For S-65C Be grade irradiated up to 1025nm~2 (~0.74dpa) at ~300 °C, the thermal conductivity was found to be similar, within experimental error, to that of the unirradiated material.146 Similarly, no effect was seen for Be S-65C after irradiation at 350 and 700 °C to a damage dose ~0.35dpa.147 Significant changes in the thermal conductivity were observed only for conditions that lead to sig­nificant changes of the beryllium structure, such as the formation of a high density of radiation defects (especially at low irradiation temperature and high dose) or high (more than tens ofpercent) swelling.144

Other physical properties (elastic modulus, coef­ficient ofthermal expansion, etc.) are not influenced by neutron irradiation (at least at the fluence and temperature ranges relevant for the beryllium armor for the ITER PFCs).

4.19.4.4.1 Swelling

It is well known that beryllium swells when irradiated by neutrons, especially during high temperature irradiation. Reviews of the available swelling data for different Be grades can be found elsewhere (see, e. g., ITER MAR,129 Billone,139 and Barabash eta/.141). The computer code ANFIBE (ANalysis of Fusion Irradiated BEryllium), has been developed and applied in the past as an interpretative and predictive tool148 for the prediction of beryllium swelling. The driving force for the swelling is the presence of He, which forms He bubbles, especially during high-temperature irradiation (more than ^400 °C) or after high-temperature annealing. The maximum values of swelling could reach approximately tens of percent at temperatures more than 600 °C and helium content more than several thousand atomic parts per million. Swelling depends on the structure of the beryllium: beryllium grades with small grain size (~8—10 pm) and high BeO content (~3-4wt%) have a higher resistance to swelling than conventional Be grades.141 As concluded in ITER MAR,129 for an irradiation temperature <^400 °C, swelling of beryl­lium containing 1500 appm He is <~1%. At higher temperature, swelling could reach the value of a few percent at the end of life.

4.20.5.3.1 Defect structure in irradiated copper and copper alloys

Copper is among the most extensively studied metals in terms of fundamental radiation damage. Several reviews on the effect of irradiation on the

microstructure of copper and copper alloys are available in the literature.60,96,97 Neutron irradiation of copper at low temperatures produces small defect clusters, dislocation loops, and SFTs. At temperatures above 150—180 °C, the density of defect clusters

starts to decrease with increasing temperature, accompanied by the formation ofvoids. This temper­ature-dependent formation of defect structures is shown in Figure 14.60 Low-temperature neutron irradiation produces a high number density of SFTs and a low number density of dislocation loops in copper. Edwards et a/.64 reported a number density of SFTs, ^2-4 x 1023m~3 and a number density of dislocation loops, 5 x 1021m~3 in OFHC copper neutron irradiated to ~0.01 dpa at 100 °C. Disloca­tion loops are believed to be of interstitial type.

(a) (b)

The size of SFTs is small, ^2-3 nm. As doses increased, the density of SFTs increased to a satura­tion level at ~0.1 dpa, while the size of SFT is inde­pendent of the dose and temperature. In general, the dislocation loop density is low, and a significant dis­location network is not formed in irradiated copper.96

Radiation hardening in copper can be adequately described by Seeger’s dispersed barrier model, and the yield strength increase is due to the formation of defect clusters.98 Singh and Zinkle96 summarized the dose dependence of the TEM-visible defect cluster density in copper irradiated near room temperature with fission neutrons, 14MeV neutrons, spallation neutrons, and 800 MeV protons (Figure 15)96 TEM — visible defect clusters were observed at a very low dose (10-5 dpa). The defect cluster density showed a linear dependence on irradiation dose at low doses. The dose dependence of the defect cluster density shifts to either a linear or a square root relation at intermediate doses (>^0.0002 dpa). The cluster den­sity reaches an apparent saturation (~1 x 1024m~3) at ~0.1 dpa. The dislocation loops range in size from ~1 to 25 nm.9 Differences in the type of irradiation (fission, fusion, spallation, etc.) have no significant effect on the defect cluster accumula­tion behavior in copper. The density of defect clus­ters in irradiated copper shows strong temperature
dependence (Figure 16).100 The defect cluster den­sity is essentially independent of the irradiation temperature between 20 and 180 °C (upper tempera­ture limit is dependent on dose rate). At higher tem­perature, the cluster density decreases rapidly with increasing irradiation temperature. At irradiation temperatures between 182 and 450 °C, the density of defect clusters was reduced by over three orders of magnitude.83,84 The saturation dose of the defect cluster density is similar, ^0.1 dpa, for all irradia­tion temperatures.96 The size distribution of visible defect clusters can be described by an exponential function10: N(d) = N0 exp(—d/d0), where N(d) is the number of defects of diameter d, N0, and d0 are constants, and their values depend on irradia­tion conditions and material purity. As the irradiation temperature decreases, a fraction of small clusters increases relative to large clusters.

Void formation occurs above 180 °C in neutron — irradiated copper.60 The peak void swelling temper­ature in copper is about 320 °C at a dose rate of 2 x 10-7dpas_1. Singh and Zinkle96 summarized the dose dependence of void density measured by TEM in copper irradiated with fission and fusion neutrons at 250-300 °C from several studies. The data showed a large variation (up to two orders of magnitude differences) of void density between

experiments. One possible source could be residual gas atoms in copper that can have a dramatic effect on void swelling in copper. Zinkle and Lee86 discussed in detail the effect of oxygen and helium on the forma­tion of voids in copper. The stacking fault tetrahedron is predicted to be the most stable configuration of vacancy clusters in copper. A small amount of oxygen (~10 appm) or helium (~ 1 appm) in copper is needed to stabilize voids. High-purity copper with low oxy­gen concentration (<5wppm) showed no significant

Irradiation temperature

void formation after 14MeV Cu ion irradiation to 40 dpa at temperatures of 100-500 °C.100

The defect microstructure (SFTs and dislocation loops) in irradiated copper alloys is essentially the same as in irradiated pure copper.22’25’64 Neutron irradiation can affect precipitate microstructure in copper alloys. When irradiated at 100 °C, the precip­itate density in CuCrZr was slightly reduced, and the mean size of the precipitates increased.13’64 Zinkle et a/.25’26 reported that when GlidCop Al25 and MAGT 0.2 were ion irradiated to 30 dpa at 180 °C, a high number density (5 x 1023 m~3) of defect clusters (primarily SFTs) with a mean size of2 nm was produced. The geometry ofoxide particles in GlidCop Al25 was transformed from triangular platelets to nearly circular platelets’ and the particle size was reduced from 10 to 6 nm after irradiation (Figure 17).25’26 The geometry and size of oxide par­ticles in MAGT 0.2 were essentially unchanged by irradiation. In general, DS copper alloys showed supe­rior particle stability under irradiation.

Limited data are available in terms of the effect of solution additions on the irradiated microstructure of copper. A study by Zinkle25 showed that solute additions (e. g., Al, Mn, Ni) to 5 at% in copper do not have significant effect on the total density of small defect clusters at low irradiation temperatures (<130 °C). However, solute additions reduce the formation of SFTs and enhance the formation of dislocation loops. The loop density and mean size in Cu-5% Mn irradiated to 1.6 dpa at 160 °C were 3 x 10[5] [6] m~3 and 23 nm, and 1.8 x 1022m~3 and 18 nm in Cu-5% Ni irradiated to 0.7 dpa at 90 °C

(Figure 18).25’26 These loop densities are more than an order of magnitude larger than the highest loop density observed in pure copper. The effect of the stacking fault energy on void formation in copper alloys was also investigated. Generally speaking, the lower the stacking fault energy, the less favorable for the formation of 3D voids. For example, swelling occurred in Cu—1—2.5% Ge alloys irradiated at 250°C, while no measurable swelling occurred in Cu-3-5% Ge that has lower stacking fault

97

energies.

4.01.3.1.1 Irradiation growth: Macroscopic behavior

One of the most specific macroscopic effects of irra­diation on materials is the dimensional change with­out applied stress. In the case of zirconium alloys, it is known that under neutron irradiation, a zirconium single crystal undergoes an elongation along the (a)

axis and a shortening along the (c) axis without significant volume evolution. Thorough reviews of this phenomenon have been given.72,150,159-163 It is observed that the elongation along the (а) axis is rapid at the beginning of the irradiation and slows down until reaching a low stationary growth rate (Figure 18). The growth strain remains small (<0.02%) and saturates at fluences less than 5 x 1024 nm~2.161,164 Eventually, at higher fluence a growth breakaway (increase of the growth rate) occurs for the annealed zirconium single crystal.161

Since the deformation of the polycrystalline clad­ding is the result of the growth of all the grains, texture has a major influence on the growth of the polycrystalline material. A weakly textured product made of zirconium alloy, with Kearns factors close to fd~ 0.33 along the three directions, such as p-quenched zirconium alloys165,166 as reviewed by Fidleris,150 exhibits a very low growth. The Kearns factor fd is the resolved fraction of basal poles along the direction d. On the other hand, strongly textured products, with most grains orientated with (c) axis along one given macroscopic direction (high Kearns factor, fd> 0.5), exhibit a negative growth in this direction and a positive growth in the direction with low Kearns factor fd< 0.2). In the case of highly textured products such as cold-worked tubing, in SRA or RXA metallurgical state, a large majority of the grains exhibit their (c) axis close to the radial direction ((c) axes oriented in the (r, в) plane with an
angle between 20° and 45° to the radial direction, the Kearns factor along the radial direction being f ~ 0.6). The directions (1120) or (1010) are along the rolling direction (low Kearns factor along the rolling, or axial direction fa ~ 0.1-0.16.167,168) Due to this strong texture, an elongation of the tube along

the rolling direction is observed159,169,168 as well as a

decrease of the thickness as shown on rolled sheet,159 the strain along the diameter of the tube remaining low.153 In the case of pressure tube for Canadian deuterium uranium (CANDU) reactors, made of cold-worked Zr-2.5Nb, since the (c) axes are mainly along the transverse direction (f ~ 0.3, fa ~ 0.05, f ~ 0.6, respectively for radial, axial, and transverse Kearns factors), the irradiation growth leads to an increase of the length in the axial direction and a decrease of the diameter.163

As for the zirconium single crystal, textured RXA Zy-4 or Zy-2 products, for instance, in the form of tubing, exhibit first a rapid elongation along the roll­ing direction, and then a decrease in the growth rate, reaching a low stationary growth rate.159 It can be noticed that the stationary growth strain of the polycrystal is higher than that for the Zr single crys — tal.161 This demonstrates the role of the grain bound­aries on the growth mechanisms. For higher fluence, higher than 3-5 x 1025 nm~2, a growth breakaway is observed, yielding a high growth rate.

It is reported150,160,166 that for polycrystalline zir­conium alloys, the grain size affects the growth rate

of RXA zirconium alloys during the initial growth transient at 553 K, the growth rate increasing when the grain size decreases. On the other hand, the stationary growth is not affected by the grain size. This phenomenon is also observed for Zircaloy-2.159 Ibrahim and Holt170 and Holt171 have also suggested that the grain shape, especially in the case of Zr-2.5% Nb material, can play a role on the growth behavior.

It is shown that for cold-worked materials (e > 10%) the growth rate increases as the cold work­ing increases150,159,160 (Figure 19). For the extreme case of SRA zirconium alloys, which could undergo up to 80% cold working followed by a SRA treat­ment, the growth rate is so high that the stationary growth rate is not observed, and from the beginning of the irradiation, the growth rate is comparable to the growth rate measured for RXA zirconium alloys after the breakaway growth. Several authors, as reviewed by Fidleris etal.159 and Holt,72 have clearly correlated the increase of the growth rate with the increase of the dislocation density due to the cold working. This also proves the importance of the ini­tial dislocations network in the growth mechanisms.

Several authors have studied the effect ofthe impu­rity and alloying elements on the growth rate and especially on the growth acceleration. At 280 °C, for a high-purity zirconium single crystal obtained by the melting zone method, no growth breakaway is observed. On the other hand, for a lower purity zirco­nium single crystal obtained by using the iodine puri­fication method161 the breakaway growth is observed.

Similarly, for polycrystalline RXA zirconium alloys, irradiated at elevated temperature (390-430 °C), the growth rate is higher than that of pure zirconium.73,160 It is particularly noticed by Griffiths eta/.73 that RXA zirconium alloys exhibit accelerated growth contrary to pure zirconium. It is believed that minor elements (Fe, Cr), and especially iron, play a major role on the breakaway.54,160 On the other hand, it appears that the tin content, in solid solution, has no effect on the stationary growth rate at high temperatures (280 °C)150,160 but that the niobium leads to a reduced growth rate compared to RXA Zy-4.168

The irradiation temperature has a complex influ­ence on the growth behavior72,150 (Figure 20). For SRA zirconium alloys, it is shown that the growth rate increases as the temperature increases. On the other hand, for RXA zirconium alloys the prebreakaway growth rate has a very low temperature sensitivity, the growth rate increasing very slowly with increasing temperature. A growth peak is even observed around 570 K, the growth rate decreasing rapidly above 620 K. However, for postbreakaway growth, the tempera­ture sensitivity is high, as high as for SRA zirconium alloys.150 It is also shown that the breakaway fluence decreases with increase in the temperature.72

4.01.3.1.2 Irradiation growth: Mechanisms

The mechanisms proposed in the literature in order to explain the growth under irradiation of zirconium and its alloys have progressively evolved as the obser­vations of the microstructure have progressed.

1/T(K)

Several comprehensive reviews of these mechanisms have been given,44’46’72’163 and a nice history of the various mechanisms for irradiation growth of zirco­nium alloys is provided by Holt.162 Some of these mechanisms are not compatible with all the observa­tions. For instance, the fact that both vacancy and interstitial (a) loops are present in the polycrystalline material, as described in the first part, shows that the model proposed by Buckley172 described in Northwood173 and Holt162 for the growth of zirconium alloys is not correct.

The most promising model that gives the best agreement with the observations is the model based on the DAD, first proposed by Woo and Gosele174 and described in detail by Woo.44 This last model is based on the assumption that the diffusion of SIAs is anisotropic, the vacancy diffusion anisotropy being low. Indeed, as reported in the first part of this chap­ter, several authors28,33,34,175 have shown, using atom­istic simulations, that the mobility of the SIAs is higher in the basal plane than along the (c) axis and that the vacancy diffusion is only slightly anisotropic.

The growth mechanism proposed by Woo44 is the most convincing model, since every feature of the growth phenomenon is understood in its frame unlike in the previous models. According to this mechanism, during the first stage of the irradiation of RXA zirconium alloys, with low initial dislocation density, the grain boundaries are the dominant sinks.

Figure 21 (a-c) The three phases of growth of

recrystallized zirconium alloys. (d) Growth mechanisms of stress relieved zirconium alloys.

Due to the rapid mobility of SIAs in the basal plane, the grain boundaries perpendicular to the basal plane are preferential sinks for SIAs. In contrast, grain boundaries parallel to the basal plane constitute preferential sinks for vacancies. This leads to a fast initial growth of polycrystalline zirconium alloys, in agreement with the model first proposed by Ball176 (Figure 21 (a)). This mechanism explains why the initial growth transient is sensitive to the grain size.

As the irradiation dose increases, the (a) loop density increases and the (a) loops become the dom­inant sink for point defects. In the absence of (c) component dislocation (as is the case in RXA zirco­nium alloys), calculations of DAD-induced bias
show that linear (a) type dislocations parallel to the (c) axis are preferential SIA sinks while (a) type loops are relatively neutral and may receive a net flow of either interstitials or vacancies, depending on the sink situation in their neighborhood. This explains why both interstitial and vacancy (a) type loops can be observed. This also explains why in the neigh­borhood of prismatic grain boundaries, or surfaces, which experience a net influx of SIAs, there will be a higher vacancy supersaturation leading to a predom­inance of vacancy loops towards interstitial loops as shown by Griffiths.46 It has to be pointed out that the simultaneous growth of interstitial and vacancy (a) type loops in the prismatic plane does not induce strain of the crystal although they are the dominant sinks (Figure 21 (b)). This explains the low station­ary growth rate observed.

For irradiation doses higher than 5 x 1025 nm~2, vacancy (c) component dislocation loops in the basal plane are observed in RXA zirconium alloys (Figure 21 (c)). The origin of the nucleation of (c) component loops remains unclear. Nevertheless, it has been shown, as described previously, that it is favored by the iron dissolution in the matrix coming from the precipitates.57,73,75,76 The appearance of (c) component defects has been clearly correlated to the breakaway growth71 (Figure 22). The fact that these vacancy (c) component basal loops are able to grow in zirconium alloys, whereas it is the (a) prismatic loops that are the most stable, is easily explained in the frame of the DAD model. Indeed, it can be shown that it is due to the DAD that vacancies are elimi­nated preferentially on the (c) component loops and on the grain boundaries parallel to the basal plane. The SIAs are eliminated on (a) type dislocations
and grain boundaries parallel to the prismatic plane. This partitioning of the point defects on these differ­ent sinks leads to the growth of the vacancy (c) component loops and therefore to the accelerated growth of RXA zirconium alloys. However, as pointed out by Griffiths et a/.,73 although there is a clear correlation between the occurrence of the breakaway and the appearance of (c) loops, the strain induced by the loops observed is much lower than the growth strain measured.

The fast and continuous growth of cold-worked or SRA zirconium alloys can also be easily explained by this model. Indeed, since in these materials the (c + a) line dislocations are already present before irradiation, under irradiation, the vacancies are prefer­entially eliminated on the dislocations with (c + a)

Burgers vector in the basal plane,72,162,163 leading to

the climb of these dislocations. On the other hand, the SIAs are eliminated on (a) type dislocations, leading to the climb of these dislocations. This parti­tioning of point defects therefore leads to the fast and continuous growth of cold-worked or SRA zirco­nium alloys (Figure 21 (d)). Here the growth created by the point-defect flux to the grain boundaries is relatively unimportant because they are not dominat­ing sinks. Irradiation growth under such circumstances is thus not sensitive to the grain size or shape.177

It has also been discussed by several authors, espe­cially by Holt,162 that due to the polycrystalline nature of the material, the growth strain of the indi­vidual grains can induce strain incompatibilities between adjacent grains that exhibit different orien­tations. Intergranular stresses can then result from these strain incompatibilities, leading to a local irradiation creep of individual grains even without

	G2	D2 D1
Many <c> component dislocations
Some <c> component dislocations
0.05

macroscopic applied stress on the material. This phe­nomenon can also affect the growth behavior of the polycrystalline material. It has also been shown that the intergranular stresses resulting from a deformation prior to irradiation can lead to a complex transient growth behavior at the beginning of the irradiation due to intergranular stress relaxation.162,178

An essential prerequisite for maximizing the ‘irra­diation creep resistance’ is to ensure38 the best combination of thermal creep behavior and long­term microstructural stability at high temperature. Hence, the present section would discuss irradiation creep in the same sequence as mentioned above.

The design principles of development of creep — resistant steels are as follows:

• Introduce high dislocation density by either trans­formations or cold work to increase the strength of the basic lattice;

• Strengthen the host lattice by either solid solution strengtheners or defects;

• Stabilize the boundaries created by phase transfor­mations by precipitating carbides along the boundaries;

• Arrest dislocation glide and climb by appropriate selection of crystal structure, solid solution, inter­faces, dislocation interactions, and crystal with low diffusivity;

• Resist sliding of grain boundaries by introducing special type of boundaries and anchoring the boundaries with precipitates;

• Ensure long-term stability of the microstructure, especially against recovery and coarsening of the fine second phase particles;

In the case of 9-12 Cr steels, the martensitic lath structure (Figure 7) decorated with only MX which should39 be stable over long-term service life is the most desired structure. Thermo-Calc evalua­tions show39 that MX can be stabilized at the expense of M23C6 only by reducing carbon to as low a value as 0.02% in 9 Cr-1Mo steel. This value is too low to ensure acceptable high temperature mechani­cal behavior of the steels. In the context of fast reactor core components, the high chromium 9-12% ferritic-martensitic steels assume relevance. Hence, an extensive database25 for a large number of commercial ferritic steels has been generated and

(a) (b)

Figure 7 The schematic39 of most undesirable (a) and desirable (b) microstructures for design of creep-resistant steels.

Figure 8(a) shows40 the continuous improvement achieved by careful modification of alloying elements, in the thermal creep behavior of successive grades of different commercial ferritic steels.

While understanding thermal creep is essential to narrow down the choice of ferritic steels for use in a fast reactor, ‘benchmarking’ the steels developed under irradiation is an essential stage before actually using the radiation-resistant steels in the reactor. The irradiation creep depends on the stress level, the temperature, and the dose. Figure 8(b) shows41 the comparison of irradiation creep of ferritic steels with competing materials like the austenitics and nickel-based alloys.

It is clear that the point defects generated during irradiation act against the design principles of devel­oping creep-resistant materials, listed earlier. The point defects accelerate the kinetics of dislocation climb, coarsen the precipitates, and generally enhance the diffusivity. In addition, the excess point defects precipitate into either interstitial or vacancy loops, but not randomly. The interaction between point defects and stress leads to the precipitation of intersti­tial loops parallel to the applied stress, while vacancy

(a)

(b)

Figure 9 The mechanisms of stress-induced preferential absorption (a) and stress-induced preferential nucleation (b) during irradiation creep.

loops form in planes perpendicular to the stress. This process (Figure 9(a)) called the stress-induced pref­erential nucleation (SIPN) results in additional creep strain solely due to irradiation. The excess point defects under temperature migrate randomly. But in the presence of an additional factor, that is, stress, the vacancies migrate preferentially to grain boundaries perpendicular to the applied stress, while the intersti­tials toward boundaries parallel to the stress. This is equivalent to removing material from planes parallel to the stress to those which are perpendicular to the applied stress, introducing additional creep strain. This process is called the stress-induced preferential absorption (SIPA) (Figure 9(b)).

The radiation-induced defects also evolve from isolated point defect to loops and voids, which have different types of influence on irradiation creep. Most often, irradiation creep occurs19,42 simultaneously with swelling and sometimes, swelling influences irradia­tion creep. At very small dose levels, swelling enhances creep rates. Beyond a certain dose levels, the creep component reduces and at high dose levels, creep dis­appears, while swelling continues. Figure 10 shows the variation in creep coefficient at various dose levels, and the regimes where swelling has an influence. The dynamics of point defects during irradiation continu­ously evolve with change in structure of dislocation network and loops. At small dose levels, there is a uniform distribution of very fine voids, which act as effective pinning centers for mobile dislocations. Thus the creep rate increases. With increase in dose levels, voids grow and multiply. The chance of interstitials and

vacancies impinging on the void surface becomes more than their reaching dislocations. The number of inter­stitials reaching a dislocation reduces. Additionally, the defect clusters, that is, the dislocation loops also undergo ‘faulting’ contributing to the density of dis­locations in the matrix. Hence, creep rate reduces, due to two factors: increased dislocation density of the matrix due to unfaulting of dislocation loops and reduced availability of interstitials to dislocations. The above process continues until complete cessation of creep, with swelling continue to take place.

At very high temperatures, the point defect migra­tion along the grain boundaries in preferential routes causes the grain boundary aided creep.

This high temperature limit of ferritic-martensitic steels restricts the application of these steels to at best, wrappers of present generation fast reactors based on oxide fuel. It is necessary to develop materi­als with better high temperature irradiation creep properties and void swelling for clad applications. The future scenario, which envisages the development of metallic fuel to ensure sustainability by breeding, could make use of ferritic steels for both clad and wrapper. This advantage arises due to the lower value ofthe anticipated clad temperatures with metallic fuels (see Chapter 3.14, Uranium Intermetallic Fuels (U-Al, U-Si, U-Mo)), whose choice is mainly to ensure sustainability using high breeding ratio.

4.03.4.2 Irradiation Embrittlement in Ferritic Steels

The stabilized ferritic steels in the normalized and tempered condition have a tempered martensitic structure with a preponderance of monocarbides that impart the necessary creep strength, while the prior austenite grain and lath boundaries are deco­rated with Cr rich M23C6 precipitates which increase the thermal stability of the steel. It is reported that thermal aging at temperatures above 773 K causes gradual but continuous degradation in upper shelf properties in addition to increase in the DBTT. The nature of embrittlement varies for different compo­nents of the reactor. For removable components such as clad, which are subjected to high temperature and pressure, with a residence time of a few years, creep embrittlement is the issue which decides their design and performance, while for permanent sup­port structures increase in hardening and loss in fracture toughness on irradiation are major issues.

The origin of embrittlement is two-fold: segrega­tion of tramp elements to prior austenite grain boundaries which make the grain boundaries deco­hesive and evolution of carbides and intermetallic phases. The latter causes progressive changes in the tempered martensitic microstructure, which deterio­rate the fracture properties of the steel, by introdu­cing irradiation hardening effects.

The increase in the ductile to brittle transition temperature, ADBTT, is known to be related to irradiation hardening, which is generally observed to saturate with fluence. Evidence for a possible maxi­mum in DBTT was observed for the 12Cr steel irra­diated in the range of 35-100 dpa in fast flux test facility (FFTF). Based on observed data in a number of cases it appears that a high fluence and/or high tempera­ture are required before a maximum is observed. This implies that the strength and impact properties are a balance between the point defect production and irradiation-induced precipitation. The precipitation during irradiation hardens the steel and irradiation accelerated recovery and aging soften the steel. The latter process is more important at high fluences and/or higher irradiation temperatures. Hence, hardening in most of these Cr-Mo steels is more than compensated for by the recovery and aging processes, leading to saturation in irradiation hardening above 723 K.

For body centered cubic materials such as ferritic martensitic steels, radiation hardening at low tempera­tures (<0.3 TM) can lead to a large increase in the DBTT and lowering of impact energy for radiation dose as low as 1 dpa (displacement per atom). The minimum operating temperature to avoid embrittle­ment in ferritic martensitic (F/M) steels is ^473- 523 K, while the upper limit is controlled by four different mechanisms: thermal creep, high temperature helium embrittlement, void swelling, and compatibility
with fuel and coolant. Void swelling is not expected to be significant in F/M steels up to damage levels of about 200 dpa.

Extensive evaluation14’15’43-58 of the embrittle­ment behavior of the ferritic steels for different chemistry is shown in Figure 11. The merit in focus­ing on chemistry around 9% chromium is very clear based on the observation of minimum shift in DBTT around this composition, under irradiation. However, higher chromium improves corrosion resistance and ease of reprocessing. Hence, chromium content has to be selected balancing these requirements. It is known44-48 that addition of phosphorous, copper, vanadium, aluminum, and silicon would increase the DBTT while sulfur reduces the upper shelf energy (USE). The 12Cr steels, HT9, show a larger shift (125 K) in DBTT as compared to modified 9Cr-1Mo steel (~54K). Hence, the balance is always between nearly nil swelling resistant 12Cr steels and 9Cr steel which is less prone to embrittlement than 12Cr steels.

Microstructural parameters, like the prior aus­tenite grain size, lath/packet size, carbides, and their distribution influence49,50 the embrittlement behav­ior. Studies on the effects of heat treatment and microstructure on the irradiation embrittlement in 9Cr-1MoVNb and HT9 steels are summarized below:

• Prior austenite grain size (PAGS) influences51 the

DBTT for the 9Cr-1MoVNb steel, but not in

12Cr-MoVW steel. This is attributed to the precipitates in the microstructure controlling the fracture behavior rather than the PAGS, in the 12Cr steel.

• The size of martensitic lath and packet, which is sensitive52 to austenitization temperature, can also affect51 the fracture behavior. Examination of the fracture surface revealed cleavage and regions of ductile tearing along prior austenite grain and lath packet boundaries. Subsurface microcracks and sec­ondary surface cracks were found associated with large boundary carbides. It was suggested that cleav­age fracture initiated in HT9 by propagation of a microcrack from a coarse carbide into the matrix. Propagation was inhibited by the intercepted boundaries, lath or grain and ductile tearing was required53 to continue propagation. The amount of tearing increased with increasing austenitization temperature.

• Tempering for the two normalization tempera­tures had very small effect on the DBTT, for the two steels.

• Irradiation of the two steels at 638 and 693 K resulted37 in an increase in DBTT and a decrease in USE for all conditions with the shift in DBTT for the 12 Cr steel being almost twice that for 9Cr steel.

•

Although the 12Cr steel with the smallest grain size had55 the lowest DBTT after 20 dpa, the effect of tempering was different. In the case

of 12Cr steel, the higher tempering temperature causes coarsening of precipitates thus accelerating fracture.

• The saturation of shift in DBTT with fluence is independent54 of tempering conditions for the 9Cr steel, while for the 12Cr steel, a maximum is observed, probably due to faster growth of preci­pitates during irradiation.

The generation of helium through (n, a) reaction in elements of structural materials is known to cause severe damage to the embrittlement behavior of core component materials. Table 6 lists the shift in DBTT, for 9 and 12CrMo steels, under reactor irra­diation, with and without helium, which demon — strates56 the harmful effect of helium. These results become more pertinent in the case of fusion reactors, where the operating conditions include the genera­tion of helium up to about ^100appmyear~

The increase in the DBTT due to irradiation is a cause of serious concern for use of ferritic steels, since it makes the postirradiation operations very difficult. Several methods have been attempted57,58 to address this problem, which includes modification of the steel through alloying additions, control of tramp elements by using pure raw materials and improved melting practices, and grain boundary engineering (GBE). However, the propensity of the problem is less if the clad thickness is low, which normally is the case to ensure best heat transfer properties. For low thickness components, the triaxial stress necessary for the embrittlement does not develop, which reduces the intensity of this otherwise serious problem of embrittlement in ferritic steels.

An approach to reduce shift in DBTT is an immediate concern in ferritic steels for core com­ponent applications and efforts to overcome this problem by selection of high purity metals, adoption of double or triple vacuum melting for steel making, strict control of tramp and volatile elements, and development of special processing methods, which would improve the nature of grain boundaries (GBE) are in progress.

4.06.1 Introduction

Refractory metals and alloys offer attractive and promising high-temperature properties, including high-temperature strength, good thermal conductiv­ity, and compatibility with most liquid metal cool­ants, many of which are suitable for applications in nuclear environments. Though many of the refrac­tory alloys have been known for more than 60 years, there are significant gaps in the materials property database for both unirradiated and irradiated conditions. In addition, significant issues related to low-temperature irradiated mechanical property degradation at even low neutron fluences restrict the use of refractory metals. Protection from oxidizing environments also restricts their use, unless suitable protection or a liquid metal coolants is used.

Much of the early research on refractory metal alloys was centered on applications in aerospace as well as cladding and structural materials for fission reactors, with particular emphasis on space reactor applications. Reviews concerning the history of these programs and the development of many of the alloys whose irradiated properties are discussed in this chapter can be found elsewhere.1-5 Due to cancellations and reintroduction of new mission cri­teria for these space reactor programs, the materials database shows similar waves in the gains of intellec­tual knowledge regarding refractory alloy and irra­diated property behavior. Unfortunately, as seen in the subsequent sections of this chapter, much of the irradiated property database for refractory metals con­sists of scoping examinations that show little overlap in either material type, metallurgical conditions (i. e., grain size, impurity concentrations, thermomechanical treatments), radiation conditions (i. e., spectra, dose and temperature), or postirradiation test conditions or methods.

The irradiation behavior of body-centered cubic (bcc) materials is known. Irradiation-induced swelling because of void formation in the material lattice is typical for temperatures between 0.3 and 0.6 Tm, where Tm is the melting temperature. Maximum swelling in refractory metals is <10% for displace­ment damage levels up to 50 dpa (displacements per atom), but typical values for fission reactor applica­tions are <4%. Alloy additions can further reduce the sensitivity to swelling, for example, rhenium additions to molybdenum or tungsten. These levels of swelling are manageable through the appropriate engineering design of components.

The generation of dislocation loops and point defects provide significant irradiation-induced strengthening or hardening of refractory metals and alloys. This in turn creates reductions in the ductility and fracture toughness of the material. This is most pronounced at temperatures <0.3 Tm, where defect mobility is reduced. The increase in the yield strength of the material because of the irradiation — induced defects can exceed the fracture strength of the material, leading to brittle behavior. These degradations in material property can begin to occur at neutron fluences as low as 1 x 102°ncm~2, or ^0.03 dpa3 and increase in severity with dose. As irradiation temperatures increase, dislocation loop and void sizes increase, whereas their densities are reduced, providing improvements in ductility, though at a reduced strength of the material. At high enough temperatures, recovery of properties to levels close to that of the unirradiated values is possible, though changes in material properties may be further influ­enced by microstructural changes such as segregation or precipitate formation of solute and transmuted spe­cies or recrystallization, which can lead to further dete­rioration of properties. Detailed information on the effects of radiation on materials is presented in Chapter 1.03, Radiation-Induced Effects on Microstructure, and in Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys. In general, the use of refractory alloys in radiation environments is not recommended at temperatures <0.3 Tm. However, new research work, particularly on molybdenum and its alloys, has shown that control over interstitial element contamination levels, grain size, and morphol­ogy, as well as the introduction of oxide dispersion strengthening, can lead to improvements in the low — temperature irradiation behavior. This is discussed in detail in this chapter.

The following sections of this chapter deal with the irradiated properties database of niobium, tanta­lum, molybdenum, and tungsten, as well as their alloys. While vanadium may sometimes be consid­ered a refractory metal, its melting temperature is considerably lower than that of the other materials mentioned. However, its radiation effects database is considerable and well advanced relative to some refractory metals and it is therefore discussed sepa­rately in Chapter 4.12, Vanadium for Nuclear Sys­tems. The irradiated properties database for refractory alloys is particularly thin, especially involving frac­ture toughness properties, irradiation creep effects, and combined radiation effects with high thermome­chanical loads such as those experienced in plasma facing components or spallation target materials. Where needed, a comparison of the unirradiated and irradiated properties of a material is given.

The internal creep rupture properties of the manu­factured 12Cr-ODS steels at 700 °C are shown in Figure 20.36 Increasing the yttria and titanium con­tents improves the internal creep rupture strength (F4 > F3 > F2 > F1). The uni-axial creep rupture strength for F4 is also plotted; there is the strength anisotropy between the uni-axial and internal hoop directions. This strength anisotropy can be associated with the slightly elongated grain structure shown in Figure 19.

The stress-strain rate relationship was investi­gated for ODS ferritic claddings to evaluate the creep deformation mode. The results of the analyses are given in the log-log plot in Figure 21.36 In general, the creep strain rate in the steady-state con­dition is expressed using applied stress o as:

" = As” [5]

where n is the stress exponent and A is the temperature-dependent coefficient.37 In the case of

1000

ra

CL

СЛ

CL

о

о

X

the uni-axial creep mode, a significantly high stress sensitivity of n = 43.7 appears. This stress exponent value is typical for an ODS alloy.37 The applied stress that initiates the strain is clearly located around 250 MPa; this stress corresponds to the so-called threshold stress for deformation. On the other hand, the stress exponent, n, is 10.4 for the internal creep
mode of F4, and a higher strain rate is found even below a stress of 200 MPa. A transverse section of this specimen shows finely equi-axed grains of 5-10 pm (Figure 19). Apart from pinning the gliding dislocations due to oxide particle-dislocation interaction, the defor­mation mechanism associated with grain morphology may be the dominant factor that induced accelerated strain in the hoop stress mode of the tubular specimen.

In order to characterize the high temperature strength of manufactured 12Cr-ODS steel cladding, its strength mechanism was evaluated from the view­point of the interaction between Y2O3 particles and dislocations. This interaction could be formu­lated by the void-hardening mechanism proposed by Srolovitz,38 in which oxide particles were re­placed by voids. The oxide particle-hardening stress sp can be evaluated by the following equation based on Scattergood and Bacon’s equation,39 which takes into account the interaction between the branches of the bowed-out dislocation around a Y2O3 particle:

sp/G = AMb/(2nX)[ln(D/r0) + B], [6]

for screw dislocation,

A =(1 + v sin2′)cos ‘/(1 — v),

B = 0.6

for edge dislocation,

A = [l — v sin2’/(1 — v)] cos ‘,

B = 0.7

where G is the Shear modulus, v is Poisson’s ratio, M is the Taylor factor, b is the magnitude of Burgers vector, and r0 is the inner cut-off radius of the dislo­cation core. The value of ‘ is the critical angle at which the dislocation detaches from the particles. This value was estimated to be ‘ = 46° for screw dislocations and ‘ = 19° for edge dislocations. Fur­ther, l is the average face-to-face distance between particles on a slip plane and is given as a function of the average particle radius rs and the average center — to-center distance ls between the particles by

l = 1.254 — 2rs, [7]

where the averages are calculated by considering the size distribution of the particles. The factor 1.25 is the conversion coefficient from regular square distri­bution to random distribution.40 The characters ls and rs represent the results of the measurement of oxide particles by means of TEM. D is the harmonic mean of 2rs and l. The values of l were calculated, and the oxide particle-hardening stress was estimated by

substituting l, M = 3.0,41 n = 0.334, b = 2.48 x 10 10 m, and G = 50 600 MPa, at 700 °C.

Figure 22 shows the results of analyses in rela­tion to the face-to-face distance between particles.36 The oxide particle-hardening stress levels estimated by using the aforementioned equations at 700 °C are represented by vertical bars, with the upper and

Face-to-face distance between particles, l (nm)

lower bars derived from an estimate of edge and screw dislocations, and with the uncertainty of r0 ranging from b to (3 x b). The measured stress in the uni-axial mode of the F4 specimen is shown by an open circle. These results imply that the higher oxide particle-hardening stress for specimen F4 is due to its shortened face-to-face particle distance l of 70 nm. The lower band represents the stress corresponding to a strain rate of 10~9s_1 in the internal hoop direc­tional mode. For the F1 specimen, as a stress level corresponding to a strain rate of 10~9s_1 approaches the oxide particle-hardening stress, the strong anisot­ropy tends to disappear. However, for the F3 and F4 specimens with a shortened distance between par­ticles, stress levels for a strain rate of 10~9s_1 in the hoop direction are degraded from the oxide particle­hardening stress. The strong anisotropy still remains in the F4 specimen. The accelerated deformation in the internal hoop direction could be the result of grain boundary sliding, since finely equi-axed grains with a small size of 5-10 pm are formed, and the grain bound­aries occupy a large fractional area in the transverse cross-section of the F4 specimen (see Figure 19). Based on these results, it seems to be difficult to control internal creep rupture strength by recrystallization processing in 12Cr-ODS steel cladding.

The mechanical and physical properties of several medium-grained and fine-grained nuclear grade gra­phites currently in production are given in Table 2 (see also Chapter 2.10, Graphite: Properties and Characteristics). The coke type, forming method, and potential uses of these grades are in Table 1. The most obvious difference between the four grades listed in Table 2 is the filler particle sizes. Grade IG-110 is an isostatically pressed, isotropic grade, whereas the others grades shown are near-isotropic and have properties reported either with-grain or against-grain. As discussed earlier (see Section 4.10.2), the orientation of the filler coke particles is a function of the forming method.

The mechanical properties of nuclear graphites are substantially altered by radiation damage. In the unirradiated condition, nuclear graphites behave in a brittle fashion and fail at relatively low strains. The stress-strain curve is nonlinear, and the fracture process occurs via the formation of subcritical cracks, which coalesce to produce a critical flaw.35,36 When graphite is irradiated, the stress-strain curve becomes more linear, the strain to failure is reduced, and the strength and elastic modulus are increased. On irra­diation, there is a rapid rise in strength, typically ^50%, that is attributed to dislocation pinning at irradiation-induced lattice defect sites. This effect is largely saturated at doses >1 dpa. Above ~1 dpa, a more gradual increase in strength occurs because of

structural changes within the graphite. For nuclear graphites, the dose at which the maximum strength is attained loosely corresponds with the volume change turnaround dose, indicating the importance of pore closure and generation in controlling the high-dose strength behavior, and may be as much as twice the unirradiated value.

The strain behavior of nuclear graphites subjected to an externally applied load is largely controlled by shear of the component crystallites. As with strength, irradiation-induced changes in Young’s modulus are the combined result of in-crystallite effects, due to low fluence dislocation pinning, and superimposed structural change external to the crystallite. The effects of these two mechanisms are generally consid­ered separable, and related by

(E/E0) irradiated = (E/Eo)pinning(E/Eo)structure I1]

where E/E0 is the ratio of the irradiated to unirradi­ated elastic modulus. The dislocation pinning contri­bution to the modulus change is due to relatively mobile small defects and is thermally annealable at ^2000 °C. The irradiation-induced elastic modulus changes for GraphNOL N3M graphite37 are shown in Figure 13. The low dose change due to dislocation pinning (dashed line) saturates at a dose < 1 dpa.

The elastic modulus and strength are related by a Griffith theory type relationship.

Strength, st =[GE/nc]l/1 [2]

where G is the fracture toughness or strain energy release rate (J m — ), E is the elastic modulus (Pa), and c is the flaw size (m). Thus, irradiation-induced changes in st and E (in the absence of changes in [G/c]) should follow st/E1/2. High-dose data reported by Ishiyama et a/.38 show significant devia­tion from this relationship for grade IG-110 graphite, indicating that changes in G and or c must occur.

Graphite is a phonon conductor of heat. Therefore, any reduction in the intrinsic defect population causes a reduction in the degree of phonon-defect scattering, an increase in the phonon mean free path, and an increase in the thermal conductivity. In graphite, such thermally induced improvements are attributable to increases in crystal perfection and a concomitant increase in the size of the regions of coherent order­ing upon graphitization. With increasing temperature, the dominant phonon interaction becomes phonon — phonon scattering (Umklapp processes). Therefore, there is a reduction of thermal conductivity with increasing temperature.39 This decrease in the thermal conductivity with increasing temperature can clearly be seen in Figure 14.

The mechanism of thermal conductivity and the degradation of thermal conductivity have been exten­sively reviewed.13,14,26,40 The increase of thermal resis­tance due to irradiation damage has been ascribed by Taylor eta/.41 to the formation of (1) submicroscopic interstitial clusters, containing 4 ± 2 carbon atoms; (2) vacant lattice sites, existing as singles, pairs, or small groups; and (3) vacancy loops, which exist in the graphite crystal basal plane and are too small to have collapsed parallel to the hexagonal axis. The contribution of collapsed lines of vacant lattice sites and interstitial loops, to the increased thermal resis­tance, is negligible.

The reduction in thermal conductivity due to irradiation damage is temperature and dose sensitive. At any irradiation temperature, the decreasing thermal conductivity will reach a ‘saturation limit.’ This limit is not exceeded until the graphite undergoes gross struc­tural changes at very high doses. The ‘saturated’ value of conductivity will be attained more rapidly, and will be lower, at lower irradiation temperatures.42 In graph­ite, the neutron irradiation-induced degradation of thermal conductivity can be very large, as illustrated in Figure 14. This reduction is particularly large at low temperatures. Bell et a/.43 have reported that the room temperature thermal conductivity of pile grade A (PGA) graphite is reduced by more than a factor of 70 when irradiated at 155 °C to a dose of —0.6 dpa. At an irradiation temperature of 355 °C, the room temperature thermal conductivity of PGA was reduced by less than a factor of 10 at doses twice that obtained at 155 °C. Above 600 °C, the reduc­tion of thermal conductivity is less significant. For example, Kelly8 reported the degradation of PGA at higher temperatures: at an irradiation temperature of 600 °C and a dose of — 13 dpa, the thermal conductiv­ity was degraded only by a factor of —6; at irradiation temperatures of 920 and 1150 °C, the degradation was minimal (less than a factor of 4 at —7 dpa). For the fine-grained, isomolded graphite shown in Figure 14, the degradation of thermal conductivity at the irradia­tion temperature (600 °C) was only by a factor of —3, but was by a factor —6 at a measurement temperature of 100 °C.

There are two principal thermal expansion coeffi­cients in the hexagonal graphite lattice; a, the ther­mal expansion coefficient parallel to the hexagonal c-axis and aa, the thermal expansion coefficient par­allel to the basal plane (a-axis). The thermal expan­sion coefficient in any direction at an angle f to the c-axis of the crystal is

a(f) = ac cos2f + aa sin2f [3]

The value of ac varies linearly with temperature from -25 x 10—6K—1 at 300 K to -35 x 10—6K—1 at 2500 K. In contrast, aa is much smaller and increases rapidly from —1.5 x 10—6K—1 at — 300 K to —1 x 10—6K—1 at 1000 K, and remains relatively constant at temperatures up to 2500 K.39

The large anisotropy in the crystal coefficient of thermal expansion (CTE) values is a direct conse­quence of the bond anisotropy and the resultant anisotropy in the crystal lattice compliances. The thermal expansion ofpolycrystalline graphites is con­trolled by the thermal closure of aligned internal porosity which forms as a result of thermal shrinkage strains on cooling after graphitization. Thus, the c-axis expansion of the graphite crystals is initially, partially accommodated by this internal porosity and a much lower bulk CTE value is observed. On further heating, the graphite crystals fill more of the available internal porosity and more of the c-axis expansion is observed. The bulk CTE thus increases with temper­ature (Figure 15).

Figure 15 Temperature dependence of the coefficient of thermal expansion for typical nuclear grade graphite.

Figure 16 The irradiation-induced changes in coefficient of thermal expansion (25-500 °C) for GraphNOL N3M graphite at two irradiation temperatures. From Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179-181, 205-208.

As the CTE of polycrystalline graphite is depen­dent on the pore structure, irradiation-induced changes in the pore structure (see discussion of structural changes in Section 4.10.4) can be expected to modify the thermal expansion behavior of carbon materials. Burchell and Eatherly37 report the behav­ior of GraphNOL N3M (which is typical of many fine-textured graphites), which undergoes an initial increase in the CTE followed by a steady reduction to a value less than half the unirradiated value of 5 x 10-6 °C-1 (Figure 16). Similar behavior is reported by Kelly8 for grade IM1-24 graphite. Heat energy is stored in the crystal lattice in the form of lattice vibrations. The Debye equation there­fore gives the specific heat, C, as follows:

Figure 17 The temperature dependence of the specific heat of graphite, a comparison of calculated values and literature data for POCO AXM-5Q graphite. Sources: ASTM C 781. Standard Practice for Testing Graphite and Boronated Graphite Materials for High-Temperature Gas-Cooled Nuclear Reactor Components, Annual Book of Standards. ASTM International: West Conshohocken, PA; Vol. 05.05; Hust, J. G. NBS Special Publication 260-89; US Department of Commerce, National Bureau of Standards, 1984; p 59.

C =9RI, 5D

[4]

where R is the gas constant (8.314J mol-1 K-1); T, the temperature; 0D, the Debye temperature; and z = hm/2kT%, where o is the frequency of vibrational oscillations; k, the Boltzmann’s constant; T, the tem­perature; and h is the Plank’s constant.

At low temperatures, where (T/0D) <0.1, z in eqn [4] is large, we can approximate eqn [4] by allowing the upper limit in the integral to go to infinity such that the integral becomes ^(я4/15), and on differen­tiating we get

C = 1941(T / yD)3 J mol-1 K-1 [5]

Thus, at low temperatures, the specific heat is pro­portional to T3 (eqn [5]). At high temperatures, z is small and the integral in eqn [4] reduces to z2dz; hence, on integrating we get the Dulong-Petit value of 3R, that is, the theoretical maximum specific heat of 24.94 J mol-1 K-1. As we are typically concerned only with the specific heat at temperatures above

(Wigner energy) will reduce the effective specific heat (see Section 4.10.4).

1

11.07Г-1-644 + 0.0003688 T0 02191

The electrical resistivity of graphite is also affected by radiation damage. The mean free path of the conduction electron in unirradiated graphite is rela­tively large, being limited only by crystallite bound­ary scattering. Neutron irradiation introduces (1) scattering centers, which reduce charge carrier mobility; (2) electron traps, which decrease the charge carrier density; and (3) additional spin reso­nance. The net effect of these changes is to increase the electrical resistivity on irradiation, initially very rapidly, with little or no subsequent change to rela­tively high fluence.14,37 A subsequent decrease at very high neutron doses is attributed to structural degradation.

Fast neutron irradiation and, in the case of car­bon dioxide-cooled reactors, radiolytic oxidation change many of the properties of graphite. The

properties of interest to the nuclear engineer are

the following:

• Stored energy — a function of fast neutron damage and temperature, due to damage to the graphite crystallites, but not affected by radiolytic oxidation other than by a reduction in mass.

• Specific heat — a function of temperature but not affected by fast neutron irradiation or radiolytic oxidation other than by stored energy, which may be considered separately.

• Dimensional changes — a function of fast neutron damage and irradiation temperature. There is also some evidence of modification by radiolytic oxida­tion. It is also modified by stress (see irradiation creep).

• CTE — a function of temperature, fast neutron damage, irradiation temperature, and stress. There is evidence that it is not modified by radio­lytic weight loss.

_ 30

28

26 c

24

(Л

22

$

20 E

18 .c

16

14

о

12

о

° 10 ————- ,——— ,———- ,——— ,———- ,——— ,——— ,———- ,

0 2 4 6 8 10 12 14 16

Fluence (1020 ncm-2 EDND)
c-axis

Fluence (1020 ncm-2 EDND)
a-axis

♦ 150 °C ■ 170 °С A 200 °С • 250 °C

Figure 22 Low-temperature changes to coefficient of thermal expansion in highly orientated pyrolytic graphite. Modified from Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260(1109), 37-49.

• Thermal conductivity — a function of temperature, fast neutron damage, and irradiation tempera­ture. It is significantly modified by radiolytic weight loss.

• Young’s modulus — a function of temperature, fast neutron damage, and irradiation temperature. It is significantly modified by radiolytic weight loss.

• Strength (tensile, compressive, flexural, and frac­ture) — a function of temperature, fast neutron damage, and irradiation temperature. It is signifi­cantly modified by radiolytic weight loss.

• Electrical resistivity — a function of temperature, fast neutron damage, and irradiation temperature. It is probably modified by radiolytic weight loss.

• Irradiation creep — a function of fast neutron dam­age, irradiation temperature, and stress.

These property changes are illustrated in Figure 30.

These dimensional changes, property changes and creep mechanisms are correlated, some more strongly than others. This has been taken advantage of in various semiempirical models for irradiation damage in graphite.57-60

In discussing the irradiation behavior ofpolycrystal — line graphite, it is useful to split these changes into low, medium, and high fluence effects. At low irradiation, fluence changes in polycrystalline graphite are strongly correlated with the crystallite changes discussed else­where; see Section 4.11.11. Typical mechanisms would

_ 30

the accumulation of stored energy, pinning (leading to a rapid increase in Young’s modulus), and the rapid decrease in thermal conductivity. At medium fluence, several of the properties saturate, such as Young’s mod­ulus and thermal conductivity. At high fluence, when crystallite growth in the ‘c’ direction has taken up much of the accommodation provided by Mrozowski cracks,11 and larger ‘cracks,’ the polycrystalline struc­ture starts to become strained, thereby generating new cracking. At extremely high fluence, beyond that expe­rienced in a modern power production reactor, the crystallite swelling becomes so large that the polycrys­talline structure completely breaks down leading to a rapid decrease in modulus and strength.

Each of the property changes is discussed in more detail below. In attempting to understand the behavior of polycrystalline graphite, reference is made to the irradiation behavior of HOPG, as previously discussed in Section 4.11.11. This is because HOPG is considered to be a representative model material for the individual crystallite structures in polycrystalline graphite.

4.11.10 Averaging Relationships

Before looking at each of the properties individually, it is first worth considering the methods developed to relate changes in the crystallites to the bulk

♦50 °C ■ 150°C » 650°C • 1000°C
0.5	1.0	1.5	2.0	2.5

Fluence (1020 ncm-2 EDND)

Figure 24 Reduction in C33 in irradiated highly orientated pyrolytic graphite. Modified from Seldin, E. J.;

Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389-3400; Summers, L.; Walker, D. C. B.; Kelly, B. T. Philos. Mag. 1966, 14(128), 317-323.

_ 5 E

4

И

o’ 3

CO

1 2

О

E

1

CO

_cc

Ш

0

50 °C

_ 0.5

E

0.3

от

0.2

о

E

0.1

to

_ra

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4

Fluence (1020 ncm-2)

650 and 1000 °C

Figure 25 Changes in C44 with irradiation in highly orientated pyrolytic graphite. Modified from Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389-3400.

properties of polycrystalline graphite. In the 1960s, Simmons derived a model based on the following assumptions: polycrystalline graphite could be con­sidered to be a single phase, porous aggregate of
identical crystals with the correct graphite symmetry, and small element volumes containing only graphite could be chosen so that their internal and external stress could be considered to be uniform.61

From first principles and by applying the laws of thermodynamics for the assumptions given above, Simmons derived the following relationship for the bulk dimensional change rate:

, dxc, , .

Ax^—h(1 _ Ax) dg

where dexx/dg, dxc/dg, and dxa/dg are the linear dimensional change rate in polycrystalline graphite in some direction V and the crystallite dimensional change rates in the ‘a’ and V directions respectively. Simmons also derived the following relationship for the bulk CTE:

axx Axac ^ (1 Ax)aa [34]

where axx, ac, and aa are the linear CTE in poly­crystalline graphite in some direction ‘x’ and the crystallite CTE in the ‘a’ and V directions, respectively.

The so-called structure factor ‘Ax’ is the summa­tion rate of change in rate of the crystallite stresses with respect to the change in bulk stress, as illustrated in the equations below.

0 50 100 150 200 250 300

Fluence (1 020 ncm-2 EDND)

Figure 26 Changes in the thermal conductivity of highly orientated pyrolytic graphite. (a) Change in thermal conductivity in the ‘a’ direction as a function of fluence, (b) Change in thermal resistance as a function of temperature, and (c) Normalized change in thermal resistivity as a function of irradiation and measurement temperature. Reproduced from Taylor, R.;

Kelly, B. T.; Gilchrist, K. E. J. Phys. Chem. Solids 1969, 30, 2251-2267.

(1 ) = Qs (°22,п + ^ll,»^

For a more detailed derivation of these equations see Hall et al61

In addition, Simmons4 provided evidence that there is a linear relationship between unirradiated CTE and initial dimensional change rate for poly­crystalline graphite (Figure 31). Brocklehurst and Bishop62 later found a similar relationship in bromi — nated graphite.

Expressions to those of Simmons have been derived by Sutton and Howard12:

«par = K1g « + K2baa

«perp K3gac ^ K4b«a [36]

where K1, K2, K3, K4, g, and b are crystal accommoda­tion and Bacon13 crystal orientation factors. A similar but more complex relationship was also derived by Jenkins.63 In the discussion of the individual
properties, the anisotropic graphite PGA and the semi-isotropic Gilsocarbon graphite are used as examples.

All commercial NPPs in the United States contain structures whose performance and function are necessary for the protection of the safety of

plant-operating personnel and the general public, as well as the environment. The basic laws that regulate the design (and construction) of NPPs are contained in Title 10ofthe Code of Federal Regulations (CFR),10 which is clarified by Regulatory Guides (e. g., R. G. 1.29),11 NUREG reports, Standard Review Plans (e. g., Con­crete and Steel Internal Structures of Steel or Concrete Containments),12 etc. In addition, R. G. 1.29 and Part 100 to Title 10 of the CFR state that NPP structures important to safety must be designed to withstand the effects of earthquakes without the loss of function or threat to public safety. These ‘safety-related’ structures are designed as Seismic Category I. Seismic Category I structures typically include those classified by the American Society of Mechanical Engineers (ASME) and the American Nuclear Society (ANS) as Classes 1, 2, and 3 (i. e., safety related).

Initially, existing building codes such as the ACI Standard 318, Building Code Requirements for Reinforced Concrete, were used in the nuclear industry as the basis for the design and construction of concrete structural members. However, because the existing building codes did not cover the entire spectrum of design requirements and because they were not always con­sidered adequate, the United States Nuclear Regu­latory Commission (USNRC) developed its own criteria for the design of Category I structures (e. g., definitions of load combinations for both operating and accident conditions). Current requirements for nuclear safety-related concrete structures, other than concrete reactor vessels and concrete containments, are also based on ACI 318, but have incorporated modifications to accommodate the unique perfor­mance requirements of NPPs. These requirements were developed by ACI Committee 349 and first published in October 1976.13 This Code has been endorsed by the USNRC as providing an adequate basis for complying with the general design criteria for structures other than reactor vessels and contain — ments.14 USNRC15 provides additional information on the design of seismic Category I structures that are required to remain functional if the Safe Shutdown Earthquake (SSE) occurs. Current requirements for concrete reactor vessels and concrete containments were developed by ACI Committee 359 and first published in 1977.16 Supplemental load combination criteria are presented in Section 3.8.1 of the USNRC Regulatory Standard Review Plan} However, since all but one of the construction permits for existing NPPs have been issued prior to 1978, it is unlikely that endorsed versions of either ACI 349 or ACI 359 were used in the design of many of the concrete structures at these plants. Older plants that used early ACI codes, however, have been reviewed by the USNRC through the Systematic Evaluation Program to determine if there were any safety concerns.18

Each boiling water reactor (BWR) or pressurized water reactor (PWR) unit in the United States is located within a much larger metal or concrete con­tainment that also houses or supports the primary coolant system components. Although the shapes and configurations of the containment can vary signifi­cantly from plant to plant, leak tightness is ensured by a continuous pressure boundary consisting of non­metallic seals and gaskets and metallic components that are either welded or bolted together. There are several CFR General Design Criteria (GDC) and ASME Code sections that establish minimum requirements for the design, fabrication, construc­tion, testing, and performance of the light-water reac­tor (LWR) containment structures. The GDC serve as fundamental underpinnings for many of the most important safety commitments in licensee design and licensing bases. General Design Criterion 2, Design Bases for Protection Against Natural Phenomena, requires the containment to remain functional under the effects of postulated natural phenomena such as earthquakes, tornadoes, hurricanes, floods, tsunami, and seiches. General Design Criterion 16, Containment Design, requires the provision of reactor containment and associated systems to establish an essentially leak — tight barrier against the uncontrolled release of radio­activity into the environment and to ensure that the containment design conditions important to safety are not exceeded for as long as required for postulated accident conditions. Criterion 53, Provisions for Con­tainment Testing and Inspection, requires that the reactor containment be designed to permit (1) appropriate periodic inspection of all important areas, such as penetrations; (2) an appropriate surveillance program; and (3) periodic testing at containment design pressure of leak tightness of penetrations that have resilient seals and expansion bellows. Current LWR contain­ments are considered as a significant element of the USNRC’s safety policy, which employs a defense- in-depth approach (i. e., successive compensatory measures are exercised to prevent accidents or miti­gate damage if a malfunction, accident, or naturally caused event occurs). The defense-in-depth philoso­phy ensures that safety will not be wholly dependent on any single element of the design, construction, maintenance, or operation at a nuclear facility (e. g., the facility in question tends to be more tolerant of failures and external challenges).

From a safety standpoint, the containment is one of the most important components of an NPP because, independent of the fuel barrier and reactor coolant pressure boundary barrier, it serves as the final barrier to the release of fission products to the outside environment under postulated accident conditions. During normal operating conditions, the containment is subject to various operational and environmental stressors (e. g., ambient pressure fluc­tuations, temperature variations, earthquakes, and wind storms). In some containment designs, the prin­cipal leak-tight barrier is surrounded by another structure (e. g., reactor or shield building) that pro­tects the containment from external events. Ensuring
that the structural capacity and leak-tight integrity of the containment has not deteriorated unacceptably because of aging or environmental stressor effects is essential to reliable continued service evaluations and informed aging management decisions. More detailed information on containments is available.1

In addition to the containment, a myriad of concrete-based structures are contained as a part of an LWR plant to provide foundation, support, shield­ing, and containment functions. Table 1 presents a listing of typical safety-related concrete structures that may be included as part of an LWR plant.20 Relative to general civil engineering reinforced con­crete structures, NPP concrete structures tend to be

Table 1 Typical safety-related concrete structures in LWR plants and their accessibility for visual examination

Concrete structure

Primary containment Containment dome/roof Containment foundation/basemat Slabs and walls Containment internal structures Slabs and walls

Reactor vessel support structure (or pedestal) Crane support structures Reactor shield wall (biological)

Ice condenser dividing wall (ice condenser plants) NSSS equipment supports/vault structures Weir and vent walls (Mark III)

Pool structures (Mark III)

Diaphragm floor (Mark II)

Drywell/wetwell slabs and walls (Mark III) Secondary containment/reactor buildings Slabs, columns, and walls Foundation

Sacrificial shield wall (metallic containments) Fuel/equipment storage pools Walls, slabs, and canals Auxiliary building Fuel storage building Control room (or building)

Diesel generator building

Piping or electrical cable ducts or tunnels

Radioactive waste storage building

Stacks

Intake structures (including concrete water intake piping and canal embankments)

Pumping stations

Cooling towers

Plant discharge structures

Emergency cooling water structures

Dams

Water wells Turbine building

Internal liner/complete external

Internal liner (not embedded) or top surface

Internal liner/external above grade

Generally accessible Typically lined or hard to access Generally accessible Typically lined Lined or hard to access Generally accessible Lined with limited access Lined

Lined with limited access Internal liner/partial external access

Accessible on multiple surfaces Top surface

Internal lined/external accessible

Internal lined/partial external

Generally accessible

Generally accessible

Generally accessible

Generally accessible

Limited accessibility

Generally accessible

Partial internal/external above grade

Internal accessible/external above grade and waterline

Partially accessible Accessible above grade

Internal accessible/external above grade and waterline

Limited accessibility

External surfaces above waterline

Limited accessibility

Generally accessible

Source: Hookham, C. J. In-Service Inspection Guidelines for Concrete Structures in Nuclear Power Plants, ORNL/NRC/LTR-95/14; Lockheed Martin Energy Systems, Oak Ridge National Laboratory: Oak Ridge, TN, 1995.

more massive and have increased steel reinforcement densities with more complex detailing. Information pertaining to a particular structure at a plant of interest can be obtained from sources such as the plant’s safety analysis report or docket file. Concrete structures that are considered to be ‘plant specific’ or unique have not been addressed in the discussion later, but some information provided for similar structures may be applicable. Additionally, the names of certain structures may vary from plant to plant depending on the nuclear steam supply system (NSSS) vendor, architect engineering firm, and owner preference. Typical safety-related concrete structures contained in LWR plants may be grouped into four categories: primary containments, contain­ment internal structures, secondary containment/ reactor buildings, and other structures.

Stainless steels undergo an evolution of phase struc­ture at reactor-relevant temperatures, even in the absence of radiation. These changes involve the for­mation of various carbides, later followed by various intermetallic phases.1,108 This evolution is accompa­nied by net changes in average lattice parameter arising from differences in partial molar volume of elements when passing from one phase to another.

By-92

52 dpa 29.8% U-796

34 dpa max 14% swelling

Figure 42 Severe embrittlement and failure in three BOR — 60 reflector assembly ducts. The ducts were made of annealed X18H10T, the Russian equivalent of 321 steel. Reproduced from Neustroev, V. S.; Ostrovsky, Z. E.; Teykovtsev, A. A.; Shamardin, V. K.; Yakolev, V. V. In Proceedings of 6th Russian Conference on Reactor Materials Science; 11-15 September 2000, Dimitrovgrad, Russia, in Russian. The maximum swelling values (from left to right) were 27.8, 29.8, and 14%. Failure was the result of high withdrawal loads arising from both swelling and bending, the latter a consequence of radial dpa gradients in the reflector.

The resulting macroscopic strains are sometimes very counterintuitive, however, especially with re­spect to their sign.

For example, formation of the less dense carbide phases leads to macroscopic densification of the alloy and shrinkage of volume,109 while the formation of denser intermetallic phases (Chi, Sigma, Laves) usu­ally leads to an increase in volume, a form of nonvoid swelling.110,111 This counterintuitive behavior is the result of the different partial molar volumes of criti­cal elements (C and Mo primarily) between the new

precipitates and the alloy matrix in which they form. Both the carbide and intermetallic phase evolution appear to be accelerated and sometimes altered under irradiation.

Other radiation-produced phases (W, G-phase) also appear to induce changes in lattice parameter

but these have not been well characterized, primarily because these phases develop concurrently with void swelling that masks their contribution.1

Garner1 provides a review of precipitation-induced strains. For the current purpose it is sufficient to note that carbide-induced densification increases with carbon content and with increasing irradiation temperature. Such volume changes for the most
common carbon levels range from 0.1% to 0.4% decrease in volume. The resulting strains may or may not be isotropically distributed, depending on whether there is a pronounced starting dislocation texture on which the carbides nucleate. This process is most pro­nounced for titanium carbides in Ti-stabilized steels. Carbide-induced strains usually develop quickly enough to be measurable before swelling strains become dominant and therefore are relatively easy to identify compared to those of slower forming phases.

The formation of intermetallic phases can gener­ate strains in the order of 1-3%. There is insufficient evidence to support anisotropy of resulting strains, but there exists some evidence that tensile stress states may accelerate the formation of these phases.110

Additionally, there is a decrease in density and a concurrent increase in volume when ferrite is formed from austenite as a result of radiation-induced segre­gation of nickel. Formation of ferrite from austenite can lead to volume increases as large as 3%, but there are no available data on potential anisotropy or stress dependence. As opposed to carbide-induced strains that develop relatively quickly, ferrite and interme­tallic strains develop rather slowly, and therefore are usually unrecognized, especially when other strain contributions arising from swelling and creep are present.

Such precipitation-induced strains are important in that while they usually saturate in magnitude, they can be a significant portion of the total net strain at low dpa levels, thereby complicating the analysis and extrapolation of void swelling and irradiation creep data. Such strains can also affect the stress distribu­tion and level in a structural component. For instance, a preloaded tie-rod or bolt will initially increase in load as a result ofcarbide-induced shrink­age even while irradiation creep proceeds to relax the load.

It should be noted that radiation-induced segre­gation can lead to overall changes in average lattice parameter without actually culminating in observable precipitation. Although there is no convincing evi­dence that segregation to void and grain boundaries produces measurable strains, it has been shown that radiation-induced spinodal-like decomposition in Fe-35Ni and Fe-Cr-35Ni alloys produces periodic oscillations in composition that are accompanied by densification in the order of ^1%.112,113 Oscillations in nickel level are almost exactly offset by out-of­phase oscillations in chromium. This demonstrates that in a single phase system the lattice parameter of a given element is not constant but is influenced by its local concentration and its association with other elements.