The only report on the solubility of hydrogen in boron carbide is that by Shirasu eta/206 They exposed crystals of boron carbide to hydrogen gas at various temperatures and pressures for 20 h, and subse­quently outgassed them during anneals in which the temperature was linearly increased at the rate of 20 K min-1 up to 1273 K. The uptake was seen to increase with the square root of pressure, and to decrease with increasing temperature (exothermic). Schnarr and Munzel207,208 measured the diffusivity of tritium in both irradiated and unirradiated boron car­bide. While the actual expression for the diffusivity for each case was not given, it can be extracted from the figures. It was noted that the apparent diffusivity decreased with increasing radiation damage until the percentage of 10B exceeded 10%. Elleman eta/.209 used the 6Li-neutron reaction to generate tritium profiles in samples of boron carbide. Diffusivity was determined by examining the rate of release of the tritium during isothermal anneals at elevated temperatures.

The diffusivity and the solubility of hydrogen in silicon carbide (a material described in Chapter 2.12, Properties and Characteristics of SiC and SiC/ SiC Composites and Chapter 4.07, Radiation Effects in SiC and SiC-SiC) have been measured twice by Causey et a/.21 ’ 1 In the first set of experi­ments,210 various grades of silicon carbide were implanted with tritium, using the neutron reaction with 6Li on the sample surfaces. The diffusivity for each material was then determined by fitting the release curves determined during isothermal anneal to those predicted by the analytical solution to the diffusion equation. The results were seen to differ strongly depending on the type and purity of the silicon carbide. As an example, the measured diffusiv — ity in hot pressed and aluminum-doped a-silicon car­bide was approximately five orders of magnitude greater than that in vapor-deposited p-silicon carbide at 1273 K. The lowest diffusivities were reported for vapor-deposited p-silicon carbide and single-crystal a-silicon carbide. In all cases, the activation energy of the diffusivity was >200 kJ mol-1 (suggesting that chemical bonding plays a strong role in the diffusion). For the diffusion of tritium in vapor-deposited silicon carbide, the diffusivity was given as D = 1.58 x 10- exp(-37 000/T) m2 s — . Deuterium solubility was also determined for the vapor-deposited silicon car­bide. The values were determined by exposing samples at elevated temperatures to deuterium gas followed by outgassing to determine the amount ofuptake. Because equilibrium retention was not obtained in the experiments, inherent in the calculations was the assumption that the diffusivity values determined in the implantation experiments were valid in the gas­eous uptake experiments. The amount of uptake was assumed to be the product of diffusivity, the solubil­ity, the sample area, and the square root of pressure. The solubility was given as K = 1.1 x 10- exp (+18 500/T) mol H2 m-2 MPa-1/2. Again, the nega­tive value of the activation energy would suggest chemical bonding of the hydrogen to the host mate­rial. In the later work by Causey et a/.,2 vapor — deposited silicon carbide was again tested. In these experiments, the implantation of energetic particles into the silicon carbide was avoided. Samples were exposed to gas containing 99% deuterium and 1% tritium at a temperature of 1573 K for 1 h. The sam­ples were subsequently outgassed at temperatures from 1373 to 1773 K. The outgassing rates were then fitted to release curves predicted by the solution to the diffusion equation to determine the diffusivity. In this case, the diffusivity was given by the expres­sion D = 9.8 x 10-8 exp(-21 870/ T) m2 s-1, one to two orders of magnitude faster than the values deter­mined earlier with energetic particles.21 The solu­bility was also determined in this study. Samples were exposed to the deuterium/tritium gas at tem­peratures from 1273 to 1873K for sufficient duration to achieve equilibrium loading. The samples were then outgassed to determine this equilibrium amount. The expression for the solubility in this case was K= 2.2 x 10-2 exp(+7060/T) mol H2 m-3 MPa-1/2. This solubility is one to two orders of magnitude lower than the one determined in the earlier experi — ments.210 If one assumes the migration of hydrogen in silicon carbide to occur along active sites on the edges of the grains, it is not unexpected that radiation damage produced by the implantation of energetic particles would increase the apparent solubility and proportionately decrease the apparent diffusivity. If hydrogen can exist only on the grain bound­aries by being attached to trap sites, higher trapping means higher apparent solubility. Conversely, higher trapping means slower diffusion. It was the apparent higher solubility on small-grained samples that led Causey et a/.211 to propose the trap-controlled grain boundary diffusion model.

The permeation of hydrogen isotopes through sili­con carbide has been measured by several groups.212-21 Verghese eta/.213 measured the permeation of a hydro — gen/tritium mixture through a KT silicon carbide tube that was manufactured by wet extrusion and sin­tering. The permeability reported for the experiments
is given by Ф = 3.8 x 108 exp(—66 000/ T) mol H2 m-1 s-1 MPa-1/2. Sinharoy and Lange212 measured the permeation of hydrogen through a tungsten tube with a CVD coating of silicon carbide. The retarding effect of the tungsten was taken into consideration in the calculation. The recorded permeation for these experiments was Ф = 2 x 10-4 exp(-6830/T) mol H2 m-1 s-1 MPa-1/2. Yao eta/.214 performed perme­ation experiments on a steel sample that had been RF sputter-coated with silicon carbide. The thickness of the coating was estimated to be 1.3 pm and contained several percent oxygen and traces ofiron. The coating was seen to decrease the permeation rate of steel by about two orders of magnitude, but did not change the activation energy. In this case, the coating was clearly porous, and the reduction in permeation was simply due to a reduction in the effective permeation surface area. The plot of the permeation values for the vapor — deposited silicon carbide by Causey et a/.21 (calcu­lated as the product of diffusivity times solubility), KT silicon carbide by Verghese et a/.,213 and CVD silicon carbide by Sinharoy and Lange212 is shown in Figure 20. The differences in the absolute values of the permeability as well as the differences in the activation energy of the process are extreme. It is difficult to even imagine that the values are for the same material. In fact, the materials are not the same. As mentioned for the original study by Causey eta/.,2 differences in impurities play a significant role in determining the behavior of hydrogen in silicon car­bide. If hydrogen does migrate along the grain bound­aries, impurity metals along those grain boundaries reduce the fraction of migrating hydrogen chemically bound to the silicon carbide. Likewise, the apparent diffusivity would be much more rapid if hydrogen trapping at the grain boundaries is reduced. In the case of the permeability measured by Sinharoy and Lange,2 2 it is difficult to believe that the measured permeation is not really controlled by permeation through the underlying tungsten with the specific surface area limited by the porous silicon carbide coating. The activation energy for the permeation in the report by Verghese eta/.213 is difficult to under­stand. The value of 555 kJ mol-1 is even greater than the chemical bond of hydrogen to carbon.215 The permeation was seen to vary by as much as an order of magnitude at the same temperature. There is no apparent explanation for the rapid change in perme­ation with temperature.

Titanium carbide has also been tested as a perme­ation barrier. Due to adhesion problems with direct deposition on steel, titanium nitride was used as an intermediate layer between the steel and titanium carbide. Forcey et a/.202 measured deuterium perme­ation through 3-pm thick layers of TiC and TiN on steel, observing a PRF of ten. For the experiments performed over the temperature range of 550-740 K, extended defects were listed as the reason for the relatively small improvement over bare steel. Checchetto eta/.201 used ion-beam assisted deposition of TiN-TiC films on steel in their permeation experiments. When the film was deposited on the downstream side, little reduction in permeation was seen. Using the deposited film on the upstream side

did yield a PRF of ^50. Shan et at216 used a CVD process to deposit their 2.5-pm thick film on steel and noted a permeation reduction of five to six orders of magnitude. It is obvious from these three studies that the deposition oftheoretically dense thin films is very difficult. There is also the question of cracking of such thin films during thermal cycling. This is dis­cussed later in this chapter.