To increase the proliferation resistance of plutonium, a coconversion method of adding plutonium nitrate and uranyl nitrate to a mixed oxide powder was devel­oped inJapan. In the MH method, about 71 of a mixed solution of uranyl nitrate and plutonium nitrate with a concentration of about 250gl~1 of heavy metal, is fed into a denitration vessel. The diameter and height of this silicon nitride vessel are about 50 and 6 cm, respectively. After microwave irradiation (2450 MHz, 16 kW), PuO2 + UO3 is formed, and then this product is calcined to PuO2 + U4O9 + U3Ogx in air for 2h at 750 °C. Subsequently, this mixture is reduced to PuO2 + UO2 (MH-MOX) powder under an atmo­sphere of N2-5% H2 mixed gas, at the same tempera­ture used for calcination.52 The obtained MH-MOX powder has sufficiently good powder characteristics to allow fabrication of MOX pellets of more than 95% TD.52,53 Full details of the MH method have been given elsewhere.53-56 With the MH method, the generation of radioactive liquid waste containing plutonium is reduced compared with other conver­sion processes.

Figure 12 shows microstructures which were observed by scanning electron microscopy (SEM) at 10 000-fold magnification, in the PuO2 powder (A) prepared by the oxalate precipitation method and MH-MOX powder (B). The microstructures of MH-MOX powder and UO2 powder (prepared by the ADU process) calcined at various temperatures have been reported in Asakura et a/.52

Examples of the characteristics of PuO2 and MH-MOX powders are shown in Table 4.

The values vary depending on the conversion condi­tions described above.

Coriou reported IGSCC (PWSCC) susceptibility for nickel-based alloys and the influence of nickel con­tent in high-temperature, high-purity water.18 It is now known, however, that this cracking susceptibility

is very dependent on the corrosion potential, as determined by the concentration of molecular hydro­gen in solution (as in PWR primary water). Subse­quently, it was reported that IGSCC is affected by the chromium content, but not by the nickel content in nickel-based alloys, as shown in Figures 23 and 24 60 Alloy 690 has higher resistance to PWSCC than Alloy 600, due to its higher chromium content.

Carbide precipitation along grain boundaries by thermal treatment (TT) at around 700 °C improves the PWSCC resistance for Alloys 600 and 690. In particular, M23C6 precipitation that is coherent with the matrix was detected along grain boundaries in the TT Alloy 690, which has excellent PWSCC resistance depending on the carbon content and the solution heat-treatment temperature. By contrast, niobium addition to Alloy 600 was found to have a poor effect on PWSCC susceptibility,60 but improves IGSCC resistance under BWR water conditions.25

2.08.4.2 In High-Temperature Gases

2.08.4.6.1 Oxidation

Nickel shows superior oxidation resistance to carbon steels and copper alloys due to the formation of a

1000

ЕЕГ

0

500

Ф

3

Q.

3

DC 100

50

70 80 90

Nickel content (%)

Figure 24 Effect of Ni content on the stress corrosion cracking fracture time under constant load test for solution-annealed nickel-based 5% chromium-iron alloys at 360 °C in simulated pressurized water reactor primary water.

nickel oxide film in air or other oxidizing environ­ments. The oxidation resistance of nickel is improved remarkably by the addition of chromium. Repeated oxidation test results for various alloys are shown in Figures 25 and 26.10,11,14 In these tests, Alloy 600 showed little weight change and was found to have better oxidation resistance than 304 or 310 stainless steels. Alloy 601 had higher oxidation resistance than Alloy 600, due to its higher chromium and aluminum contents. Alloy 690 also had higher oxidation resis­tance than Alloy 600 due to its higher chromium content.

2.08.4.6.2 Nitriding

Low-alloy steels are highly susceptible to nitriding in active atmospheres such as high-temperature ammo­nia gas. To obtain resistance to nitriding, the addition of nickel is effective. Austenitic stainless steels have higher resistance to nitriding than low-alloy steels, for example. Nickel-based alloys have significantly better resistance to nitriding. Alloy 600 shows excel­lent resistance to nitriding in ammonia production plant environments.

2.08.4.6.3 Sulfidation

Nickel-based alloys are highly susceptible to sulfi­dation. Nickel forms a eutectic with sulfur at

C — heavy oil

(200 ppmV, 50 ppmNa, 2.5% S)
Temperature : 816 °C

a : 4500 h b : 9429h

c : 6450 h d : 1200h [11]

temperatures above 645 °C and the scales on nickel lose their protective properties at higher tempera­tures. The addition of chromium to nickel-based alloys is effective for improving sulfidation resistance, and alloys containing higher than 20% chromium show good sulfidation resistance.

Nickel undergoes severe corrosion in combustion gases ofcrude petroleum. When vanadium is present in these gases, corrosion occurs due to the formation of low-temperature-melting compounds with vana­dium oxides (so-called vanadium attack). When sul­fur is present in crude petroleum, sulfide corrosion occurs.

50% chromium-50% nickel and 60% chromium — 40% nickel alloys are rare nickel-based materials with excellent resistance to vanadium attack and sulfide corrosion. Figure 2761 shows exposure test results for the supports of the super heater tubes of a fossil-fuel electric power plant. The data indicate that both alloys are indeed highly resistant to corrosion.

2.08.3 Summary

The excellent corrosion resistance and mechanical characteristics of various nickel-based alloys have been described in this chapter. In particular, typical corrosion, mechanical, other physical properties data, together with general fabrication information, have been reviewed.

Copper-based alloys have more than 5000 years of history, and iron-based alloys more than 4000 years. However, nickel-based alloys were developed only in the last 100 years or so. This very short history for nickel-based alloys means that some unknown, or uncertain, or unexpected scientific properties will be still remaining to be discovered for these alloys. Consequently, continuing and assiduous studies of nickel based alloys are plainly required (see Chapter

4.4, Radiation Effects in Nickel-Based Alloys and Chapter 5.04, Corrosion and Stress Corrosion Cracking of Ni-Base Alloys).

It is worth addressing the processing method first because this information is useful for a better under­standing of the structure of SiC/SiC. The manufac­ture of long fiber-reinforced composites requires three main steps14’15’28’29:

1. preparation of fibrous preform,

2. fiber coating’ which provides an interface material (interphase), and

3. infiltration of the matrix.

2.12.3.1 Fibrous Preform

The preforms of SiC/SiC composites are made of refractory SiC-based continuous fibers. The latest near-stoichiometric SiC fibers (such as Hi-Nicalon type S and Tyranno-SA3 fibers) are the most appro­priate for those CVI SiC/SiC foreseen for nuclear applications. These fibers exhibit high strength, high stiffness, low density, and high thermal and chemical stability to withstand long exposures at high tempera — tures.30 Finally, the fiber diameter must be small (<20 pm) so that the fibers can be woven easily.

The fiber preforms may consist of

1. A simple stack of unidirectional fiber layers or fabrics (1D or 2D preforms).

2. A multidirectional fiber architecture (3D preforms). Weaving in four or five directions can also be used.

The 2D layers are stacked and kept together using a tool or using fibers in the orthogonal direction (3D preforms).

The main thermodynamic properties of the consid­ered LM (such as density, thermal expansion, heat capacity, enthalpy, surface tension, sound velocity,

and compressibility) are measured with satisfactory precision in the region close to normal melting tem­perature. (An exception is the saturated vapor pres­sure, which is well measured for Na, but not so well for the other two LM: Pb and especially for Pb-Bi(e) at low temperatures — see Section 2.14.3.2.)

Oxide fuels are one of the most popular selections for fast reactor fuel systems, metallic fuels being the other.78 The basis of this popularity can be largely attributed to the great successes achieved in fabrica­tion and operation of LWR oxide fuels.42 Nowadays, LWR operators are seeking ever higher burn-ups of their fuel to attain an economical advantage for LWRs compared to other power plants burning coal
and natural gas. However, the current fuel design has reached its limit at an estimated burn-up of ^80 GWd tU_ . 9 In addition, LWRs produce outlet cool­ant water at a maximum temperature of 320 °C; this limits the efficiency of converting heat to elec­tricity to ^33% and precludes its use as process heat for H2 production.29 The above disadvantages in LWRs based on UO2 fuel may possibly be over­come by the very high temperature reactor (VHTR). The VHTR is fueled by tiny fuel particles embedded in graphite and are cooled by helium (see Chapter 3.06, TRISO Fuel Production). Certain R&D pro­jects still remain to introduce the VHTR commer­cially, in place of LWRs.

For next generation fuel systems that need to burn MAs and process the fuel in a manner that never yields pure plutonium, modifications will be required to minimize waste generation, maximize safety, and maintain operation economics.42 At present, oxide fuels have a higher potential for use in next genera­tion reactor systems than other fuels because a wealth of data has been accumulated for oxide fuels such as fuel fabrication, irradiation behavior, and reproces­sing. As time is still needed to switch from LWRs to FBRs, other fuel systems still have a chance to be the next generation fuel systems through development of innovative technologies.

Detailed accounts of the manufacture of polygranular synthetic graphite may be found elsewhere.2,4,7 Figure 2 summarizes the major processing steps in

the manufacture of synthetic graphite. Synthetic graphite consists of two phases: a filler material and a binder phase. The predominant filler materials are petroleum cokes made by the delayed coking process or coal-tar pitch-derived cokes. The structure, shape, and size distribution of the filler particles are major variables in the manufacturing process. Thus, the properties are greatly influenced by coke morphol­ogy. For example, the needle coke used in arc furnace electrode graphite imparts low electrical resistivity and low coefficient of thermal expansion (CTE), resulting in anisotropic graphite with high thermal shock resistance and high electrical conductivity, which is ideally suited for the application. Such needle-coke materials would, however, be wholly unsuited for nuclear graphite applications, where a premium is placed upon isotropic behavior (see Chapter 4.10, Radiation Effects in Graphite). The coke is usually calcined (thermally processed) at 1300 °C prior to being crushed and blended.

The calcined filler, once it has been crushed, milled, and sized, is mixed with the binder (typically a coal-tar pitch) in heated mixers, along with certain additives to improve processing (e. g., extrusion oils). The formulations (i. e., the amounts of specific ingre­dients to make a specified grade) are carefully fol­lowed to ensure that the desired properties are attained in the final products. The warm mix is trans­ferred to the mix cylinder of an extrusion press, and

the mix is extruded to the desired diameter and length. Alternately, the green mix may be molded into the desired form using large steel molds on a vertical press. Vibrational molding and isostatic pressing may also be used to form the green body. The green body is air — or water-cooled and then baked to completely pyrolyze the binder.

Baking is considered the most important step in the manufacture of carbon and graphite. The pitch binder softens upon heating and goes through a liquid phase before irreversibly converting into a solid car­bon. Consequently, the green articles can distort or slump in baking if they are not properly packed in the furnace. If the furnace-heating rate is too rapid, the volatile gases evolved during pyrolysis cannot easily diffuse out of the green body, and it may crack. If a sufficiently high temperature is not achieved, the baked carbon will not attain the desired density and physical properties. Finally, if the baked artifact is cooled too rapidly after baking, thermal gradients may cause the carbon blocks to crack. For all of these reasons, utmost care is taken over the baking process.

Bake furnaces are usually directly heated (electric elements or gas burning) and are of the pit design. The furnaces may be in the form of a ring so that the waste heat from one furnace may be used to preheat the adjacent furnace. The basic operational steps include (1) loading, (2) preheating (on waste gas), (3) gas heating (on fire), (4) cooling (on air), and (5) unloading. Typical cycle times are of the order of hundreds of hours (Figure 3). The green bodies are stacked into the furnace and the interstices filled with pack materials (coke and/or sand). Thermocouples are placed at set locations within the furnace to allow

direct monitoring and control of the furnace temper­ature. More modern furnaces may be of the car — bottom type, in which the green bodies are packed into saggers (steel containers) with ‘pack’ filling the space between the green body and the saggers. The saggers are loaded onto an insulated rail car and rolled into a furnace. The rail car is essentially the bottom of the furnace. Thermocouples are placed within the furnace to allow direct monitoring and control of the baking temperature.

The furnaces are unpacked when the product has cooled to a sufficiently low temperature to prevent damage. Following unloading, the baked carbons are cleaned, inspected, and certain physical properties determined. The carbon products are inspected, usually on a sampling basis, and their dimensions, bulk density, and specific resistivity are determined. Measurement of the specific electrical resistivity is of special significance since the electrical resistivity correlates with the maximum temperature attained during baking. Minimum values of bulk density and maximum values of electrical resistivity are specified for each grade of carbon/graphite that is manufactured.

Certain baked carbon products (those to be fur­ther processed to produce synthetic graphite) will be densified by impregnation with a petroleum pitch, followed by rebaking to pyrolyze the impregnant pitch. Depending upon the desired final density, pro­ducts may be reimpregnated several times. Useful increases in density and strength are obtained with up to six impregnations, but two or three are more common. The final step in the production of graphite is a thermal treatment that involves heating the carbons to temperatures in excess of 2500 °C. Graphitization is achieved in an Acheson furnace in which heating occurs by passing an electric current throughout the baked products and the coke pack that surrounds them. The entire furnace is covered with sand to exclude air during operation. Longitu­dinal graphitization is increasingly used in the indus­try today. In this process, the baked forms are laid end to end and covered with sand to exclude air. The current is carried in the product itself rather than through the furnace coke pack. During the pro­cess of graphitization (2500-3000 °C), in simplistic terms, carbon atoms in the baked material migrate to form the thermodynamically more stable graphite lattice.

Certain graphite require high chemical purity. This is achieved by selecting very pure cokes, utiliz­ing a high graphitization temperature (>2800 °C), or by including a halogen purification stage in the manufacture of the cokes or graphite, either during graphitization or as a postprocessing step. Graphite manufacture is a lengthy process, typically 6-9 months in duration.

Graphite structure is largely dependent upon the manufacturing process. Graphites are classified according to their ‘grain’ size8 from coarse-grained (containing grains in the starting mix that are gener­ally >4 mm) to microfine-grained (containing grains in the starting mix that are generally <2 pm). The forming process will tend to align the grains to impart ‘texture’ to the green body. The extrusion process will align the grains with their long axis parallel to the forming axis, whereas molding and vibrational molding will tend to align the long axis of the particles in the plane perpendicular to the forming axis. Thus, molded graphite has two perpendicular with-grain (WG) orientations and one against-grain (AG) orientation, whereas extruded graphite has one WG orientation (parallel to the billets long axis) and two AG orientations. Isostatically pressed graphite does not exhibit a preferred orientation. Examples of various graphite microstructures are present in Figures 4-10. The graphite grades shown in Figures 4-10 have all either been used in nuclear applications or been candidates for nuclear reactor use.9 Grade AGOT (Figure 4) was used as the moderator in the earliest nuclear reactors in the United States. Pile grade A (PGA) graphite (Figure 5) was used as the moderator in the early air­cooled reactors and Magnox reactors in the United Kingdom.9 Grade NBG-18 is a candidate for the next generation of high-temperature reactors. Grade IG-110 (Figures 7 and 8) is the moderator material

in the high-temperature test reactor in Japan, and grade 2020 graphite (Figures 9 and 10) was a candidate for the core support structure of the modular high — temperature gas-cooled reactor in the United States.

A comparison of Figures 4 and 7 indicates the range of nuclear graphite textures. Figure 4 shows the structure of AGOT graphite, an extruded medium-grained, needle-coke graphite (maximum filler size ~-0.7 5 mm) and Figure 7 shows the struc­ture of IG-110 graphite, an isostatically pressed, fine-grained graphite (maximum filler size ~10 pm). Similarly, grade 2020 (Figures 9 and 10) is also a fine-grained, isostatically pressed graphite. The UK graphite PGA is extruded needle-coke graphite with a relatively coarse texture (Figure 5). The large individual needle-coke filler particles (named needle coke because of their acicular structure) can clearly be distinguished in this graphite. Another dominant feature of graphite texture can easily be distinguished in Figure 5, namely porosity. Graphite single crys­tal density is 2.26 gcm~3, while the bulk density is 1.75 gcm~ . The difference can be attributed to the porosity that is distributed throughout the graph­ite structure.10 About half the total porosity is open to the surface, while the remainder is closed. In the case of PGX graphite, large pores in the structure result in relatively low strength. The formation of pores and cracks in the graphite during manufacture adds to the texture arising from grain orientation and causes anisotropy in the graphite physical properties. Three classes of porosity may be identified in syn­thetic graphite:

• Those formed by incomplete filling of voids in the green body by the impregnant pitch; the voids originally form during mixing and forming.

• Gas entrapment pores formed from binder phase pyrolysis gases during the baking stage of manufacture.

• Thermal cracks formed by the anisotropic shrink­age of the crystals in the filler coke and binder.

Isotropic behavior is a very desirable property in nuclear graphite (see Chapter 4.10, Radiation Effects in Graphite) and is achieved in modern nuclear graphite through the use of cokes11 with an isotropic structure in the initial formulation.

Coke isotropy results in large measure from the optical domain structure of the calcined coke. The optical domain size is a measure of the extended — preferred orientation of the crystallographic basal planes. Essentially, the optical domain size and struc­ture (domains are the isochromatic regions in the coke and binder revealed when the structure is viewed at high magnification on an optical micro­scope under polarized light) controls the isotropy of the filler coke. Anisotropic ‘needle’ cokes have relatively large extended optical domains, whereas ‘isotropic’ cokes exhibit smaller, randomly orientated domains. The domain structure of a coke is devel­oped during delayed coking through pitch pyrolysis chemistry (mesophase formation) and coking trans­port phenomena. At the atomic scale, orientation of the crystallographic structure is characterized using X-ray diffraction analysis. The crystal spacing within the graphitized artifact may be determined (the dimensions a and c in Figure 1). Moreover, the extent to which the basal planes are parallel to one another, or crystal coherence length (la), and the mean height over which the layers are stacked in a coherent fashion (4) may be defined. These two para­meters, 4 and 4 (the crystal coherence lengths), define the perfection of the crystal (contained within the graphitized coke and binder) and the degree of graphitization.

An important feature of artificial graphite struc­ture, which has a controlling influence upon the material properties, is that the structural feature dimensions span several orders of magnitude. The crystal lattice parameters are fractions of a nanometer (a = 0.246 nm, c = 0.67 nm). The crystallite ‘coherent domains’ or extent of three-dimensional order, la and 4, are typically tens of nanometers (4 = stack height = 15-60 nm and la = stack width = 25-60 nm). The thermal microcracks between planes are typically the size of crystallites. Within the graphite, the opti­cal domain (extended orientation of crystallites) may typically range from 5 to 200 pm and largely controls the isotropy of synthetic graphite. As discussed ear­lier, graphite grain size (usually refers to largest filler particles) is a manufacturing variable and is typically in the range 1 pm to 5 mm. Finally, the pore size, depending upon the category and location (pores could be within filler or binder phases) is commensu­rate with grain size. The largest pores (excluding ther­mal cracks between the crystal layers) are typically 10 pm to a few millimeters.

Properties of CVD SiC, CVI SiC/SiC, and NITE — SiC/SiC have been reported and discussed. The CVI SiC/SiC composite combines the advantageous properties of CVD SiC and the benefits of reinforce­ment by SiC-based continuous fibers. In particular, CVI SiC/SiC exhibits damage tolerance, limited sen­sitivity to flaws and notch, high load-carrying capac­ity, and improved reliability. As opposed to earlier generations of SiC/SiC that were reinforced with Nicalon fibers, the database for SiC/SiC reinforced with the advanced near-stoichiometric fibers is incomplete. The properties of Nicalon-reinforced SiC/SiC should provide a useful baseline. The main features of the mechanical behavior of SiC/SiC composites have been described, and the relationships between the microstructure and properties have been discussed. They have been established on CVI SiC/SiC composites reinforced with Nicalon fibers. They should be reproduced on those CVI SiC/ SiC composites reinforced with near-stoichiometric Hi-Nicalon S fibers. The inherent surface roughness of Tyranno-SA3 fibers is an issue.

Precise knowledge of the mechanisms that govern the mechanical behavior is useful for proper use or design with SiC/SiC. Damage and ultimate fracture of CVI SiC/SiC involve load transfer from matrix to fibers at various length scales defined by the 2D woven structure and the tow microstructure. At high tempera­tures, additional load transfers are driven by the local stress relaxation induced by temperature and/or envi­ronment. The ultimate fracture and delayed failure are dictated by the tows. Scale effects and scatter in strength data are limited when compared to monolithic ceramics. Fiber-matrix interfaces and interphases play a significant role in damage tolerance and load­carrying capacity. Interfaces resistant to crack extension are beneficial to composite performances. However, this scheme is invalid when the fiber surface is rough.

Properties of CVI SiC/SiC can be tailored via engi­neering of the interfaces and the use of advanced fibers.

Plutonium is an isotopically composition-variable material and the variation is attributable to its genera­tion reaction in LWR fuel, the initial uranium enrich­ment and burn-up of the LWR fuel, and so forth. It needs various methodologies and much prudence in its handling because its nuclear properties differ notice­ably from one isotope (nuclide) to another. Table 124,25 summarizes the principal nuclear properties of typical nuclides in MOX fuel, including uranium isotopes. A material with high content of 238Pu is more calorific owing to its decay mode (a) and short life. Therefore, the content of 238Pu would be the limiting factor for handling batch sizes in a fabrication process. 241Pu, which also has a short life, causes alteration in the isotopic composition even during a relatively short period, for example, during storage after fuel

fabrication but before loading into a reactor. Besides the above, neutron reaction cross-sections are com­pletely different in isotopes and reactor types. Taking such variations in the cross-sections into consideration, MOX fuel is prepared, in view of plutonium content, to secure sufficient in-core reactivity.26

The nuclear characteristics of uranium and pluto­nium are needed for the evaluation of radiation exposure during the fuel fabrication process. In par­ticular, the short life of a nuclide merits attention with regard to exposure to radiation. All isotopes listed in Table 1 are a-emitters, especially 238Pu, which has highly significant a-radioactivity. 241Am, which is adjunct to 241Pu, is also a strong a-emitter. These two nuclides also give off strong g-ray emissions fol­lowing their a-decay. The major sources of neutrons are the even-A (mass number) plutonium isotopes such as 238Pu, 240Pu, and 242Pu because of their high probability for spontaneous fission. In addition, espe­cially in oxide fuels such as MOX fuel, a-particle bombardment of oxygen isotopes is an important factor that determines neutron emission. 238Pu and 241Am have a higher specific (per unit mass) influence on this reaction than other nuclides because of their large a-ray emission rates, as mentioned above. In addition, these two nuclides have a somewhat higher Q-value (a-ray energy) for decay and this increas­ingly affects the neutron production rate.

Turning to the topic of safeguards, the large neu­tron yield by spontaneous fission from the MOX fuel is utilized for a neutron coincidence counting method for inventory verification. This method uses the fact that neutrons from spontaneous fission or induced fission are essentially emitted simultaneously. This measurement can be made in the presence of neu­trons from room background or (a, n) reactions because these neutrons are noncoincident, or ran­dom, in their arrival times. The detection signals of these neutrons are analyzed and plutonium isotopes are determined by their quantity.

Burnable poison suppresses initial fuel reactiv­ity during fuel life and compensates fuel reactivity with the gradual reduction in burnable poison with burn-up. Consequently, the fuel burn-up reactivity is lowered and this lowered reactivity leads to an extended operation cycle period. Burnable poison is often mixed into oxide fuel. Gadolinium is a typical one; it has a variety of stable and substable isotopes and some of them (155,157Gd) have large thermal capture cross-sections. They are used in the form of a sesqui-oxide compound, gadolinia, in oxide fuel.

The composition, atomic structure, and microstruc­ture of an irradiated fuel are complex and evolve during irradiation. Because the thermal conductivity depends on each of these characteristics, a large number of parameters are required to determine its value. At the mesoscopic scale, the main relevant parameters are the microstructure (porosity, grains size, fission gas bubbles, and plutonium or additives distribution) and the eventual precipitates. The para­meters acting at the microscopic (atomic) scale are plutonium and the additives that are dissolved in the lattice, the soluble and insoluble fission products (fission products) and fission gases dispersed as atoms, radiation damage (interstitials, vacancies), stoichiometry, and fission density. A scheme sum­marizing some of these parameters is presented in Figure 3. The radiation effects in UO2 are described with more detail in Chapter 2.18, Radiation Effects in UO2.

2.11.3.1 Introduction

Zirconium hydride is used as a material for neutron reflectors in fast reactors. The evaluation of the ther­mal conductivity, elastic modulus, and other basic properties of zirconium hydride is extremely impor­tant for assessing the safety and cost-effectiveness of nuclear reactors. Metal hydrides, ofwhich zirconium hydride is a typical example, are also very interesting because they exhibit unique properties and shed light
on some fundamental aspects of physics. As part of work on metals such as zirconium, Yanamana et al. have successfully created crack-free, bulk-scale metal hydrides, and systematically investigated their funda­mental properties — particularly at high tempera­tures. Here, we present an outline of the results on the fundamental properties of zirconium hydride. Figure 6 shows the zirconium-hydrogen binary phase diagram.19

2.11.3.2 Production of Zirconium Hydride20

We used polycrystalline (grain size: 20-50 pm) ingots of high-purity zirconium as the starting material for producing hydrides. The main impurities present in the zirconium were O (0.25 wt%), H (0.0006 wt%), N (0.0024 wt%), C (0.003 wt%), Fe (0.006 wt%), and Cr (0.008 wt%). The hydride was generated with high-purity hydrogen gas (7 N) at a prescribed pres­sure, using an advanced ultra-high vacuum Sieverts instrument. Details of the instrument configuration are given in Figure 7.

The procedure for synthesizing hydrides varies according to the type of metal. This is due to the phase transition, from metal to hydride that is accom­panied by a massive increase in volume due to hydro­genation, and to differences in the strength of the hydride. Figure 8 shows the external appearance of zirconium hydride substances produced by the author’s group.

Low temperature incubator (inner temperature: 298 K)

Turbo-molecular pump

Oil rotary vacuum pump

Ionization vacuum gauge

Liquid nitrogen trap

Compressed hydrogen gas cylinder

Reactor for high temperature (quartz glass)

Electric resistance furnace (<1273 K)

Figure 7 Schematic diagram of advanced Sieverts instrument.

2.11.3.3 Lattice Parameter20

Zirconium hydride or deuteride described here was all fcc_C1 (8) ZrH2-x or ZrD2-x single-phase crys­tals with a fluorite structure. The lattice parameters at ambient temperature of zirconium hydride or deu­teride are plotted in Figure 9, as a function of hydro­gen content (CH). The lattice parameter of zirconium hydride increases slightly with increasing hydrogen content, according to the following formula:

a(nm) = 0.4706 + 4.382 x 10-3 x CH(H/Zr).

2.11.3.4 Elastic Modulus and Hardness20,21

Figure 10 illustrates the hydrogen content depen­dence of the elastic modulus of zirconium hydride or deuteride, determined using an ultrasonic pulse echo method. The elastic modulus of zirconium hydride is higher than that of the pure metal, and decreases slightly with increasing hydrogen content. The hydrogen content dependence of the elastic modulus of zirconium hydride is expressed by the following equations (E: Young’s modulus, G: Shear modulus, and B: Bulk modulus):

E(GPa) = 187.7 — 33.28 x Ch(H/Zr)

G(GPa) = 73.59 — 14.19 x CH(H/Zr)

B(GPa) = 130.0 — 2.329 x CH(H/Zr)

Figure 11 illustrates the hydrogen content depen­dence of the Vickers hardness of zirconium hydride and deuteride. The graph clearly shows that the Vickers hardness of the hydride is higher than that of pure zirconium, and that it decreases slightly with increasing hydrogen content. Generalizing these results, we can conclude that increasing the hydrogen content has the effect of making zirconium hydride and deuteride plastically ‘softer.’ The relationship between the hardness and hydrogen content depen­dence for zirconium hydride is expressed by the following formula:

0.479

0.478

ф

ф

E

ra

ra

CL

0.477 —

Figure 9 Hydrogen content dependence of the lattice parameter of zirconium hydride and deuteride.

Hy(GPa) = 7.190 — 2.773 x CH(H/Zr)