Materials Science and Engineering, Al14 (1989) 179-187
Hydrogen Permeation Behaviour in Austenitic Stainless Steels SUN XIUKUI, XU JIAN and LI YIYI
Institute of Metal Research, Academia Sinica, 2-6 Wenhua Road, Shenyang (China) (Received March 16, 1988; in revised form January 1989)
The permeability, diffusivity and solubility of hydrogen in six types of austenitic stainless steel-316L, 316LN, 21-6-9, 21-9-9, 304 and 321--have been measured by a gaseous permeation technique in the temperature range 200-430 °C. The effect of the cold-working and heat treatment conditions and the alloy composition of the materials on hydrogen permeation has been investigated. The results indicate that the permeability and diffusivity of hydrogen in various alloys obey Arrhenius relationships over the experimental temperature range and the hydrogen permeation behaviour is not significantly influenced by cold-working and heat treatment conditions of the materials but is slightly influenced by alloy composition. The difference between the hydrogen permeation behaviour of pure iron, general alloying steels and austenitic stainless steels has been discussed; and a comparison between the present work and data in the literature has been made. 1. Introduction
The investigation of hydrogen permeation behaviour in austenitic stainless steels is very important in order to understand the hydrogen embrittlement mechanism and practical application of austenitic steels. Because of experimental difficulties, however, the hydrogen permeation parameters of these materials are scattered throughout the literature [1-8]; few reports of the effects of metallurgical factors on hydrogen permeation have been given and the results obtained are not completely consistent. The results of Louthan and Derrick  have shown that the hydrogen permeability and diffusivity are not significantly affected by the alloying elements, cold working and heat treatment. Gromov and Kovneristyi  have reported that the permeation parameters are somewhat affected by the 0921-5093/89/$3.50
alloying constituents. Recently Perng and Altstetter  have demonstrated that the alloy composition does not influence the permeation rate but that cold working reduces the permeability and diffusivity. In this paper, the hydrogen permeation parameters of six types of anstenitic stainless steel have been measured using an ultrahigh vacuum-gas phase permeation technique. The effects of alloying elements, cold working and heat treatment on hydrogen permeation behaviour are reported.
2. Experimental method
2.1. Simple principle The permeation of gaseous hydrogen through metals is a complicated physicochemical process. The process includes adsorption and dissociation of molecular hydrogen on the metal surface, solution and diffusion of the atomic hydrogen, recombination of the atomic hydrogen and desorption of the molecular hydrogen on the metal surface. In general, the diffusion step is the rate-controlling step in the hydrogen permeation process, if the surface is clean and the surface reaction is sufficiently fast. If the permeation surface area of the specimen is sufficiently large compared with the thickness, only one-dimensional diffusion perpendicular to the permeation surface is considered. According to Fick's first law, the hydrogen diffusion rate J, i.e. the diffusion flux in a unit time and on a unit area, is expressed as dc J= -D-dx
where D is thehydrogen diffusivity, dc/dx is the concentration gradient along the x axis which is perpendicular to the permeation surface of the © Elsevier Sequoia/Printed in The Netherlands
membrane specimen. When steady state diffusion is achieved, i.e. a constant concentration gradient is established in the specimen (time t ~ oo ), j is expressed as J~: J= -
- D ( C2 - Cl)
It can be seen from eqn. (6) that the physical concept of the permeability is the steady state permeation rate of hydrogen through a unit thickness and a unit area of the specimen at a unit pressure. The permeability is a physical constant of material and is independent of hydrogen pressure and specimen thickness. The hydrogen diffusivity can be calculated by the time lag method:
where C~ and C z are the hydrogen concentrations at the inlet and outlet surfaces respectively and l is the specimen thickness. Sievert's law has already demonstrated the relationship between the hydrogen concentration and pressure P: C=SP
where tL is the lag time, L e. the time when the transient permeation rate Jt=0.617J~o [11, 12]. The hydrogen solubility S is readily calculated from the known • and D according to eqn. (5).
where S is the hydrogen solubility. From eqns. (2) and (3), Joo can be rewritten as DS(Pll/2 - p21/2)
2.2. Specimen preparation The specimens were taken from cold-rolled sheets of six types of commercial austenitic stainless steel--316L, 316LN, 21-6-9, 21-9-9, 321 and 304--the chemical compositions of which are listed in Table 1. Three treatment conditions of the specimen were studied as shown in Table 2. The effect of a'-martensite transformation on hydrogen permeation in metastable alloys 321 and 304 has already been studied in detail elsewhere . Therefore, annealed specimens of alloys 321 and 304 with a single austenite phase were only studied in this work.
where P1 and P2 are the equilibrium hydrogen pressures at the inlet and outlet surfaces respectively. In general, P1 ">P2, and DS is defined as the permeability O: (5)
(P = DS
Thus, eqn. (4) is simplified as follows: •
Stainless steel 316L 316LN 21-6 -9 21 -9 -9 304 321
Chemical constituents of experimental materials Amount (wt.%) of the following elements Cr
17.02 16.96 20.27 19.58 18.63 17.96
14.27 12.30 7.56 9.45 8.82 8.08
1.05 1.48 9.26 9.30 1.22 0.49
0.14 0.39 0.40 0.33 0.47 0.84
0.027 0.025 0.031 0.032 0.062 0.053
2.19 2.71 -----
-0.23 0.26 0.27 ---
<0.003 <0.005 < 0.003 <0.003 0.022 0.010
0.009 0.007 0.027 0.034 0.035 0.032
--0.03 0.02 ---
--0.05 0.05 -0.51
--0.08 0.08 ---
Treatment conditions and microstructures of specimens
Microstructure of the following types of stainless steel 316L
C old rolled 3 6 - 8 5 %
Solid solution, 1050 °C, 1 h, water-quenched
A n n e a l e d , 9 50 °C, 4 h, furnace cooled
The specimens were cut into disks 43 mm in diameter and 0.15-0.4 mm thick. Before the hydrogen permeation experiment, the specimens were polished with No. 5 metallographic emery paper, cleaned with acetone and then immediately electroplated with palladium on both surfaces. The electroplating was performed in a bath with aqueous solution (pH 7) of palladium chloride (2.5 g l -t) dibasic ammonium phosphate (20 g 1-1), dibasic sodium phosphate (100 g 1-1) and ammonium chloride (25 g 1-1) at 1.25 mA cm- 2 for 4 rain at room temperature. The coating thickness is 0.1 ,um.
2.3. Apparatus and procedure The gas phase permeation dynamic monitoring method was used in the present work. The experimental apparatus and operation procedure have been reported elsewhere [14, 15]. As shown in Fig. 1, the apparatus is an ultrahigh vacuum system made of stainless steel and consists of a specimen chamber, charging hydrogen chamber,
monitoring chamber and other auxiliary parts. The specimen was sealed in the specimen chamber using two gold O-rings, dividing the system into an upstream side and a downstream side, both evacuated to less than 10 -6 Pa. The permeation experiment was carried out in the temperature range from 200 to 430 °C. When a constant vacuum was achieved at a given temperature, high purity (99.9999%) hydrogen gas was introduced into the upstream side to 0.1 MPa. The hydrogen flux permeated through the specimen into the downstream side was monitored with a quadrupole mass spectrometer, which has been calibrated against a standard leak, under dynamic conditions with continuous pumping, and the dynamic plot of permeation rate vs. time was recorded as shown in Fig. 2. The steady state permeation rate J~o and the lag time t L can be obtained from Fig. 2. According to eqns. (6) and (7), the hydrogen permeability and diffusivity can be calculated. 3. Results and discussion
3 14 --N 2
3.1. Effect of the surface condition The oxide film which consists of C r 2 0 3 and seriously affects the normal hydrogen permeation is readily formed on the surface of austenitic stainless steels [8, 9]. The effect of this oxide film on hydrogen permeation has also been found in present work. The experiment with an uncoated palladium specimen showed that the permeation rate was significantly reduced after two hydrogen permeation runs and yellow oxide was observed on the specimen surface.
Fig. 1. Schematic diagram of the permeation apparatus: 1, adsorption pump; 2, outgassing valve; 3, vacuum gauge; 4, ion pump; 5, Bayard-Alpert gauge; 6, ion pump control; 7, quadrupole mass spectrometer control; 8, quadrupole mass spectrometer; 9, furnace; 10, specimen; 11, standard leak; 12, pressure transducer; 13, vacuum pressure gauge; 14, high purity hydrogen generator; 15, Pirani gauge; 16, temperature control; t7, blockade oil trap; 18, liquidnitrogen trap; 19, helium leak detector.
' ' material 21-6-9 1=0.014cm.
I /Jt=O.617J= 8~'ltb~*- 4
12 ' 16 ' 20 24 28 32 time (mln) Fig. 2. Hydrogen permeation rate vs. time for type 21-9-9 steel. (T=410 °C; l=0.217 mm; S = 12.57 cm2; tt = 1 5 9 s; D = 4 . 9 4 x 1 0 - 1 1 m 2 s l ; j = 2 . 0 3 x l 0 - 4 m l ( S T P ) s 1).
Fig. 3. Steady state permeation rate input pressure.
the square root of
450 400 350
4C ~, 30
¢'C 10 -1°
1°~ \ rol I. sol. 316LN • 316L '~ •
-11 m 10 Z" E
i¢2 o 0
40 ;0 111(cm -1)
Fig. 4. Steady state permeation rate vs. reciprocal specimen thickness. Temperature
450 400 ,
1.4 ' ' 1.5
16 i ' '1.7 1.8 1/T(xlO-ZK "1)
Fig. 6. Hydrogen permeability and diffusivity vs. reciprocal temperaturefor type 21-6-9 steel.
0 q f ~
(°C) 300 250 i
=o :E d
e cold-rolled o solution • annealed 1
1.5 16 17 1.8 1/T(xlO'3K -1 )
~ • \
Fig. 5. Hydrogen permeability and diffusivity vs. reciprocal temperaturefor type 316L steel. To eliminate the oxidization of the specimen surface during hydrogen permeation, the specimen was electroplated with palladium on both surfaces. The experiments have indicated that the steady state permeation rate is not only directly proportional to the square root of input hydrogen pressure but also inversely proportional to the thickness of specimen as shown in Figs. 3 and 4. It demonstrates that, after the specimen has been electroplated with palladium, the effect of the oxidization of the specimen surface is eliminated and the hydrogen permeation process is controlled by lattice diffusion.
i i ,16 ,'7 48 ,9 2'0 ~(xlO'3
K -1 )
Fig. 7. Hydrogen permeability and diffusivity vs. reciprocal temperaturefor type 21-9-9 steel.
3.2. Effect o f c o l d working a n d heat treatment
The hydrogen permeabilities and diffusivities of the austenitic stainless steels 316L, 21-6-9 and 21-9-9 with different cold-working and heat treatment conditions are given in Figs. 5-7. It can be seen that, even ff the cold-working and heat treatment conditions of the specimen for each alloy are different, the hydrogen diffusivity and perm-
eability are the same and obey Arrhenius relationships in the experimental temperature range. The various permeation parameters for each alloy have been obtained by a least-squares fit of the permeability and diffusivity data and are listed in Table 3. It can be seen from the results in Figs. 5-7 that the hydrogen permeability and diffusivity in austenitic stainless steels are not significantly affected by the cold-working and heat treatment conditions of the alloys. This observation is different from the hydrogen permeation behaviour of a-Fe and general low alloy steels. This dissimilarity could be caused by the difference in the crystal structures of the materials. Because a large number of the lattice defects that can act as trapping sites for hydrogen exist in general a-Fe and low alloy steels, the usually obtained diffusion activation energy is virtually the apparent hydrogen diffusion activation energy, which contains the lattice diffusion activation energy and the binding energy between the hydrogen and trapping sites. In a-Fe with a b.c.c. structure, the lattice hydrogen diffusion activation energy is lower, generally 6-8 kJ mo1-1 [12, 16, 17]. However, the binding energy between the hydrogen and lattice defects, e.g. dislocations and microvoid, is larger, generally 24-30 kJ mo1-1 [12, 16, 18, 19]. Obviously, these defects can be thought of as "deep traps" with regard to the lattice diffusion. The existence of the "deep traps" necessarily leads to an increase in the apparent
diffusion activation energy and to a decrease in the diffusivity. Therefore the difference in the cold-working and heat treatment conditions of this type of material will cause an obvious disparity in hydrogen permeation parameters. The other ferritic and martensitic steels are all b.c.c. alloys and similar to a-Fe in hydrogen permeation behaviour. Austenitic stainless steels have an f.c.c, structure, a high content of alloying elements and a complex chemical composition. The hydrogen diffusion activation energy in this type of material is greater than in the b.c.c, alloys, e.g. 51-55 kJ mol-1 is given in the present work and is independent of the experimental temperature and the cold-working and heat treatment conditions of specimen. Consequently, it can be considered that this value basically represents the lattice diffusion activation energy of hydrogen in the austenitic stainless steels. On the contrary, the binding energy between the hydrogen and lattice defects, such as dislocations, is lower in austenitic stainless steels, e.g. 18.4-18.8 kJ mol-~ [19, 20]. These defects seem to be "shallow traps" with regard to the lattice diffusion. Therefore trapping by the lattice defects does not seem to have an effect on hydrogen in austenitic stainless steels in the experimental temperature range.
3.3. Effect of alloy composition The hydrogen permeabilities and diffusivities of six types of austenitic stainless steel are pre-
Hydrogen permeation data in austenitic stainless steels
Stainless Temperature Hydrogen pressure steel (°C) (MPa)
= ~,, exp( - H~/RT) (l) 0
304 310 304L 316L 321 316 302 304 347 304 316 316L 316LN 21-6-9 21-9-9 304 321
100-600 199-506 350-700 350-700 224-660 143-745 200-300 539-917 400-900 227-927 227-927 200-430 200-430 200-430 200-430 200-430 200-430
0.01-3.0 4.4 x 10-4-4.4 x 10 -2 1.33x10 5-2.66x10 4 1.33x10 5-2.66x10 4 (0.6-1.2) x 10 2 1.4x10 9-1.4x10-2 6x10-4_1 xl0 2 10 3-10 i 0.001-0.1 0.001-0.1 0.1 0.1 0.1 0.1 0.1 0.1
(mol H 2 m J s 1 MPa i/z)
1.04 x 10 3 2 . 8 4 x 1 0 -4 7.12x10 5 1.17x10 4 8.53 x 10-5 2.36x10 4 5.54 x 10-~ 2.59x10 4
64 56.56 58.52 62.70 59.4 63.5 65.47 67.36
6.25 x 2.70 x 3.90x 1.85 x 2.52 x 2.83 x 2.85 x 2.44 x
65.4 61.7 64.06 60.66 61.29 62.29 62.43 62.70
10 10 10 10 10 10 10 10
4 4 ~ 4 4 4 4 4
D = DII e x p ( - H D / R T )
2 . 7 2 x 1 0 6 54.4 5.15 x 10 -7 48.82 3 . 1 4 x 1 0 7 53.92 4 . 1 5 x 1 0 7 53.71
[11 [2j  [3 [4't 
1 . 6 5 x 1 0 -7 47.85 9 . 9 6 x 1 0 7 52.12 7.01 x 10 -7 48.0 8 . 2 5 x 1 0 7 49.7 6.32 x 10 -7 47.8 4 . 7 9 x 1 0 7 51.59 4.33 x 10 7 50.99 2.49 x 10 -7 51.15 6.33 x 10 -7 54.99 6.43 x 10 -7 53.40 4 . 6 1 x 1 0 -7 53.79
[51 [611 [711 [81 
Present work Present work Present work Present work Present work Present work
184 Temperature (oc) 450
coldrol. sol. 316L e ®
"~o • ~ •\~N N,~
1.6 1.7 1.8 l I T (xlO'3K "1)
Fig. 8. Hydrogen permeabilities v s . reciprocal temperature for various austenitic stainless steels.
Temperature (°C) 350 300 250 , r r
316LN • 10-12 _-21-6-9 o 21-9-9 • 304 321 l
~qL ~ • • ® • I
1.6 1.7 1.8 l I T (xl0-3K -1 )
Fig. 9. Hydrogen diffusivities vs. reciprocal temperature for various austenitic stainless steels.
sented in Fig. 8 and Fig. 9 respectively. Arrhenius behaviour is observed over the experimental temperature range. The hydrogen permeation parameters calculated by means of a least-squares fit of the experimental results for each alloy are listed in Table 3. It can be seen from Fig. 8 that the hydrogen permeabilities of the six types of austenitic stainless steel with different alloy compositions are indistinguishable. This demonstrates that the hydrogen permeability is not affected by the alloy composition. Figure 9
indicates that the diffusivity is slightly affected by the alloy composition. The diffusivities of type 316L, 316LN and 304 stainless steels are very similar as all the diffusivities of type 21-6-9, 219-9 and 321 stainless steels. The data for the former group of alloys are larger than those of the latter group by about 70%. This difference is probably due to the variation in alloy composition. It can be seen from Table 1 that the chemical compositions of type 316L, 316LN and 304 stainless steels are similar, except for nickel, molybdenum and nitrogen. This illustrates that the hydrogen permeation parameters are not significantly affected by these three elements. The chemical compositions of type 21-6-9 and 21-9-9 stainless steels are also very similar, and they contain more chromium and manganese than do type 316L stainless steel etc. and small quantities of aluminium, titanium and copper. Tanabe et al.  have already reported that the alloying elements aluminium, silicon, titanium, chromium, manganese, niobium and vanadium may reduce the hydrogen permeability and diffusivity in stainless steels or other high temperature alloys; the reason probably originates in the oxidization of these minority elements. Tanabe et al. also assumed that the activation energies of permeation and diffusion for an alloy are the weighted sum of the activation energies of the various constituent metals, and the activation energies of permeation and diffusion for the metal chromium have been roughly estimated to be 120 + 60 kJ mol- 1 and 60 + 40 kJ mol- l  respectively. Thus the activation energies of permeation and diffusion for the alloys have shown a tendency to increase with increasing chromium content. In addition, it is well known that the carbide and nitride particles of chromium and titanium are stronger traps than other lattice defects for hydrogen. Consequently, it can be seen from present results that titanium and chromium tend to decrease the hydrogen diffusivity; in particular, the effect of titanium is more obvious and the difference in the diffusivities of type 304 and 321 stainless steels is mainly caused by the titanium content in type 321. The analysis for present results shows that hydrogen diffusivity seems to be affected by titanium and chromium in the form of trapping rather than in the form of oxidization, because the hydrogen permeabilities of austenitic stainless steels with different compositions are essentially the same. This observation indicates
that these elements are not oxidized, as discussed above. Trap sites in materials can depress the hydrogen diffusivity but do not influence the steady state permeation rate or permeability of hydrogen [21-23]. Also, the high manganese content and aluminium and copper appear to be unfavourable to hydrogen diffusion. However, by and large, the effect of alloy composition on hydrogen diffusion in austenitic stainless steels is not as significant as in b.c.c. materials. The reason is similar to that given above. If the difference in hydrogen diffusivities caused by alloy composition can be neglected, a least-squares fit of the experimental results for the six types of austenitic stainless steel yields • = 2.81 x 10 -4 exp
(-62.27kJmo1-1) RT m o l m I s-I MPa-1/2
X 1 0 -7
- 53.62 kJ mo1-1) RT m 2 s -1
S =4.88 x 102 exp
- 8.65_ kJ mol I RT ] m o l m -3 M P a -1/2
3.4. Analysis and comparison In the hydrogen permeation measurement for ct-Fe and low alloy steels by the electrochemical or gaseous permeation method, the permeation plot for the first run is evidently different from that for the second run when the permeation experiment is continuously carried on for the same specimen [23-26]. The steady state permeation in the second run is built up more rapidly than the first run, and the breakthrough time tb, as shown in Fig. 2, is smaller than that in the first run. Therefore the diffusivity obtained from the first-run plot with time lag method is smaller than that from the second-run plot. The third- and fourth-run plots are similar to the second-run plot. This phenomenon is thought to be caused by the irreversible traps existing in the materials. After the first-run permeation, all the irreversible traps are fully filled by hydrogen; thus the steady state permeation for the second run is quickly built up. The above phenomenon is not observed in the present experiment. The second-run permeation plot is just the same as the first. In austenitic stain-
less steels, as discussed in Sections 3.2 and 3.3, the effect of various traps on hydrogen permeation and diffusion is not very significant or is slight. These traps can also be thought of as reversible. The calculation and experimental results of Caskey and Sisson  have also indicated that the hydrogen concentration trapped by dislocation is only 3% of the hydrogen concentration dissolved in the lattice for the austenitic stainless steels with a large dislocation density. Such a small variation in hydrogen concentration is within the limits of measurement sensitivity. In addition, because the hydrogen permeability and diffusivity are extremely small in the austenitic stainless steels, the quantity of hydrogen entering the specimen during the palladium electroplating for 4 min at room temperature is also very small and is concentrated in the surface layer of the specimen. This hydrogen can be completely released during the heating process (5-6 h) at the experiment temperature before the permeation experiment to obtain a constant vacuum. Consequently, the diffusivities obtained from the firstand second-run permeation plots do not differ. In other words, the diffusivity is independent of the time for austenitic stainless steels. Generally speaking, the diffusivity for materials with a high hydrogen concentration, such as niobium and tantalum, depends on the hydrogen concentration . In austenitic stainless steels, because the hydrogen content is quite small (small concentration), e.g. only 1.2 ppm for the six types of stainless steel at 298 K and 0.1 MPa hydrogen pressure in the present work, and almost all the hydrogen is mobile, the diffusivity and steady state permeation rate are independent of the hydrogen concentration in alloys. In the present work, the hydrogen solubility is calculated by dividing the permeability by the diffusivity. It can be seen from the discussion in Sections 3.2 and 3.3 that the coldworking and heat treatment conditions do not influence the solubility in austenitic stainless steels and, in contrast with the effect on diffusivity, titanium and chromium tend to increase the solubility. Caskey and Sisson  have determined the hydrogen concentration of type 304L and 21-6-9 stainless steels at 470 K and 69 MPa hydrogen pressure by the gas analysis method. The hydrogen concentrations calculated from the permeabilities and diffusivities at 470 K and 69 MPa pressure in the present work and the results of Caskey and Sisson  and Louthan
186 TABLE 4
Hydrogen concentration at 470 K and 69 MPa
H concentration (wt.ppm)
Present work Present work
Annealed Cold worked" (36%)
Present work Present work
Solid solution Cold worked (70%)
Present work Present work Present work
Six austenitic Six austenitic Six austenitic
Annealed Solid solution Cold worked
Permeation Permeation Permeation
113.4 113.4 113.4
  
304L 304L 304L
Annealed Cold worked (100%) High energy rate forged
Gas analysis Gas analysis Gas analysis
72 71 71
  
21-6-9 21-6-9 21-6-9
Annealed Cold worked High energy rate forged
Gas analysis Gas analysis Gas analysis
118 126 119
Six austenitic Six austenitic
Annealed Cold worked
85.1 b 85.1 b
aContaining part of a ' martensite phase. bDeuterium concentration.
and Derrick  are all listed in Table 4. In the work of Caskey and Sisson, the specimen surfaces were electropolished or ground up to 600 grit abrasive. The hydrogen concentrations for electropolished specimens are approximately 8% larger than that for ground specimens when the specimens with the same cold-working and heat treatment conditions are compared. The results of Caskey and Sisson listed in Table 4 are the concentrations for the ground specimens. It can be found from Table 4 that (a) the effects of cold-working and heat treatment conditions on the hydrogen solubility in present work are consistent with the work of Caskey and Sisson  and Louthan and Derrick  on stainless steels with a single austenite phase, (b) in the present work and the work of Caskey and Sisson, the hydrogen concentrations for type 21-6-9 stainless steel are all about 70% larger than that for type 304, (c) the hydrogen concentrations for type 304 and 21-6-9 stainless steels in the present work are larger than those in the work of Caskey and Sissons by 29% and 27% respectively and (d) the results for the six austenitic steels in the present work are 33% larger than those in the work of Louthan and Derrick . The above differences in hydrogen concentration seem to arise from the different experimental methods. In the work of Caskey and Sisson  the specimens are charged to saturation in 69 MPa
hydrogen gas a t 4 7 0 K and then the total hydrogen content is directly measured immediately with a Leco RH-1 hydrogen determinator. In this measurement process, if the operation is not very careful, some of the hydrogen could be lost. Although the hydrogen measured in the present work is mobile hydrogen, it is almost equal to the total hydrogen in the austenitic steels as mentioned above, and there is not any hydrogen loss in the measurement process. The result of Louthan and Derrick  is actually the deuterium concentration. There would be some difference between the solubilities of hydrogen and deuterium. However, in a general way, the above differences in hydrogen solubilities are within the experimental errors for austenitic stainless steels. It could be said that these results are approximately consistent with each other.
(1) After austenitic stainless steels are electroplated with palladium on both surfaces of the specimens, the steady state hydrogen permeation rate is directly proportional to the square root of the inlet hydrogen pressure and inversely proportional to the thickness of specimen. This is consistent with the conclusions for other types of material in the literature [6, 26, 29, 30]. T h i s
demonstrates that the hydrogen permeation process is controlled by lattice diffusion in the present experiment. (2) The hydrogen permeation behaviour of austenitic stainless steels is remarkably different" from those of a-Fe and general low alloy steels. These disparities are mainly caused by the difference between the crystal structures of the two types of material. (3) In austenitic stainless steels, there is no difference between the first-run and second-run permeation plots when the permeation experiment is continuously carried on for the same specimen. The hydrogen diffusivity and permeability are independent of the test time and original hydrogen concentration in alloys. (4) The hydrogen permeation behaviour of the austenitic stainless steels is not significantly affected by the cold-working and heat treatment conditions of the materials. The trapping by lattice defects, such as dislocations, for hydrogen is not obvious in the experimental temperature range. (5) In austenitic stainless steels, the alloy compositions do not influence the hydrogen permeability but have a limited effect on the hydrogen diffusion process. Titanium and chromium tend to reduce hydrogen diffusivity and to increase solubility owing to the trapping of hydrogen, but the effect is not as large as in b.c.c. materials [12, 31, 32]. (6) The hydrogen solubilities obtained by permeation measurement for austenitic stainless steels in the present work are approximately consistent with those determined by the gas analysis method for total hydrogen by Caskey and Sisson .
Acknowledgments The authors are grateful to Professor Fan Cungan and Ma Luming who donated the experimental materials.
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