1. PURPOSE
1.1 PURPOSE AND BACKGROUND 
One potential failure mechanism for titanium and its alloys under repository conditions is via the 
absorption of atomic hydrogen in the metal crystal lattice. The resulting decreased ductility and 
fracture toughness may lead to brittle mechanical fracture called hydrogen-induced cracking 
(HIC) or hydrogen embrittlement. For the current design of the engineered barrier without 
backfill, HIC may be a problem since the titanium drip shield can be galvanically coupled to 
rock bolts (or wire mesh), which may fall onto the drip shield, thereby creating conditions for 
hydrogen production by electrochemical reaction. 
The purpose of this scientific analysis and modeling activity is to evaluate whether the drip 
shield will fail by HIC or not under repository conditions within 10,000 years of emplacement. 
This Analysis and Model Report (AMR) addresses features, events, and processes related to 
hydrogen induced cracking of the drip shield. REV 00 of this AMR served as a feed to Waste 
Package Degradation Process Model Report and was developed in accordance with the activity 
section Hydrogen Induced Cracking of Drip Shield of the development plan entitled Analysis 
and Model Reports to Support Waste Package PMR (CRWMS M&O 1999a). This AMR, 
prepared according to Technical Work Plan for: Waste Package Materials Data Analyses and 
Modeling (BSC 2002), is to feed the License Application. 
Titanium Grade 7 (Ti-7) (UNS R52400) is now favored for construction of the drip shield for the 
waste package due to its excellent corrosion resistance (CRWMS M&O 2000a, Section 1.2). 
This alloy consists of 0.3% Fe, 0.25% O, 0.12-0.25% Pd, 0.1% C, 0.03% N, 0.015% H, 0.4% 
total residuals with the balance being Ti (CRWMS M&O 1999b, p. 45; DTN: 
MO9906RIB00052.000). Other titanium alloys that will be discussed in this report include Ti-2, 
Ti-5, Ti-12, Ti-16, Ti-17 and Ti-24. Ti-2 is essentially commercially pure titanium. It is 
believed that the most effective way of reducing hydrogen absorption is to choose a more crevice 
resistant alloy (Shoesmith et al. 1997, p. 15). This improvement in resistance to crevice 
corrosion can be attributed to the addition of alloying elements, which reinforce passivity. 
Therefore, susceptibility to crevice corrosion is reduced through the alloying series Ti-
2.Ti-12.Ti-16 (Shoesmith et al. 1997, p. 15). The major additions are Ni (0.6-0.9%) for Ti-12 
and Pd (0.04-0.08%) for Ti-16 (CRWMS M&O 1999b, pp. 49, 51; DTN: 
MO9906RIB00052.000). A decrease in the corrosion rate by the addition of a noble metal such 
as Pd in Ti-16 is noted in Shoesmith et al. (1997, p. 21). The composition of Ti-7 is almost 
identical to that of Ti-16, but the Pd content is higher in Ti-7 (0.12-0.25%) than in Ti-16 (0.04-
0.08%) (CRWMS M&O 1999b, pp. 45, 51; DTN: MO9906RIB00052.000). The similarity 
between Ti-7 and Ti-16 was also noted by CRWMS M&O (2000a, Section 1.2). 
The drip shield will experience a wide range of conditions during its service life (CRWMS M&O 
2000a, Section 1.3). Initially, the underlying waste packages will be hot, and drip shield surfaces 
will be dry due to the heat generated from radioactive decay. The temperature will eventually 
drop to levels where both humid air and aqueous phase corrosion will be possible. Crevice 
corrosion and slow general passive corrosion will produce hydrogen, which can be absorbed into 
the metal. In a drip shield design without backfill, hydrogen generation may be caused by the 

galvanic couple between the titanium drip shield surface and ground support components (such 
as rock bolts, wire mesh, and steel liners used in the drift), which may fall onto the drip shield 
surface. 
HIC is a mechanism of premature failure of the drip shield, which is one of the most important 
principal factors, and should be considered in the repository performance assessment. In this 
AMR, a realistic model will be used to account for the degradation of the drip shield due to the 
effects of HIC induced by crevice corrosion and general corrosion, based upon data generated by 
the project. The model will follow the approach of Shoesmith et al. (1997). Both qualitative and 
quantitative assessments will be made to evaluate the effect due to the galvanic couple. 
1.2 RESOLUTION OF COMMENTS IN THE NUCLEAR REGULATORY 
COMMISSION LETTER OF DECEMBER 21, 2001 
1.2.1 Resolution of Container Life and Source Term Agreement 6.02 
The NRC has requested the following information to support a potential licensing review 
(Reamer 2001, pp. 11-13): 
1. Provide better justification and needs to verify that the critical hydrogen concentration 
chosen is a realistic and representative value for the onset of cracking in Ti Grade 7. An 
evaluation of the critical hydrogen concentration for Ti Grade 24 should also be given. 
ApproachAs mentioned in section 1.1, Ti-7 is similar to Ti-16 in terms of chemical 
composition; the only difference is that Ti-7 contains more palladium than Ti-16. In 
addition, Table 1 of Been and Grauman (2000) shows that the key mechanical properties 
of Ti-16 and Ti-7 are identical (minimum yield strength = 275 MPa: minimum tensile 
strength = 345 MPa). Since both of these alloys are a-alloys containing minimal 
amounts of -phase, it is expected that they will exhibit very similar responses to applied 
stresses in an acidic environment. Based on this argument, the value of HC for Ti-7 
would be expected to be at least that for Ti-16 (i.e., ~ 1000 g/g). 
A convincing argument can also be made that the rates of hydrogen absorption will be 
independent of the Pd content of the alloy. The rate of hydrogen absorption is influenced 
by many parameters including pH and temperature. The hydrogen absorption efficiency 
decreases dramatically with increasing pH (Foroulis 1980, Okada 1983; Phillips et al. 
1974) and also with time at a constant pH (Noel et al. 1996, Figure 18). The results of 
Okada (1983) show a very significant drop in hydrogen absorption efficiency around pH 
2. Over a 50-hour period Phillips et al. (1974) found no hydrogen absorption for pH = 
10.3 and pH = 14. These measurements of hydrogen absorption (Okada 1983, Phillips et 
al. 1974) are based on accurate electrochemical measurements of hydrogen production 
rates and accurate measurements of the amount of hydrogen absorbed using the well-
characterized vacuum degassing technique. 
This influence of the presence of a surface oxide on the absorption efficiency is very 
marked. The results of Okada (1983) clearly suggest that any tendency for the Pd content 
in Ti to accelerate hydrogen absorption will be suppressed when a passive oxide is 

present (i.e., for pH = 4). It can be concluded that the primary factor controlling H 
absorption is the presence and properties of the passive oxide, not the noble metal content 
of the alloy. The results of Tomari et al. (1999) support this conclusion. The rate of 
hydrogen absorption was found to decrease parabolically with exposure time even under 
cathodically polarized conditions (equivalent to the presence of galvanic coupling with 
carbon steel) when hydrogen absorption would be expected to be accelerated through a 
reduced oxide film. The same parabolic relationship was obtained for Ti-2 and Ti-17 
(equivalent to Ti-16 but with a slightly lower oxygen content). The observation of a 
parabolically decreasing rate is evidence that the absorption efficiency is decreasing with 
time. This suggests that surface absorption sites are becoming saturated and that the rate 
of absorption will eventually become controlled by the rate of diffusive transport within 
the titanium. For the a-alloys (Ti-2, Ti-16, and Ti-7) this transport rate will be very low. 
The results of Kim and Oriani (1987a) show that a similar parabolic decrease in H 
absorption rate is observed on Ti-12, and the rate is almost independent of temperature in 
saturated brines close in composition to the most aggressive conditions expected in 
Yucca Mountain. 
Further supporting evidence comes from observations made on Ti alloys covered by 
thermally grown oxides. The presence of such oxides have been shown to suppress 
hydrogen absorption in aggressive brine environments (6% NaCl) up to temperatures 
~ 120C (Mon 2002, Attachment IV; Been and Grauman 2000; and Covington 1979), 
and the hydrogen absorption observed in deaerated HCl (2 = pH = 4; 50C to 250C) was 
completely stopped by aerated conditions (Shimogori et al. 1985). 
The above mentioned evidence clearly suggests that the HIC behavior of Ti-7 should be 
at least as resilient as that of Ti-16. 
The testing of Ti-5 and its Pd-modified version Ti-24 show that the addition of Pd to Ti-5 
improves the alloys corrosion resistance in an analogous manner to that observed when 
Pd is added to the Ti-2 alloy to produce Ti-16 (Schutz 1995; Schutz and Xiao 1993; and 
Kitayama et al. 1992). While the absolute corrosion performance of Ti-24 doesnt quite 
match that of Ti-16, it has many similar corrosion properties. The addition of Pd severely 
reduces the active corrosion rate in hot acidic solutions and eliminates susceptibility to 
crevice corrosion except in extremely aggressive conditions. Thus the alloy does not 
crevice corrode in naturally aerated 20% NaCl up to 260C. The mechanism by which 
Pd improves the corrosion resistance is the same as in Ti-7 and Ti-16, i.e., the 
accumulation of Pd in the corroding surface ennobles the alloy. Also, hydrogen-charging 
studies (Kitayama et al. 1992) revealed no measurable influence of Pd addition to Ti-5 (to 
produce Ti-24) on the alloys susceptibility to hydrogen absorption or embrittlement. 
Although no specific measurements exist, there is no reason to expect that the general 
passive corrosion rates for Ti-24 will be significantly different from those observed for 
Ti-16, Ti-12, and Ti-7 in Long Term Corrosion Test Facility (LTCTF) measurements at 
Lawrence Livermore National Laboratory (LLNL). Ti-24 is an a- alloy and, while the 
composition and physical structure of the passive oxide film will be slightly different on 

this alloy compared to that on the a-alloys (Ti-2, Ti-7, Ti-16) (Ask et al. 1988-89), this 
would not be expected to influence corrosion of the alloy under repository conditions. 
Any difference that might exist would only be expected under extreme conditions of 
acidity or alkalinity, especially in the former case if significant F-concentrations were 
present in the exposure environment (Schutz and Thomas 1987). The mechanism by 
which F- attacks Ti is through defects and/or flaws in the oxide. A greater probability of 
the existence of such flaws would be expected in the passive film on Ti-24 due to the 
differential oxide growth rates expected on the a- and -phases in the substrate. 
However, since Ti-24 will be located under the shell of the drip shield, it should not be 
exposed to the seepage drips, which could produce such conditions (i.e., differential 
oxide growth rates) by evaporative concentration. 
As for the a-alloys, whether or not the alloy is subject to HIC will be determined by the 
critical hydrogen concentration (HC) for slow crack growth and the rate of hydrogen 
absorption to achieve this critical concentration. The critical hydrogen concentration has 
been established for Ti-5 using similar experimental techniques employed to determine 
the critical hydrogen concentrations for Ti-2, Ti-12, and Ti-16 (Hardie and Ouyang 
1999). Using the slow strain rate technique on precracked compact tension specimens 
precharged with known amounts of hydrogen, it was shown that the fracture toughness of 
Ti-5 was not significantly altered until the hydrogen level in the alloy exceeded 200 g/g. 
This compares to values of 400 g/g (Ti-2), 400 g/g (Ti-12) and ~ 1000 g/g (Ti-16) 
(Shoesmith et al. 1997, 2000; Ikeda and Quinn 1998a). It seems likely that the presence 
of Pd will lead to a higher value of the critical hydrogen concentration in Ti-24 in a 
similar manner to the increase observed when adding the same amount of Pd to Ti-2 to 
produce Ti-16. In the absence of experimental measurements it is not possible to specify 
how much this improvement might be. A modest improvement to 400 g/g seems a 
reasonable HC value for Ti-24. 
A considerable amount of evidence exists to show that the rate of absorption of hydrogen 
is predominantly controlled by the properties of the oxide film, especially its thickness 
and defect properties (Pan et al. 1997; Blackwood et al. 1988; Hurlen and Hjornkol 
1990). Since the alloy will experience a long period of slow thermal oxidation prior to 
exposure to aqueous conditions, it is the properties of this oxide that will determine the 
rate of hydrogen absorption. Thermally grown oxides on Ti-2 have been shown to 
suppress hydrogen absorption on subsequent exposure to aggressive brine solutions (up 
to 6% NaCl) up to temperatures ~ 120C (Mon 2002, Attachment IV; Been and Grauman 
2000; and Covington 1979), and the hydrogen absorption observed in de-aerated HCl 
(2 = pH = 4; 50C to 250C) was completely stopped under aerated conditions 
(Shimogori et al. 1985). In the absence of exposure of the oxide film to seepage drips, 
Ti-24 would be expected to perform just as well. 
2. Provide an evaluation of possible detrimental effects of fluoride on hydrogen uptake and 
its effects on the critical hydrogen concentration and subsequent cracking. 
ApproachDetailed description of the detrimental effects of fluoride on titanium has 
been provided in Section 1.2.2. As mentioned therein, the influence of fluoride on 

hydrogen absorption by titanium is expected to be proportional to the active corrosion 
rate and the period for which it can sustain it. The primary damage mode for titanium 
under these circumstances will be the loss of wall thickness due to titanium dissolution 
(via complex formation with F-), not the increase in hydrogen content of the alloy leading 
to early failure by HIC. 
Again, as explained in Section 1.2.2, under repository conditions at Yucca Mountain, the 
Ti-7 drip shield surface will be covered with an oxide grown relatively slowly under 
thermal conditions. Such a film is expected to have a low defect density (Pan et al. 1997; 
Blackwood et al. 1988) and, hence, should be highly resistant to dissolution including 
that induced by reaction with F-. Moreover, the repository environment at Yucca 
Mountain contains lots of silica along with a variety of metal ions, which have the 
potential to consume F- ions, and thereby reduce the deleterious effects of fluoride on the 
Ti-7 drip shield to a very low level. Support of this view is provided by the observation 
that, during evaporation of J-13 water, the concentration of F-continued to decrease (due 
to precipitation) and ultimately dropped below the level of detection (CRWMS M&O 
2000a, Section 4.1.2). 
Finally, it is to be noted that HC is a material property and, hence, it should not be 
affected by an increase in corrosion rate of the material concerned (i.e., Ti-7) in the 
presence of F-ions. Most probably, this is why the effect of fluoride on HC for Ti-7 has 
not received attention in the past, and it is for this reason that a recent literature search did 
not find a published scientific article focused on this topic. 
3. Provide the results for the Ti Grade 24 structural drip shield components. 
ApproachSee approach for item number 1 mentioned above. 
1.2.2 Resolution of Container Life and Source Term Agreement 6.03 
The NRC has requested the following information to support a potential licensing review 
(Reamer 2001, pp. 11-13): 
Provide an evaluation of the possible detrimental effects of fluoride on possible hydrogen 
uptake rates, as well as enhanced corrosion resulting in higher than currently estimated 
hydrogen generation rates. 
ApproachFluoride is reported to be corrosive to titanium and its alloys over the full 
range of concentrations (Schutz and Thomas 1987). In general, the susceptibility of 
titanium to corrosion in environments containing F- is attributed to the formation of 
complex anions, such as TiF62-and TiF63-, which are highly stable and soluble in electrolyte 
solutions (Reclaru and Meyer 1998; Brossia and Cragnolino 2001). A few previous reports 
have noted that, in the absence of fluoride inhibitors (e.g., Al3+, Fe3+, etc.), solutions 
containing over 20 ppm of F- may attack titanium when the solution pH drops below 67 
(Schutz and Thomas 1987; Pulvirenti et al. 2002). However, measurements of corrosion 
potential by a number of other researchers suggest that a combination of a significant F-
concentration (> 500 g/g) and a low pH is critical for the loss of passivity of titanium 

alloys (Nakagawa et al. 1999; Nakagawa et al. 2001; Reclaru and Meyer 1998; Lorenzo de 
Mele and Cortizo 2000; Wilhelmsen and Grande 1987). In artificial saliva (at 37C) 
passivity was maintained in solutions containing 5001000 ppm F-providing the pH was 
above 4.3. For lower pH values the corrosion potential fell from the range 0 to -500 mV 
(vs. saturated calomel electrode [SCE]) to -850 mV. The latter potentials are a clear 
indication that aggressive active corrosion conditions had been established. Analyses of Ti 
concentrations in solution, using Inductively Coupled Plasma Emission Spectroscopy-Mass 
Spectrometry, confirmed this loss of passivity (Nakagawa et al. 1999). Polarization curves 
confirmed the introduction of an active region and also a significant increase in passive 
current densities at more positive potentials (Nakagawa et al. 2001). Thus, the passive 
current density in the presence of 500 ppm F-increased as follows; 10 A/cm2 (pH = 5), 
2030 A/cm2 (pH = 4.5), 4050 A/cm2 (pH = 4.3), 6090 A/cm2 (pH = 4.2), 
400A/cm2 (pH = 4.0), and 700A/cm2 (pH = 3.0) (Nakagawa et al. 2001). It should be 
noted that while these values illustrate the dependence on pH the absolute values should be 
treated with caution since they were recorded under dynamic polarization conditions and 
the films are not representative of steady-state defect-annealed oxide films. 
The importance of the condition of the passive film was demonstrated by Lorenzo de Mele 
and Cortizo (2000). Their measurements, conducted in synthetic saliva (pH = 6.5) 
containing 0.2 mol/L of F-, showed that if the F-was added shortly after electrode 
immersion while the oxide film was still growing, then the corrosion potential fell rapidly 
indicating attack of the oxide. However, if the oxide was allowed to grow for 4 days, a 
sufficient period for the annealing out of most of the defects in the oxide, then the 
subsequent addition of this concentration of F-had no observable effect over the 
subsequent 2 days of exposure. They concluded that the susceptibility to F-was associated 
with defects and flaws in the oxide. 
These results (Lorenzo de Mele and Cortizo 2000) strongly suggest that the influence of F-
is via the same mechanism as that involved for Cl-and requires that F- utilize defects within 
the oxide in order to initiate oxide breakdown rather than chemically-attacking the outer 
oxide surface. Polarization curves recorded by Lorenzo de Mele and Cortizo (2000) at 
different [F-] showed that a conventional breakdown potential could be observed. This 
breakdown potential decreased from +0.4 V (vs. SCE) in 0.02 mol/L NaF to 0.25V for 0.5 
mol/L NaF. The presence of blisters in the oxide film after polarization experiments 
(observed by Scanning Electron Microscopy) confirmed the local, as opposed to general, 
attack of the oxide in the presence of F-. 
Pan et al. (1997) showed that the oxide defect density, determined from film capacitance 
measurements using MottSchottky plots, for films produced by ion-beam assisted 
sputtering deposition had both a lower defect density and a lower passive current density 
than those of oxides grown under natural corrosion conditions in aqueous solutions. These 
results clearly connect the passive dissolution rate to the density of defects in the oxide and 
are consistent with the results of Blackwood et al. (1988) who also found a rate of 
dissolution dependent on defect density (e.g., degree of crystallinity) of the oxide film and 
with the claims of Hurlen and Hjornkol (1990) that the kinetics of passive corrosion of 

titanium in neutral solutions is controlled by the migration of the predominant defect, an 
oxygen vacancy across the oxide film. 
In the presence of fluoride, the rate of hydrogen absorption would be expected to be 
proportional to the rate of the corrosion process, as it would be under active corrosion 
conditions within a crevice (Noel et al. 2001). Even under these conditions there is 
accurate electrochemical evidence to show that the hydrogen absorption efficiency 
decreases with time (Noel et al. 1996). This can be attributed to the formation of a surface 
hydride layer leading to control of the hydrogen absorption process by transport across this 
layer. This conclusion is consistent with the extensive results of Phillips et al. (1974). 
Once this layer is formed, its removal by anodic dissolution and reformation by the 
cathodic reduction of protons achieve steady state and the rate of hydrogen absorption 
becomes controlled by the rate of diffusion of hydrogen into the alloy from the 
hydride/alloy interface (Shoesmith et al. 2000). 
Thus, the influence of fluoride on hydrogen absorption will be proportional to the active 
corrosion rate it can sustain and the period for which it can sustain it. The primary damage 
mode under these circumstances will be the loss of wall thickness, not the increase in 
hydrogen content of the alloy leading to early failure by HIC. 
A number of species, in particular certain compounds containing elements from groups V 
and VI of the periodic table (e.g., oxyanions of S, Sb, Se, Te, and CN-), are known to be 
catalysts for hydrogen absorption (Jones 1996, p. 335). These promoters act by retarding 
the formation of molecular hydrogen gas and enhancing the residence time of hydrogen 
atoms on the surface, thereby facilitating its absorption into the metal. A specific 
interaction with surface-adsorbed H-states is required for a chemical species to act as a 
catalyst to increase the hydrogen absorption efficiency. There is no evidence that halide 
ions can interfere with the hydrogen absorption process in this manner. Consequently, 
fluoride is not expected to function as a catalyst for hydrogen absorption. 
Finally, based on the above discussion it can be concluded that F-attack of passive titanium 
requires acidic conditions (pH ~ 4) as well as the presence of flaws in the oxide which 
allow access of F- to the metal surface and the establishment of localized acidic conditions. 
Once these conditions are established, attack of the metal by F- is rapid. This raises a 
question as to why Brossia and Cragnolino (2001) observed aggressive active corrosion of 
Ti in neutral 1 mol/L NaCl solution. Two factors appear to be important in understanding 
their work: 
1. 
The values were recorded in polarization experiments on freshly prepared Ti 
specimens covered only with a defective anodic film. As discussed above, attack of 
the metal by F- is to be expected on such surfaces. 
2. 
The temperature was 95C, a value at which flaws could develop in the solution-
grown oxide film on Ti (Shibata and Zhu 1995). 
By contrast, under repository conditions at Yucca Mountain, the Ti-7 drip shield surface 
will be covered with an oxide grown relatively slowly under thermal conditions, and 

scientific evidence exists (see Section 1.2.1) to show such a film will have a low defect 
density and, hence, should be highly resistant to dissolution including that induced by 
reaction with F-. Moreover, the repository environment at Yucca Mountain contains silica 
and a variety of metal ions, which have the potential to consume F-ions, and thereby keep 
the deleterious effects of fluoride on the Ti-7 drip shield at a very low level. This view is 
supported by the observation that, during evaporation of J-13 water, the concentration of F-
decreased (due to precipitation) and ultimately dropped below the level of detection 
(CRWMS M&O 2000a, Section 4.1.2). 
2. QUALITY ASSURANCE 
The Office of Civilian Radioactive Waste Management Quality Assurance program applies to 
this AMR. All types of waste packages were classified per QAP-2-3 as Quality Level-1. 
CRWMS M&O (1999c, p. 7), Classification of the MGR Uncanistered Spent Nuclear Fuel 
Disposal Container System, is cited as an example of a waste package type. This classification is 
still in effect. The applicability of Quality Assurance Requirements and Description 
requirements (DOE 2002) to Revision 00, ICN 01, and ICN 02 activities was established in 
activity evaluation, 1101213PM7 Waste Package Analysis & Models - PMR (CRWMS M&O 
1999d), which was prepared per QAP-2-0. These earlier versions of this AMR were developed 
in accordance with the development plan (CRWMS M&O 1999a) and AP-3.10Q, Analyses and 
Models. Consistent with this development plan, no special controls on the electronic 
management of data were necessary in these activities. 
This revision of the AMR has been developed per AP-SIII.10Q procedure under Technical Work 
Plan for: Waste Package Materials Data Analyses and Modeling (BSC 2002), which, in turn, 
was developed in accordance with AP-2.27Q, Planning for Science Activities. The applicability 
of Quality Assurance Requirements and Description requirements (DOE 2002) is documented in 
the technical work plan (TWP) (BSC 2002, Section 8). A process control evaluation was 
performed in accordance with AP-SV.1Q, Control of the Electronic Management of Information. 
Results documented in Addendum A of the technical work plan (BSC 2002) specifies measures 
to be adopted for the electronic management of data/information. 
3. USE OF SOFTWARE 
No computer software or models have been used to support the development of the analysis and 
modeling activities described in this AMR. 

4. INPUTS
4.1 DATA AND PARAMETERS 
The following input data and/or parameters are obtained from Weast (1978, p. B-177): 
.Ti = density of titanium 
MTi = Molecular weight of titanium 
The following input data and/or parameters for Ti-7 are obtained from CRWMS M&O (2000a, 
Section 6.5.4): 
Ruc = passive general corrosion rate (maximum). 
The following design parameters are obtained from BSC (2003) and 10 CFR 63.114(d), 
respectively: 
d0 = Half of drip shield thickness 
t = time of emplacement 
For project consistency, these parameters have been tabulated as shown in Table 1. All of these 
input data and parameters contained in Table 1 will be used to evaluate the HIC effects on the 
drip shield. 
Table 1. Input Parameters 
Parameter 
Name 
Parameter 
Source DTN 
Parameter 
Value(s) Units 
Distribution (or 
single value if fixed) 
.Ti = Density of Weast N/A 4.5 g/cm3 Fixed value 
titanium 1978, p. B177 
MTi = Molecular Weast N/A 47.9 N/A Fixed value 
weight of 1978, p. B-
titanium 177 
Ruc = Rate of N/A DTNs: 
general passive MO0003SPASUP02.003 3.88x10-4 mm/year Maximum value 
corrosion LL990610605924.079 1.76x10-4 mm/year Maximum value 
(maximum) 
d0 = Half of drip BSC 2003 N/A 7.5 mm Fixed value 
shield thickness 
t = time of 
emplacement 
10 CFR 
63.114(d) 
N/A 10,000 years Fixed value 
In general, for the input parameters of the HIC Model, a conservative approach will be used in 
that input values at (or near) lower or upper bounds (as appropriate) will be used to ascertain that 
ANL-EBS-MD-000006 REV 01 15 February 2003 

the estimated output value is conservative. However, the primary sources of uncertainty in the 
HIC Model are associated with the selection of values for the rate of general passive corrosion, 
Ruc, and the fractional hydrogen absorption efficiency, fh, discussed in Section 5.5. In evaluating 
the model, the maximum values for these parameters will be used for a conservative estimate of 
hydrogen concentration in the drip shield. 
For the purposes of this analysis, the maximum possible general corrosion rate of the drip shield 
is considered to be 3.88x10-4 mm/year (DTN: MO0003SPASUP02.003). This value results from 
measurements of the corrosion rate of samples with both weight-loss and creviced geometries 
(and corrected for potential effects of silica deposition). Conceptually, this general corrosion 
rate would be appropriate for use only on the drip shield outer surface since only the drip shield 
outer surface could be contacted by dripping water. Dripping water could potentially lead to the 
formation of scale and/or salt deposits under which crevice conditions could form. On the 
underside of the drip shield, no crevice conditions are possible as only condensed (hence 
relatively impurity-free) water could contact these regions. Therefore, the maximum possible 
general corrosion rate for the underside of the drip shield is considered to be 1.76x10-4 mm/year 
(DTN: LL990610605924.079). This value results from measurements of the corrosion rate of 
samples with weight-loss geometries only (and corrected for potential effects of silica 
deposition). Since the rate of hydrogen generation on the drip shield surface, and hence the rate 
of hydrogen absorption into the drip shield, will be proportional to the rate of drip shield 
corrosion, and the two surfaces of the drip shield are to corrode at two different rates, the average 
of these two maximum corrosion rates, 2.82x10-4 mm/year, is appropriate for its use in the HIC 
model. 
4.2 CRITERIA 
The Waste Package Technical Work Plan (BSC 2002, Table C5) has identified the following 
acceptance criteria (AC) from Project Requirements Document (Curry and Loros 2002) and 
Yucca Mountain Review Plan, Draft Report for Comment (NRC 2002). According to the Waste 
Package TWP (BSC 2002, Section 3), this AMR is to address only these applicable acceptance 
criteria: 
1. 
System Description and Demonstration of Multiple Barriers (Curry and Loros 2002, 
PRD-002/T-014, PRD-002/T-016; NRC 2002, Section 4.2.1.1.3) 
 
AC1: Identification of barriers is adequate. 
 
AC2: Description of the capability of identified barriers is acceptable. 
2. 
Degradation of Engineered Barriers (Curry and Loros 2002, PRD-002/T-015; NRC 2002, 
Section 4.2.1.3.1.3) 
 
AC1: System description and model integration are adequate. 
 
AC2: Data are sufficient for model justification. 

4.3 CODES AND STANDARDS 
The acceptance criteria listed in Section 4.2 are consistent with the methodology described in the 
ASTM Standard Practice C 1174 for prediction of the long-term behavior of engineered barrier 
system components in a geologic repository (ASTM C 1174-97 1998). 
5. ASSUMPTIONS 
The following assumptions are used for the HIC model of this AMR. The technical basis for 
each assumption is also provided so that these assumptions require no further verification. 
5.1 
Due to the close similarity in chemical composition between Ti-7 and Ti-16, available 
material property data for Ti-16 are assumed to be applicable to Ti-7 if the said material 
property data are not available for Ti-7. This assumption is used in Sections 5.2, 6.1, and 
5.2 
The HC value for Ti-7 is assumed to be at least 1000 g/g. This assumption is based on 
Assumption 5.1 and results reported by Ikeda and Quinn (1998a, p. 7), which indicated that 
the HC value for Ti-16 is between 1000 and 2000 g/g. This assumption is necessary 
because HC data are not available for Ti-7. This value is conservative on a number of 
counts: (1) the value of 400 g/g is the lower limit of the measured range of 400 to 1000 
g/g measured at room temperature (~25C) for Ti-2 (Shoesmith et al. 1997, pp. 7-11); (2) 
the temperature of the drip shield will be well above room temperature for the majority of 
its anticipated lifetime and Clarke et al. (1995) have shown HC for both Ti-2 and Ti-12 rises 
to ~1000g/g at a temperature of 92C to 100C; and (3) the value of 400 g/g was 
measured for Ti-2, whereas the value for Ti-16, a palladium containing a-alloy with similar 
properties to Ti-7, appears to be > 1000 g/g (Ikeda and Quinn 1998a, p.7). This 
assumption is used in Sections 6.1.3 and 6.2.2. 
6.2. 
It is noted that 400 g/g, the lower bound value observed for other titanium alloys (e.g., 
Ti-2 and Ti-12) by Shoesmith et al. (1997, pp. 7-11), was used as the critical hydrogen 
concentration for Ti-7 in REV 00 of this AMR. As discussed above, this value appears to 
be very conservative based on data reported by Ikeda and Quinn (1998a, p. 7). 
5.3 
Mechanisms for hydrogen generation and absorption on a titanium surface may include a 
galvanic couple, hydrogen produced in atomic form by the corrosion process, and the direct 
absorption of hydrogen produced by water radiolysis. As indicated by Shoesmith et al. 
(1997, p. 17), the direct absorption of radiolytically produced hydrogen is insignificant 
except at high dose rate (> 102 Gy/h) and high temperature (> 150C), clearly unattainable 
under Yucca Mountain drip shield conditions. Therefore, only generation and absorption 
of hydrogen produced by general passive corrosion and by corrosion in the presence of a 
galvanic couple are considered in this AMR. 
5.4 
The passive oxide film growth rate and, hence, the corrosion rate are assumed to be 
constant in time for the formula used to calculate the concentration of hydrogen in the 
metal (Shoesmith et al. 1997, p. 22). The rationale for this assumption is that the 
assumption of constant corrosion rate is conservative, and less conservative corrosion 

models assume that the rate decays with time (CRWMS M&O 2000a, Section 5.3). It is 
also implicit in this assumption and this formula that any hydrogen absorbed into the 
titanium remains within the remaining wall thickness and is not subsequently removed as 
the general corrosion process progresses with time. This is a very conservative assumption, 
since any titanium hydride formed within the alloy will be unstable with respect to titanium 
oxide (Beck 1973) and should, therefore, be removed by the general corrosion process. 
This assumption is used in Sections 6.1.5, 6.1.6, and 6.2.3. 
5.5 High and low values for fractional efficiency of hydrogen absorption, fh, for titanium alloys 
have been assumed to be 0.015 and 0.005, respectively. The rationale for this assumption 
is that these values were adopted from the results of experimental measurements by Okada 
(1983) performed under extremely conservative environmental conditions. These 
aggressive electrochemically polarized conditions are considered unachievable in the 
proposed repository at Yucca Mountain. For example, they were measured on Ti-2 under 
constant applied current conditions with an applied current of 0.5 mA/cm2 in acidic sodium 
sulfate solutions of pH = 4 (at 25C). Under these conditions, the electrode potential 
achieved is < -1.0V (vs. SCE), 250 to 400 mV more negative than would be anticipated 
under even galvanic coupling conditions to carbon steel (Hodgkeiss et al. 1987). At these 
potentials redox transformations in the oxide will render it permeable to hydrogen 
(Shoesmith et al. 1997, Figure 2 and p. 2). However, this permeability will not be present 
under non-galvanic conditions. Furthermore, the impermeability of the passive oxide has 
been shown to be improved by thermal oxidation prior to immersion (Mon 2002, 
Attachment III). Such conditions are expected to prevail for an extensive period of 
emplacement in the repository prior to wetting. 
Although fh decreases with time as noted by Noel et al. (1996, Figure 18) and Tomari et al. 
(1999, Table 1), a value of 0.015 has been conservatively assumed constant throughout 
10,000 years. This value is exceedingly conservative in comparison with the fh value of 
0.00014 reported by Tomari et al. (1999, Table 1) on the basis of experiments conducted 
for 1,440 hours under electrochemically polarized conditions. This conservative 
assumption has been used in Sections 6.1.5, 6.1.6, 6.2.3, 6.3.3, and 6.3.5. 
5.6 
The extent of galvanic corrosion due to coupling between carbon steel and Ti-7 will be 
limited by the small area of wet contact between these two metals, and the low cathode-to-
anode ratio (Ti/Fe) which is likely to approach one. This is a rational assumption because 
the point of galvanic contact also simultaneously experiencing seepage drips is likely to be 
an incident of very low probability. The resulting inability of the site to remain wet due to 
run-off and the accumulation of iron oxide corrosion products will significantly impede 
galvanic corrosion. The passivating effects of iron oxides on the Ti surface and the 
presence of ferric ions in the corrosive solutions are well known in the scientific literature 
(Covington and Schutz 1981; Schutz and Thomas 1987). In view of these limitations, the 
mass of steel that could be consumed by galvanic corrosion and the initial contact area are 
assumed to be 22.727 kg (50 lbs) and 75 cm2, respectively. These assumptions have been 
used in Section 6.3.4 and 6.3.5. 

5.7 
A mathematical model and an alternative conceptual model (ACM) are proposed in 
Sections 6.3.3 and 6.3.5, respectively, to predict the hydrogen concentration in the drip 
shield due to a galvanic coupling between the drip shield and a carbon steel segment. 
Simplifying assumptions are employed to develop the mathematical model. Because the 
assumptions are to simplify the system for the mathematical model, they do not require a 
technical basis or verification. These assumptions are: 
 
The drip shield is treated as an infinite flat plate. This is a conservative assumption 
since it allows a larger surface area to be available for galvanic coupling between the 
drip shield and a carbon steel segment than would be the case with a curved drip shield 
structure. This assumption is also to simplify the geometry for the mathematical model. 
 
The wetted contact plane between the drip shield surface and the carbon steel segment 
is circular in shape. This assumption is also to simplify the geometry for the 
mathematical model. 
 
Hydrogen is absorbed into the drip shield through the wetted contact plane and 
immediately reaches the other surface. The initial region of the drip shield 
experiencing hydrogen absorption has the shape of a circular disk. This assumption is 
also to simplify the geometry for the mathematical model. 
 
As hydrogen starts to diffuse in the drip shield, the radius of the circular region 
containing hydrogen will expand at a constant speed. This assumption is also to 
simplify the geometry for the mathematical model. 
6. ANALYSIS/MODEL 
6.1 
DESCRIPTION OF MODEL FOR HYDROGEN INDUCED CRACKING 
6.1.1 Introduction 
The purpose of this scientific analysis and modeling activity is to evaluate whether the drip 
shield will fail by HIC or not under repository conditions within 10,000 years of emplacement. 
This AMR addresses features, events, and processes related to hydrogen induced cracking of the 
drip shield. 
The drip shield will experience a wide range of conditions during its expected service life of 
10,000 years (CRWMS M&O 2000a, Section 1.3). Initially, the underlying waste packages will 
be relatively hot, and drip shield surfaces will be dry due to the heat generated from radioactive 
decay. The temperature will eventually drop to levels where both humid air and aqueous phase 
corrosion will be possible. Crevice corrosion and slow general passive corrosion may be 
initiated and propagate in the metal. Both types of corrosion will produce hydrogen, which can 
be absorbed into the metal (Shoesmith et al. 1997, p. 2). In a drip shield design without backfill, 
hydrogen generation may be caused by the galvanic couple between the titanium drip shield 
surface and the ground support (such as rock bolts, wire mesh, and steel liners used in the drift), 
which may fall onto the drip shield surface. 

Failure occurs if (1) the wall penetration by corrosion exceeds the corrosion allowance or if (2) 
the amount of hydrogen absorbed exceeds the critical hydrogen concentration, HC, for failure due 
to hydrogen induced cracking (Clarke et al. 1994, Figure 8). The first type of failure mechanism, 
i.e., wall penetration by crevice and general corrosion, is treated in another AMR (CRWMS 
M&O 2000a). This AMR deals with the failure mechanism associated with HIC. 
HIC (also called hydrogen embrittlement) is characterized by decreased fracture toughness or 
ductility of the metal due to absorbed atomic hydrogen. The usual mechanical failure mode for a 
ductile material is the ductile tearing observed during slow crack growth. Decreased fracture 
toughness causes fast crack growth (brittle fracture) of a normally ductile material under 
sustained load. During slow crack growth, the material will fail as the stress intensity factor K 
reaches a value KS. During fast crack growth, the same material will fail as the stress intensity 
factor K reaches a value KH, which is less than KS (Shoesmith et al. 1997). Figure 1 (Shoesmith 
et al. 1997, page 9, Figure 7) represents, schematically, the combinations of stress intensity 
factor and hydrogen concentration leading to (1) fast crack growth (brittle fracture) controlled by 
KH, (2) slow crack growth controlled by KS due to either sustained load cracking or ductile 
rupture, or (3) no failure. 
For corrosion induced HIC, an approach has been adopted to predict when HIC might become a 
potential failure process for the drip shield, following a Canadian precedent (Shoesmith et al. 
1997). The basic premise of the model is that failure will occur once the hydrogen content 
exceeds a certain limit or critical value. Both qualitative and quantitative assessments will be 
conducted to evaluate the effect due to the galvanic coupling. 
This report will address: 
Processes by which Ti-7 can absorb hydrogen and, hence, eventually become susceptible to 
HIC 
Analysis procedures for predicting the consequences of hydrogen absorption 
Acquisition of input data for the analysis parameters 
Results. 

KH 
Fast 
KS 
No 
[H]C 
Hydrogen concentration 
Slow 
crack fracture 
failure 
Stress intensity factor, Kgrowth 
Source: Shoesmith et al. 1997. 
Figure 1. Schematic Showing the Combinations of Stress Intensity Factor and Hydrogen Concentration 
Leading either to Fast Crack Growth (Brittle Failures) or Slow Crack Growth due to either 
Sustained Load Cracking or Ductile Rupture or to No Failure 
6.1.2 Processes by which Hydrogen is Absorbed 
In accordance with Schutz and Thomas (1987, p. 673), factors have been identified that can lead 
to HIC, i.e., hydrogen uptake and possible embrittlement of a and near-a titanium alloys (such 
as titanium grade 2, grade 7 and grade 12) in aqueous media. The three general conditions that 
must exist simultaneously for the hydrogen embrittlement of a alloys are: 
A mechanism for generating hydrogen on a titanium surface. The mechanisms may include a 
galvanic couple, hydrogen produced in atomic form by the corrosion process, and the direct 
absorption of hydrogen produced by water radiolysis. As indicated in Section 5 (Assumption 
5.3), the direct absorption of radiolytically produced hydrogen is insignificant except at high 
dose rate (> 102 Gy/h) and high temperature (> 150C), clearly unattainable under Yucca 
Mountain drip shield conditions. Therefore, only hydrogen produced by a galvanic couple and 
by general corrosion are considered in this AMR. 
A metal temperature above approximately 80C (175F) when the diffusion rate of hydrogen into 
a titanium becomes significant. 

A solution pH less than 3 or greater than 12 or impressed potentials more negative than -0.7 V 
(SCE). 
Assessment of HIC due to the effects of a galvanic couple will be presented in Section 6.3. The 
remainder of Section 6.1 and Section 6.2 will be devoted to HIC resulting from hydrogen 
produced by corrosion of titanium. 
Both crevice corrosion and general passive corrosion will be accompanied by hydrogen 
production and, hence, possibly by the absorption of hydrogen into the metal. 
For crevice corrosion the hydrolysis of dissolved metal cations leads to acidification within the 
occluded area and the development of active conditions in which the metal is unprotected by an 
oxide film. Once initiated, crevice corrosion is supported by both the reduction of oxygen on 
passive surfaces external to the crevice and the reduction of protons (Ti + 4H+ . Ti4+ + 2H2) on 
exposed metal surfaces inside the crevice. The latter process can lead to the absorption of atomic 
hydrogen into the metal in sufficient quantities to produce extensive hydride formation 
(Shoesmith et al. 1997, Figure 3, p. 4). 
For the passive non-creviced or inert crevice conditions expected to prevail, the corrosion of the 
titanium alloy will be sustained by reaction with water under neutral conditions 
(Ti + O H 2 . TiO2 + H 2 2) and will proceed at an extremely slow rate. This process will 
2 
generate hydrogen, which must pass through the TiO2 film before absorption into the underlying 
Ti alloy can produce HIC. Redox transformations (TiIV-> TiIII) in the film are required before 
the oxide becomes significantly transparent to hydrogen. Significant cathodic polarization of the 
metal (generally only achievable by galvanic coupling to carbon steel or the application of a 
cathodic protection potential) is required for these transformations to occur. Measurements of 
absorbed hydrogen suggest a threshold potential of -0.6 V versus SCE above which no 
absorption occurs (Shoesmith et al. 1997, p. 4). Only impressed current cathodic protection or 
galvanic coupling to active alloys such as Fe, Zn, or Mg could produce such cathodic potentials. 
Thus, hydrogen generation rates, and, hence, hydrogen absorption rates, are expected to be very 
low on Ti alloy surfaces. 
6.1.3 Critical Hydrogen Concentration, HC 
Provided that wall penetration by corrosion does not exceed the corrosion allowance, HIC failure 
is assumed to occur when the material has absorbed enough hydrogen to exceed the critical 
hydrogen content, HC. Observations regarding the critical hydrogen content, HC, for Ti-2 and 
Ti-12 are summarized as follows: 
Using the slow strain rate technique on precracked compact tension specimens precharged with 
known amounts of hydrogen, it has been shown that the fracture toughness of Ti-2 and Ti-12 is 
not significantly affected until the hydrogen content exceeds HC. Once above HC, there is no 
slow crack growth, and only fast crack growth is observed. 
HC is sensitive to the microstructure and texture of the material with respect to the orientation of 
the crack and the applied stress. Preferential pathways for cracking are formed along -phase 
stringers introduced through the manufacturing process. An HC of 500 g/g has been measured 

in Ti-12 containing cracks propagating in the directions defined by these stringers. Crack 
propagation perpendicular to these features is not observed up to HC = 2000 g/g. Heat 
treatment to remove this laminar structure by randomly reorienting the residual -phase can lead 
to a decrease in HC to ~400 g/g. 
Even for manufactured plate materials that do not have high -phase content, e.g., Ti-2, the 
laminar structure introduced by rolling appears to dominate the cracking behavior. As a 
consequence, depending on crack orientation, HC varies between ~400 and 1000 g/g. Since 
Ti-2 does not contain as much -phase as Ti-12, heat treatment does not exert a significant 
influence on HC. Welding produces a larger change in the microstructure than does heat 
treatment. The high weld temperature results in significant microstructural changes in the 
weldment. This results in a small decrease in strength. The heat affected zone does not appear 
to be sufficiently large to influence the cracking behavior because, for both Ti-2 and Ti-12, HC is 
slightly decreased near the weld metal compared to the base metal. It did not decrease below 
500 g/g. 
It can also be assumed that failure is less likely at elevated temperature. This assumption is 
justified, in part, by recognizing that the hydrogen solubility would be higher at elevated 
temperature. However, it has also been shown experimentally that HC increases markedly with 
temperature (Clarke et al. 1995). While the maximum critical stress intensity factor for Ti-2 
decreases slightly from ~50 MPam1/2 to ~40 MPam1/2 at 95C, only slow crack growth was 
observed up to hydrogen concentrations of ~2000 g/g, clearly indicating an enhanced resistance 
to growth of brittle-like cracks as the temperature is increased. A similar increase in resistance 
to brittle fracture was observed for Ti-12 at 95C, the maximum stress intensity factor decreasing 
from ~60MPam1/2 to ~45 MPam1/2, while HC increased to ~1000g/g (Clarke et al. 1995). 
Preliminary creep measurements clearly indicate that increased creep deformation at higher 
temperatures was a major factor in preventing the development of a sufficiently high stress 
concentration to initiate fast fracture (Clarke et al. 1995). 
HC data are not available for Ti-7 but more recent data reported by Ikeda and Quinn (1998a, p. 7) 
indicated that the HC value for Ti-16 is between 1000 and 2000 g/g. As noted in Section 1, Ti-7 
and Ti-16 are similar alloys because of their similar chemical compositions. 
The HC value for Ti-7 is assumed to be at least 1000 g/g. This assumption is based on 
Assumption 5.1 and data reported by Ikeda and Quinn (1998a, p. 7), which, as indicated earlier, 
concluded that the HC value for Ti-16 is between 1000 and 2000 g/g. This assumption is 
necessary because HC data are not available for Ti-7. 
It is noted that 400 g/g, the lower bound value observed for other titanium alloys (e.g., Ti-2 and 
Ti-12) by Shoesmith et al. (1997, pp. 7-11) was used as the critical hydrogen concentration for 
Ti-7 in REV 00 of this AMR. This value appears to be excessively conservative based on data 
reported by Ikeda and Quinn (1998a, p. 7). 

6.1.4 Hydrogen Absorption During Crevice Corrosion 
Crevice corrosion, if it initiates, will propagate at a rate much higher than that of general 
corrosion. Crevice propagation, as mentioned in Section 6.1.1, can lead to failure by wall 
penetration as well as failure due to HIC because of hydrogen transported into the metal. 
The most effective way to reduce hydrogen absorption is to choose a more crevice corrosion 
resistant alloy (Shoesmith et al. 1997, p. 15). Improvements in resistance to crevice corrosion 
have been attributed to alloying elements, which reinforce passivity. Susceptibility to crevice 
corrosion is eliminated through the alloying series Ti-2.Ti-12.Ti-16 (Shoesmith et al. 1997, 
p. 15). In accordance with CRWMS M&O (1999b, pp. 49, 51), the major additions are Mo (0.2 
0.4%) and Ni (0.6-0.9%) for Ti-12 and Pd (0.04-0.08%) for Ti-16. Decreased crevice corrosion 
rates by adding alloying elements such as Mo and Ni to Ti-12 and Pd to Ti-16 are noted in 
Shoesmith et al. (1997, p.16), and discussed in recent reviews of Ti crevice corrosion behavior 
(Schutz 1988). The effect of deliberate alloying of Ti-12 with Ni and Mo, which segregate at 
grain boundaries and intermetallics, is to improve the susceptibility to crevice corrosion. Thus, 
while some crevice corrosion occurs in Ti-2, only minor crevice corrosion damage prior to 
repassivation occurs in Ti-12. The effect of adding Pd to ennoble titanium while avoiding 
segregation in Ti-16 further improves the susceptibility to crevice corrosion, and no crevice 
corrosion damage is observed in Ti-16. It can be seen from CRWMS M&O (1999b, pp. 45, 51) 
that the composition of Ti-7 is almost identical to that of Ti-16, but the Pd content is even higher 
in Ti-7 (0.12-0.25%) than in Ti-16 (0.04-0.08%). Therefore, Ti-7, like Ti-16, is not expected to 
suffer any crevice damage. 
Crevice corrosion is one form of localized corrosion of a metal surface (ASM 1987, p. 4). The 
model developed by CRWMS M&O (2000a, Section 6.4.1) for the Ti-7 drip shield assumes that 
localized attack occurs only if the open circuit corrosion potential, Ecorr, exceeds the threshold 
potential for breakdown of the passive film, Ecritical. This is the universally accepted criterion for 
processes such as pitting. Experimental measurements reported by CRWMS M&O (2000a, 
Section 6.4.3) for Ecorr and Ecritical were obtained from various test environments expected in the 
repository. These test environments include simulated dilute water (SDW), simulated 
concentrated water (SCW), and simulated acidic concentrated water (SAW) at 30, 60, and 90C, 
as well as simulated saturated water (SSW) at 100 and 120C. SCW is about one thousand times 
more concentrated than J-13 well water and is slightly alkaline (pH~8). SAW is about one 
thousand times more concentrated than J-13 well water and is acidic (pH~2.7). J-13 well water 
(Harrar et al. 1990) is the groundwater from Yucca Mountain. The experimental measurements 
show that the threshold Ecritical is consistently greater than Ecorr. CRWMS M&O (2000a, Section 
6.4.3) indicated that the difference, Ecritical - Ecorr, falls in the range from 995 mV to 1652 mV for 
the various test environments of SDW, SCW, SAW, and SSW at temperatures from 20150C. 
It is therefore concluded that hydrogen absorption during crevice corrosion or any other form of 
localized corrosion such as pitting may be ignored in the HIC model. 

6.1.5 Hydrogen Absorption During General Passive Corrosion 
According to Shoesmith et al. (1997, p. 15), there are two processes by which hydrogen could be 
produced, and possibly absorbed, under passive conditions: (1) direct absorption of hydrogen 
produced by water radiolysis and (2) absorption of atomic hydrogen produced by the corrosion 
process to produce oxide. The direct absorption of radiolytically produced hydrogen does not 
appear to be significant except at high dose rate (> 102 Gy/h) and high temperature (> 150C) 
(Shoesmith et al. 1997, p. 17). This condition is clearly unattainable under Yucca Mountain drip 
shield conditions and will not be considered, leaving the corrosion process as the only feasible 
source of hydrogen for absorption. 
Under anoxic conditions, when passive corrosion should prevail, the corrosion potential for 
passive titanium must reside at a value at which water reduction can couple to titanium 
oxidation, 
Ti + O H 2 . TiO2 + H 2 2
2 
and, hence, must be at, or more negative than, the thermodynamic stability line for water. At 
such potentials, titanium hydrides are thermodynamically stable with respect to the metal. 
Consequently the passive film can be considered only as a transport barrier and not as an 
absolute barrier. The rate of hydrogen absorption at the corrosion potential will be controlled by 
the rate of the corrosion reaction, which dictates the rate of production of absorbable hydrogen. 
Since titanium oxide, TiO2, is extremely stable and protective in the drip shield environment, the 
corrosion reaction will be effectively limited to an oxide film growth reaction. 
While the rate of hydrogen production and, hence, absorption may be assumed directly 
proportional to the rate of film growth, the fraction of hydrogen absorption needs to be 
determined. Based on experimental measurements of Okada (1983), a value in the range of 
0.005 to 0.015 may be adopted for fh. These values represent the minimum and maximum values 
measured in the pH range 4 to 5 and can be considered conservative since they were measured 
on Ti-2 under constant applied current conditions with an applied current of 0.5 mA/cm2 at 25C 
in sodium sulfate solutions at pH = 4. The electrode potential achieved during these experiments 
was -1.14 V (vs. SCE), about 500 mV more negative than the threshold value of -0.6 V (vs. 
SCE) for hydrogen absorption. Similar measurements on Pt and Ni-coated Ti-2 specimens gave 
a value of fh only marginally higher under these conditions, a clear indication that the Pd-content 
of Ti-7 would not be expected to increase this value. The applied current density used was 500 
A/cm2 compared to a value of 1.45 nA/cm2 calculated as the current that would flow at a 
general corrosion rate of 50 nm/year [the 70th percentile value of general corrosion rates 
measured for Ti-16 in the LTCTF at LLNL (CRWMS M&O 2000a, Section 6.5.4)]. 
The above mentioned values of fh are appropriate for Ti-2 since the passive film is a good 
transport barrier to hydrogen absorption. Furthermore, the impermeability to hydrogen 
absorption will be improved by the period of dry thermal oxidation expected under repository 
conditions. The improved resistance to hydrogen absorption conferred by thermal oxidation has 
been documented by Mon (2002, Attachment III). For Ti-12, since Ni, in either Ti2Ni 
intermetallics or in  phase, is present, a higher value would be more appropriate. For Ti-16 and 

Ti-7 in which intermetallic formation is avoided, the values of 0.005 to 0.015 are most 
appropriate. 
It may be pointed out that in the previous versions of this AMR (REV 00, ICN 01 and 02) a 
higher value of 0.1 was mentioned as the upper bound value of fh. In this revision (REV 01), this 
value has been ruled out because of the following reasons: 
1. 
It was adopted for calculations on Ti-12 (Shoesmith et al. 2000). However, Ti-12 is a 
-alloy unlike Ti-16 and Ti-7, which are a-alloys. 
2. 
It was obtained from crevice corrosion studies on Ti-2. Since Ti-7 and Ti-16 are immune 
to crevice corrosion, the results of such studies conducted on Ti-2 are not applicable in 
the case of Ti-7 and Ti-16. 
Based on a constant film growth rate and, hence, corrosion rate, Shoesmith et al. (1997, p. 22) 
indicated that the concentration of hydrogen in the metal, HA, in g/mm3 can be calculated as a 
function of time of emplacement (t in years) from the expression: 
HA = 4(.Ti /103)R f M t Ti(d - Ruct)]-1 (Eq. 1)
h uc [ o 
where 
HA = hydrogen (g/mm content 3 ) 
.Ti = (g/cm Ti of density 3) = 5.4 (Weast 1978, p. B-177) 
fh = absorption for efficiency fractional 
R = rate (mm/year) corrosion passive general of 
uc 
t = years in t emplacemen of time 
MTi =Ti of mass atomic = 9.47 (Weast 1978, p. B-177) 
d0 = half of drip shield thickness 
The fractional efficiency for absorption, based on previous discussion, is fh = 0.015 for Ti-7. 
Considering a Ti-7 plate with 1 mm2 surface area, it is noted in Equation 1 that (1) the amount of 
hydrogen in grams produced by the general corrosion after t years of emplacement is 
4(.Ti/103)R t/MTi based on the reaction Ti + O H 2 . TiO2 + H 2 2 , (2) multiplication by 
uc 2 
the factor fh converts the produced hydrogen into absorbed hydrogen, and (3) the remaining 
volume of Ti-7 alloy in mm3 is represented by (d - Ruct) . The derivation of Equation 1 is
o 
based on a constant general corrosion rate. It was noted by CRWMS M&O (2000a, Sections 5.2, 
6.3, and 6.5.2) that the assumption of constant corrosion rate is conservative, and less 

conservative corrosion cases assume that the rate decays with time. It was also noted above that 
it is conservative to assume that all the hydrogen absorbed during general corrosion is retained 
within the remaining drip shield wall thickness as the general corrosion process proceeds. It is 
most likely that the majority of it will be removed by conversion to the more thermodynamically 
stable titanium oxide as corrosion progresses. 
Using the relationship H (g of H/g of Ti) = H (g of H/mm3 of Ti)(1000)/(10-6.Ti), Equation 1 
can be rewritten for HA in g/g as follows: 
HA = R f x10 4 ucM t Ti(d - Ruct)]
6h [-1 (Eq. 2)
o 
where 
HA = content hydrogen (g/g) 
The rate of general passive corrosion, Ruc, can be calculated from the rate of oxide film 
thickness, Roxby the following formula: 
Ruc = Rox (. /M )(. /MTi )-1 (Eq. 3)
ox oxTi 
where 
.ox = g/cm in oxide the of density 3 
Mox = oxide the of mass molecular 
Since the value of (. /M )(. /MTi )-1is always greater than unity, it is conservative to 
ox oxTi 
assume that Ruc = Rox . 
The maximum general corrosion rate (Ruc) reported in Section 4.1 has been used for the 
estimation of hydrogen concentration of the drip shield (see Section 6.2.3). 
6.1.6 Alternative Conceptual Model for Hydrogen-Induced Cracking Under General 
Passive Corrosion Conditions 
The present model, as described in Section 6.1.5, is based on the use of a general passive 
corrosion rate of titanium, which is constant over the lifetime of the drip shield. It is assumed 
that this process will be driven by the reaction of titanium with water to produce hydrogen and 
that a constant fraction of this hydrogen will be absorbed into the alloy and will eventually be 
precipitated as hydride (Assumptions 5.3 and 5.5). 
Since measured corrosion rates (taken from LTCTF measurements at LLNL) are low, absorbed 
hydrogen will diffuse deeply into the alloy rather than accumulate at the corroding surface to be 
released as the corrosion front progresses into the metal. As a consequence, hydride formation 

would occur uniformly throughout the wall thickness. Furthermore, it is assumed (Assumption 
5.4) that, once absorbed, hydrogen would not be released, leading to a predicted long term 
acceleration in the rate of hydrogen accumulation as the volume of available metal is reduced by 
conversion to oxide. Since failure by HIC could occur once a critical hydrogen concentration 
(HC) is achieved, this acceleration leads to a very rapid achievement of this critical value and, 
hence, a very conservative prediction of failure times by HIC. 
Key conservatisms adopted in this model are the following: 
1. 
Corrosion of titanium will be driven by interaction with water to produce absorbable 
hydrogen (Assumption 5.3). Since conditions will be generally oxidizing, reaction will 
be predominantly with dissolved oxygen, a process that does not produce hydrogen. 
2. 
Corrosion rates will be maintained constant over the lifetime of the drip shield 
(Assumption 5.4). 
3. 
Hydrogen release by the conversion of the metal containing hydrogen to oxide as the 
corrosion progresses is negligible (Assumption 5.4). 
An alternative, less conservative, and hence, more realistic conceptual model can be defined by 
relaxing conservatisms (2) and (3). The retention of the first conservatism is judicious, since 
titanium corrosion could be, at least partially, supported by water reduction in the concentrated 
saline environments possible as a consequence of evaporative concentration of seepage waters 
contacting the alloy surface. In these environments, the dissolved oxygen concentration would 
be low until temperatures cooled and salinity levels declined. This assumption is conservative 
since, with time, support of the general passive corrosion of titanium would be expected to shift 
from water reduction to oxygen reduction as the concentration of dissolved oxygen in the water 
increased with decreasing temperatures. As a consequence, the total amount of hydrogen 
produced, and hence the fraction absorbed, would decrease. 
In this alternative model, the corrosion rate is given by the passive oxide film growth rate as 
before but is taken to decrease with time as the oxide thickens and becomes more impermeable 
to hydrogen. The rate of hydrogen absorption is directly proportional to the corrosion rate and 
the amount of hydrogen absorbed, calculated in the identical manner to that used in the 
performance assessment model. This means that this new model incorporates the same 
efficiency factor for hydrogen absorption. Consequently, the new model will predict a decrease 
in hydrogen absorption rate with a time-dependency equal to that for the decrease in corrosion 
rate. 
As in the performance assessment (PA) model, the rate of diffusion of hydrogen in the titanium 
is taken to be rapid compared to the corrosion rate. Consequently, absorbed hydrogen will be 
dispersed throughout the full wall thickness; i.e., the concentration of hydrogen in the metal will 
be the same throughout. However, some hydrogen will be released as corrosion of the hydrogen-
containing alloy continues. The rate of release will be proportional to the corrosion rate and to 
the concentration of hydrogen that exists in the metal at that time. 

The criterion for failure remains the same as in the adopted model; i.e., failure could occur 
instantly once the amount of hydrogen in the material (HA) exceeds the critical value (HC). The 
value of HC remains unchanged. 
A schematic comparison of the original and the new conceptual models is given in Figures 2 and 
3. The terms are defined as follows; 
RCorrcorrosion rate of titanium
RHArate of hydrogen absorption into titanium
fhfractional efficiency of hydrogen absorption
RHRrate of absorbed hydrogen release due to continuing corrosion
n1time exponent for the decrease in corrosion rate with time, t
koxide film growth constant
JHflux of hydrogen in the metal at time t.
(HTi)tconcentration of hydrogen in the metal at time t.
HCcritical hydrogen concentration in the metal for failure by HIC.
Figure 2. A Schematic Representation of the Current Performance Assessment Conceptual Model for 
Hydrogen Absorption During Passive General Corrosion of Titantium (refers to the 
mathematical model described in Section 6.1.5) 

Figure 3. A Schematic Representation of the Alternate Performance Assessment Conceptual Model for 
Hydrogen Absorption During Passive General Corrosion of Titantium 
The expected decrease in corrosion rate with time is the accepted mechanism for passive film 
growth in exposure environments in which the metal cation has a very low solubility as is the 
case for titanium in neutral to alkaline solutions (Baes and Mesmer 1986). Such a decrease in 
corrosion rate has been observed for Alloy 22 in tests conducted in the LTCTF at LLNL 
(CRWMS M&O 2000b, Section 6.9.1), and a similar behavior is anticipated for titanium. 
Considerable electrochemical evidence exists to demonstrate that the rate of film growth 
decreases with time (Leitner et al. 1986; McAleer and Peter 1982; Nishimura and Kudo 1982; 
Beck 1982) according to the field-assisted ion-transport mechanism when subjected to a constant 
applied potential. While in these studies the potential driving oxide growth was applied 
electrochemically, a similar decrease in film growth rate with time would be expected to occur 

due to the open-circuit polarization of titanium by a soluble redox reagent in an aqueous solution 
(e.g., O2). 
The universally-accepted structure for passive films, an inner protective barrier layer covered by 
an outer more porous, and hence less protective, hydrated layer, has been demonstrated for 
passive films on Ti (Pan et al. 1994), and a number of features of passive oxide film growth, 
besides an increase in thickness, have been shown to lead to a decrease in corrosion rate with 
time. Since defect transport, or ion transport via defects, is the primary mechanism for both film 
growth (Macdonald 1999) and film dissolution (Blackwood et al. 1988), a defect annealing 
process (Leitner et al. 1986) will lead to a decrease in corrosion rate. Such a decrease has been 
observed in neutral (pH ~5 to 6) hydrogen peroxide solutions (Fonseca and Barbosa 2001). 
Studies in phosphate-buffered saline solutions (Pan et al. 1994) have shown that the 
incorporation of ionic species into the outer porous hydrated layer leads to a sealing, at least 
partially, of the porosity and an increase in film impedance, equivalent to a decrease in corrosion 
rate. Even under the aggressive conditions experienced in pulp and paper bleaching plants, the 
incorporation of Ca2+, Mg2+, and particularly SiO42-, into the oxide film have been shown to 
inhibit corrosion (Schutz and Xiao 1994; Whyllie et al. 1994). Based on these observations, it is 
reasonable to expect that the accumulation of silica on the surface will lead to a suppression of 
corrosion rate with time. 
The use of a hydrogen absorption rate that is directly proportional to the corrosion rate, and, 
hence, can be determined by multiplying the corrosion rate by a time-independent absorption 
efficiency, is very conservative. When a coherent oxide is present on titanium, hydrogen 
absorption would not be expected until the passive film became a degenerate semiconductor, a 
condition requiring the application of a potential = -0.6V (vs SCE) (Shoesmith and Ikeda 1997). 
Conditions that would allow hydrogen absorption, therefore, require that one of the following 
conditions apply: 
1. 
There is a source of polarization to a potential = -0.6V 
2. 
Destruction of the oxide occurs such that active conditions are achieved 
3. 
The presence in the alloy of intermetallics can act as hydrogen absorption windows in the 
otherwise impermeable oxide (CRWMS M&O 2000c, Appendix A). 
Since passive corrosion conditions prevail, condition (2) will not occur, and the cathodic 
polarization of titanium achieved by galvanic coupling to carbon steel sections within the 
repository drifts is modeled separately. Corrosion potential measurements on Ti-7 conducted in 
simulated Yucca Mountain groundwaters at General Electric Corporate Research and 
Development yielded a value of approximately -120 mV (vs. SCE) for immersion times greater 
than 2 days (BSC 2001, Section 3.3.3). Providing the potential remains within the band gap 
region of the TiO2 passive film (-0.6V to ~2.4V [vs. SCE]), then there is considerable published 
evidence to show that the oxide is an excellent barrier to hydrogen absorption (Been and 
Grauman 2000; Covington 1979; Shimogori et al. 1985). Clearly, a substantial period of dry 

oxidation (under ventilated conditions) would be expected to improve the shields resistance to 
hydrogen absorption. 
This influence of the passive film in suppressing hydrogen absorption is also clearly 
demonstrated in the electrochemical measurements of Okada (1983) from whose data the value 
of the absorption efficiency used in the model was adopted. In this model, a value of 1.5% has 
been adopted for the efficiency. In reality, the actual value should be considerably less, since 
Okadas value was measured under cathodically polarized conditions, when the specimen 
potential would be very much lower than the threshold value of ~-0.6 V (vs. SCE), and hence, 
beyond the band gap region of the oxide. 
Available evidence suggests that the presence of impurity Fe may increase the absorption of 
hydrogen by Ti-2 (Cotton 1970; Covington and Schutz 1981). Observed increases in hydrogen 
absorption by Ti-12 (Kim and Oriani 1987a) may also be attributed to the alloying additions, 
especially Ni and Mo. By contrast, no increase in hydrogen absorption was observed for Pd-
containing alloys providing the pH was greater than ~4 (Okada 1983). The key difference 
between these two categories is that the Ni in Ti-12 and the Fe (present as an impurity) in Ti-2 
can form reactive (i.e., capable of sustaining both anodic and cathodic reactions) intermetallics 
(Ti2Ni, TixFe) and/or -phase, whereas intermetallic formation is rare in Ti-7, and -phase 
formation does not occur. 
The electrochemical results of Okada (1983) show that there was no difference in the hydrogen 
absorption efficiency for Ti-2 specimens and specimens coated with noble metal (Pt) providing 
pH = 4 when a passive oxide will be present. This demonstrates that any tendency for the noble 
metal content of alloys Ti-7 and Ti-16 to catalyze hydrogen absorption (as observed under 
active, acidic conditions) (Fukuzuka et al. 1980) is masked when a passive oxide film is present. 
Based on this discussion, the assumption that the efficiency for hydrogen absorption is constant 
and directly proportional to the corrosion rate is clearly conservative. The adoption of a value 
measured under polarization conditions equivalent to galvanic coupling makes the value for the 
efficiency conservative, since there is no reason to fear that the Pd content of Ti-7 will influence 
this value when the material is passive; i.e. in the neutral to alkaline pH range anticipated under 
Yucca Mountain conditions. 
6.1.7 Comparison Between the Current Performance Assessment Conceptual Model 
and the Alternative Conceptual Model for Hydrogen Absorption by Titanium 
During General Passive Corrosion 
A comparison of the current PA conceptual model as discussed in Section 6.1.5 with the ACM 
discussed in Section 6.1.6 identifies the following key differences and commonalities: 
 
In the PA model, the passive corrosion rate is assumed constant with time, whereas in the 
ACM it is allowed to decrease with exposure time. 
 
In the PA model, all the hydrogen absorbed is retained and not re-released as the 
corrosion front progresses into the alloy. By contrast, in the ACM some of the absorbed 
hydrogen is re-released as the corrosion front progresses into the alloy. 

 
In both models, the rate of hydrogen absorption is assumed directly proportional to the 
passive corrosion rate and the proportionality constant is a time-independent hydrogen 
absorption efficiency. 
 
In both models, the rate of transport of hydrogen in the alloy is assumed to be very fast 
compared to its rate of absorption. As a consequence, hydrogen in the alloy is uniformly 
distributed throughout the full wall thickness. 
6.2 APPLICATION OF HIC MODEL TO DRIP SHIELD 
6.2.1 Material 
Titanium Grade 7 (Ti-7) (UNS R52400) is now being considered for construction of the drip 
shield for the waste package. As indicated in Section 1, this alloy consists of 0.3% Fe, 0.25% O, 
0.120.25% Pd, 0.1% C, 0.03% N, 0.015% H, and 0.4% total residuals with the balance being 
Ti. 
Both crevice corrosion and general passive corrosion will be accompanied by hydrogen 
production and, hence, possibly by the absorption of hydrogen into the metal. It has been 
concluded in Section 6.1.4 that hydrogen absorption during crevice corrosion will not be 
considered in the HIC model since crevice corrosion is essentially eliminated for Ti-7. 
General corrosion rate for Ti-16 has been reported in Section 4.1. As may be noted in Table 1, 
the calculated maximum value for Ruc is 2.82x10-4 mm/year. 
6.2.2 Determination of the Critical Hydrogen Concentration, HC 
As discussed in Section 6.1.3, Shoesmith et al. (1997, p. 11) concluded that a conservative value 
of HC = 500 g/g can be adopted as the critical hydrogen concentration in rolled plate material 
using Ti-2 and Ti-12 for predicting container lifetime. The lowest HC value observed for Ti-2 
and Ti-12 is ~400 g/g. Ikeda and Quinn (1998a) indicated that HC for Ti-16 is at least 1000 
g/g and may be much greater. 
The HC value for Ti-7 is assumed to be at least 1000 g/g. This assumption is based on 
Assumption 5.1 and data reported by Ikeda and Quinn (1998a, p. 7), which, as indicated earlier, 
concluded that the HC value for Ti-16 is between 1000 and 2000 g/g. This assumption is 
necessary because HC data are not available for Ti-7. 
It is noted that 400 g/g, the lower bound value observed for other titanium alloys (e.g., Ti-2 and 
Ti-12) by Shoesmith et al. (1997, pp. 7-11), was used as the critical hydrogen concentration for 
Ti-7 in REV 00 of this AMR. This value appears to be excessively conservative based on data 
reported by Ikeda and Quinn (1998a, p. 7). 

6.2.3 Determination of Hydrogen Concentration 
Based on Equation 2, the hydrogen concentration in the metal can be rewritten as follows: 
6 -1 
h [
HA = R f x10 4 ucM t Ti(d - Ruct)]
o 
where 
HA = content hydrogen (g/g) 
.
Ti= of density g/cm in Ti 3 = 5.4 (Weast 1978, p. B-177) 
fh = absorption for efficiency fractional
R= rate (mm/year) corrosion passive general of
uc 
t = years in t emplacemen of time
MTi =Ti of mass atomic = 9.47 (Weast 1978, p. B-177)
do = half of drip shield thickness
More information about the model parameters are provided in Table 2. 
Table 2. Model Inputs Used in Current Performance Assessment Model for Hydrogen Absorption During 
Passive General Corrosion of Titanium 
Input Name Input Description Input Source (DTN if 
applicable) 
Value or Distribution 
(Units) 
fh Fractional efficiency for 
hydrogen absorption 
Assumption 5.5 0.005  0.015 
.Ti Density of Ti Weast 1978, p. B-177 4.5 g/cm3 
Ruc Rate of general passive 
corrosion (maximum) 
DTNs: 
MO0003SPASUP02.003 
LL990610605924.079 
3.88x10-4 mm/year 
1.76x10-4 mm/year 
do Half of drip shield 
thickness 
BSC 2003 7.5 mm 
t Time of emplacement 10 CFR 63.114(d) 10,000 years 
MTi Atomic mass of Ti Weast 1978, p. B-177 47.9 
As seen in Table 2, the fractional efficiency for hydrogen absorption is fh = 0.015 for Ti-7. The 
rate of general passive corrosion is Ruc= 2.82x10-4 mm/year (maximum value) based on 
Sections 4.1 and 6.2.1. The time of emplacement is t = 10,000 years. Half of the minimum wall 
thickness (15 mm) is used for do, i.e., do = 7.5 mm. Using half of the thickness of the drip shield 
ANL-EBS-MD-000006 REV 01 34 February 2003 

in calculations is equivalent to assuming that there is twice the surface area of Ti-7 available to 
absorb hydrogen because absorption can occur under as well as on top of the drip shield. 
Conservative Estimate: 
Ruc = 2.82x10-4 mm/year 
fh = 0.015 
do = mm 5.7 
t = years 000,10 
From Equation 2, HA = 755 g/g < HC = 1000 g/g. 
6.2.4 Results 
The analytical estimate presented in Section 6.2.3 based on Equation 2 indicated that there exists 
a big margin of safety for the drip shield against the effects of HIC. The hydrogen concentration 
in the drip shield at 10,000 years after emplacement is 755 g/g resulting from a conservative 
estimate. The estimated hydrogen concentration is less than the critical hydrogen concentration 
of 1000 g/g for Ti-7. 
As discussed in Section 6.1.5, the current model is highly conservative in that it assumes all the 
hydrogen absorbed during general corrosion is retained within the remaining drip shield wall 
thickness as the general corrosion process proceeds. It is most likely that the majority of it will 
be removed by conversion to the more thermodynamically stable titanium oxide as corrosion 
progresses. With this conservative assumption, the hydrogen concentration in the drip shield 
increases continuously with time as the drip shield thickness (thus the volume of the remaining 
drip shield) is reduced by general corrosion. This hydrogen concentration increase from the 
thinning of the drip shield is in addition to the hydrogen pick-up from the general corrosion. 
With the current conservative model, the time for the hydrogen concentration in the drip shield to 
exceed the critical hydrogen concentration will always be before the time to failure of the drip 
shield by general corrosion. 
As noted in Section 6.1.3, crack growth in the drip shield by HIC requires embrittlement of the 
drip shield (measured by the hydrogen concentration exceeding the critical hydrogen 
concentration) and applied stress. When the drip shield is subject to HIC, the failure will be by 
through-wall cracks, and the cracks are likely to be those generated by SCC. However, those 
cracks are self-limited and will be eventually plugged by corrosion products and scale deposits 
(CRWMS M&O 2000d, Section 6.5.5), especially calcite from available CO2 and silicates from 
ground waters (Section 6.3.2). Therefore, even when the drip shield is breached by through-wall 
cracks induced by HIC, the intended design function of the drip shield (i.e., preventing dripping 
water from directly contacting the underlying waste package) will not be compromised 
significantly from the HIC failure. 

6.2.5 Parameter Uncertainties 
Uncertainties and Impacts on Model outputKey input parameters in the model for hydrogen 
absorption during passive corrosion of titanium are Ruc, fh and HC. 
Ruc was measured at the LTCTF of the LLNL using calibrated instruments and following quality 
assurance procedures. Therefore, the data used to determine Ruc is qualified from origin. The 
maximum value of Ruc (see Sections 4.1) was used for a conservative estimate of hydrogen 
concentration in the drip shield. 
The constant value of fh = 0.015 adopted in the model was determined under extremely 
conservative conditions unachievable in the repository, and, therefore, constitutes a conservative 
upper limit for this parameter (Assumption 5.5). Consequently, the model predictions obtained 
using this value will also be conservative. Adoption of a value of fh > 0.015 (i.e., the value of 0.1 
as discussed in Section 6.1.5) is not merited because it applies only to Ti-12, an alloy containing 
 phase. Since Ti-7 is an a-alloy, which does not undergo crevice corrosion, the adoption of this 
higher value of 0.1 is not merited. 
The value of 1000 g/g adopted for HC is a conservative lower limit as discussed in Section 
6.1.3. Measurements show the value could be as high as 2000 g/g. Adoption of a higher value, 
or a value distributed between 1000 g/g and 2000 g/g, would further decrease the probability 
of failure by HIC before 10,000 years. 
The design parameters, d0 and t, are certain. Also, MTi and .Ti are physical constants and 
therefore can be considered certain. 
Sources of UncertaintiesA principal source of uncertainty is measurement error, particularly 
in the parameters Ruc and HC. Another significant uncertainty is the lack of measurements on 
Ti-7, leading to the adoption of parameter values measured on the similar alloy Ti-16. The small 
database of measurements on the fracture toughness of Ti-7 (Ti-16) and the absence of 
information on the time-dependent decrease in the value of fh have lead to the adoption of 
limiting conservative values for these parameters. This last uncertainty includes the uncertainty 
associated with variability in materials properties. 
6.3 HYDROGEN INDUCED CRACKING OF TI-7 DUE TO GALVANIC COUPLING 
6.3.1 Introduction 
As indicated in Section 6.1.2, the three general conditions that must exist simultaneously for the 
hydrogen embrittlement of a alloys are: 
 
A mechanism for generating hydrogen on a titanium surface 
 
A metal temperature above approximately 80C (175F) 
 
A solution pH less than 3 or greater than 12, or impressed potentials more negative than 
-0.7 V (SCE). 

In accordance with Schutz and Thomas (1987, p. 673), one hydrogen-generation mechanism is 
the coupling of active metals such as zinc, magnesium, and aluminum to titanium leading to 
hydrogen uptake and eventual embrittlement of the titanium provided the other two conditions 
described above are met. A similar possible problem occurs when titanium is in galvanic contact 
with carbon steels. In a drip shield design without backfill, hydrogen generation may be caused 
by the galvanic couple between the titanium drip shield surface and ground supports (such as 
rock bolts, wire mesh, and steel liners used in the drift), which may fall onto the drip shield 
surface. Therefore, HIC or hydrogen embrittlement of the drip shield due to galvanic couple can 
not be ruled out since the conditions associated with metal temperature and water or moisture pH 
(or corrosion potential) as described above are certainly attainable under Yucca Mountain drip 
shield conditions. 
6.3.2 Qualitative Assessment 
Given the expected evolution of ground waters potentially contacting the drip shield and the 
temperature regime within the repository (CRWMS M&O 2000a, Section 1.3), the required 
conditions for hydrogen absorption by the titanium drip shield are clearly present when 
galvanically coupled to sections of the steel components. If occurring while temperatures are 
high (= 80C) and concentrated ground waters are present, then local hydrided hot spots are 
possible. These hot spots will only occur at those sites where both contact to carbon steel and 
the establishment of lasting (at least periodically) aqueous conditions are achieved. Their 
formation is likely to be promoted if the contact points coincide with abraded or scratched areas 
of the shield caused by the impact of collapsing sections of the ground support structure. 
At higher temperature, hydrogen will be more rapidly absorbed and transported into the bulk 
structure to produce the hydride distribution required for extensive crack propagation. It is less 
likely that a hydride layer will be retained at the surface. As the temperature falls, the formation 
of surface hydrides becomes more likely as the rate of hydrogen transport into the bulk of the 
metal decreases. The formation of surface hydrides tends to reduce the efficiency of subsequent 
hydrogen absorption (Noel et al. 1996), and their presence has little effect on structural integrity. 
The efficiency of Fe-Ti galvanic couples to cause hydrogen induced cracking of the titanium drip 
shield will be limited for a number of additional reasons: 
1. 
The contact area is likely to be small and the anode to cathode area (area of steel and 
titanium, respectively) low. Since relatively low volumes of groundwater are likely to 
contact both metals simultaneously, anode to cathode areas close to one seem likely. If a 
couple with a small anode to cathode area ratio was established (i.e., a small piece of 
steel in contact with a large area of drip shield), then one would expect the couple to be 
more rapidly exhausted as the steel is consumed. Under these conditions, the amount of 
hydrogen absorbed would be limited. Future quantitative studies may be conducted to 
determine how much hydrogen would be liberated per unit mass of Fe, whether that 
much hydrogen could cause HIC in the adjacent Ti, and how fast hydrogen diffuses into 
the surrounding Ti to keep hydrogen low. 

While temperatures are high (> 80C), the intermittent nature of seepage dripping onto 
the drip shield should lead to only limited periods of the aqueous conditions required to 
sustain an active galvanic couple, thereby limiting hydrogen absorption while 
temperatures are high enough to drive hydrogen transport into the metal. For 
temperatures below 80C, even galvanic polarization below the potential threshold of 
-0.6 V (vs. SCE) produces only innocuous surface hydride films (Schutz and Thomas 
1987). Also, at these lower temperatures, the small amounts of dissolved O2 in the 
solutions forming the galvanic couples will make it very difficult to achieve polarization 
of the galvanic potential to less than -0.6 V (vs. SCE). Additionally, the intermittent 
wetting and drying cycles anticipated on the drip shield will lead to the ready formation 
of calcareous and mineral deposits, which are well known to dramatically suppress 
galvanic currents, thereby stifling hydrogen absorption (Lunde and Nyborg 1993). 
2. 
Conditions in the repository will be oxidizing, making it less likely that the couple will 
sustain water reduction, and hence hydrogen absorption. However, at high temperatures 
in concentrated saline solutions, the amount of O2 dissolved in the solution forming the 
couple will probably be too low to displace water reduction as the primary cathodic 
reaction. However, the ferrous ion product of steel dissolution will be homogeneously 
oxidized to ferric species by dissolved O2. If conditions remain neutral, this should lead 
to the formation of insoluble Fe(III) oxides/hydroxides, and little influence would be 
exerted on the galvanic couple. However, any tendency for acidification or the 
development of alkaline conditions will increase the ferric ion solubility. Under 
evaporative conditions, this could lead to quite high dissolved ferric ion concentrations 
and the establishment of a galvanic potential sufficiently positive to avoid hydrogen 
absorption into the titanium. It is well documented that only parts per million 
concentrations of multivalent transition metal cations such as Fe(III) are required to 
polarize titanium to passive conditions (Covington and Schutz 1981; Schutz and Thomas 
1987) and ferric oxide deposits are passivating to Ti (Covington and Schutz 1981). Also, 
experimental evidence exists to show that galvanic currents (Hodgkiess et al. 1987) and 
the rate of hydrogen absorption (Lund and Nyborg 1993) will decrease with time as 
deposits (calcite from available CO2 and silicates from groundwaters) accumulate (Lund 
and Nyborg 1993; Hodgkiess et al. 1987). 
3. 
In the proposed repository environment, both the Ti drip shield and steel component 
surfaces will experience a considerable period of dry high temperature (= 85C 
depending on whether the higher- or lower-temperature operating mode is adopted). This 
will leave both Ti and steel passivated (especially Ti) and avoid galvanic contact. While 
passivity on the steel may be subsequently disrupted, loss of passivity of the Ti drip 
shield will not be so readily achieved. Experimental evidence exists to show that air 
oxidation prevents hydrogen absorption by Ti even in aggressive (0.5% to 6% HCl) 
solutions at elevated temperatures (70C to 250C) (Mon 2002, Attachment III). 
4. 
Titanium has a large tolerance for hydrogen (see Section 6.1.3 for critical hydrogen 
concentration), and substantial concentrations must be achieved before any degradation 
in fracture toughness is observed. This concentration level, as indicated in Section 6.1.3, 
has been measured to be in the range of 400 to 1000 g.g-1 for Ti-2 and Ti-12. Recent 

measurements suggest that the tolerance for hydrogen of the Ti-16 (a titanium alloy very 
similar to Ti-7 in chemical composition) may be in the 1000 to 2000  g.g-1 range (Ikeda 
and Quinn 1998a, p. 7). According to Section 6.1.3, the critical hydrogen concentration 
for Ti-7 is 1000  g.g-1. 
Given the high critical hydrogen concentration, the large volume of available titanium in the drip 
shield into which absorbed hydrogen can diffuse, and other reasons stated above, hydrogen 
embrittlement of the Ti drip shield is highly unlikely. 
6.3.3 Mathematical Model for Hydrogen Absorption and Diffusion in Drip Shield 
As described in Section 6.3.2, hydrogen absorbed in the Ti-7 drip shield due to galvanic coupling 
with carbon steel may diffuse to the rest of the drip shield. A mathematical model is proposed in 
this section to predict the hydrogen concentration in the drip shield due to a galvanic couple 
between the drip shield and a carbon steel segment. The absorption and diffusion of hydrogen in 
the drip shield is a complex process. A mathematical model is possible only if a number of 
necessary assumptions are adopted. These assumptions (also mentioned in Section 5.7) are as 
follows: 
 
The drip shield is treated as an infinite plate. The thickness of the plate is denoted by H. 
 
The area of galvanically coupled wetted surface is considered circular in shape. While 
the actual geometry of the wetted area will undoubtedly have a more complex geometry, 
its representation by an equivalent circular area will not significantly alter the 
calculations described below. The radius of this circular wetted area is denoted by ro. 
 
Hydrogen is absorbed into the drip shield through the contact plane and immediately 
reaches the other surface. The initial region of the drip shield with hydrogen absorption 
has the shape of a circular disk with a radius ro and thickness H. 
 
As hydrogen starts to diffuse in the drip shield, the radius of the circular region will 
expand at a constant speed v. As a result, the radius r(t) of the circular disk at a time t is 
expressed by the following equation: 
r(t) = ro + v t (Eq. 4) 
The amount of hydrogen (in g or mg, for example), Q(t), absorbed in the drip shield at a time t, 
therefore, is: 
= t t()=. . p ro ) Ht + ... [ 2p r()] t H t - ) dr(t) (Eq. 5)
t Q ( Ti2 
Ti t(
= t 0
where . is the hydrogen absorption rate, in ppm (or  g/g) /day, for example, and . Ti is the mass 
density of the Ti-7 drip shield (see Section 6.1.5). 

Based on the reaction: 2Fe + 3H2O . Fe2O3 + 3H2, the maximum amount of hydrogen, formed 
by corrosion of the carbon steel, that can be absorbed by titanium is: 
Q = fFe W3 /MFe (Eq. 6)
max 
where W is the mass of the carbon steel segment, fFe is the fraction of hydrogen produced by
oxidation of the total mass of carbon steel, and MFe is the atomic mass of Fe.
If tmax is the time when Q(t) reaches Qmax, tmax can be obtained by solving for the following
equation: 
= t tmax
2
Q . . = Ti pro )Ht + ... [2pr()]t H t - )dr(t) (Eq. 7)
max ( max Ti t( max 
= t 0 
Using Eq. 4, Eq. 7 can be written as: 
23
.
. 2
Q .. = Ti pr H o2t + rovt + 
v t max .. 
(Eq. 8)
max . max max 
3 .
.. 
The maximum hydrogen concentration HA(ro) in ppm (or g/g), for example, will be in the 
region of the drip shield beneath the initial contact area and at the time tmax, i.e., 
HA () . = t (Eq. 9)
ro 
max 
The hydrogen concentration HA(r) tends to be reduced at locations with a radius r from the center 
of the initial contact area, i.e., 
r - ro .(
HA () . =.. 
t -. for ro = r = t r max ) (Eq. 10)
r max
. v . 
At t > tmax, HA(r) will start to decline as the hydrogen diffusion continues and the source of
hydrogen exhausts.
As a numerical example, the following input data are considered:
H = Ti-7 plate thickness = 15 mm 
ro = radius of the initial contact area = 50.8 mm (2 in.) 
.Ti = density of Ti-7 = 4.5x106 g/cm3 
fFe = fraction of hydrogen available for absorption by Ti-7 drip shield = 0.015 (refers to 
Assumption 5.5) 

W = mass of carbon steel = 22.727x103 g (50 lb.) 
MFe = atomic mass of Fe = 55.847 (Weast 1978, p. B-177) 
The corrosion potential of carbon steel is estimated to be about -0.6 V (SCE), and a galvanic 
couple would polarize Ti down to that level. Based on Shoesmith et al. (1995, Figure 19) and 
Murai et al. (1977, Figure 5), the hydrogen absorption rate, ., for Ti-7 would be about 0.5 
g/g/day. Inspection of Figure 5 in Murai et al. (1977) shows that the maximum value of . 
measured in the potential range -0.6 V to 1.0 V (SCE) is ~0.7 g/g/day. Therefore, this value is 
adopted as the maximum . value in this sensitivity study. No data are available for the hydrogen 
diffusion rate, v, of Ti-7. For the base case of the numerical example, the values used for . and v 
are, respectively, 0.5 g/g/day and 1mm/day. For the parametric study, . = 0.5, 0.7 g/g/day, 
and v = 1, 2, 3 mm/day are considered, and plots of hydrogen concentration versus distance from 
the original contact area are shown in Figure 4. It can be seen that, in all cases of the parametric 
study, the hydrogen concentration does not exceed the critical hydrogen concentration of 1000 
g/g for Ti-7 by a rather comfortable margin. The results of the parametric study appear to be 
consistent with the qualitative assessment presented in Section 6.3.2. The choice of the 
mathematical model and the value of . appear to be reasonable and conservative. The value of v 
is high but would also appear to be reasonable, since it ignores the probability that the hydrogen 
will be retained at the Ti surface and the efficiency of absorption would then decrease. The 
calculation can be considered bounding since all the features likely to suppress hydrogen 
absorption are ignored in the calculation. These features are summarized above in section 6.3.2. 

500
450
400
350
300
250
200
150
100
50 
0 
0 200 400 600 800 1000 1200 
H Concentration (ppm) 
/
/
/
/
/
/
0.5 ppmday & 1 mm/day 
0.5 ppmday & 2 mm/day 
0.5 ppmday & 3 mm/day 
0.7 ppmday & 1 mm/day 
0.7 ppmday & 2 mm/day 
0.7 ppmday & 3 mm/day 
Distance (mm) 
Figure 4. Hydrogen Concentration in Drip Shield Versus Distance from Contact Area 
6.3.4 Parameter Uncertainties and Sensitivity Analyses 
Uncertainties and Impacts of Uncertainties on Model OutputsThe key parameters in this 
model are the mass of steel (W), the radius of the initial contact area (ro), the fraction of 
hydrogen absorbed (fFe) , the critical hydrogen concentration, HC, the hydrogen absorption rate 
(.) and the hydrogen diffusion rate (v). 
The values of W and ro adopted will both influence the analysis but include uncertainties that 
cannot be quantified. The value of W = 22.727 kg (50 lbs) for the mass of steel available for 
consumption in a single galvanic couple is a conservative limit, since it ignores the stifling of 
coupled conditions by the evolving local environment at the coupled site, as discussed in Section 
6.3.2. 
The choice of a value of ro ~ 50.8 mm (2 in.) is a realistic but uncertain value. The size of the 
contact area will be decided by the amount and distribution of aqueous solution trapped at the 
carbon steel/Ti-7 contact site, not the actual metal to metal contact area. The loss of wet contact 
by evaporation, gravitational flow and capillary absorption into accumulating corrosion products 

will be the primary influence limiting ro. Although it is difficult to quantify it is unlikely that a 
large uncertainty is associated with the chosen value. 
The value of ., the hydrogen absorption rate, is a very conservative upper limit since it ignores 
the observed decrease in absorption rate with time as discussed in Section 6.3.5. As a 
consequence, the model will predict a conservative rapid accumulation of hydrogen. To be 
consistent with this constant high rate of hydrogen absorption it was necessary to also assume a 
fast rate of hydrogen transport (v) within the Ti-7 (Assumption 5.7). The values of v chosen 
were arbitrary but much higher than the transport rates that would be calculated from measured 
diffusion coefficients of hydrogen in Ti. This conservatism is partially offset by allowing lateral 
hydrogen transport away from the contact area. 
The conservative natures of the parameters HC and fFe (equivalent to fh) have been discussed in 
Section 6.2.5. 
Sources of UncertaintiesAny event sequence(s) causing variations in the mass of the steel 
anode, and the anode and cathode areas will alter the model output and serve as sources of 
uncertainty. 
6.3.5 Alternative Conceptual Model For HIC Under Galvanically-Coupled Conditions 
An additional source of hydrogen that could lead to its absorption by the drip shield and, hence, 
possibly to shield failure by HIC, is wet contact between the drip shield and carbon steel ground 
supports components (e.g., rock bolts, wire mesh and steel drift liners) within the repository 
drifts. 
In the PA model, as noted in Section 6.3.3, the galvanic coupling supports corrosion of the 
carbon steel according to the reaction 
2Fe + 3H2O  Fe2O3 + 3H2 
and a certain fraction of the hydrogen, produced on the drip shield surface, is then absorbed into 
the titanium. This fraction, expressed as an absorption efficiency, is the same as that absorbed 
under non-galvanically-coupled conditions and used in the model for hydrogen absorption during 
general passive corrosion. This is appropriate since the absorption efficiency was measured 
under cathodically polarized conditions akin to those which will prevail when the titanium is 
galvanically coupled. While this makes the use of this efficiency very conservative for general 
passive corrosion, it is more realistic for galvanically-coupled conditions. 
Once absorbed, the hydrogen rapidly diffuses through the full wall thickness of the shield and, 
with time, diffuses laterally within the drip shield wall to adjacent locations not in direct galvanic 
contact with the steel. Corrosion of the carbon steel, and hence hydrogen absorption by titanium, 
stop once all the available contacting steel has been consumed. The criterion for failure is the 
same as that used in the general passive corrosion model, i.e., failure by HIC is instantaneous 
once the concentration of hydrogen in the titanium exceeds the critical value, HC. 

This model contains a number of key conservatisms mentioned below: 
1. 
The possible extent of galvanic corrosion will be limited by the small area of contact 
between the two metals, the low cathode-to-anode ratio (Ti/Fe) which is likely to 
approach one, the low probability that the point of galvanic contact will simultaneously 
experience seepage drips, the inability of the site to remain wet due to run-off, and the 
stifling of steel corrosion by the accumulation of iron oxide corrosion products. In view 
of these limitations, the assumption that a considerable mass of steel (up to 50 lbs) could 
be consumed by galvanic corrosion is very conservative (Assumption 5.6). 
2. 
Once absorbed, hydrogen will diffuse rapidly within the metal (Assumption 5.7, third 
bullet). However, practically measured diffusion coefficients for hydrogen in a-titanium 
are low and thermally-activated in the normal Arrhenius manner (Phillips et al. 1974). 
3. 
The absorption efficiency for hydrogen remains constant throughout the duration of the 
galvanic corrosion process (Assumption 5.5). Considerable evidence exists to show that, 
under galvanically coupled conditions, a surface hydride layer is formed (Noel et al. 
1996; Mon 2002, Attachment II; Tomari et al. 1999) and, as a consequence, the 
efficiency for hydrogen absorption decreases substantially with time. 
An alternative, less conservative, and hence, more realistic model can be defined by 
incorporating a flux term for hydrogen absorption in the alloy and using a time-dependent 
absorption efficiency. This is equivalent to relaxing the conservatisms in (2) and (3) above. 
The rate of hydrogen production is taken as constant and equal to the rate of carbon steel 
corrosion until the amount of available steel is consumed, after which no further hydrogen 
absorption due to galvanic coupling occurs. However, since a considerable literature exists to 
show that the rate of hydrogen absorption follows a parabolic relationship (e.g., Tomari et al. 
1999), the absorption efficiency is taken to decrease with time. This acknowledges that the 
amount of hydrogen that can absorb is limited to the amount required to form a surface hydride 
layer. 
The flux of hydrogen within the material is calculated using Ficks laws with a diffusion 
coefficient, which is dependent on temperature according to Arrhenius Law as shown in Figure 
6 as JH. As a consequence, a steady-state is eventually established in which the rate of hydrogen 
absorption into the surface of the titanium is controlled by the temperature-dependent flux of 
hydrogen from the surface into the bulk of the material. Lateral diffusion of hydrogen to sites 
not in direct galvanic contact with the titanium is ignored. 
A schematic comparison of these two models is given in Figures 5 and 6. This comparison 
emphasizes the conservatisms in the adopted PA model compared to the alternative conceptual 
model. The additional terms are defined as follows: 
(RCorr)CScorrosion rate of carbon steel galvanically-coupled to titanium 
(fh)tthe time-dependent fractional absorption efficiency for hydrogen in to titanium 
n2time exponent for the decrease in (fh)t as a function of time 
DHtemperature-dependent diffusion coefficient of hydrogen in a-titanium 

D0diffusion coefficient of hydrogen in a-titanium at 25C 
.c/.xconcentration gradient of hydrogen in the titanium defined by Ficks laws 
EA#activation energy for hydrogen diffusion in a-titanium 
Vmvolume of titanium containing hydrogen 
Rgas constant 
Ttemperature. 
NOTE: CS = carbon steel. 
Figure 5. A Schematic Representation of the Current Performance Assessment Conceptual Model for 
Hydrogen Absorption (HA) During Galvanic Coupling of Titanium (refers to the mathematical 
model described in Section 6.3.3) 

NOTE: CS = carbon steel. 
Figure 6. A Schematic Representation of the Alternate Performance Assessment Conceptual Model for 
Hydrogen Absorption During Galvanic Coupling of Titanium 
ANL-EBS-MD-000006 REV 01 46 February 2003 

A number of studies have been published which clearly show the rate of hydrogen absorption 
decreases parabolically with time (Kim and Oriani 1987a; Noel et al. 1996; Phillips et al. 1974; 
Tomari et al. 1999; Foroulis 1980). Since all of these studies were conducted under cathodically 
polarized conditions akin to, or much more polarized than, those anticipated as a consequence of 
galvanic coupling to carbon steel, the use of a decreasing absorption efficiency in our model is 
appropriate. In all cases the storage of hydrogen in the surface of the titanium was clearly 
demonstrated by the identification of a surface hydride layer. Since the majority of these studies 
were performed in acidic solutions with large applied current densities, the hydride layers were 
much thicker (many microns) than would be anticipated for galvanic coupling conditions under 
Yucca Mountain conditions. 
Tomari et al. (1999) and Phillips et al. (1974) showed that the parabolic relationship for 
hydrogen absorption possessed a t1/2 time-dependence consistent with the form adopted in the 
model; i.e., the hydrogen absorption process is controlled by the transport of hydrogen from the 
surface layer into the bulk of the alloy. The presence of the surface hydride, or the storage of 
hydrogen in a concentrated surface layer, was confirmed optically (Noel et al. 1996; Phillips 
et al. 1974), by glow discharge spectroscopy (Mon 2002, Attachment II) and by SIMS depth 
profiling (Tomari et al. 1999). Glow discharge spectroscopy was also used to show the change 
in distribution of hydrogen in the alloy as a consequence of transport over time. 
Tomari et al (1999) showed that Ti-2 and Ti-17 (Ti-16 with a slightly higher oxygen 
concentration) exhibited effectively identical parabolic relationships for hydrogen absorption, 
indicating that the Pd alloy content had no measurable influence on either hydrogen absorption 
or its subsequent transport within the alloy. This is consistent with the conclusion (above) that 
any possible influence of the Pd content of the titanium is masked under passive conditions. 
Based on similar studies performed as a function of temperature by Phillips et al. (1974) and 
Mon 2002, Attachment II), a diffusion coefficient for hydrogen transport in titanium (2 4 x 10-12 
cm2/s) and an activation energy of 60  3.4 kJ/mol were determined. 
Based on this discussion, the use of a time-dependent hydrogen absorption efficiency, which 
decreases with time, is appropriate. The use of a constant absorption efficiency in the PA model 
makes that model conservative by comparison to this alternative model. The adoption of a t-1/2 
time dependency for this efficiency is also justified. Since well-documented values exist for the 
diffusion coefficient of hydrogen in titanium and the activation energy for this diffusion process, 
the form of the model is appropriate. 
6.3.6 Comparison between the Current Performance Assessment Conceptual Model 
and the Alternative Conceptual Model for Hydrogen Absorption During 
Galvanic Coupling 
A comparison of the current PA conceptual model for hydrogen absorption as discussed in 
Section 6.3.3 with the ACM discussed in Section 6.3.5 reveals the following commonalities and 
differences: 
 
In both models, the rate of hydrogen absorption is directly proportional to the rate of 
carbon steel corrosion. 

 
In both models, the extent of carbon steel corrosion is limited by assuming a limited mass 
of steel is available by galvanic coupling. 
 
In both models, an arbitrary, but reasonable, area of wetted contact between the steel and 
the Ti-7 is assumed. 
 
In the PA model, the efficiency for hydrogen absorption is assumed to be constant with 
time, whereas, in the ACM model it is allowed to decrease with time as hydrogen 
accumulates in the surface of the Ti. 
 
In the PA model, the absorbed hydrogen is rapidly transported through the full wall 
thickness. In the ACM model, hydrogen diffuses through the Ti-7 at a rate controlled by 
the measured temperature dependent diffusion coefficient. 
 
In the PA model, hydrogen is allowed to diffuse laterally along the wall of the drip shield. 
In the ACM model lateral diffusion is not allowed. 
6.3.7 Worst Case Considerations 
As a worst case scenario, it is considered that the majority portion of the hydrogen generated by 
the section of carbon steel will be absorbed by the drip shield and that the absorbed hydrogen 
will be retained in the vicinity of the contact area and not be diffused to the rest of the drip 
shield. In this case, there may be a local embrittlement of the drip shield in the contact area. It is 
further considered that the contact area is impacted by a rock fall, and residual stress and a crack 
are initiated. As a result, a fast brittle fracture is possible if the stress intensity factor calculated 
from the residual stress and the crack size exceeds the fast fracture stress intensity factor KH, 
which tends to have a lower value than the critical slow crack growth stress intensity factor KS 
(also called fracture toughness) as indicated in Figure 1. 
Calculated stress intensity factors due to rock fall in the drip shield are listed in CRWMS M&O 
(2000d, Attachment II) for various crack sizes (depths). Based on CRWMS M&O (2000d), an 
initial crack depth of 50 m (i.e., the maximum incipient crack size, based on CRWMS M&O 
2000d, Section 6.5.2) can be considered. The equivalent stress intensity factor, on the basis of 
CRWMS M&O (2000d, Attachment II), is < 5 MPam1/2. From CRWMS M&O (2000d, Figure 
27), it can be seen that the fracture toughness KS for Ti-7 should be at least 30 MPam1/2. Noting 
KH ~ 1/3 KS from Figure 8 of Clarke et al. (1994), brittle fracture of the drip shield would require 
a value of KH = 10 MPam1/2. Since this value of KH is greater than 5 MPam1/2, brittle fracture 
due to hydrogen embrittlement would not be a problem for the drip shield subjected to galvanic 
coupling with a section of the carbon steel ground support. 
Even though fast brittle fracture does occur and the surface crack turns into a through-wall crack, 
the length of the crack will still be limited because the surrounding area of the drip shield is not 
embrittled by hydrogen. The crack opening will eventually be plugged by corrosion product 
(CRWMS M&O 2000d, Section 6.5.5). 

7. MODEL VALIDATION
Models described in this AMR are expected to predict accurately and precisely the hydrogen 
concentration in the Ti-7 drip shield for a period of at least 10,000 years. This extraordinarily 
long time factor makes it extremely difficult to validate these models in the usual way, i.e., by 
comparison of model predicted values with those observed experimentally for the whole range of 
time (ASTM C 1174, Sections 19.3 and 20.4). Consequently, a different approach was adopted 
in the TWP (BSC 2002) for the validation of these models. According to this new approach, 
these models were validated by validating the input parameter values used and comparing these 
parameters and model predictions to available peer reviewed and qualified project data. The 
medium level of confidence in the titanium HIC models, as mentioned in the TWP, was obtained 
by validating three key model parameters, e.g., the passive corrosion rate, the critical hydrogen 
concentration, and the hydrogen absorption efficiency. In order to build further confidence in 
these models, the model output values were corroborated with those available in the peer 
reviewed scientific literature. In addition, four criteria were developed and documented in the 
TWP (BSC 2002) to ensure that the required level of confidence in these models for the models' 
stated purposes has been achieved. They are: (1) the critical hydrogen concentration of the 
grade-7 titanium alloy is consistent with the literature; (2) the general corrosion rates used in the 
model are consistent with the literature; (3) a comparison of the rates of hydrogen absorption 
calculated by the model to the calculated fluxes of absorbed hydrogen within the alloy 
demonstrates that the assumption that all hydrogen absorbed remains in the alloy is conservative; 
and (4) the hydrogen contents for grade-7 titanium and similar alloys predicted by the model are 
consistent with the literature. A detailed description of the validation of HIC models in light of 
these criteria is given below. 
7.1 VALIDATION OF MODEL FOR HYDROGEN ABSORPTION DURING 
GENERAL PASSIVE CORROSION OF TITANIUM 
According to TWP (BSC 2002), model validation will involve meeting the selected criteria as 
discussed below: 
Criterion 1 
Is the critical hydrogen content (HC) of the Grade-7 titanium alloy (Ti-7) used by the model 
consistent with the results of mechanical tests published in the peer-reviewed literature? 
Using the slow strain rate technique on precracked compact tension specimens precharged with 
known amounts of hydrogen, it has been shown that the fracture toughness of Grades -2, -12, and 
16 (Ti-2, Ti-12, Ti-16) are not significantly altered until their hydrogen levels exceed critical 
values 400 g/g, 400 g/g, and >1000 g/g, respectively. It should be noted that these values are 
lower limiting values, not mean values and, hence, are conservative. These values, and the 
details of their measurement, are published in the peer-reviewed literature (Clarke et al. 1994, 
1997; Shoesmith et al. 2000; Ikeda et al. 2000). The model assumes that, for Ti-7, the minimum 
HC value of 1000 g/g is realistic. While this realistic value was measured on Ti-16, there is no 
reason to expect that Ti-7, with a higher Pd content than Ti-16, will possess a value of HC lower 
than1000 g/g. 

The value of HC is the hydrogen concentration above which slow crack growth is no longer 
observed in these tests and only fast crack growth occurs. It is sensitive to the microstructure 
and texture of the material with respect to the orientation of the crack and the applied stress. The 
hydride precipitation in the a-phase, which for HA > HC can lead to the decrease in fracture 
toughness, is favored close to the basal planes (Clarke et al 1997, p. 1550). For Ti-12, -phase 
laminations are introduced into the material through the manufacturing process (rolling) leading 
to a preferential pathway for cracking along these laminations. These can be easily seen in the 
metallographic cross sections presented in Clarke et al. (1994). Thus, the measured value of HC 
for cracks propagating along these laminations is ~400 to 600 g/g compared to a measured 
value of = 1800 g/g for cracks propagating perpendicular to them (compare Figures 10 and 12 
to Figure 14 in Clarke et al. 1994). Heat treatments to remove these laminations lead to a value 
of HC ~400 g/g, irrespective of the direction of crack propagation, since a preferential direction 
of orientation for hydride precipitates no longer exists. The HC values for Ti-2 varied between 
400 and 1000 g/g depending on the crack orientation. Heat treatments associated with welding 
did not significantly reduce these values (Clarke et al. 1995). 
Measurements with the Pd-containing alloy Ti-16, a material with very similar properties to Ti-7, 
show that the value of HC is between 1000 and 2000 g/g (Ikeda et al. 2000; Ikeda and Quinn 
1998a, 1998b). This is thought to be due to the increased solubility of hydrogen in this material 
(Ikeda et al. 2000; Ikeda and Quinn 1998b). Since no -phase laminations are present in Ti-16, 
its hydrogen-cracking behavior is not expected to be directional in the manner shown for Ti-12. 
Since Ti-7 possesses similar mechanical properties to Ti-16, the adoption of a model value of HC 
of 1000 g/g is reasonable. This similarity in properties is illustrated by the equal minimum yield 
and ultimate tensile strengths for Ti-2, Ti-16 and Ti-7 (Been and Grauman 2000; Schutz 1995). 
The values of HC adopted in the model are conservative for other reasons also. Slow cracking in 
the a-alloy Ti-2 is dominated by creep and plastic deformation at the crack tip leading to 
blunting (Clarke et al. 1997); i.e., there is no sustained load cracking as could be the case for the 
stronger Ti-12 (Clarke et al. 1994). This inability to achieve a sufficient stress intensity factor 
for fast fracture is most evident at low strain rates indicating that crack blunting should dominate 
under the very slow straining conditions expected under Yucca Mountain conditions. The value 
of HC has also been shown to increase markedly with temperature. At 90C the value of HC for 
Ti-2 has been measured to be > 1000 g/g (Clarke et al. 1995). Whether or not sustained load 
cracking or creep/plastic deformation will occur in a slow strain rate tests prior to fast fracture 
depends on the strength of the alloy. Since both Ti-16 and Ti-7 are a-alloys with very similar 
strengths (minimum yield strengths are all around 275 MPa (Been and Grauman 2000), they 
would be expected to exhibit similar cracking behavior. By contrast Ti-12 has a minimum yield 
strength of 345 MPa. To date, a critical hydrogen concentration has not been measured for Ti 
Grade-24, proposed for use in the support structure for the drip shield. This alloy is a much 
stronger material than the Ti-7 (minimum yield strength of 828 MPa [Schutz 1995]) and would 
be expected to experience sustained load cracking at lower hydrogen concentrations. A value of 
HC for an alloy identical to Ti-24 except for the absence of Pd (i.e., Ti-5) has been measured to 
be ~200 g/g using the identical experimental procedure to that employed for the other Ti alloys 
discussed above (Hardie and Ouyang 1999). If the influence of Pd addition is taken to have a 
similar effect on Ti-5 (to yield Ti-24) as its addition to Ti-2 (to yield Ti-16) did, then it is not 
unreasonable to expect that HC for Ti-24 will be increased to 400 g/g or better. 

Criterion 2 
Is the database of general corrosion rates used in the model consistent with the rates published 
in the peer-reviewed literature? 
The rates used in the model are those based on weight loss measurements on specimens of Ti-7 
exposed in the LTCTF at LLNL to exposure environments chosen to simulate those anticipated 
within Yucca Mountain (CRWMS M&O 2000a, Section 6.5.4). The rates fit a Weibull 
distribution with 50, 90, and 100 percentile values of 25, 100, and 350 nm/year, respectively, and 
comprise the most extensive database of general corrosion rates available. The rates were 
independent of solution composition, pH over the range 2.7 to 8.0, and temperature over the 
range 60C to 90C. 
These corrosion rates can be validated by comparison to other published values. Mattsson and 
Olefjord (1990) measured rates in the range 0.5 to 4 nm/year on Ti-2 and Ti-7 in compacted 
clays saturated with saline solutions (95C). Rates of 40 to 60 nm/year can be obtained from the 
plots of oxide film thickness as a function of time published by Kim and Oriani (1987b). These 
authors also found only a marginal increase in rate with increasing temperature between 25C 
and 108C (the boiling point of the brine). A similar absence of a temperature dependence (90C 
to 200C) was observed by Smailos and Koster (1987) and Smailos et al. (1986) for Ti-7 in 
aggressive Mg-dominated brines1 over an exposure period of ~3.5 years. For this last set of 
results the rates were < 100 nm/year. Blackwood et al. (1988) measured rates in acidic (pH = -1 
to 1) and alkaline solutions (pH =14). Since they also measured a pH dependence over the acidic 
pH range, it was possible to estimate (by extrapolation to more neutral pH values) that the rates 
in more neutral environments should be in the range 15 to 20 nm/year for Ti-2. Since these rates 
were measured using a range of techniques from electrochemical (Blackwood et al. 1988) to 
surface analytical (Mattsson and Olefjord 1990; Kim and Oriani 1987b) to weight change 
(Smailos and Koster 1987; Smailos et al. 1986), they validate not only the values but also the 
weight change method used to determine the rates on the LTCTF specimens. 
Criterion 3 
Does a comparison of the rates of hydrogen absorption calculated by the model to the calculated 
fluxes of absorbed hydrogen within the alloy demonstrate that the assumption that all hydrogen 
absorbed remains in the alloy is conservative? 
The amount of hydrogen absorbed is determined in the model using the general corrosion rates 
(discussed above) multiplied by an electrochemically measured absorption efficiency taken from 
the published literature (Okada 1983). This value was measured for commercially pure titanium 
(equivalent to Ti-2) at pH in the range 3.5 to 4.5, but only a marginally higher value was 
obtained under the same conditions for titanium plated with the noble metal platinum (Okada 
1983). This clearly indicates that the primary controlling factor is the presence of oxide films, 
not the noble metal content of the electrode surface. Consequently, the Pd-containing Ti-7 alloy 
would be expected to exhibit similar absorption efficiency. This similarity between the 
absorption behavior of commercially pure Ti (Ti-2) and Ti-7 is supported by the measurements 
1 Q-brine (wt %)  1.4 NaCl; 4.7 KCl; 26.8 MgCl2; 1.4 MgSO4; 65.7 H2O (pH = 4.9 at 25C) 

of Fukuzuka et al. (1980), who showed that, even under non-passive conditions (boiling HCl), 
the rate of hydrogen absorption by Ti-0.15%Pd (equivalent to Ti-7) was only a factor of two 
greater than for commercially pure Ti (Ti-2). Since surface analytical evidence exists to show 
that the oxide film formed on Ti-0.15%Pd (1.25% HCl at 125C) becomes more passive (i.e., 
thickens more slowly) with time than that on Ti (Shimogori et al. 1982), its ability to prevent 
hydrogen formation and absorption will improve more with time than that of Ti. This justifies 
the adoption of a hydrogen absorption efficiency measured on Ti for use in the model with Ti-7. 
There are published grounds to believe that the absorption efficiency used in the model is 
conservative. Thermally grown oxides (and oxides grown under pickling conditions) have been 
shown to increase the resistance to hydrogen absorption (Been and Grauman 2000; Covington 
1979; Mon 2002, Attachment IV), and the hydrogen absorption observed in deaerated HCl 
(2 = pH = 4; 50C to 250C) was exceedingly low under aerated conditions (Shimogori et al. 
1985). 
Calculations of the diffusive fluxes within the a-modification of titanium using published values 
of diffusion coefficients (Phillips et al. 1974) indicate that any hydrogen absorbed should be 
transported into the bulk of the alloy. However, the accumulation of experimental evidence 
shows that the amount of hydrogen absorbed increases parabolically (i.e., the rate of absorption 
decreases) with time. This suggests that the hydrogen is accumulating in the surface region and 
not being transported deep into the alloy. This is particularly clear in the results of Kim and 
Oriani (1987b) for Ti-12. These results are particularly pertinent since they were measured in 
saturated brines over the temperature range 25C to 108C, conditions close to those anticipated 
in Yucca Mountain. Although the corrosion rate (measured as an oxide film growth rate) 
increased linearly with time, at a rate within the range measured in the LTCTF tests (CRWMS 
M&O 2000a, Section 6.5.2), the hydrogen absorption rate was parabolic and almost independent 
of temperature. It should be noted that the -phase in Ti-12 leads to higher hydrogen absorption 
rates than would be expected in Ti-2 (Lunde and Nyborg 1993) and other a-alloys. Westerman 
(1990) showed that the susceptible Ti-12 absorbed only small amounts of hydrogen over an 
18-month exposure period to concentrated brines at 150C and that absorption stopped after ~6 
months. 
These results show that hydrogen will accumulate in the surface of titanium alloys. This is likely 
to be more marked in Ti-7 than in Ti-12, since the former alloy does not contain the -phase 
ligaments that facilitate diffusion into the alloy bulk (CRWMS M&O 2000c, Section 3.4). 
Accumulation of hydrogen in the surface of the alloy in this manner would mean that it was 
re-released as corrosion progressed. As a consequence, the hydrogen content of the alloy would 
increase more slowly than predicted by the model. This makes it very conservative to assume in 
the model that all the hydrogen absorbed due to general corrosion accumulates until the critical 
hydrogen concentration for a decrease in fracture toughness is achieved. 

Criterion 4 
Are the hydrogen contents for grade-7 titanium and similar alloys predicted by the model 
consistent with values published in the peer-reviewed literature? 
The majority of hydrogen absorption measurements have been performed under aggressive 
conditions not relevant to Yucca Mountain; i.e., under acidic conditions or with a large cathodic 
potential applied to simulate cathodic protection conditions. For more relevant conditions 
(concentrated brines at 150C), Westerman (1990) measured only marginal hydrogen absorption 
(10 to < 30 g/g) by Ti-12 ( an alloy more susceptible to hydrogen absorption than Ti-16 and 
Ti-7) over 12 to 18 months of exposure. The majority of the absorption appeared to occur in the 
first 6 months and then stop. These measurements were made in the presence of a radiation field 
(3 x 104 Rad/h) and it is possible that the hydrogen permeability of the oxide films was reduced. 
An improvement in resistance of the passive film on Ti-12, to hydrogen absorption in the 
presence of a radiation field, was also observed by Kim and Oriani (1987a). Under cathodic 
polarization conditions in NaCl (pH = 1) negligible amounts of hydrogen were absorbed into 
Ti-16 for applied potentials = -600 mV (vs SCE) (Ikeda et al. 2000; Ikeda and Quinn 1998b). In 
the absence of galvanic coupling such negative potentials should be unattainable under Yucca 
Mountain conditions. 
Of particular relevance are the results of Shimogori et al. (1985) who measured the amount of 
hydrogen absorbed in dilute HCl as a function of pH (2 to 4), temperature (50C to 250C) and 
degree of aeration (deaerated and aerated). Over an exposure period of 10 days the amount of 
hydrogen absorbed increased with a decrease in pH and an increase in temperature for deaerated 
conditions. For the most relevant condition (pH = 4) absorption was immeasurable at 50C and 
only 1-2 g/g at 250C. For aerated conditions essentially no measurable hydrogen was 
absorbed irrespective of the pH or temperature over a similar exposure period. There is also 
experimental evidence that aeration suppresses hydrogen absorption by Ti-0.15Pd (equivalent to 
Ti-7) more than it does on by Ti (1% HCl, 70C) (Fukuzuka et al. 1980). Since conditions 
within Yucca Mountain are expected to be oxidizing these results are consistent with the 
expectation that hydrogen absorption by the drip shield will be minor. 
Two attempts to predict the hydrogen absorption rates of Ti nuclear waste containers have been 
published. The model adopted for the drip shield was used to predict the hydrogen absorption 
rates for Canadian waste containers (Shoesmith et al. 2000). In these calculations the very low 
corrosion rates expected in the compacted clay buffer environment proposed for this repository 
(0.4 to 3.3 nm/year) were used with a little higher hydrogen absorption efficiency than that used 
in the Yucca Mountain Project drip shield model. It was predicted that the critical hydrogen 
concentration of 500 g/g would not be exceeded in < 106 years. These calculations are to be 
compared with the current PA model prediction of 755 g/g after 104 years. Based on cathodic 
polarization measurements as a function of applied potential, pH and duration of polarization (60 
days), a parabolic hydrogen absorption relationship was adopted, and it was calculated that only 
17 g/g of hydrogen would be absorbed in 103 years (Tomari et al. 1999). Given that this model 
is based on a parabolic relationship compared to the linear relationship used in the current PA 
model, it strongly supports the model prediction that critical hydrogen concentrations will not be 
exceeded in < 104 years. 

Conclusion: In light of the above discussion, it may be concluded that the input parameter 
values used in the HIC model of titanium during passive general corrosion, and the model output 
corroborate extremely well with those reported mostly in the peer-reviewed scientific literature. 
It is also clear from the above discussion that the validation activities performed for building 
confidence in the model have sufficiently strong scientific bases, and that all of the four criteria, 
used to determine that the required level of confidence in the model has been achieved, have 
been met. Recent publication of this model in a peer-reviewed scientific journal (Shoesmith et 
al. 2000) has further increased the level of confidence in this model. 
7.2 VALIDATION OF THE MODEL FOR HYDROGEN ABSORPTION BY 
TITANIUM UNDER GALVANICALLY COUPLED CONDITIONS 
According to the TWP (BSC 2002), model validation for HIC of titanium under galvanically 
coupled conditions is to be accomplished by meeting the selected criteria described below: 
Criterion 1 
Is the critical hydrogen concentration (HC) of the Grade-7 Titanium alloy (Ti-7) used by the 
model consistent with the results of mechanical tests published in the peer-reviewed literature? 
The critical hydrogen concentration used in this model is the same as that used in the general 
passive corrosion model for HIC. Since the difference between that model and this one for HIC 
under galvanically-coupled conditions is in the processes leading to hydrogen absorption, not in 
the materials properties defining the resistance to cracking, the use of an identical value of HC is 
appropriate. Consequently, the arguments used to validate the values of HC used in the general 
passive corrosion model are equally applicable to this model. 
It should be recalled that the value of HC used is the minimum absorbed hydrogen concentration 
for which a reduction in fracture toughness is observable. For general passive corrosion this 
concentration would be achieved universally throughout the drip shield making it vulnerable to 
fracture over a wide area. By contrast, in the case of galvanic coupling, hydrogen absorption will 
be confined to a relatively small area of the shield (~1000 to 2000 cm2 affected area for the Fe/Ti 
contact area of ~75 cm2 assumed [see Assumption 5.6] in the model). Hence, fracture would be 
expected to occur over a much smaller area of the shield, making the consequences of failure in 
the galvanically coupled case much less severe. 
Criterion 2 
Is the database of general corrosion rates used in the model consistent with the rates published 
in the peer-reviewed literature? 
In the galvanically coupled case, the rate of hydrogen absorption by the Ti-7 is the key rate used 
in the model. This absorption rate is supported by the corrosion rate of the coupled steel. The 
value used for the rate of hydrogen absorption is taken from the measurements of Murai et al 
(1977) (discussed in Shoesmith et al. (1995). A rate of 0.5 g/g/day was used. This value was 
measured in artificial seawater (30C) at a potential of -800 mV (vs. SCE). This is an 
appropriate value since the corrosion potential of a galvanic couple between carbon steel and Ti 

in neutral seawater has been measured to be in the ranges -662mV to -766mV at 18C and 
-736mV to -763mV at 60C (Hodgkeiss et al. 1987). 
This value of absorption rate is conservative for a number of reasons. The rate is calculated from 
the total amount of hydrogen absorbed over the 180-day exposure time employed and, hence, is 
an average linear rate over that period. A simple linear extrapolation of this rate indicates that 
the galvanic couple would have to persist, without any decrease in rate, for 2000 days to achieve 
the critical HC value of 1000 g/g. Actually, since the model allows a considerable amount of 
hydrogen to diffuse laterally within the shield (i.e., to adjacent sites not in galvanic contact with 
the steel), the period required to exceed HC would be greater. However, considerable published 
evidence exists to show that the rate of hydrogen absorption under galvanically coupled 
conditions decreases parabolically with time (Phillips et al. 1974; Noel et al. 1996; Tomari et al. 
1999). This is equivalent to a decrease in the hydrogen absorption efficiency since the carbon 
steel corrosion rate and, hence, the rate of hydrogen production would not necessarily also 
decrease with time. This parabolic decrease means the actual absorption rate beyond 180 days 
used by Murai et al. (1977) (i.e., for the large majority of the exposure period required to exceed 
the HC at the rate used in the model) would be less than the rate used in the model which assumes 
the constant value measured by Murai et al (1977) applies well beyond the 180 day period over 
which it was measured. As a consequence the amount of hydrogen absorbed for the assumed 
amount of carbon steel corroded over the duration of the galvanic couple will be less than 
calculated by the model. 
Hodgkeiss et al. (1987) have measured the current flowing in a carbon steel/Ti galvanic couple 
with a 1:1 contact area (1 cm2) in seawater (pH = 8.1; 60C; 3645 days exposure) to be in the 
range 14 to 22 A. The results of Mon (2002) show that the steel corrosion rate in a galvanic 
couple is approximately proportional to the contact area. If this proportionality is accepted and 
the current measured by Hodgkeiss et al. (1987) is taken to be representative of that anticipated 
under repository conditions, then for the assumed contact area in the model (~75 cm2), the 
galvanic current would be 1,650 A. Application of Faradays Law then demonstrates that such a 
current would not produce, by a large margin, sufficient hydrogen to maintain an absorption rate 
of 0.5g/g/day at a fractional absorption efficiency of 0.015, the value used in the model. 
The key point in this conservative argument is the applicability of the data of Hodgkeiss et al. 
(1987) to repository conditions. Higher temperatures and more aggressive exposure 
environments, especially those containing FeII corrosion products (Mon 2002, Attachments I & 
IV), could lead to the production of considerably larger amounts of hydrogen and, hence, to 
higher absorption rates. However, such rates would be prevented by the production of FeIII 
species, under oxidizing repository conditions, which are known to re-enforce the passivation of 
titanium (Schutz and Thomas 1987; Covington and Schutz 1981) and/or the presence of thermal 
oxides grown during the ventilation period or while temperatures are high prior to the formation 
of aqueous conditions. Substantial published evidence exists to show that thicker oxide films 
grown thermally (Mon 2002, Attachment III) or in a pickling solution (Mon 2002, Attachment 
IV), prevent hydrogen absorption up to temperatures approaching 120C. The results reported in 
Mon (2002, Attachment III) show that a 40-nm thick thermally grown oxide maintains its 
thickness and prevents hydrogen absorption even in 2% HCl at 70C. 

Criterion 3 
Does the comparison of the rates of hydrogen absorption calculated by the model to the 
calculated fluxes of absorbed hydrogen within the alloy demonstrate that the assumption all 
hydrogen absorbed remains in the alloy is conservative? 
Unlike the case for general passive corrosion the titanium drip shield wall is not being thinned by 
material loss during the galvanically coupled corrosion process. In this latter case the titanium 
maintains its wall thickness while cathodically absorbing hydrogen coupled to anodically 
corroding carbon steel. Thus, the assumption that all absorbed hydrogen remains in the alloy is 
realistic, not conservative. 
Criterion 4 
Are the hydrogen contents for Grade-7 titanium and similar alloys predicted by the model 
consistent with values published in the peer-reviewed literature? 
The majority of hydrogen absorption measurements under galvanically coupled conditions (or in 
the presence of cathodic polarization to simulate such conditions) have been made in aggressive 
environments not relevant to Yucca Mountain, i.e., in acidic environments or with a large 
potential that simulates cathodic protection conditions. 
By comparison to the hydrogen absorption rates measured by Tomari et al. (1999) (neutral 
carbonate solutions of pH 8 containing 1000 g/g of chloride at 25C) at -750 mV (vs. SCE), the 
absorption rate used in the model is very conservative. Based on the parabolic relationships 
measured by Tomari et al. (1999), only a few tens of g/g of hydrogen would be absorbed by 
Ti-2 (or Ti-17; Ti-16 with a little extra oxygen content) over thousands of years. By contrast, the 
rate used in the model would predict that the critical hydrogen concentration would be exceeded 
in ~5 to 10 years, providing support by the galvanic couple is maintained. The model rates are 
also conservative compared to those measured by Lee et al. (1986) in seawater (pH = 8, 25C) 
which are approximately one order of magnitude lower. 
A number of additional features should also be noted: 
1. 
While the rate of hydrogen absorption is higher at higher temperatures (in the absence of 
a thermally grown oxide), a very large majority of the hydrogen is stored in a surface 
hydride layer (Noel et al. 1996; Mon 2002, Attachment II) which drastically limits the 
absorption rate (Noel et al. 1996). This makes the use in the model of a constant 
absorption efficiency of 0.015 conservative. Even for large applied potentials or applied 
currents (equivalent to galvanic currents well in excess of those sustainable on the drip 
shield [> 105 A through the galvanic contact area assumed in the model]) the thickness 
of the surface hydride layer would be only tens of microns thick. This is insufficient to 
threaten the full wall integrity of the drip shield. 
2. 
Experimental evidence exists to show that the rate of hydrogen absorption (Lunde and 
Nyborg 1993) will decrease with the accumulation of deposits, calcite from available 

CO2, and silicates from groundwater. The accumulation of silicate deposits has been well 
characterized in LTCTF specimens. 
Conclusion: In light of the above discussion, it may be concluded that the input parameter 
values used in the HIC model of titanium under galvanically coupled condition and the model 
output corroborate extremely well with those reported mostly in the peer-reviewed scientific 
literature. It is also clear from the above discussion that the validation activities performed for 
building confidence in the model have sufficiently strong scientific bases and that all of the four 
criteria used to determine that the required level of confidence in the model has been achieved, 
have been met. 
8. CONCLUSIONS 
A simple and conservative model has been developed to evaluate the effects of hydrogen induced 
cracking on the drip shield. The basic premise of the model is that failure will occur once the 
hydrogen content exceeds a certain limit or critical value, HC. Quantitative evaluation based on 
the HIC model described in Section 6.1 indicates that the drip shield material (Ti-7) is able to 
sustain the effects of HIC. 
The quantitative evaluation includes analytical estimates presented in Section 6.2.3 based on 
Equation 2 indicating that there exists a big margin of safety for the drip shield against the 
effects of HIC. The hydrogen concentration in the drip shield at 10,000 years after emplacement 
is 755 g/g resulting from a conservative estimate (see Table 3). The estimated hydrogen 
concentration is much less than the critical hydrogen concentration of 1000 g/g for Ti-7, a 
conservative value based on Section 6.1.3. 
Table 3. Output of Model for Hydrogen Absorption During Passive General Corrosion of Titanium 
Output Name Output 
Description 
DTN 
Output Uncertainty 
Sources of 
Uncertainty 
Uncertainty 
Distribution (if 
Characteristic 
Values (if 
applicable) applicable) 
HA Hydrogen N/A Uncertainties N/A 755 g/g 
content in the 
drip shield 
associated with 
Ruc and fh 
(conservative 
value) 
(estimated) 
It may be noted that the model predictions shown in Table 3 are based on a very conservative 
and constant upper bound value of 0.015 for fh. However, on the basis of available literature 
evidence for the absorption of hydrogen by titanium, the value of fh, the absorption efficiency 
would be expected to be less than this (Tomari et al. 1999, Table 1). It should be noted that this 
value was measured under very aggressive electrochemically polarized conditions considered 
unachievable under repository conditions. Consequently, fh may be taken to have a value 
between 0.005 and 0.015. The adoption of this range of values means that even the maximum 
value for the general passive corrosion rate will not lead to failure of the drip shield by HIC in 
10,000 years. Since good published evidence exists to show the hydrogen absorption rate will 

have decreased substantially long before the critical hydrogen concentration could be exceeded, 
it is reasonable to conclude that failure of titanium drip shield by HIC will not occur. 
In a drip shield design without backfill, hydrogen generation may be caused by the galvanic 
coupling between the titanium drip shield surface and the ground support (such as rock bolts, 
wire mesh, and steel liners used in the drift), which may fall onto the drip shield surface. A 
qualitative assessment of HIC due to effects of galvanic coupling, as presented in Section 6.3.2, 
concludes that, given the high critical hydrogen concentration, the large volume of available 
titanium in the drip shield into which absorbed hydrogen can diffuse, and other reasons stated in 
Section 6.3, hydrogen embrittlement of the Ti drip shield is highly unlikely. 
A parametric study based on a proposed mathematical model (Section 6.3.3) indicates that 
hydrogen concentration in the drip shield due to a galvanic couple will most likely not exceed the 
critical value. Consideration of the worst case scenario leads to the conclusion that hydrogen 
embrittlement is unlikely even though the hydrogen concentration may exceed the critical value. 
This is because of the fact that the stress intensity factor induced by a rock fall appears to be 
below the fracture toughness of titanium although that may be degraded by hydrogen absorption. 
Furthermore, if cracks do become through wall, they will be self-limited and eventually plugged 
by corrosion product (Section 6.3.7). 

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Shibata, T. and Zhu, Y.-C. 1995. "The Effect of Film Formation Conditions on the Structure and 
Composition of Anodic Oxide Films on Titanium." Corrosion Science, 37, (2), 253-270. [New 
York, New York]: Pergamon Press. TIC: 248556. 
Shimogori, K.; Satoh, H.; and Kamikubo, F. 1985. "Investigation of Hydrogen Absorption-
Embrittlement of Titanium Used in the Actual Equipment." Titanium, Science and Technology, 
Proceedings of the Fifth International Conference on Titanium, Congress-Center, Munich, FRG, 
September 10-14, 1984. Ltjering, G; Zwicker, U.; and Bunk, W.; eds. 2, 1111-1118. Oberursel, 
Germany: Deutsche Gesellschaft fr Metallkunde. TIC: 252879. 
Shimogori, K.; Satoh, H.; Tomari, H.; and Ooki, A. 1982. "Analysis of Passive Film on Titanium 
Formed in Dilute HCl Solutions at Elevated Temperatures." Titanium and Titanium Alloys, 
Scientific and Technological Aspects, International Conference on Titanium, 3d, Moscow State 
University, 1976. Williams, J.C. and Belov, A.F., eds. 2, 881-890. New York, New York: 
Plenum Press. TIC: 252876. 
Shoesmith, D.W. and Ikeda, B.M. 1997. The Resistance of Titanium to Pitting, Microbially 
Induced Corrosion and Corrosion in Unsaturated Conditions. AECL-11709. Pinawa, 
Manitoba, Canada: Whiteshell Laboratories. TIC: 236226. 
Shoesmith, D.W.; Hardie, D.; Ikeda, B.M.; and Noel, J.J. 1997. Hydrogen Absorption and the 
Lifetime Performance of Titanium Waste Containers. AECL-11770. Pinawa, Manitoba, 
Canada: Atomic Energy of Canada Limited. TIC: 236220. 

Shoesmith, D.W.; Ikeda, B.M.; Bailey, M.G.; Quinn, M.J.; and LeNeveu, D.M. 1995. A Model 
for Predicting the Lifetimes of Grade-2 Titanium Nuclear Waste Containers. AECL-10973. 
Pinawa, Manitoba,Canada: Atomic Energy of Canada Limited. TIC: 226419. 
Shoesmith, D.W.; Nol, J.J.; Hardie, D.; and Ikeda, B.M. 2000. "Hydrogen Absorption and the 
Lifetime Performance of Titanium Nuclear Waste Containers." Corrosion Reviews, 18, (4-5), 
331-359. London, England: Freund Publishing House. TIC: 252989. 
Smailos, E. and Kster, R. 1987. "Corrosion Studies on Selected Packaging Materials for 
Disposal of High Level Wastes." Materials Reliabity in the Back End of the Nuclear Fuel Cycle, 
Proceedings of a Technical Committee Meeting, Vienna, 2-5 September 1986. IAEA TECHDOC421, 
7-24. Vienna, Austria: International Atomic Energy Agency. TIC: 252877. 
Smailos, E.; Schwarzkopf, W.; and Koster, R. 1986. "Corrosion Behaviour of Container 
Materials for the Disposal of High-Level Wastes in Rock Salt Formations." Nuclear Science and 
Technology. EUR 10400. Luxembourg, Luxembourg: Commission of the European 
Communities. TIC: 248245. 
Tomari, H.; Masugata, T.; Shimogori, K.; Nishimura, T.; Wada, R.; Honda, A.; and Taniguchi, 
N. 1999. "Hydrogen Absorption of Titaniumfor Nuclear Waste Container in Reducing 
Condition." Zairyo-to-Kankyo, 48, 807-814. [Tokyo, Japan: Japan Society of Corrosion 
Engineering]. TIC: 252965. 
Weast, R.C., ed. 1978. CRC Handbook of Chemistry and Physics. 59th Edition. West Palm 
Beach, Florida: CRC Press. TIC: 246814. 
Westerman, R.E. 1990. "Hydrogen Absorption and Crevice Corrosion Behaviour of Titanium 
Grade-12 During Exposure to Irradiated Brine at 150C." Corrosion of Nuclear Fuel Waste 
Containers, Proceedings of a Workshop, Winnipeg, Manitoba, 1988 February 9-10. Shoesmith, 
D.W., ed. AECL-10121. Pages 67-84. Pinawa, Manitoba, Canada: Atomic Energy of Canada 
Limited. TIC: 227040. 
Whyllie, W.E., II.; Brown, B.E.; and Duquette, D.J. 1994. "The Corrosion of Titanium in 
Alkaline Peroxide Bleach Liquors." Corrosion 94, The Annual Conference and Corrosion Show. 
Paper No. 421. Houston, Texas: NACE International. TIC: 253214. 
Wilhelmsen, W. and Grande, A.P. 1987. "The Influence of Hydrofluoric Acid and Fluoride Ion 
on the Corrosion and Passive Behaviour of Titanium." Electrochimica Acta, 32, (10), 14691474. 
[New York, New York]: Pergamon Press. TIC: 248574. 
9.2 CODES, STANDARDS, REGULATIONS, AND PROCEDURES 
10 CFR 63. 2002. Energy: Disposal of High-Level Radioactive Wastes in a Geologic 
Repository at Yucca Mountain, Nevada. Readily available. 

AP-3.10Q, Rev. 2, ICN 4. Analyses and Models. Washington, D.C.: U.S. Department of Energy, 
Office of Civilian Radioactive Waste Management. ACC: MOL.20010405.0009. 
AP-2.27Q, Rev. 0, ICN 0. Planning for Science Activities. Washington, D.C.: U.S. Department 
of Energy, Office of Civilian Radioactive Waste Management. ACC: MOL.20020701.0184. 
AP-SIII.10Q, Rev. 0, ICN 2. Models. Washington, D.C.: U.S. Department of Energy, Office of 
Civilian Radioactive Waste Management. ACC: MOL.20020506.0911. 
AP-SV.1Q, Rev. 0, ICN 3. Control of the Electronic Management of Information. Washington, 
D.C.: U.S. Department of Energy, Office of Civilian Radioactive Waste Management. ACC: 
MOL.20020917.0133. 
ASTM C 1174-97. 1998. Standard Practice for Prediction of the Long-Term Behavior of 
Materials, Including Waste Forms, Used in Engineered Barrier Systems (EBS) for Geological 
Disposal of High-Level Radioactive Waste. West Conshohocken, Pennsylvania: American 
Society for Testing and Materials. TIC: 246015. 
QAP-2-3, Rev. 10. Classification of Permanent Items. Las Vegas, Nevada: CRWMS M&O. 
ACC: MOL.19990316.0006. 
QAP-2-0, Rev. 5. Conduct of Activities. Las Vegas, Nevada: CRWMS M&O. ACC: 
MOL.19980826.0209. 
10. ATTACHMENTS 
No attachments.