Values between 2.4 and 5.1 MPa Vm have been measured for CVD p-SiC, depending on the test technique employed and grain size. Fracture tough­ness of CVD SiC increases slightly at elevated tem­peratures. It does not exceed 6 MPaVm.

2.12.2.1.4 Fracture strength

As is usual with brittle ceramics, fracture data exhibit a significant scatter, as flaws that have a random distribution induce fracture. An important conse­quence is that the fracture stress is not an intrinsic characteristic. It is, instead, a statistical variable, which depends on several factors including the test method, the size of test specimens, and the number of test specimens.24 Therefore, a universal reference value of fracture strength cannot be recommended.

It is widely accepted that the Weibull model satis­factorily describes the statistical distribution of fail­ure strengths: (s/s0)mdV /V„ where P is the probability of failure, s is the stress, s0 is the scale factor, m is the Weibull modulus, V is the volume of specimen, and V0 is a reference volume (1 m3 is generally used); m reflects the scatter in data, and s0 is related to the mean value of the strength.

The strength data for a given geometry and stress state can be determined using eqn [4]. However, m, s0, and V0 must be available. It is important to note that the estimate of s0 depends on V0.24 It will be substantially different if V0 = 1 m3 or 1 mm3. This dependence is ignored in most publications, even in the work by Snead and coworkers23 in which a number of s0 values are reported. When V0 is not given, the estimate of s0 is meaningless. The strength cannot be determined safely. Unfortunately, reliable s0 values (characteristic strength in a few papers) cannot be recommended here until the authors have completed their papers. The values of Weibull modulus of CVD SiC at room temperature reported in Snead etal23 span a large range, from 2 to 12. The following values were measured using tensile tests on CVI SiC/SiC mini­composites: m = 6.1, s0 = 10.5 MPa (V0 = 1 m3).25,26

2.12.2.1.5 Thermal creep23

Primary and secondary creep deformations have been reported in the literature for CVD SiC (high-purity and polycrystalline p-SiC). Creep in SiC is highly depen­dent on the crystallographic orientation. The loading orientation of 45° from the CVD growth axis is the direction in which the most prominent creep strain is observed. A review of creep behaviors of stoichiometric CVD SiC has been provided by Davis and Carter.27

Primary creep of CVD SiC occurs immediately upon loading and tends to saturate with time. The primary creep strain generally obeys the following relationship:

ec = Ap (s/G)n(t/t)p [5]

where Ap, p, and t are creep parameters, and t is the time elapsed. n = 1.63, Ap = 29, p = 0.081, and t = 0.0095 s for the temperature of 1923 K. These para­meters are for the loading orientation of 45° from the CVD growth axis. In severe conditions, primary creep strain in the CVD SiC can reach as high as 1%.

Steady-state creep rates for polycrystalline mate­rials have been measured only above 1673 K, when the stress axis is 45° inclined from the deposition direction; temperatures as high as 2023 K are required when the stress axis is parallel to the depo­sition direction. The strain rate is given by a power — law creep equation:

de/dt = As(s/G)nexp(—Q / k^T) [6]

where As = 2.0 x 103, n = 2.3, Q= 174kJmol—1 (acti­vation energy), s is the applied stress, G is the shear modulus, and kb is the Boltzmann constant.

The saturated vapor pressure is a very important parameter of LM coolants for safety estimations. It is directly related to the latent heat of evaporation. The boiling temperature increases (with the decreas­ing rate) when pressure increases. At low pressures, where the vapor behaves as a perfect gas and the evaporation enthalpy (AHe) is practically constant, it follows from the Clausius-Clapeyron equation that temperature dependence of the saturated vapor pres­sure is about exponential:

ps(T) =ps, i exp(-AHe/RT) [2]

R = 8.314J mol-1 K-1 is the universal gas constant.

Equation [2] can provide approximate values for equilibrium-saturated vapor pressures over a wide range of temperature due to the relatively small variation of AHe with temperature at low pressures. At high pressures, closer fits of the experimental results can be obtained by adding supplementary temperature-dependent terms:

ln(ps(T))= A + T + C ln(T) + DT [3]

Very often this correlation is used without the last term.

Evaluation of data on the temperature depen­dence of the saturation vapor pressure of sodium in a wide range of temperatures was performed many

times3,8,43,44 and correlations of type [3] were pro­posed that can describe the temperature dependence of the Na saturation pressure in a wide temperature range with an uncertainty of 1-25%. In the recent compilation of IAEA,26 a polynomial of the fifth order was proposed for the saturation pressure of Na constructed on the bases of the recommended data presented in Bystrov et al, Browning and Potter,44 and Vargaftik and Voljak.45 In the review of Fink and Leibowitz,22 a detailed analysis of previous compilations was performed and a correlation of type [3] proposed in Browning and Potter44 was selected, which allows to calculate the Na saturated vapor pressure with an uncertainty of <3% in the temper­ature range from 864 K to the region close to T:(Na). The uncertainty increases at lower temperatures and reaches 24% at T = 400 K. This correlation is recommended for high temperatures in the current work and its coefficients are given in Table 4.

Not many data exist on the saturation vapor pres­sure of lead. In 1973, they were reviewed by Hultgren eta/.,5 who analyzed different sources and presented a table with the recommended data in the temperature range of 298.15-2100 K. Later, Kubaschewski and Alcock46 reanalyzed the available data on the Pb satu­ration vapor pressure and proposed a correlation of type [3]. In more recent compilations23,27,33,47,48 and in the present work, the simpler correlation of type [2] or the correlation of type [3] with the coefficients from

Kubaschewski and Alcock46 are recommended for the saturated vapor pressure of lead (see Table 4).

Data on the saturated vapor pressure of Pb-Bi(e) are very limited in number49-51 and in the tempera­ture range (508-1023 K), and a larger dispersion exists between them, especially at lower tempera­tures. Therefore, they can be described with correla­tion [2]. Nevertheless, in order to take into account the formation of Bi2 molecules in the Pb-Bi vapor phase, Morita et al5 proposed to use a four-term correlation [3] for temperature interval 700-2000 K. The coefficients are included in Table 4; the coeffi­cient D was corrected to obtain TB = 1927 K at ps = 0.101 325 MPa (1 atm).

At low temperatures, the saturated vapor pressure of the considered LM can be estimated with the same or better uncertainty with eqn [2]; the coefficients recommended in the compilations33,34 for Na, Pb,

and Pb-Bi(e) are given in Table 5.

Figure 3, in which the saturated vapor pressure is plotted as a function of temperature, illustrates the

correlation [2] for Na, Pb-Bi(e), and Pb in the tem­perature range of TM0 to TB0.

It should be noticed that for Na and Pb, the uncer­tainty of this correlation is about 7-10% when the temperature is 50-100 °C higher than normal melt­ing temperature and lower than normal boiling tem­perature. The uncertainty increases rapidly beyond this interval, at lower and higher temperatures. For Pb-Bi(e), the uncertainty of this correlation is about 10% at temperatures higher than 1500 K; however, it becomes more than 50% when temperature decreases below 1000 K.24,34 At lower temperatures, the saturated pressure is too low to be measured correctly. The values of AHe presented in Table 5 are slightly higher but in satisfactory agreement with AHB 0 presented in Table 2.

The PWR fuel assembly consists of fuel rods, grids, the top nozzle, the bottom nozzle, the instru­mentation tube, and guide tubes. First, the skeleton assembly is made, which is an assembly of the instru­mentation tube and the grids. Then, the fuel rods and the guide tubes are inserted into the skeleton assembly. Finally, the top nozzle and the bottom nozzle are mounted on the guide tubes by screws.

The BWR fuel assembly consists of fuel rods, water rods, grid spacers, the upper tie plate and the lower tie plate. First, the water rods, grid spacers, and the lower tie plate are assembled. Then, fuel rods are inserted into grid spacers and tie rods are connected to the lower tie plate. Finally, the upper tie plate is mounted and connected to the tie rods with screws.

The dimensions and appearance of the fuel assem­blies are inspected and the BWR fuel assembly is attached to the channel box before loading it into a reactor.

Austenitic stainless steels are also a key component for MFR systems because of many of the properties and vast experience in fission nuclear systems described above. An important example of rapid alloy development is presented, which is part of the US contribution to the international fusion demon­stration project in France, called ITER, and in­cludes ^20% of the first wall (FW) and shield components. The ITER project could include nearly

100 modules from austenitic stainless steel (316LN — ITER Grade or — IG) each weighing ^3.5 T and 366 FW panels (SS/CuCrZr/Be). An example of the shield wall module is shown in Figure 18. Traditional machining of the cooling channels shown in Figure 18 results in a loss of ^30% of the raw mate­rial during fabrication. A US industry manufacturing assessment indicates that casting the shield modules (including the cooling channels) results in major cost savings when compared to fabrication via welding together quarter modules machined from large for­gings. However, because casting produces a large grain size, low dislocation density, and extensive seg­regation of alloying elements, the strength properties of such cast components are frequently inferior to those of conventionally forged and annealed compo­nents. Additional R&D has been performed27 in recent years to ensure that the properties of cast 316L(N)-IG equivalent grades meet ITER Structural Design Criteria,28-30 which require cast steel performance that is similar to or no worse than wrought equivalent material.

On the basis of past development experience, archive material analysis, and simulations, several improvement strategies were identified as part of this effort to modify and upgrade the properties of the standard CF3MN cast stainless steel grade (which is described in more detail in Busby et a/.27) (Table 3). The primary strategy identified for boost­ing the YS was increased strengthening by additions of N and Mn; N is the most powerful solid solution strengthener (0.1 wt% should increase strength by 50 MPa). However, Mn increases are also required to raise the solubility limit of N. In addition, Mn is also an austenite stabilizer, and increases both strength and strain-hardening rate. Industrial part­ners were involved in the fabrication of test alloys

to help speed scale-up to larger test articles. Alloys with the most minor alloying additions were studied most extensively, with one alloy showing the greatest performance, which is designated CF3MN-US (Table 3). Mechanical testing (tensile, impact, and fracture toughness) was performed along with examinations of physical properties, porosity, weldability, and resistance to stress-corrosion cracking.

To accelerate the transition to heavy-section cast­ings, tensile tests were conducted on both cast keel blocks and specimens cut from the larger cross­section as-cast ingots (from both the surface and center regions). These different specimen locations help illustrate the potential differences in mechanical performance. The results from room temperature test­ing demonstrated no systematic difference between types of specimens, locations of testing, or locations or types of specimen used.

Elevated temperature tests were also performed. The yield stress results for samples in the as-cast condition are illustrated in Figure 19 and are com­pared to the minimum requirements for use in ITER applications. At all temperatures, the CF3MN-US exceeds the minimum required strength and meets the ITER acceptance criteria.

An evaluation of impact properties on the CF3MN-US was also conducted. Initial testing was performed using a drop-weight machine setup with a maximum capacity of 325J potential energy for ini­tial screening tests. Two tests of CM3F-US with the drop weight machine set at 325J were performed.

Only one specimen at —196 °C (liquid nitrogen tem­perature) broke. To demonstrate the excellent tough­ness of the materials in the temperature range of interest for the ITER shield module applications, additional testing was performed at higher tempera­tures. Tests were performed at room temperature (27 °C), 100, 200, and 300 °C, again using a drop — weight machine. All tests at all temperatures were fully ductile and very tough, and none of the speci­mens tested fractured. Figure 20 shows a photograph of a Charpy specimen tested in the drop-weight machine at 300 °C. It is clear that the specimen did not fracture when tested with a maximum potential energy of 270J, so its actual impact toughness is higher than that. As indicated earlier in this section, wrought 316 and 347 steel typically have Charpy impact toughness of 100—150J, so the CF3M-US cast stainless steels exhibit excellent impact tough­ness, even at the liquid nitrogen temperature (—196 °C). The stated minimum impact toughness for the ITER shield module materials is 60 J, whereas the tested specimens exhibited impact toughness values ranging from 140 to 262 J at —196 °C.

Finally, testing of fracture toughness properties on the CF3MN-US was performed. The 12.5-mm thick compact tension (0.5T C(T)) specimens were tested at room temperature and at 90 and 190 °C. At least two specimens of each alloy were tested at each temperature in general accordance with the ASTM E 1820-06.31 As expected from the simpler Charpy
impact data reported previously, all alloys exhibited very high fracture toughness at all test temperatures. Moreover, none of the specimens exhibited crack extension regardless of test temperature. The value of critical J-integral, J1C, was above 800 kJ m—2 at all tested temperatures, comparable or better than that of the equivalent wrought austenitic steel, meeting the ITER acceptance requirements.

While not shown in detail here, testing and evalu­ation of the most promising alloy under development, CF3MN-US, has been completed for several proper­ties. Composition, ferrite content, microstructure, porosity, mechanical properties (tensile, impact, and fracture toughness), irradiation performance, stress — corrosion cracking performance, and weldability have all been found to meet ITER acceptance cri­teria. This combination of past experiences, expertise, and new tools demonstrates new opportunities for rapidly developing improved austenitic steels for advanced reactor applications such as ITER. It is also reasonable to expect that the new CF3MN-US steel may have attractive properties in either the cast or wrought condition for advanced LWR core or structural support designs and applications.

During cyclic fatigue at room temperature, matrix damage is determined by the maximum stress. It is created during the first cycles. Fatigue resistance is governed by the damage of fibers and the fiber — matrix bonds. Two different fatigue behaviors have been observed: after 1000 cycles, either the elastic modulus remains constant and the specimen is run­ning out, or it decreases until the specimen ultimately fails.48 The modulus degradation reflects either wear at cracked interfaces48 or growth of interface cracks. Under stresses smaller than 100 MPa, ultimate fail­ures are generally not observed after 106 cycles under tension-tension fatigue.

At high temperatures, additional phenomena acti­vated by environment (oxidation, creep, or slow crack growth) may operate and cause the extension of initial stress-induced matrix damage and interface cracks as well as the weakening of fibers by degradation or reloading. Reloading of fibers involves changes in load sharing. If the strength of tows is exceeded by the applied stresses according to the mechanism described previously, ultimate failure occurs. The rup­ture of tows dictates the lifetime.

The matrix cracks created upon loading become the pathways for the ingress of oxygen into the mate­rial. The PyC interphase is consumed, which causes fiber reloading. Creep of the SiC matrix (at very high temperatures, above 1200 °C) makes the stresses on the fiber to increase, which enhances matrix creep and further fiber reloading. Creep of fibers (at tem­peratures above 1200 °C) causes matrix reloading and possible matrix and interface cracking or crack prop­agation, leading to fiber reloading by decrease ofload carried by the matrix.

Finally, the slow crack growth in fibers (at tem­peratures below 1100 °C) is activated by oxidation of carbon grain boundaries, leading to delayed failure.49,50 This phenomenon, which was observed at

intermediate temperatures between 500 and 900 ° C, was first referred to as the ‘pest phenomenon’ by a few authors. The SiC/SiC were claimed to be susceptible to degradation by oxidation embrittlement.

In order to protect the PyC interphase against oxidation, multiple coating concepts have been explored and multilayered interphases and matrices have been developed.21 Such multilayered matrices contain phases that produce sealants at high temperatures, causing healing of the cracks and preventing oxygen from reaching the cracks and the interphases.2,22,51 Lifetime is also improved with oxidation-resistant interphases such as BN or multilayers.52

2.12.6.3 Thermal Shock

CVI SiC/SiC composites have been tested under thermal shock with excellent results.2,53 CVI SiC/ SiC generally had good strength retention after ther­mal shock cycles involving heating up to the desired temperature and then cooling down in water at 20 °C.

2.15.2.1.2.1 Flowability

In pellet fabrication, powder flowability is one of the most important characteristics that determine the productivity of the fabrication process. It is well known that blended powders have very poor powder flowability, just after milling.5 Therefore, the milled powder is granulated or mixed with a powder having good flowability to ensure uniform die filling and good compaction behavior.5-7 Carr indices are a well-known method to evaluate powder flowability of dry solids.8,9 The powder flowabilities of micro­wave heating denitrated MOX (MH-MOX) powder and ammonium diuranate (ADU) powder have been evaluated on the basis of Carr indices both before and after granulation.1 ,

2.15.2.1.2.2 Effective thermal conductivity

The temperature of MOX powder increases by self heat generation of plutonium by a-decay when the powder is kept in the fuel fabrication process. In a MOX fuel fabrication plant, the temperature increase in MOX powder should be prevented because the excessive temperature increase of MOX powder may possibly cause changes in powder characteristics (e. g., O/M ratio variation), degradation of additives (e. g., lubricant agents), and overheating of equipment in the fabri­cation process. An example of a preventive measure against the temperature increase of MOX powder is the use of a storage vessel that has radiator plates.

The effective thermal conductivity ofMOX powder is important for estimating its temperature distribution. The effective thermal conductivity of a powder can be defined as the combination of thermal conductivities of powder particles and the atmospheric gas because the volume fraction of the atmosphere gas in the total volume is large. In addition, particle shapes, mean particle size, specific surface area, and O/M ratio of powder particles influence the effective thermal con­ductivity of the powder.12 Figure 1 shows the effective thermal conductivities of various MOX

under a controlled atmosphere to improve their mechanical strength. The powder compact is com­posed of individual grains separated by 35—50 vol.% porosity. During sintering, the following major changes commonly occur: an increase in grain size, and changes in pore shape, pore size, and pore num­ber. In the early stages of sintering, the powder par­ticles begin to mutually bond. In the middle stage, grain growth, disappearance of pores, and formation of closed pores occur. The pellet densification pro­ceeds according to the shape change from a point contact to a face contact between grains. In the last stage, disappearance of the closed pores occurs. The diffusion of uranium, plutonium, and oxygen, the evaporation-condensation process of their com­pounds, the grain growth process, the pore migration process, and the pore disappearance processes are important for understanding the process of sintering. To obtain pellets with high mechanical strength and density, it is desirable to eliminate as much porosity as possible.

Diffusion coefficients ofthese elements are needed for evaluating the sintering behavior (e. g., volume shrinkage in the fuel fabrication technology). Section 9.1.6.1 shows that the oxygen self diffusion coeffi­cients of actinide oxides increase with increasing deviation from stoichiometry near the stoichiometric region and that the diffusion coefficients of cations in hyperstoichiometric actinide oxides increase dras­tically with deviation from stoichiometry. It was shown that the diffusion coefficient of plutonium in (U0.8 Pu0.2)O2±x has the lowest value near the stoi­chiometric region and it increases significantly with an increase in deviation from stoichiometry (see Figure 2).

Vapor species of oxide fuel and its vapor pressure are required to assess the redistribution of elements, pore migration, and fuel restructuring. The O/M ratio dependencies of vapor pressures in the vapor species of uranium oxide, plutonium oxide, and MOX are shown in Figures 26 and 27 of Section 9.1.5. The vapor pressures of each of these species have a large dependency on the O/M ratio and their behavior is different in each vapor species.

Temperatures used during dewaxing and sintering are very important factors in the fabrication process. The Huttig and Tamman temperatures, which are defined as the start temperatures for surface diffusion and volume diffusion of powder particles, respec­tively, are provided for establishing temperatures for dewaxing and sintering. These temperatures can be easily calculated using melting point temperature.

Utility requirements are tending to become ever more demanding as they attempt to extend cycle lengths further and further and at the same time demand increased fuel discharge burnup. If this trend continues, it is likely that burnable poisons will be used even more extensively and with ever-increasing sophistication (which usually means more heterogeneous distribution of the poison material radially and axially within the fuel assemblies). The use of enriched 10B, already available commercially, is likely to increase and there may also arise a demand for isotopically enriched gadolinia and erbia. These developments are likely to increase the pressure on fuel manufacturers to develop automated fabrication lines that can more easily deal with more complicated burnable poison loadings, and on the developers of nuclear design codes to more easily accommodate very heterogeneous burnable poison loadings through a higher level of automation. Since there are only around 90 naturally occurring elements in the Periodic Table, the number of candidate burnable poison materials is virtually exhausted and it is unlikely that any new ones will emerge (see Chapter 2.09, Properties of Austenitic Steels for Nuclear Reactor Applications, Chapter 2.15, Uranium Oxide and MOX Production and Chapter 2.19, Fuel Performance of Light Water Reactors (Uranium Oxide and MOX)).

2.11.2.1 Introduction

Apart from its use as a neutron reflector and modera­tor in nuclear reactors, beryllium is in strong demand for use in X-ray windows of medical and industrial equipment, acoustic speaker diaphragms, galvano mir­rors for laser drilling, reflected electron guard plates in semiconductor production equipment, and various other applications. It is also widely used in the electri­cal and electronic industry, particularly in beryllium — copper alloys for wrought metal production and for molds and other forging tools and dies. In electronics, in particular, the need for beryllium has been growing rapidly in recent years with the trend toward lighter, thinner, and smaller electronic components. In the following sections, we outline the methods of its pro­duction and processing and discuss its basic properties.

2.11.2.2 Production and Processing Methods1

Among the 30 or so naturally occurring ores, the most economically important is beryl, which contains 10-14% beryllium oxide (BeO). At present, the two main industrial processes used to extract BeO from beryl are the fluoride method and the sulfuric acid method. Both of these yield BeO of industrial-grade purity, which is used as a raw material for Be-Cu mother alloys, electronics manufacture, refractories, and other fields of application. For use in nuclear reactors, BeO is further purified by recrystallization or precipitation.

Metallic beryllium (Be) is produced from BeO or Be(OH)2 by either of two industrial processes. One involves the formation of BeF2 followed by its ther­mal reduction with Mg to produce Be pebbles, and the other involves the formation of BeCl2 followed by its electrolysis to produce Be flakes. The resulting pebbles and flakes are high in Mg and Cl2 content, respectively, and these impurities are removed by vacuum melting.

The principal techniques of Be processing are molding by powder metallurgy, warm or hot working, and joining or welding. In hot-press sintering, which has been widely developed for Be molding, the start­ing material is commonly —200 mesh Be powder, which is inserted between graphite dies and then pressure molded in vacuum at high temperatures (1323 K). The resulting moldings are commonly called ‘hot-press blocks,’ and can be obtained with high integrity and near theoretical density. Other molding methods that may be employed include spark plasma sintering and cold-press sintering.

Cold working of Be at room temperature is extremely difficult because of its low elongation, and it is accordingly formed into plates, rods, or tubes by ‘warm working’ at 773-1173 Kor ‘hot work­ing’ at 1273-1373 K. In either case, the Be must be covered with mild steel or some other material and the intervening air withdrawn before it is heated, as it readily oxidizes at high temperatures.

Various methods have been developed for Be join­ing and welding. These include mechanical joining and resin bonding, electron-beam and diffusion weld­ing, and brazing and soldering. Because of its high oxygen affinity, however, any process in which the Be is heated must be performed under an appropriate inert gas or vacuum.

Though strength and fracture properties are con­trolled by sample processing, the elastic and plastic deformation behavior inherent to ZrC may be under­stood in terms of its chemical bonding. Transition metal carbides are known to be brittle at ambient temperatures but ductile at high temperatures.

The metal-like conductivity and the fcc structure of ZrC suggest that deformation along the fcc metal slip systems is possible. In fcc, the {111} (110) system corresponds to the slip of close-packed planes along close-packed directions and requires the lowest stress to form and move a dislocation. The same is assumed for the rocksalt structure.

However, many crystals with the NaCl structure are ionic, and slip along the above system is inhibited due to the energy required to overcome strong Cou — lombic repulsion when in the half-glided position.3 Instead, ionic rocksalt compounds prefer to slip along {110} planes, maintaining attractive Coulombic forces. If ZrC is known to slip along {111} planes, the degree of ionic bonding must not be large.

Even if there is no ionic prohibition to slip in ZrC, the directed nature of its covalent bonding is still an impediment to slip. Strong metal-carbon bonds in an octahedral coordination inhibit slip on close-packed {111} planes, leading to high shear stresses required for dislocation mobility at low temperatures. The preferred mechanism of deformation is then brittle fracture, which occurs by cleavage on {100} planes.12,114 Brittle fracture persists at least to 1000 K.3

At elevated temperatures, however, a ductile-brittle transition has been observed. Based on compressive loading of single-crystal ZrC09,115 ZrC0 945116, and arc- melted ZrC094,117 plastic yield was reported at 1172, 1353, and 1473 K. Microplasticity was reported at 1273 K, with gross plastic yield above 1773 K, as observed by fractography of tensile specimens.118 Trans­mission electron microscopy (TEM) of dislocations in ZrC0.98 after elevated temperature compression re­vealed microplasticity above 1420 K.119

The crystallography of the slip system has been investigated. Lee and Haggerty116 measured the crit­ical resolved shear stress (CRSS) in the compression of single crystal ZrC0945 samples grown along (100) and (111) directions, and one sample whose axis corresponded to the ‘0.5’ orientation for {111} (110) slip, indicating that one of these 12 equivalent slip systems was oriented at 45° to the crystal axis. When loaded uniaxially along the crystal axis, a maximum resolved shear stress (t) of half of the applied stress (or 0.5s) would be achieved, according to Schmid’s law,

t = a cos f cos l [11]

where f and l are, respectively, the angles separating the slip plane normal and the slip direction from the load axis. CRSS as a function of temperature for the various samples are plotted in Figure 27. The results of Williams115 for compression of (100) oriented single crystal ZrC0.875 are consistent with those of Lee and Haggerty, plotted in the same figure.

The slip planes were identified by slip traces on the samples, while the Burgers vector was confirmed to be 1/2(110) by diffraction contrast of dislocations, a result consistent with the TEM analysis of Britun
eta/.119 Slip was induced on either {100}, {110}, or {111} planes, depending on the crystal axis and its slip system of maximum resolved shear stress. For loading along the (100) direction, the maximum resolved shear stress was for the {110} plane; for (111) loading, slip was along {100} planes; and for the ‘0.5’ oriented sample, slip was along {111} planes. CRSS for slip along {100} was highest, and along {110} and {111} was approximately equal over the temperature range where they overlapped.

Hannink et a/’22 characterized the anisotropy of room-temperature Knoop hardness on the {100} surface of single crystal ZrC0.94, rotating the long axis of the Knoop indenter azimuthally; hardness as a function of rotation angle is plotted in Figure 28. Hardness varied sinusoidally with rotation, with a minimum occurring when the indenter was aligned along (100) directions, and maximum for indenter alignment with (110) directions. This indicated that the slip system was {110} (110), normally asso­ciated with ionic crystals. The authors also measured Vickers hardness, which exhibited anisotropy in the same sense as the Knoop measurements, but with lower amplitudes, in agreement with the results of Kumashiro eta/}2

Hannink et at}22 suggested that the active slip system is dependent on temperature, as they observed for TiC096 and VC0 83. At ‘low’ tempera­tures (room temperature for TiC and ZrC, 87 K for VC), the {110} (110) slip system was active, as seen by the hardness anisotropy described earlier. At higher temperatures (883 K for TiC, 623 K for VC), maxi­mum and minimum hardness occurred respectively for indenter alignment with (100) and (110), which is the opposite to that observed for {110} (110) slip. ‘High’ temperature slip in TiC and VC was on the {111} (110) or {100} (110) system. Kohlstedt101 pro­posed that covalent, directional bonding dominates at low temperatures, prohibiting slip on {111} planes and resulting in high hardness, while at high tem­peratures, the degree of covalent bonding is decreased as the {111} (110) slip characteristic of fcc metals is favored, resulting in the observed hardness drop with temperature.

Capacity for plastic deformation has also been observed to vary with the C/Zr ratio. A monotonic decrease in hardness with decreasing C/Zr ratio is seen in Figure 23. Although not determined for ZrC, CRSS of TiCx was observed to decrease with decreasing C/Ti ratio.115 The author explains this intuitively in terms of fewer C—Ti bonds that must be broken during dislocation motion. Hollox12 attrib­uted these results to a decrease in the contribution made by carbon atoms to cohesion in TiC as the C/Ti ratio is reduced, further citing the band structure calculations of Lye and Logothetis11 which
indicated that carbon donates electrons to and strengthens metal-metal bonds. Hollox also inferred that the DBTT might decrease as the carbon-to — metal ratio is decreased, but this has not been demon­strated conclusively.

Plutonium is extracted from spent fuels in the repro­cessing plants in the form of plutonium nitrate. In order to utilize extracted plutonium for MOX fuel production, plutonium nitrate is converted to oxide powder by three methods: one is an oxalate pre­cipitation method; the other two methods involve coconversion with uranium, the ammonium uranyl plutonyl carbonate (AUPuC) conversion method, and the microwave heating denitration method (MH method). The AUPuC conversion method is described in Section 39.5.2.3 as part of the AUPuC fuel fabrication process.

2.15.5.1.1 Oxalate precipitation method

In the oxalate precipitation method, the plutonium oxide powder is prepared from plutonium nitrate by the following two reactions.50 Pu(NO3)4 + 2(COOH)2 ! Pu(COO)4 + 4HNO3

Pu(COO)4 ! PuO2 + 2CO2 + 2CO

Oxalate acid, H2(COOH)2, is added to plutonium nitrate solution at about 60 ° C, and the temperature maintained until the precipitation reaction (1) is completed. The plutonium oxalate precipitate is filtered and then dried in air. Dried plutonium oxalate is calcined in a furnace at temperatures from 350 to 650 °C. It has been reported that reac­tion (2) begins below 100 °C and is completed at around 350 °C.50 The characteristics of the obtained PuO2 powder vary depending upon the precipita­tion and calcination conditions, that is, the precipi­tation temperature, addition rate of oxalate acid to plutonium nitrate, oxalate acid concentration, and calcination temperature. This PuO2 powder is commonly utilized as a feed material for MOX fuel production in the world. The microstructure and characteristics of PuO2 powder prepared by the oxalate precipitation method have also been explained elsewhere.51