chromia-based dopants are improved relative to standard fuel in cladding without PCI-resistant liners (Delafoy et al., 2007).

In order to establish measures for managing materials degradation, the deg­radation mechanisms must first be fully understood. Inspection techniques, mitigation methods and repair technologies depend on knowledge grounded in experimental studies of degradation mechanisms or in field operating experience within power plants (IAEA, 2011). There are three stages to managing materials ageing in nuclear power: preventive action; monitoring and inspection; and repair and replacement. In preventive action, improve­ment of the materials, reduction of stress and improvement of water chem­istry can be used as measures to prevent cracks of Ni alloys (IAEA, 2011). Surveillance of pressure vessels can be carried out through monitoring and inspection to check the soundness of parts. For example, by checking for leakage of primary coolant through wall cracks in J-welds of the upper ves­sel head penetration (VHP) or lower bottom mounted instrumentation (BMI) nozzles, pressure boundary performance can be maintained. In the case of coolant leakage, boric acid residues on the outside of the pressure vessel or carbon steel corrosion products can be detected through visual inspection. Cracking can occur in operating power plants due to material properties or residual stress, therefore the timing of cracking can differ from the experimental result. Regardless of the cause, it is important to detect cracks as early as possible. Besides visual inspection, methods such as pene­trant testing or eddy-current testing (ECT), and ultrasonic testing (straight beam and longitudinal wave angle beam UT) can also be used. In repair and replacement, the damaged parts should be isolated from the corrosive envi­ronment, or the tensile stress upon them reduced. In the case that these two measures are inappropriate, the parts should be replaced with others made from more corrosion-resistant materials. In order to systematically manage PWR structural materials, a common objective has been established, and much research has been done as a collaborative effort between many coun­tries. The joint research programme for 2011 is as shown in Table 7.1.

Ashby and Verall46 proposed a model to describe the process of GBS accom­modated by diffusional flow. According to this model [3.23]

where

and a0 is the threshold stress. As shown in Fig. 3.7, a natural outcome of diffusional flow during GBS is grain switching. The grains change their neighbors during the process of sliding and such a change is assisted by dif — fusional flow. The threshold stress term present in Equation [3.23] appears

due to an increase in grain boundary area, a resultant of the grain switching process. Though this model was successful in explaining the experimentally observed switching of grains during deformation, it failed to predict the stress dependence of strain rate. Moreover, Spingarn and Nix47 suggested that grain switching cannot be entirely attributed to diffusional flow as the diffusion paths are physically incorrect.

As outlined in Section 4.3.1, irradiation produces damage in the form of small dislocation loops (<a> component loops) which harden the material. The result is an increase in strength and decrease in ductility.

At reactor start-up, the tensile properties are the unirradiated properties reported by the fuel supplier. Mechanical properties begin changing imme­diately upon startup, and by an exposure of 5 MWd/KgU or a fluence of about 1 x 1025 n/m2 (E > 1 MeV) an increase in strength and decrease in ductility reach fluence-saturated values. Figure 4.20 illustrates this point for Zircaloy-4 irradiated and tested at 315°C (588K) (after Morize et al., 1987). Note also that the UTSs of cold worked stress relieved (CWSR) and recrys­tallized (RX) materials become similar at low exposures. This is a general trend which depends on the balance of hardening by pre-existing disloca­tions (cold work) and irradiation-produced defects.

Fuel cladding requires sufficient strength to prevent inward plastic defor­mation of the cladding at beginning-of-service conditions. PWR strength must be higher than for BWRs due to the higher water pressure needed to suppress boiling; therefore, PWR Zircaloy cladding has traditionally been in the cold work stress relieved annealed (SRA) condition. The discussion above points out that the difference in strength between SRA and RXA materials is short-lived under reactor conditions.

Properties of zirconium alloys and their applications in LWRs 177