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Measurement of the low temperature diffusivity and solubility of tritium in an iron-base superalloy
Materials Science and Engineering, A 103 (1988) 257 262
257
Measurement of the Low Temperature Diffusivity and Solubility of Tritium in an Iron-base Superalloy S. L. R O B I N S O N
Sandia National Laboratories, Livermore, CA (U.S.A.) S. M. MYERS
Sandia National Laboratories, Albuquerque, NM (U.S.A.) W. B. ESTII.L
Sandia zValional Laboratories, Livermore, CA (U.S.A.) (Received March 27. 1987; in revised form December 21, 1987)
Abstract
Tritium diffusivity and solubility measurements" were made in an iron-base superalloy at 273 K by analyzing the depth distribution of 3He which is the immobile product of the radioactive decay of tritium. 1he measurements were made by nuclear reaction analysis using the reaction 3He(d,p)4He. The diffusivity and solubility agreed, within the precision of the measurement technique, with predicted values from higher temperature studies. Surface preparation effects were characterized by comparing subsurface :He concentrations, allowing relative tritium concentrations to be compared. 1. Introduction
The extrapolation of high temperature permeability data to a lower temperature range is subject to substantial uncertainties. As temperature decreases, surface barriers may become dominant [1, 2], alternative paths of diffusion may contribute significantly [3] and the effects of trapping may become more pronounced [4]. Permeation experiments become progressively more difficult with decreasing temperature because of the lower diffusivities and solubilities. In this work we extended hydrogen diffusivity and solubility measurements to 273 K in a low diffusivity material by a novel approach which circumvents many of the limitations of standard methods. The steel was exposed for 50 days to high-pressure tritium at 273 K, thereby producing a shallow diffusion profile of 20-30 ~m depth. During this time the tritium decayed continuously to 3He, the half-life being 12.38 years. Subsequently, this helium was (t921-5(193/88/$3.5()
depth profiled using the nuclear reaction 3He(d,p)4He, and the tritium profile was calculated from the results [5]. The helium provides an essentially immobile marker for the quantity of tritium present in the steel because it is strongly trapped by lattice defects [6-8] and by clustering (self-trapping) after migrating short distances [6-8]. On the depth scale of interest here, the helium distribution is even stable at temperatures of several hundred degrees Celsius [6, 7]. Consequently, the tritium can be thermally offgassed before ion-beam analysis without significantly perturbing the results. The above approach has two important features which complement the capabilities of thinfoil permeation experiments. First, the high depth resolution of the nuclear reaction analysis, about 0.5/,m, permits measurements to be taken when the diffusion distance is microscopic. Diffusivities may now be measured at lower temperatures. The second desirable feature is that it yields detailed information on the tritium concentration profile, thereby permitting the effects of diffusivity, solubility and surface permeability to be distinguished in a more straightforward manner. 2. Sample preparation and tritium exposure
The alloy chosen for this investigation is a modified A-286 alloy (known as JBK-75) of the nominal composition given in Table 1. The alloy is strengthened by g' precipitation which is developed through a dual age treatment of 8 h at 873 K followed by 8 h at 948 K. The volume fraction of y' is about 10 vol.% with an average 7 © Elsevier Sequoia/Printed in The Netherlands
258
TABLE 1
JBK-75 (Modified A-286) composition (wt.%)
Ni
Cr
Mn
C
V
AI
B
Mo
Ti
29-31
13.5-16.0
0.20
0.1-0.3
2.0-2.3
0.15-0.35
0.002 (Max)
1.0-1.5
2.0
High energy rate forged and aged for 8 h at 873 K followed by 8 h at 948 K.
precipitate diameter of 7-7.5 nm. Strengthening is also imparted by high energy rate forging (HERF) which produces dislocation densities to 5 × 1015 m -2. The samples tested were exposed to tritium in the HERF, 7' strengthened condition. Disc-shaped permeation specimens 9.5 mm in diameter and 0.5 mm thick were prepared from a high energy rate forging by cutting with a low speed diamond saw. A variety of surface conditions were prepared to study hydrogen permeation through different surfaces. Mechanically polished specimens were prepared using successively finer abrasives, ending with an 0.05 /~m alumina abrasive. Metallographic analysis showed the disturbed surface layer to be much less than 0.5/~m thick. An electropolishing solution of 180 ml butyl alcohol, 300 ml methyl alcohol and 30 ml perchloric acid was used at 75 V and at a temperature of 233 K. Electrical leads were spot welded to the side of the sample to perform this step and the spot weld area avoided during nuclear reaction analysis. A number of specimens were palladium coated by sputter deposition to a thickness of 36 nm, as subsequently determined by Rutherford back scattering analysis. Other specimens were subjected to a Nitradd ® cleaning procedure after mechanical polishing. This process uses a proprietary cleaning solution containing nitric acid. Weld fusionzone specimens were similarly prepared from an autogenous gas tungsten arc weld performed on forged JBK-75 but were not aged following the welding process. A stainless steel pressure vessel was used to expose the specimens to 69 MPa tritium gas at 273 K. Stainless steel washers were used to separate the specimens and the areas occluded by the washers avoided during analysis. All parts were cleaned in ethanol and acetone and dried to avoid hydrocarbon contamination of the tritium gas and surface contamination of the specimens. Prior to tritium filling, a vacuum of 1 × 10 -6 Torr was achieved. Ice packs maintained the temperature between 273 and 275 K during the two hour filling procedure and the later transfer stages. A
TABLE 2 Species
T2 DT HT CH3T 3He
Gas composition Vol. % at end of exposure
90.48 1.53 4.06 0.13 3.80 100.00
freezer, held at 271 K to compensate for radioactive decay heat inside the specimen chamber, maintained the temperature for the following 50 days. The tritium filling was 93 vol.% pure, with the analysis listed in Table 2. In the calculations to follow in Sections 3 and 4, the effective tritium pressure is taken to be 69 MPa x 0.93 = 64 MPa. After the 50 day charge period, the tritium gas was removed. Tritium was outgassed from the specimens by immediately evacuating the vessel to 1 × 10 -6 Torr and heating it to 473 K for 14 days. The specimens were then removed for nuclear reaction analysis of the helium.
3. Measurement and interpretation of 3He profiles Aged and tritium-offgassed specimens were bombarded with a beam of deuterons at a series of energy values from 2.0 MeV down to 0.3 MeV. These deuterons induced the nuclear reaction 3He(d,p)4He, yielding high energy protons which were counted for each energy value. A profiling experiment begins at high energy and proceeds to the lower energy so that deposition of deuterium in the material is beyond the range being measured and D - D reactions are minimized. Details of the profiling technique and analysis are discussed in ref. 5 where it is shown that the depth sensitivity is about 0.5/zm and the helium sensitivity is about 0.1 appm. Proton yield data for the tritium-exposed specimen with a mechanically polished surface are given by the open squares in Fig. 1. These results were obtained with 40/zC of 2D+ per point, pro-
259
ducing peak proton yields above 1000. Deconvolution procedures discussed in ref. 5 produced the helium concentration points plotted in Fig. 2. The validity of the deconvolution was checked by reconvohiting the concentration profile to regain nuclear-reaction yield ws. energy. This procedure produced the filled symbols in Fig. 1 where the consistency with the starting data is seen to be good. The 3He concentration profile is not related to the tritium profile by a simple constant Of proportionality but instead, at each depth, the profile is
proportional to the time-integral of the varying tritium concentration. As a result, in deducing the tritium diffusivity and solubility of the measured 3He distribution, it is necessary to utilize a detailed description of the evolution of the system during the period of exposure. This was accomplished by numerically solving the coupled system.
a
o at 40 E
_ Gross Yield
o
° o
:=L
•
30
~Reconvoluted
Data
E 0
u
O:
OtC.(x,t)=DTOx 2 CT--
Cv r
(1
CT t)= + T
(2)
where Cr and CH~ are, respectively, the tritium and helium concentrations in atom fraction, DT is the tritium diffusivity, and r is the tritium half-life divided by In 2. For a fully permeable surface the boundary condition on eqn. ( 1 ) is
20 El
CT(X~O,t)=CToexp
..I Ill
z oI..on.-
10
a.
0
!
i
1 2 I n c i d e n t D e u t e r o n Energy, MeV
Fig. l. Proton yield vs. energy data for the case of mechanically polished specimens. Both raw data and the results of the deconvolution-reconvolution procedure are shown, indicating proper functioning of the deconvolution. The filled squares represent the overlap of the starting data and the reconvoluted data.
T=273 K, 62.4 MPa Tritium, 1204 Hours Fitted Tritium Parameters:
E
#
15 L
~
~
D y = 3,49 x 10 -17 m2/s C T = 2050 appm
=.£ 10 o
"r"
, 2
, 4
, 6
, 8
, 10
Depth, ~m
Fig. 2. Measured helium concentration-depth profile for the mechanically polished specimen. The solid line was fitted using the tritium diffusivity and solubility parameters listed in the figure.
--0.5 tr)
(3)
where CT,,is the tritium solubility at zero time and the exponential factor is due to the decay of tritium gas during the exposure. The tritium parameters D v and (),, are then adjusted to produce agreement between the calculated and measured 3He concentration profiles. The results of this calculation are given by the curve in Fig. 2, where the consistency with experimental results is seen to be good. The deduced tritium diffusivity is 3.49 x 10-~7 m 2 s-~, and the solubility is 2050 appm. An additional check on the nuclear reaction analysis beyond those discussed in ref. 5 is available from the palladium-coated specimens which were included in the experiment. A large surface peak of high helium concentration (6750 appm) was observed on the palladium-coated specimens. The area of this surface peak corresponds to a 3He areal density of 1.43 x 1() ~5 atoms cm 2, which given a tritium-exposure period of 50 days, implies that the tritium areal density before outgassing was 1.85x 1017 atoms cm-2. Furthermore, from Rutherford back scattering with 2.0 MeV 4He, we know that the areal density of palladium was 2.46 x 10 ~v atoms cm 2. Consequently, the atomic ratio of tritium to palladium was 0.75. By comparison, we predict an atomic ratio of 0.82 from the reported equilibrium phase diagram [9], this difference being well within the esti-
260 mated absolute uncertainty of our measurements of about 15%. The 3He profile in the JBK-75 matrix beneath the palladium film of the above specimen could be clearly resolved, notwithstanding the large concentration peak associated with the palladium layer, although the presence of this peak did reduce the accuracy of the concentration points appreciably. The measured 3He concentrations are in good agreement (15%) with those of Fig. 2 from which we infer that surface barriers did not substantially affect the tritium behavior in the mechanically polished, uncoated material. 4. Results and discussion
The tritium diffusivity from Fig. 2 is plotted in Fig. 3 after multiplying by (3/2) 1/2 to correct for the mass difference. The error in our measured D is estimated at a factor of 1.5. The value is seen to agree well with the extrapolated Louthan-Derrick results [10] for austenitic stainless steels. Intermediate temperature results from Swansiger [11, 12] for the JBK-75 alloy are also included. (Swansiger's permeability was divided
1 0 "9 Dataof Louthanand Derrickfor stainlesssteels, X annealedand cold worked
10-1o
1 0 11
and
¢M .
10 "12
\
1 0 "13 O ~:
_=
Data from Swansiger:
1 0 -14
JBK-75. Heft. 8+8 Aged
Deric(ekxtrapolated)
\ \ /
.~#\¼
\
¢~ 1 0 I s C3
\
JBK-75 Heft. 8+8 Aged.MechanicalPolish ~' 1 0 "is
Uriceftain~tii'~Sextr°ad' )polared L-D .... Its- ~ - - - ~ 1 0 -17
1 0 "18 0.001
0.002
0.003
0.004
l/T, l/Kelvins Fig. 3. Diffusivity values from the literature are plotted as a function of temperature. T h e fitted diffusion value agrees well with high temperature austenitic stainless steel diffusion values extrapolated over two orders of magnitude of diffusi o vity.
by the solubility [10] to obtain the diffusivity.) The consistency among the results is good. The solubility of 2050 appm from Fig. 2 is in good agreement with an extrapolation of the higher temperature results of Louthan and Derrick [10]. Using a fugacity of 94 MPa for tritium at 273 K and at a pressure of 64 MPa [ 13], we obtain the solubility 1840 appm from ref. 10. This agreement is well within the experimental uncertainties involved. These results are, however, in conflict with the results of Caskey and Sisson [14] in which the HERF JBK-75 alloy had a lower solubility for hydrogen than single phase austenitic stainless steels did. The source of this discrepancy is unknown. Unlike a thin-film permeation experiment in which steady state permeation is not sensitive to the presence of saturable traps, this method infers the matrix tritium diffusivity from the depth profile of 3He. The presence of trapped tritium would yield higher 3He concentrations than the matrix alone, which would result in an erroneous solution to eqns. ((1)-(3)). Within the sensitivity of the experiment, no evidence was found for trapping. Because of the high tritium pressure used, stress state effects on diffusivity are possible. However, these effects are predicted [15-17] to be within the sensitivity of the method and were not observed. The effects of surface preparation may be examined by comparing the relative proton yields of differently prepared specimens at specific deuteron beam energies which correspond approximately to specific depths. (While specific tritium concentrations are not measured, similar proton yields imply similar helium and therefore similar tritium concentrations over the depth interval in which the nuclear reaction occurs.) In order for the surface preparation comparisons to be valid, several conditions must be met. The subsurface tritium diffusivity must be the same in all specimens regardless of the presence of surface barriers, and the surface barrier must be thin. The effect of the barrier must be simply to reduce the tritium concentration immediately below the barrier. These conditions appear to be met. Table 3 illustrates the effects of mechanically polishing, electropolishing, Nitradd * cleaning, palladium coating and combinations of these treatments on the relative (normalized) tritium concentrations at two deuteron energies, corresponding roughly to two depths. The multiple numbers indicate measurements taken on different samples.
261
TABLE 3 Effect of surface preparation on the subsurface tritium concentration Preparation
Mechanically polished Electropolished Mechanically polished and Nitradd * cleaned Mechanically polished and palladium coated Electropolished and palladium coated
Relative tritium concentrations 1.0 Me V 4- 51~m a
2.0 Me V 8-10 /zm ~
1.24, 1.0 0.088, 0.01,0.0008 0.0032, 0.002
1.0, 0.84
1.04, 0.96 ~' 1.12, 0.28 b
aApproximate depth of measurement. bThe surface peak contribution is subtracted from the gross measurement. See text for explanation of surface peak.
The mechanically polished surfaces gave the highest values, although some variability was observed, possibly reflecting variations in the surface oxide. The electropolished specimen values are erratic and one to three orders of magnitude lower while the Nitradd ® cleaned values are consistent and almost three orders of magnitude below the mechanically polished values. Previous Auger sputtering measurements [11, 18] and permeation studies [11, 12] have shown that Nitradd ® cleaning promotes a thin uniform oxide layer while electropolishing promotes a thicker nonuniform oxide layer. The thick oxide layers may be cracked or defective [11, 18, 19], which leads to variable subsurface concentrations. The palladium-coated specimens, having a significant surface peak, cannot be directly compared to the other surfaces. As was discussed previously, the palladium surface contributions may be subtracted from the total, allowing the net values to be compared. We see then that the palladium-coated samples give the same yield at 2.0 MeV as the mechanically polished specimen. Palladium coating over the electropolished surface has overcome the permeation-reducing effect of the oxide in three out of the four measurements. 5. Conclusions (i) Nuclear reaction analysis has extended the temperature range of hydrogen diffusivity and solubility measurements to 273 K in a low diffusivity iron-base austenitic superalloy. (ii) The measured 273 K diffusivity agreed well with extrapolated higher temperature austenitic stainless steel diffusivity and the measured
solubility agreed well with literature solubility values. (iii) Surface barrier effects from various cleaning and preparation treatments were compared. Thin coherent oxide films formed on Nitradd ®cleaned surfaces showed the most consistently low proton yields, while electropolishing produced an effective but uneven and possibly cracked oxide film. Palladium coating over electropolished surfaces overcame the protective oxide effect in most cases. Acknowledgments Significant contributions were made by B. G. Brown in specimen preparation and by J. Banks in performing beam analysis measurements. The support of the U.S. Department of Energy under Contract DE-AC04-76DP00789 is gratefully acknowledged. References 1 T.S. Elleman and K. Verghese, J. Nucl. Mater., 53 (1974) 299-306. 2 W. G. Perkins, in A. W. Thompson and 1. M. Bernstein (eds.), Effect of Hydrogen On Behavior of Materials', AIME, New York, NY, 1975, p. 355. 3 P. G. Shewmon, Diffusion in Solids, McGraw-Hill, New York, NY, 1963, p. 164. 4 H. H. Johnson and R. W. Lin, in A. W. Thompson and 1. M. Bernstein (eds.), Hydrogen Effects In Metals, AIME, New York, NY, 1981, p. 3. 5 S. M. Myers, G. R. Caskey, Jr., D. E. Rawl, Jr. and R. D. Sisson, Jr., Metall. Trans. A, 14 (1983) 2261-2267; S. M. Myers, in G. E. Murch, H. K. Birnbaum and J. R. Cost (eds.), Nontraditional Methods" in Diffusion, The Metallurgical Society of AIME, Warrendale, PA, 1984, p. 137. 6 G.J. Thomas, W. A. Swansiger and M. I. Baskes, J. Appl. Phys., 50 (1979) 6942-6947. 7 G. J. Thomas and R. Bastasz, J. Appl. Phys., 52 (1981) 6426-6428. 8 W. D. Wilson, C. L. Bisson and M. h Baskes, Phys. Rev. B, 24(1981) 5616-5624. 9 W. M. Mueller, J. P. Blackledge and G. G. Libowitz, Metal Hydrides, Academic Press, New York, NY, 1968, pp. 635-643. 10 M. R. Louthan and R. G. Derrick, Corros. Sci., 15 (1975) 565-577. 11 W. A. Swansiger and R. Bastasz, J. Nucl. Mater., 85 and 86 (1979) 335-339. 12 W. A. Swansiger, in J. S. Watson and F. W. Wiffen (eds.), Proc. Int. Conf. on Radiation Effects and Tritium Technology for Fusion Reactors, Gatlinburg, Tennessee, October 1-3, 1975, Conf. 750989, USERDA, Vol. 4, p. IV-401. 13 B. Baranowski, Ber. Bunsenges. Phys. ('hem. 76 (1972) 714-724.
262 14 G. R. Caskey, Jr. and R. D. Sisson, Jr., Scr. Metall., 15 (1981) 1187-1190. 15 E. C. Aifantes and W. W. Gerberich, in A. W. Thompson and I. M. Bernstein (eds), The Effect of Hydrogen on Behavior of Metals, AIME, New York, NY, 1975, p. 350; Acta Mech., 29(1978) 169-184. 16 K. Farrell and M. B. Lewis, Scr. Metall., 15 (1981)
661-664. 17 K. Shinohara and J. P. Hirth, Philos. Mag., 27(1973) 883. 18 W.E. Estill, Results presented at The Eighth Conference on Surface Studies, Savannah River Laboratory, Aiken, SC, May 1978. 19 W. A. Swansiger, R. G. Musket, L. J. Weirick and W. Bauer, J. Nucl. Mater., 53 (1974) 307.
257
Measurement of the Low Temperature Diffusivity and Solubility of Tritium in an Iron-base Superalloy S. L. R O B I N S O N
Sandia National Laboratories, Livermore, CA (U.S.A.) S. M. MYERS
Sandia National Laboratories, Albuquerque, NM (U.S.A.) W. B. ESTII.L
Sandia zValional Laboratories, Livermore, CA (U.S.A.) (Received March 27. 1987; in revised form December 21, 1987)
Abstract
Tritium diffusivity and solubility measurements" were made in an iron-base superalloy at 273 K by analyzing the depth distribution of 3He which is the immobile product of the radioactive decay of tritium. 1he measurements were made by nuclear reaction analysis using the reaction 3He(d,p)4He. The diffusivity and solubility agreed, within the precision of the measurement technique, with predicted values from higher temperature studies. Surface preparation effects were characterized by comparing subsurface :He concentrations, allowing relative tritium concentrations to be compared. 1. Introduction
The extrapolation of high temperature permeability data to a lower temperature range is subject to substantial uncertainties. As temperature decreases, surface barriers may become dominant [1, 2], alternative paths of diffusion may contribute significantly [3] and the effects of trapping may become more pronounced [4]. Permeation experiments become progressively more difficult with decreasing temperature because of the lower diffusivities and solubilities. In this work we extended hydrogen diffusivity and solubility measurements to 273 K in a low diffusivity material by a novel approach which circumvents many of the limitations of standard methods. The steel was exposed for 50 days to high-pressure tritium at 273 K, thereby producing a shallow diffusion profile of 20-30 ~m depth. During this time the tritium decayed continuously to 3He, the half-life being 12.38 years. Subsequently, this helium was (t921-5(193/88/$3.5()
depth profiled using the nuclear reaction 3He(d,p)4He, and the tritium profile was calculated from the results [5]. The helium provides an essentially immobile marker for the quantity of tritium present in the steel because it is strongly trapped by lattice defects [6-8] and by clustering (self-trapping) after migrating short distances [6-8]. On the depth scale of interest here, the helium distribution is even stable at temperatures of several hundred degrees Celsius [6, 7]. Consequently, the tritium can be thermally offgassed before ion-beam analysis without significantly perturbing the results. The above approach has two important features which complement the capabilities of thinfoil permeation experiments. First, the high depth resolution of the nuclear reaction analysis, about 0.5/,m, permits measurements to be taken when the diffusion distance is microscopic. Diffusivities may now be measured at lower temperatures. The second desirable feature is that it yields detailed information on the tritium concentration profile, thereby permitting the effects of diffusivity, solubility and surface permeability to be distinguished in a more straightforward manner. 2. Sample preparation and tritium exposure
The alloy chosen for this investigation is a modified A-286 alloy (known as JBK-75) of the nominal composition given in Table 1. The alloy is strengthened by g' precipitation which is developed through a dual age treatment of 8 h at 873 K followed by 8 h at 948 K. The volume fraction of y' is about 10 vol.% with an average 7 © Elsevier Sequoia/Printed in The Netherlands
258
TABLE 1
JBK-75 (Modified A-286) composition (wt.%)
Ni
Cr
Mn
C
V
AI
B
Mo
Ti
29-31
13.5-16.0
0.20
0.1-0.3
2.0-2.3
0.15-0.35
0.002 (Max)
1.0-1.5
2.0
High energy rate forged and aged for 8 h at 873 K followed by 8 h at 948 K.
precipitate diameter of 7-7.5 nm. Strengthening is also imparted by high energy rate forging (HERF) which produces dislocation densities to 5 × 1015 m -2. The samples tested were exposed to tritium in the HERF, 7' strengthened condition. Disc-shaped permeation specimens 9.5 mm in diameter and 0.5 mm thick were prepared from a high energy rate forging by cutting with a low speed diamond saw. A variety of surface conditions were prepared to study hydrogen permeation through different surfaces. Mechanically polished specimens were prepared using successively finer abrasives, ending with an 0.05 /~m alumina abrasive. Metallographic analysis showed the disturbed surface layer to be much less than 0.5/~m thick. An electropolishing solution of 180 ml butyl alcohol, 300 ml methyl alcohol and 30 ml perchloric acid was used at 75 V and at a temperature of 233 K. Electrical leads were spot welded to the side of the sample to perform this step and the spot weld area avoided during nuclear reaction analysis. A number of specimens were palladium coated by sputter deposition to a thickness of 36 nm, as subsequently determined by Rutherford back scattering analysis. Other specimens were subjected to a Nitradd ® cleaning procedure after mechanical polishing. This process uses a proprietary cleaning solution containing nitric acid. Weld fusionzone specimens were similarly prepared from an autogenous gas tungsten arc weld performed on forged JBK-75 but were not aged following the welding process. A stainless steel pressure vessel was used to expose the specimens to 69 MPa tritium gas at 273 K. Stainless steel washers were used to separate the specimens and the areas occluded by the washers avoided during analysis. All parts were cleaned in ethanol and acetone and dried to avoid hydrocarbon contamination of the tritium gas and surface contamination of the specimens. Prior to tritium filling, a vacuum of 1 × 10 -6 Torr was achieved. Ice packs maintained the temperature between 273 and 275 K during the two hour filling procedure and the later transfer stages. A
TABLE 2 Species
T2 DT HT CH3T 3He
Gas composition Vol. % at end of exposure
90.48 1.53 4.06 0.13 3.80 100.00
freezer, held at 271 K to compensate for radioactive decay heat inside the specimen chamber, maintained the temperature for the following 50 days. The tritium filling was 93 vol.% pure, with the analysis listed in Table 2. In the calculations to follow in Sections 3 and 4, the effective tritium pressure is taken to be 69 MPa x 0.93 = 64 MPa. After the 50 day charge period, the tritium gas was removed. Tritium was outgassed from the specimens by immediately evacuating the vessel to 1 × 10 -6 Torr and heating it to 473 K for 14 days. The specimens were then removed for nuclear reaction analysis of the helium.
3. Measurement and interpretation of 3He profiles Aged and tritium-offgassed specimens were bombarded with a beam of deuterons at a series of energy values from 2.0 MeV down to 0.3 MeV. These deuterons induced the nuclear reaction 3He(d,p)4He, yielding high energy protons which were counted for each energy value. A profiling experiment begins at high energy and proceeds to the lower energy so that deposition of deuterium in the material is beyond the range being measured and D - D reactions are minimized. Details of the profiling technique and analysis are discussed in ref. 5 where it is shown that the depth sensitivity is about 0.5/zm and the helium sensitivity is about 0.1 appm. Proton yield data for the tritium-exposed specimen with a mechanically polished surface are given by the open squares in Fig. 1. These results were obtained with 40/zC of 2D+ per point, pro-
259
ducing peak proton yields above 1000. Deconvolution procedures discussed in ref. 5 produced the helium concentration points plotted in Fig. 2. The validity of the deconvolution was checked by reconvohiting the concentration profile to regain nuclear-reaction yield ws. energy. This procedure produced the filled symbols in Fig. 1 where the consistency with the starting data is seen to be good. The 3He concentration profile is not related to the tritium profile by a simple constant Of proportionality but instead, at each depth, the profile is
proportional to the time-integral of the varying tritium concentration. As a result, in deducing the tritium diffusivity and solubility of the measured 3He distribution, it is necessary to utilize a detailed description of the evolution of the system during the period of exposure. This was accomplished by numerically solving the coupled system.
a
o at 40 E
_ Gross Yield
o
° o
:=L
•
30
~Reconvoluted
Data
E 0
u
O:
OtC.(x,t)=DTOx 2 CT--
Cv r
(1
CT t)= + T
(2)
where Cr and CH~ are, respectively, the tritium and helium concentrations in atom fraction, DT is the tritium diffusivity, and r is the tritium half-life divided by In 2. For a fully permeable surface the boundary condition on eqn. ( 1 ) is
20 El
CT(X~O,t)=CToexp
..I Ill
z oI..on.-
10
a.
0
!
i
1 2 I n c i d e n t D e u t e r o n Energy, MeV
Fig. l. Proton yield vs. energy data for the case of mechanically polished specimens. Both raw data and the results of the deconvolution-reconvolution procedure are shown, indicating proper functioning of the deconvolution. The filled squares represent the overlap of the starting data and the reconvoluted data.
T=273 K, 62.4 MPa Tritium, 1204 Hours Fitted Tritium Parameters:
E
#
15 L
~
~
D y = 3,49 x 10 -17 m2/s C T = 2050 appm
=.£ 10 o
"r"
, 2
, 4
, 6
, 8
, 10
Depth, ~m
Fig. 2. Measured helium concentration-depth profile for the mechanically polished specimen. The solid line was fitted using the tritium diffusivity and solubility parameters listed in the figure.
--0.5 tr)
(3)
where CT,,is the tritium solubility at zero time and the exponential factor is due to the decay of tritium gas during the exposure. The tritium parameters D v and (),, are then adjusted to produce agreement between the calculated and measured 3He concentration profiles. The results of this calculation are given by the curve in Fig. 2, where the consistency with experimental results is seen to be good. The deduced tritium diffusivity is 3.49 x 10-~7 m 2 s-~, and the solubility is 2050 appm. An additional check on the nuclear reaction analysis beyond those discussed in ref. 5 is available from the palladium-coated specimens which were included in the experiment. A large surface peak of high helium concentration (6750 appm) was observed on the palladium-coated specimens. The area of this surface peak corresponds to a 3He areal density of 1.43 x 1() ~5 atoms cm 2, which given a tritium-exposure period of 50 days, implies that the tritium areal density before outgassing was 1.85x 1017 atoms cm-2. Furthermore, from Rutherford back scattering with 2.0 MeV 4He, we know that the areal density of palladium was 2.46 x 10 ~v atoms cm 2. Consequently, the atomic ratio of tritium to palladium was 0.75. By comparison, we predict an atomic ratio of 0.82 from the reported equilibrium phase diagram [9], this difference being well within the esti-
260 mated absolute uncertainty of our measurements of about 15%. The 3He profile in the JBK-75 matrix beneath the palladium film of the above specimen could be clearly resolved, notwithstanding the large concentration peak associated with the palladium layer, although the presence of this peak did reduce the accuracy of the concentration points appreciably. The measured 3He concentrations are in good agreement (15%) with those of Fig. 2 from which we infer that surface barriers did not substantially affect the tritium behavior in the mechanically polished, uncoated material. 4. Results and discussion
The tritium diffusivity from Fig. 2 is plotted in Fig. 3 after multiplying by (3/2) 1/2 to correct for the mass difference. The error in our measured D is estimated at a factor of 1.5. The value is seen to agree well with the extrapolated Louthan-Derrick results [10] for austenitic stainless steels. Intermediate temperature results from Swansiger [11, 12] for the JBK-75 alloy are also included. (Swansiger's permeability was divided
1 0 "9 Dataof Louthanand Derrickfor stainlesssteels, X annealedand cold worked
10-1o
1 0 11
and
¢M .
10 "12
\
1 0 "13 O ~:
_=
Data from Swansiger:
1 0 -14
JBK-75. Heft. 8+8 Aged
Deric(ekxtrapolated)
\ \ /
.~#\¼
\
¢~ 1 0 I s C3
\
JBK-75 Heft. 8+8 Aged.MechanicalPolish ~' 1 0 "is
Uriceftain~tii'~Sextr°ad' )polared L-D .... Its- ~ - - - ~ 1 0 -17
1 0 "18 0.001
0.002
0.003
0.004
l/T, l/Kelvins Fig. 3. Diffusivity values from the literature are plotted as a function of temperature. T h e fitted diffusion value agrees well with high temperature austenitic stainless steel diffusion values extrapolated over two orders of magnitude of diffusi o vity.
by the solubility [10] to obtain the diffusivity.) The consistency among the results is good. The solubility of 2050 appm from Fig. 2 is in good agreement with an extrapolation of the higher temperature results of Louthan and Derrick [10]. Using a fugacity of 94 MPa for tritium at 273 K and at a pressure of 64 MPa [ 13], we obtain the solubility 1840 appm from ref. 10. This agreement is well within the experimental uncertainties involved. These results are, however, in conflict with the results of Caskey and Sisson [14] in which the HERF JBK-75 alloy had a lower solubility for hydrogen than single phase austenitic stainless steels did. The source of this discrepancy is unknown. Unlike a thin-film permeation experiment in which steady state permeation is not sensitive to the presence of saturable traps, this method infers the matrix tritium diffusivity from the depth profile of 3He. The presence of trapped tritium would yield higher 3He concentrations than the matrix alone, which would result in an erroneous solution to eqns. ((1)-(3)). Within the sensitivity of the experiment, no evidence was found for trapping. Because of the high tritium pressure used, stress state effects on diffusivity are possible. However, these effects are predicted [15-17] to be within the sensitivity of the method and were not observed. The effects of surface preparation may be examined by comparing the relative proton yields of differently prepared specimens at specific deuteron beam energies which correspond approximately to specific depths. (While specific tritium concentrations are not measured, similar proton yields imply similar helium and therefore similar tritium concentrations over the depth interval in which the nuclear reaction occurs.) In order for the surface preparation comparisons to be valid, several conditions must be met. The subsurface tritium diffusivity must be the same in all specimens regardless of the presence of surface barriers, and the surface barrier must be thin. The effect of the barrier must be simply to reduce the tritium concentration immediately below the barrier. These conditions appear to be met. Table 3 illustrates the effects of mechanically polishing, electropolishing, Nitradd * cleaning, palladium coating and combinations of these treatments on the relative (normalized) tritium concentrations at two deuteron energies, corresponding roughly to two depths. The multiple numbers indicate measurements taken on different samples.
261
TABLE 3 Effect of surface preparation on the subsurface tritium concentration Preparation
Mechanically polished Electropolished Mechanically polished and Nitradd * cleaned Mechanically polished and palladium coated Electropolished and palladium coated
Relative tritium concentrations 1.0 Me V 4- 51~m a
2.0 Me V 8-10 /zm ~
1.24, 1.0 0.088, 0.01,0.0008 0.0032, 0.002
1.0, 0.84
1.04, 0.96 ~' 1.12, 0.28 b
aApproximate depth of measurement. bThe surface peak contribution is subtracted from the gross measurement. See text for explanation of surface peak.
The mechanically polished surfaces gave the highest values, although some variability was observed, possibly reflecting variations in the surface oxide. The electropolished specimen values are erratic and one to three orders of magnitude lower while the Nitradd ® cleaned values are consistent and almost three orders of magnitude below the mechanically polished values. Previous Auger sputtering measurements [11, 18] and permeation studies [11, 12] have shown that Nitradd ® cleaning promotes a thin uniform oxide layer while electropolishing promotes a thicker nonuniform oxide layer. The thick oxide layers may be cracked or defective [11, 18, 19], which leads to variable subsurface concentrations. The palladium-coated specimens, having a significant surface peak, cannot be directly compared to the other surfaces. As was discussed previously, the palladium surface contributions may be subtracted from the total, allowing the net values to be compared. We see then that the palladium-coated samples give the same yield at 2.0 MeV as the mechanically polished specimen. Palladium coating over the electropolished surface has overcome the permeation-reducing effect of the oxide in three out of the four measurements. 5. Conclusions (i) Nuclear reaction analysis has extended the temperature range of hydrogen diffusivity and solubility measurements to 273 K in a low diffusivity iron-base austenitic superalloy. (ii) The measured 273 K diffusivity agreed well with extrapolated higher temperature austenitic stainless steel diffusivity and the measured
solubility agreed well with literature solubility values. (iii) Surface barrier effects from various cleaning and preparation treatments were compared. Thin coherent oxide films formed on Nitradd ®cleaned surfaces showed the most consistently low proton yields, while electropolishing produced an effective but uneven and possibly cracked oxide film. Palladium coating over electropolished surfaces overcame the protective oxide effect in most cases. Acknowledgments Significant contributions were made by B. G. Brown in specimen preparation and by J. Banks in performing beam analysis measurements. The support of the U.S. Department of Energy under Contract DE-AC04-76DP00789 is gratefully acknowledged. References 1 T.S. Elleman and K. Verghese, J. Nucl. Mater., 53 (1974) 299-306. 2 W. G. Perkins, in A. W. Thompson and 1. M. Bernstein (eds.), Effect of Hydrogen On Behavior of Materials', AIME, New York, NY, 1975, p. 355. 3 P. G. Shewmon, Diffusion in Solids, McGraw-Hill, New York, NY, 1963, p. 164. 4 H. H. Johnson and R. W. Lin, in A. W. Thompson and 1. M. Bernstein (eds.), Hydrogen Effects In Metals, AIME, New York, NY, 1981, p. 3. 5 S. M. Myers, G. R. Caskey, Jr., D. E. Rawl, Jr. and R. D. Sisson, Jr., Metall. Trans. A, 14 (1983) 2261-2267; S. M. Myers, in G. E. Murch, H. K. Birnbaum and J. R. Cost (eds.), Nontraditional Methods" in Diffusion, The Metallurgical Society of AIME, Warrendale, PA, 1984, p. 137. 6 G.J. Thomas, W. A. Swansiger and M. I. Baskes, J. Appl. Phys., 50 (1979) 6942-6947. 7 G. J. Thomas and R. Bastasz, J. Appl. Phys., 52 (1981) 6426-6428. 8 W. D. Wilson, C. L. Bisson and M. h Baskes, Phys. Rev. B, 24(1981) 5616-5624. 9 W. M. Mueller, J. P. Blackledge and G. G. Libowitz, Metal Hydrides, Academic Press, New York, NY, 1968, pp. 635-643. 10 M. R. Louthan and R. G. Derrick, Corros. Sci., 15 (1975) 565-577. 11 W. A. Swansiger and R. Bastasz, J. Nucl. Mater., 85 and 86 (1979) 335-339. 12 W. A. Swansiger, in J. S. Watson and F. W. Wiffen (eds.), Proc. Int. Conf. on Radiation Effects and Tritium Technology for Fusion Reactors, Gatlinburg, Tennessee, October 1-3, 1975, Conf. 750989, USERDA, Vol. 4, p. IV-401. 13 B. Baranowski, Ber. Bunsenges. Phys. ('hem. 76 (1972) 714-724.
262 14 G. R. Caskey, Jr. and R. D. Sisson, Jr., Scr. Metall., 15 (1981) 1187-1190. 15 E. C. Aifantes and W. W. Gerberich, in A. W. Thompson and I. M. Bernstein (eds), The Effect of Hydrogen on Behavior of Metals, AIME, New York, NY, 1975, p. 350; Acta Mech., 29(1978) 169-184. 16 K. Farrell and M. B. Lewis, Scr. Metall., 15 (1981)
661-664. 17 K. Shinohara and J. P. Hirth, Philos. Mag., 27(1973) 883. 18 W.E. Estill, Results presented at The Eighth Conference on Surface Studies, Savannah River Laboratory, Aiken, SC, May 1978. 19 W. A. Swansiger, R. G. Musket, L. J. Weirick and W. Bauer, J. Nucl. Mater., 53 (1974) 307.