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Effects of deformation on hydrogen permeation in austenitic stainless steels
Acra metall. Vol. 34, No. 9, pp. 1771-1781,1986 Printed in Great Britain. All rights reserved
EFFECTS
oool-6160/86 $3.00+ 0.00 Copyright 0 1986 Pergamon Journals Ltd
OF DEFORMATION ON HYDROGEN PERMEATION IN AUSTENITIC STAINLESS STEELS
TSONGPYNG
PERNG and C, J. ALTSTE-M’ER
Department of Metallurgy and Mining Engineering, University of Illinois at Urbana-Champaign, Urbana, IL 61801, U.S.A. (Received 20 June 1985; in reoised form 3 January 1986) A~-Transient and steady state guxes of hydrogen were measured for annealed and deformed AISI 301, 304 and 310 austenitic and annealed AL 2942 ferritic stainless steel membranes using a gas phase permeation technique at T = lOO-350°C. Permeability and effective diffusivity and solubility constants were calculated from these data. Up to 80% deformation of the stable AISI 310 alloy made only a relatively small change in the transport parameters. Deformation of AISI 301 and 304 resulted in various amounts of stress-induced a’ martensite, which greatly enhanced the effective hydrogen diffusivity and permeability. The relationship between phase changes and hydrogen transport parameters was modeled using various assumptions about the microstructure. Effective solubility and diffusivity values am discussed in terms of dislocation trapping and transport. R6a11&--Nous avons mesure les flux transitoires et stationnaires d’hydrogene dans des membranes d’aciers inoxydables austenitiques recuits et deform&s AISI 301, 304 et 310 et d’acier inoxydable ferritique recuit Al 29-4-2, I l’aide dune technique de permeation en phase gazeuse entre 100 et 350°C. A partir de ces resultats, nous avons calculi: les constantes de permeabilite, de diffusivid effective et de solubilitt effective. Jusqu’ B 80%, la deformation de l’alliage stable AISI 310 ne prod& qu’un changement relativ~ent petit des parametres de transport. La d~fo~ation des alliages AISI 301 et 304 produisait diverses quantitCs de martensite a‘ induite par la contrainte, ce qui augmentait beaucoup la diffusivite et la permeabilite de l’hydrogene. Nous avons mod&id la relation entre les changements de phases et les parametres de transport de l’hydrogene en faisant different= hypotheses sur la microstructure. Nous discutons les valeurs de la solubilite et de la diffusivite effectives en considerant le piegeage et le transport par les dislocations.
ZItppmmeaiassung-Der
Wasserstoff-FluB im Ubergang und im station&en Zustand wurde an metallischen Membranen mit einer G~pha~n-Du~hd~n~n~t~hnik im Te~~ratur~~ich zwischen 100 und 350°C gemessen. Als Membranmaterial dienten ausgeheilte und verformte austenitische rostfreie Stiihle AISI 301, 304, 310 und der ausgeheilte ferritische rostfreie Stahl AL 29-4-2. Aus diesen Megdaten wurden die Permeabilitiit und die effectiven Diffusivittits- und Liislichkeitskonstanten ermittelt. Bis zu einer Verformung von 80% fand sich in dem stabilen Stahl AISI 310 nur eine relativ geringe binderung der Transportkonstanten. Die Verformung von AISI 301 und 304 filhrte zu verschiedenen Mengen an spannungs-induxiertem a’-Martensit, der die effektive Wasserstoffdiffusividit und permeabilitat stark erhohte. Der Zusa~~hang xwischen P~n~nde~ngen und Parametem des Was~~to~rans~rtes wurden untet verschiedenen Annahmen iiber die Mikrostruktur modelliert. Die Werte fir egektive Liislichkeit und Dilfusivitiit werden anhand des Einfangs an Versetxungen und des Transports entlang von Versetzungen diskutiert.
1. ~RODU~ON Knowledge of hydrogen transport behavior in a metal is important to the understanding of the kinetics of hydrogen embrittlement of that metal. Austenitic stainless steels are generally considered less susceptible to hydrogen embrittlement than ferritic alloys due, in part, to the lower diffusivity of hydrogen in the austenite than in the ferrite phase. Troiano and his coworkers, however, found that the embrittlement of austenitic stainless steel had essentially the same dependence on strain rate and test temperature as bee metals and the embrittlement phenomenon was probably ~ffusion~ontroll~ [l]. To embrittle f.c.c. metals, one to two orders of magnitude
greater hydrogen content was required than in b,c.c. metals. Generally speaking, metastable austenitic stainless steel alloys such as AI.91 30 1 or 304 are more susceptible to hydrogen embrittlement than stable alloys such as AISI 310. The metastable alloys are those which are relatively easy to transform to martensite when stressed. The pro~nsity for this transfo~ation depends primarily on the alloy composition and temperature. Alloys with more austenite stabilizers or higher stability factor [2] are found to be more resistant to martensitic transformation as a result of deformation or formation of a as a result of casting or annealing. It is generatly acknowledged that tr~sfo~ed CC’(b.c.c.) martensite is a sufficient, but not necessary condition to allow hydrogen 1771
1772
TSONG-PYNG
PERNG and ALTSTETTER:
embrittlement. Alloys which do not form a’ also suffer hydrogen embrittlement, but with less severity. Previous studies have shown that when notched unstable austenitic stainless steel alloys were stressed in gaseous hydrogen or after cathodic charging, stable (su~~tical) slow crack growth (SCG) under constant load was observed, and a‘ martensite was associated with the fracture surface [3-71. Various theories have been proposed to explain the kinetics of hydrogen-induced slow crack growth behavior in steels and other structural materials. It is frequently postulated that in order for a crack to initiate and propagate, hydrogen has to be transported to and accumulated in an “embrittlement region” at or near the crack tip. The rate of crack advance is controlled by the rate of supply and accumulation of hydrogen in that region [g-14]. Knowledge of hydrogen entry and transport behavior in plastically deformed material is, therefore, important to understand the kinetics of SCG of that material. Since the SCG of austenitic stainless steel may also involve phase transformation and a’ martensite may act as a suitable medium for entry into and transport of hydrogen within the matrix, info~ation on the effects of a’ phase on hydrogen transport in austenitic stainless steel is needed to understand the SCG behavior. In this paper we use annealed and homogeneously deformed specimens to simulate the state of material at a crack tip. Hydrogen transport parameters (diffusivity and permeability) and solubility were ~lculat~ from hydro~n flux measurements in a number of stainless steels in both annealed and deformed conditions using a gas-phase permeation technique. The effects of deformation and a’ martensite on hydrogen transport are discussed.
A schematic diagram of the permeation apparatus is shown in Fig. 1. The ion pumped ultrahigh vacuum (UHV) system and operation procedures have been described elsewhere [15]. Briefly, the technique employed a circular membrane specimen (net surface area 13.85cm2) clamped between two UHV chambers, both evacuated to less than 10G6Pa. Hydrogen gas was purified by diffusion through palladium and then introduced into one UHV chamber. The flux of hydrogen passing through the specimen was measured as an ion current in an ion pump in the other UHV chamber, The detection ion pump, Fig. 1, was calibrated by periodically opening the valve to a known capillary leak. The sensitivity of the detection ion pump was 2.0 x lo-*cm3 H,(STP)/s. The inlet hydrogen pressure ranged from 0.1 to 30 kPa. The furnace was capable of heating to 6OO”C, and the specimen tem~rature was constant to within +0.2”C. Most experiments were carried out at IOOtAL 29-4-2 is a product of the Allegheny Ludlum Steel Corporation.
HYDROGEN PERMEATION IN STAINLESS STEELS IDN FLANGE
CONYAINING
GUN
SPECIMEN
YITRANCE !ON PWP-
CALIBW\lED CAPILLARY
HYDROGEN LEAK
Fig. 1. Permeation apparatus schematic.
350°C. Under these conditions, the pressure build-up in the detection chamber was always below 1.33 mPa and was thus negligible with respect to the input pressure. The steady state flux, J,, and lag time, tL, were measured to calculate the ~~eability constant, 4, and an effective diffusivity, I), using the following equations, respectively
(1)
where h is the thickness of specimen and PH, is the inlet pressure of hydrogen gas. Flux transients were continuously recorded after increments in hydrogen pressure in the inlet chamber. Lag times, t,, were taken as the time to a~omplish 0.630 of the change to a new steady state after an inlet pressure increment. Lag time was independent of inlet hydrogen pressure. The ratio of permeability constant to diffusivity is the effective solubility constant, S. Specimen materials were commercial AISI 301, 304, 310 austenitic and AL 29”4-2t ferritic stainless steel alloys. The compositions are listed in Table 1. Some specimens of AISI 301 and 304 which had undergone various degrees of deformation and therefore contained various amounts of a’ (b.c.c.) martensite phase were prepared by a schedule of rolling, annealing, and electro-polishing to the desired thickness. Deformed specimens of AISI 310 were also prepared. The content of x‘ phase was measured magnetically with a commercial ferrite meter using stacks of sheet to simulate a thicker specimen. A sheet specimen of AL 29-4-2 was used in
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
PERMEATION
IN STAINLESS
1773
STEELS
Table 1. Chemical comwsitions fwt%f of the alloys used for permeation tests Material
C
AISI 301 AN 304 AISI 310
0.052 0.062 0.054
AL 29-4-2
0.0029 0.10
P
S
Si
0.023 0.022
0.009 0.020 0.004
0.48 0.57 0.65
17.1 18.35 24.76
0.01
0.009
0.10
29.5
Mn 1.28 1.31 1.90
Cr
the annealed condition. All specimens were electropolished in a chilled (- 2°C) bath of perchloric and glacial acetic acids (1:4), ion sputtered in Ar at 0.3 mA/cm2 for 120min, and then immediately coated with Pd by vapor deposition. Both entry and exit surfaces were sputtered and Pd-coated. In a bulk diffusion-controlled process, the permeated flux of hydrogen is expected to be not only proportional to the square root of input hydrogen pressure, but also inversely proportional to the thickness of specimens. To check this, three or four measurements of diffusivity and permeability under various pressures of hydrogen at each temperature were made for each specimen, In addition, three annealed 301 specimens with different thicknesses were also studied. The reproducibility was excellent for either case. The average values of permeability and diffusivity were used. The results for deformed 301 and 304 alloys which contained various amounts of a’ phase were compared with those. for annealed 301 and 304 austenitic and AL 29-4-2 ferritic alloys. The effects of Tmz
lo+ +
Ni
MO
N
7.25 9.19 19.55
0.24 0.14 0.11
0.038
2.23
3.93
0.012
0.049
Ti
Cu
Co
Sn
0.27 0.17
0.16
0.016 0.012
a’ phase on hydrogen diffusion and permeation were also studied by comparing the results of deformed 301 and 310 alloys. 3. RESULTS
3.1. AL 29-4-2 ferritic stainfess steel The results for a 195 pm thick specimen of AL 29-4-2 ferritic stainless steel are shown in Figs 2 and 3. A distinct break in the diffusivity curve at - 170°C was observed. The diffusivity in the high temperature region (T > 170°C) was approximately 20-30 times less than that of pure a-Fe, but had an identical activation energy of 7 kJ/mol [16]. No break was observed in the permeability curve in the same temperature range. It has been suggested that trapping effects may account for the downturn of diffusivity in the low t~~rature range [l&-20]. It is generally acknowledged that alloying elements in a-Fe may reduce the diffusivity by various degrees, depending upon the types and amounts of alloying elements [ 171. Nelson et al. reported that, depending on the microstructure, the medium alloy steel AISI 4130 had almost the same as or somewhat smaller diffusivity than pure a-Fe [Zl]. It is expected that AL 29-4-2,
AlSi 301 : A
," 24.0
,"
134
*,
38.3
183
10-g
u
t: -
AL 29 -4
1o-‘2[ 1.5
2.0
1000/r
2.5 (K-l)
Fig. 2. Hydrogen diffuaivity in annealed and deformed AISI 301 and in annealed AL 29-4-2.
1.5
I
I
I
1
I
,
I
‘r 8 i
2.0
1000/T
-2
,
I
>
2.5 (K-l)
Fig. 3. Hydrogen permeability in annealed and deformed AISI 301 and in annealed AL 29-4-2.
PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS
with 35% of alloying element in a-Fe matrix, would exhibit much lower diffusivity. The permeability in AL 29-4-2 was a factor of about 10-l 5 less than that in pure a-Fe. Since the effective solubility is the quotient of permeability over diffusivity, this means that the solubility of hydrogen in AL 29-4-2 was higher than in a-Fe. This, again, might be due to the high amount of alloying. The diffusivity and ~~eability constant of this specimen are described by the following equations
of a’ martensite phase (38.5, 63.5, 65.0 and 86.5%, respectively) were then tested. The content of a’ for each specimen did not change as a result of the test. A separate experiment showed that the stress-induced a’ in both AISI 301 and 304 was stable during annealing at 435°C for a period of 4 days. The results for these specimens along with the annealed alloy are summarized in Table 2 and piotted in Figs 2 and 3. With increasing defo~ation the diffusi~ty increased sharply from that of annealed AISI 301 toward that of AL 29-4-2. Furthermore, a distinct break in the diffusivity curve became increasingly apparent at the higher a’ contents. The peak was similar to that in AL 29-4-2, but at a somewhat higher temperature. The activation energy for diffusion in the high temperature region decreased with increasing a’ content but not as markedly as one might expect from the presence of so much b.c.c. phase (Table 2). The permeability of deformed specimens also increased with the degree of deformation. In fact, the specimen with 86.5% 01’was found to have higher permeability than AL 29-4-2. This has also been observed in a cryoformed 301 alloy with approximately 90% a’, which had higher permeability than both AL 29-4-2 and 301 with 86.5% a’ [22]. The f.c.c. phase dissolved more hydrogen than the b.c.c. phase, but hydrogen diffused more slowly in f.c.c. than in b.c.c. These opposing effects cause the permeability of severely deformed AISI 301 to exceed that of annealed AL 29-4-2. However, comparisons between 301 with a large amount of a’ and the 100% a in AL 29-4-2 must be made with caution because of the different solute contents (N 26% vs N 36%), dislocation densities and grain sizes. A comparison between severely deformed unstable alloy and a ferritic steel with the same amount of solute might be more appropriate, though the nature of the solute elements must also be considered, not just the total solute content. The effects of a’ phase on hydrogen effusion and permeation can be seen better in Figs 4 and 5, where the diffusivity and permeability are plotted as functions of a’ content.
1774
TSONG-PYNG
D = 6.40 x 10-sexp
- 7.0 (kJ/mol) RT m*/s, (T > 170°C)
= 8.45 x lo-* exp
(3)
- 33.7 (kJ/mol) RT Cm2Js, (T < 170°C)
Cp= 2.20 x 10m6exp
- 384(kJ/mol) RT
(4)
cm3 (SIP) cm-s-Pa”2’ (5)
3.2. AISI 301 austenitic stainless steel Type 301 alloy is the most unstable of the standard austenitic stainless steels. Even in the fully annealed condition, a small amount of residual ferritic phase might be present in the f.c.c. matrix, depending on the exact composition and annealing temperature. Both annealed and deformed specimens were studied. Three annealed 301 specimens with different thicknesses (81, 109, and 134 pm) were tested first. They all yielded the same diffusivity and ~~eability constants. The diffusivity and permeability constant of hydrogen in the annealed 301 stainless steel follow the Arrhenius relations
D = 1.30 x 10e3exp
-49.2 (kJ/mol) cm2/s RT
4 = 4.96 x 10e6 exp
- 55.0 (kJ/mol) cm3 (STP) RT cm.s*Pa’/*’
(6)
(7) Specimens with various degrees of deformation (24.0, 37.0, 37.5 and 67.5%) and therefore various amounts
3.3. AISI 304 austenitic stainless steel To examine further the effects of defo~ation and the attendant phase transfo~ation two specimens of
Table 2. Pm-exponential factors and activation energies for effective diffusivity and permeability of hydrogen in anneakd and defamed AK1 301 stainless steel % Deformation (% a’) 0 (0) 24 (38.5) 37 163.5) 37.j (65j
(T > 2Oo?z) (T < 200°C)
67.5 (86.5) (T (T < > 2OO*c) 2OWC)
4 (cm’is) 1.30 x 10-S 8.67 x 1O-4 2.70 x IO-’ 3.10 x 10-3 1.25 x IO-’ 4.60 x 10-S 1.64 x 10-J
(kg& 49.2 42.5 39.3 40.1 45.6
(cm3 (STPgn.s.
Pal@) (k.$L)
4.96 x 1O-6 2.01 x 10-b 6.00 x lo-”
55.0 47.4 45.4
4.46 x 1O-6
44.6
5.50 x 10-G
40.2
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
PERMEATION
IN STAINLESS STEELS
1775
r (“C) lo+
350
300
250
I
I
200 I def. WI
AL 29-4-2. AISI
304
-*-
o
-
o’l% 0
50.0
9.5
W7 10-g
-
10-a
2.0
1.6
1.6
1000/T
Fig. 6. Hydrogen diffusivity in annealed and deformed AISI 304.
lo+ 0
20
40
60
(K-‘1
60
a’%
Fig. 4. Effect of a’ martensite on enhancement of hydrogen diffusivity in deformed AISI 301.
a more stable, type 304, stainless steel were tested. One specimen with a thickness of 97 pm was in the annealed condition. Another one was cold rolled to a 50% reduction of thickness, resulting in 9.5% a’
to a thickness and permeation
martensite, and then electropolished of 112 pm. The hydrogen diffusion
results at T = 200-350°C for the two specimens are plotted in Figs 6 and 7. The data for the annealed one can be described best by the relations D = 7.69 x lo-’ exp
- 53.3 (kJ/mol) ~*/s RT
(8)
10-e
T (‘C) 10-g
i
350
300
250 I
I
200 I def.(%I
a’(%
lo- ‘2
1 lo- l2 1 0
I 20
I 40
I 60
I 60
100
1o-‘31 16
a’%
Fig. 5. Effect of a’ martensite on enhancement of hydrogen Permeability in deformed AISI 301.
1.6
2.0 lOOO/
Fig.
7.
T (K-l)
Hydrogen permeability in annealed and deformed AISI 304.
1776
TSONG-PYNG
PERNG and ALTSTE’ITER:
HYDROGEN PERMEATION
IN STAINLESS STEELS
Table 3. Effective hydrogen transport parameters in annealed and deformed AISI 310 stainless stee1
Deformation cm
4 m is)
38
7.45 x 10-J 7.16
52.8 53.0
3.30 x 10-s 3.50 10-J
58.5 57.3
2.15 3.04
4.52 5.47
60 80
3.70 x 10-j 3.48 x IO-’
52.6 52.8
3.09 x 10-S 3.30 x 10-s
58.7 59.1
5.19 5.89
6.15 6.33
4 = 2.60 x lO-‘exp
(cm3(~T~~~~.s.~a’/‘~
- 57.8 (kJ/mol) cm3 (STP) RT cm*s*Pa1’2 (9)
and the following equations are used to describe the deformed one -40.0 (kJ/mol) cm2/s D = 4.94 x 10m4exp RT
(~,~,~
(at. ~~,~a’l’~
observed after up to 80% deformation. The results for these specimens are shown in Table 3 and plotted in Figs 8 and 9. The effective diffusivity and permeability constant of hydrogen in the annealed 310 alloy can be described by the following equations D = 7.16 x 10e3exp
- 53.0 (kJ/moi) RT cm*/s
Cp= 3.50 x low5 exp
- 58.5 (kJ/mol) cm3 (STP) RT cm.s*Pa”2’
(10)
-49.1 (kJ/mol) cm3 (SIP) 4 = 7.27 x lOa exp RT cm*s.Pa”2’
~~,~
(13)
(11) It was found that both hydrogen diffusivity and permeability in the deformed specimen were higher than those for the annealed one, probably due to the presence of a’ phase. The deformed specimen also had lower activation energies of diffusivity and permeability.
3.4. AISI 310
austenitic stainless Steel
Finally, a completely stable austenite, AISI 310, was used to assess the effects of deformation alone on hydrogen transport. Four specimens of this alloy which had undergone 0,30,60 and 80% deformation (with thicknesses of 175, 188, 187 and 82 pm, respectively) were tested. No phase transformation was
(12)
Both quantities were slightIy lower (about a factor of two at 300°C) than those reported by Quick and Johnson [24]). The permeabilities for the deformed specimens were found to differ by less than 25% from those for the annealed specimen. The ~ffusi~ties in the 60 and 80% deformed specimens were about a factor of 2 lower than those in the annealed and 30% deformed alloys. No break in the diffusivity curve in the low temperature region was observed.
4. DISCUSSION
Three annealed austenitic stainless steel alloys have been studied. It has been reported that hydrogen transport parameters in austenitic stainless steels are Tf%l
T PC) ,o_9y0
3yo
150 ,
200 ,
270
,
o Annealed
AlSl 310
1.6
18
2.0
1000/T
0
30’hCW 6O%CW
.
eo%CW
l
2.2
AISI 310
.
\
2.4
(K-l)
Fig. 8. Hydrogen diffusivity in anneaied and deformed AISl 310.
1.6
1.8
. 30% 13 60Y.
cw CW
aOY.CW
2.0
2.2
24
1000/T (K-‘1
Fig. 9. Hydrogen liability in annealed and deformed AISi 310.
TSONG-PYNG PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS 4 = 1.20 x IOU5exp
1771
56.1 (kJjmo1) cm3 (STP) RT
cm*s.Pa”*’
(15) Type 301 0
l
Equations (14) and (15) are shown as straight lines in Figs 10 and 11, respectively, and all data points fall within a factor of two of these lines. A general expression for effective hydrogen solubility constant in austenitic stainless steels can be obtained as the quotient of equations (14) and (15).
0
302 A&
S = 5.97 x 10s3 exp -6.8~~‘mo1
lo-lot/ , , , , i.6
2.0
1.8
1000/T
I
1
I
I
2.4
2.2
(K-‘I
Fig. 10. Comparison of hydrogen diffusivity in four types of
austenitic stainless steel.
relatively independent of the austenite composition [25]. In this experiment we have also found that the differences of diffusivity and permeability among these alloys were relatively small. The hydrogen diffusivity and permeability in the annealed AISI 301, 304 and 310 alloys are summary in Figs 10 and Il. The data for AISI 302 from a separate study [23] using the same apparatus are also included. A least-squares fit of the diffusivity and permeability constant data for the four alloys yielded I) = 2.01 x low3 exp
-49.3 (kJ/mol) cm2js (14) RT
TIT) 350 10-91
300
250
200
150
I
I
I
I
I
Aurtenitic stainless steels
1000/T (K-‘I
Fig. 11. Comparison of hydrogen permeability in four types of austenitic stainless steel.
~~~~~.
(16)
The results of this work are in good agreement with previously reported values using similar experimental conditions [21,24-40]. The effects of defo~ation on hydrogen permeation in austenitic stainless steels were studied using AISI 301,304 and 3 10 alloys. Hydrogen permeability and diffusivity for each temperature were repeatedly measured and found to be the same each time, even when using different inlet hydrogen pressures. This implies that the filling of pre-existing deep traps was not an important factor under our test conditions. Deformation of type 301 alloy yielded up to 86% of a’ martensite and enhanced both hydrogen diffusivity and permeability by up to two orders of magnitude. In type 310 alloy, no a’ was observed after up to 80% deformation. The permeability was found to be almost the same for the annealed and deformed specimens, whereas the diffusivity was slightly smaller in the severely deformed ones. The difference in behavior between type 301 and 310 alloys is almost certainly due to the different phases present after deformation and not due to different dislocation confi~rations and densities. The effect of dislocations on hydrogen diffusion is somewhat controversial. In annealed f.c.c. metals, it has been concluded that dislocation trapping is frequently not important [41,42]. In deformed f.c.c. alloys, Louthan et al. reported some data which indicated that hydrogen diffusi~ty was not affected by cofd work 1251.In the present work, however, a slight decrease of hydrogen diffusivity was observed in the severely deformed AISI 310 specimens. It is possible that the lower hydrogen diffusivities were caused by dislocation trapping, and this hypothesis will now be examined. When hydrogen enters a specimen with traps the time for the transient to appear on the exit surface is increased, partly because the hydrogen must first fill the deep traps. This is reported, then, as a lower effective diffusivity. However, once the trap sites are saturated the steady state hydrogen flux is essentially unchanged, and thus the permeation is unchanged. This is the effect observed in the case of deformed AISI 310 steel. In order to minimize systematic errors in thickness measurement, thickness uniformity, surface condition, etc. we have taken the ratio of permeability to effective diffusivity, i.e. the
1778
PERNG and ALTSTETTER:
TSONG-PYNG
effective solubility, to characterize the trap population introduced by cold work. From equations (12) and (I 3), the effective hydrogen solubility constant in the annealed 310 alloy can be obtained S =4.89 x lo-‘exp
= 2.63 x 10”exp
= 3.04exp
-5.47 (kJ/mol) RT
- 5.47 (~/mol) RT
- 5.47 (kJ/mol) R1”
cm3H,(STP) cm3metal. Pa”2
atom ~ cm3 *Pall*
at ppm/Pa”*.
(17)
The formulae for the deformed specimens can also be constructed from the constants included in Table 3. The total concentration of hydrogen in the annealed alloy in PH2= 1 atm and at 573 and 423 K (within the temperature range of this experiment) are calculated to be 307 and 205 at. ppm, respectively. These values along with those for the deformed specimens are shown in Table 4. The increments of hydrogen concentration in the deformed alloys can be seen. We assume that in the annealed specimen dislocation trapping is negligible, so that the effective (or total) ~n~ntration, C,, is equal to the lattice concentration, C,. If we express concentrations in units of H atoms/cm3 and use the symbols C; and C;, then C; is calculated to be 2.66 x lOI and 1.77 x lOl9 H/cm3 at 573 and 423 K, respectively. The increments in hydrogen concentration, C; - C;, due to deformation are taken to be the amounts of hydrogen in trap sites. These are tabulated in Table 4, along with the fraction of the hydrogen in trap sites, (C+-C;)/Ci. This fraction increases rapidly at first and then less rapidly, which is not unexpected. It might be commented here that one would expect this behavior to be rather different in ferritic materials, where the lattice solubility is several orders of mag~tude lower. In this case one would expect trapping effects to dominate. Experiments are underway to examine this in a ferritic stainless steel. Thomas [41] reported that the hydrogen binding to dislocations in f.c.c. metals was essentially due to
Table 4. Cakulated
HYDROGEN
PERMEATION
IN STAINLESS STEELS
elastic interactions, with a binding energy of 9.6-13.5 kJ/moI. This is different from that in iron or ferritic steels, where core interaction constitutes a significant contribution to hydrogen trapping [19]. From the binding energy the fraction of trap sites occupied by hydrogen atoms, C,, can be calculated at any temperature. Assuming that all trap sites have the same energy of 12 kJ/mol and that this energy is independent of dislocation density or hydrogen content, a very great simplification for the case of a dislocation stress field, we estimate that the density of trap sites, N, (= C;CC;/C,), ranges from 0.5 to 5.0 x 102’ cm-j. This is the same order of magnitude as that estimated by Oriani for deformed iron and ferritic steels (431. The estimated increase in N, was not large in going from 30 to 80% deformation, considering the orders of magnitude increase in dislocation density. At a relatively high average dislocation density of lO’*cn-* there would be several H atoms per lattice plane along a dislocation line. We note that taking 12 kJ/mol as the binding energy, C, at 423 K should have been higher by 60% than that at 573 K. The fraction of hydrogen in trap sites, (C; - Ci)/C;, seemed to be insensitive to this difference. Nevertheless, this simple calculation did illustrate that in the deformed 310 stainless steel, it is reasonable to assume that increased dislocation density could be effective in trapping a certain amount of hydrogen, increasing the effective hydrogen concentration, and slightly reducing the effective hydrogen diffusitivity. In the deformed type 301 alloy, both hydrogen diffusivity and permeability were increased. Dislocation trapping certainly could not explain this result. On the other hand, if disl~ations provided preferential short-circuit paths for fast diffusion, one might expect to see this behavior. However, the experiments with two alloys of different austenite stability showed that the extent of enhancement was a function of u’ content rather than the amount of deformation. Type 304 was more stable than type 301 alloy, since only 9.5% a‘ phase was induced after 50% deformation, whereas 65.0% LX’phase was formed in a 37.5% deformed 301 alloy. Had dislocations enhanced hydrogen diffusion, the enhancement in the 50% deformed 304 alloy should have been higher than in
effective solubility’ and trapping parameters for hydrogen in AISI 310 stainless steel Total concentration
Trapped concentration c;-c; (lOI H/cm’)
Fraction of H in traps c; - c;jc;
Deformation (W
(at.zptn)
0 30 60 80
307 340 456 498
0 :
205 293 244
2.53 2.11
0.76 0.34
0.30 0.16
80
311
2.69
0.92
0.34
(10’e&m3)
At T = 573 K (300°C) 2.66 2.95 0.3 3.95 1.3 4.31 1.6 At T-423K
‘1 atm pressure of H,.
1.77
0.10 0.33 0.38
(150°C)
TSONG-PYNG PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS
Table 5. Effective hydrogen solubility” in AISI 301 alloy at 200 and 300°C Deformation (X) 0 24.0 37.0 37.5 67.5
& 0 38.5 63.5 65.0 86.5
Concentration of H (at. ppm) 200°C 300°C 173 132 93 91 60
223 I64 122 111 76
‘1 atm pressure of H,.
the 37.5% deformed type 301. This was not the case (see Figs 2 and 6). Other evidence for this argument is the decrease of effective hydrogen solubilities resulting from the deformation in 301. The effective hydrogen solubilities at 200 and 300°C in the annealed and deformed 301 specimens calculated from the ~~eability and ~ffusivity are shown in Table 5. Unlike the deformed AISI 310 alloy, the effective hydrogen solubility decreased with increasing amount of deformation. This was ascribed to the presence of tl’ phase which accommodated much less hydrogen in solution than the y phase. The results of this series of experiments led us to conclude that the enhan~ment of hydrogen diffusivity and permeability in deformed 301 was not due to the enhanced dislocation density, but the presence of a’ martensite. Only in the deformed 310 alloy, where no tl’ was formed was there an effect which could be ascribed to dislocation trapping. It is noted that the activation energy for hydrogen diffusion at low tem~ratures in the 67.5% deformed 301 steel (48.8 Wfmol) was much larger than that in annealed AL 29-4-2 (33.7 kJ/mol). This is not unexpected if we consider that the highly dislocated Co martensite may yield significant trapping. Deformed AL 29-4-2 may well exhibit higher activation energies than the annealed alloy. Experiments are underway to examine this possibility. It is also interesting to note that the hydrogen transport behavior in deformed 304 was different from both 301 and 310 alloys. For deformed 304, the hydrogen ~~eability increased more than the diffusivity, resulting in a higher apparent solubility of hydrogen. Since Type 304 is more stable than type 301, only 9.5% tl’ was induced after 50% deformation. Interpolation from Table 3 shows that with this amount of deformation, there was -35% increase of hydrogen solubility in AISI 310 at T = 200-3~°C. On the other hand, inte~olation from Table 5 shows that the presence of 9.5% a’ in AISI 301 may result in only -5% decrease of hydrogen solubility in the same temperature range, neglecting the effect of dislocation trapping of hydrogen. These two opposite effects may lead to a higher effective hydrogen solubility in deformed 304. From Figs 6 and 7, the ratios of hydrogen solubility in deformed and annealed 304 were found to be 1.34 and 1.66 at 200 and 3OO”C, respectively. The motivation for studying hydrogen transport
1779
behavior in the deformed austenitic stainless steel originated from the fact that ~1’ martensite was associated with the fracture surface and the region near the crack tip when unstable alloys are stressed in hydrogen. The much higher hydrogen diffusivity and ~~eability in martensite than in the f.c.c. matrix rationalizes the hydrogen embrittlement susceptibility and hydrogen-induced SCG behavior of AISI 301. While it is very difficult to directly measure hydrogen transport at a crack tip, it is desirable to obtain these parameters for various martensite contents in a unifo~ly deformed alloy, so that fluxes into and out of the crack tip region may be estimated. Several models were used to quantitatively analyze the enhanced diffusivity and permeability in alloys containing various amounts of a’ in a y matrix. The quantitative predictions for the phase mixtures depend very greatly on the ~~rostructure. First, upper and lower bound estimates of the diffusivity arose from assuming that the a’ and y were present as continuous lamellae and the hydrogen flux passed either parallel to the lamellae (upper bound) or through the lamellae in series (lower bound). The diffusivity in the parallel case can be written as a linear combination of diffusivities of the a’ and y phases. The series case uses a Iinear combination of resistivities (reciprocal diffusivity) of the two phases. These two estimates are compared in Fig. 12 to the values of diffusivity measured to 200°C for cold rolled AISI 301 (shown as points in Fig. 12). For this figure the computed curves used the measured value of diffusivity (4.79 x 10-9cm2/s) for austenite and an assumed value of 1.5 x 10m6cm’/s for the martensite. Figure 12 shows that this assumed lamellar micro-
I
’
’
0
20
40
60
80
I
100
Volume % 0’
Fig. 12. Estimation of the enhanced hydrogen diffusivity in deformed AISI 301 . at 2OOT.. Points are experimentally .
determined values.
1780
TSONG-PYNG
PERNG and ALTSTE’ITER:
structure gives values which can differ by more than an order of magnitude from the measured diffusivities. A model used by Louthan et al. [33] gave somewhat better results than the series curve in Fig. 12, but values were still almost an order of magnitude too low. The interpolation formula, equation (18), was derived by McLachlan [451 for the conductivity of a mixture containing a volume
HYDROGEN
f
(Pin- PJ
For a thr~~imensional body Pi= (l)Jvbi3 where i stands for ~1,y or the mixture and A is a fitting parameter, which can be thought of as a fractal dimension. The plot of equation (18) shown in Fig. 12 for A = 1.487 and D,. = 1.5 x 10V6cm*/s, fits the experimental points remarkably well. The values of A and DES were optimized to give this close fit. It is seen that the optimized value of I), for a 100% a’ AISI 301 alloy is about an order of magnitude less than the measured value for AL 29-4-2. This difference is probably due to differences between, (a) the model micro-structure and the true microstructure of AISI 301, (b) the two alloy compositions and (c) the defect contents of 100% (Y’and annealed ferrite. The application of the above models to prediction of permeation constants gave results similar to Fig. 12. The good fit to the experimental data was obtained for A = 1.487. and & = 5 x lo-” cm3/ cm*s.Pa’lz. This is a factor of 4 greater than that for AL 2942. Since saturated deep traps are not expected to affect the permeability, we are led to the conclusion that the alloy content is one of the principal reasons for this difference. The estimated values of D and # for 301 containing 100% tl’ are consistent with a linear dependence of hydrogen solubility on a’ content. 5. CONCLUSIONS 1. A gas-phase permeation technique fully used to study hydrogen transport
was success-
in a number of stainless steels. The results for the annealed austenitic stainless steels were in good agreement with reported values determined by similar techniques. 2. The highly alloyed AL 29-4-2 ferritic stainless steel had lower diffusivity and permeability than lower alloy ferritic steels by a factor of 10-20. 3. Deformation of the metastable austenitic AISI 301 and 304 alloys resulted in various amounts of a’ martensite and enhanced both hydrogen diffusivity and permeability. 4. No phase transformation was observed in AISI 310 after up to 80% deformation. The effective hydrogen diffusivities in the deformed 310 were slightly decreased, probably due to dislocation trapping.
IN STAINLESS STEELS
Acknowledgements-This research was supported by the National Science Foundation under Contract NSF DMR 83-03421. The AL 2942 ferritic stainless steel was supplied by the Allegheny Ludlum Steel Corporation. We wish to thank Professor D. S. McLachlan of the University of the Witwatersrand, Johannesburg, South Africa for his help in fitting the hydrogen diffusivity data.
fraction, f, of spheres of highly conducting material (cc’) immersed in a lower conductivity medium (y).
P,,, + (A - l)P,
PERMEATION
REFERENCE 1. M. B. Whiteman and A. R. Troiano, Corrosion 21, 53 (1965). 2. C. J. Novak, Handbook of Stainless Steels (edited by D. Peckner and I. M. Bernstein). DD. 4-29. McGraw-Hill. New York (1977). ‘. __ 3. D. Eliezer, D. G..Chakrapani, C. J. Altstetter and E. N. Push. Metall. Trans. A 1OA. 935 (1979). 4. R.iitu, N. Narita, C. Altstetter, H.‘Bimbaum and E. N. Pugh, Merall. Truns. A llA, 1563 (1980). 5. N. Narita and H. K. Bimbaum, Scrip& metall. 14.1355 (1980). 6. S. Singh and C. Altstetter, Melall. Trans. A 13A, 1799 (1982). 7. G. Schuster and C. Altstetter. Metal% Trans. A I4A. 2085 (1983). Acta metall. 22. 1965 8. R. A. Oriani and P. H. Joseohic. . (1974). 9. R. A. Oriani and P. H. Josephic, Acta metall. 25, 979 (1977). 10. W. W. Gerberich, Y. T. Chen and C. St. John, Metall. Trans. A 6A, 1485 (1975). 11. W. W. Gerberich and Y. T. Chen, Metall. Trans. A 6& 271 (1975). 12. W. W. Gerberich, J. Garry and J. F. Lessar, E#ecfs of Hydrogen on Behavior of Materials (edited by A. W. Thompson and I. M. Bernstein), p. 70. T.M.S.-A.I.M.E, Moran, Wyoming (1976). H. P. van Leeuwen, Corrosion 29, 197 (1973). ::: H. P. van Leeuwen, Corrosion 31, 154 (1975). 15. P. H. Studebaker, M.S. thesis, Univ. of Illinois at Urbana-Champaign, Ill. (1981). 16. N. R. Quick, Ph.D. thesis, Cornell Univ., N.Y. (1975). 17. J. VBlkl and G. Alefeld, Diffusion in Solids (edited by A. S. Nowick and J. J. Burton), p. 252. Academic Press, New York (1975). 18. C. A. Wert, Hydrogen in iwetals, Topics in Appiied Physics (edited by G. Alefeld and J. Vitlkl), Vol. 29, p. 305 (1978). 19. J. P. Hirth, MetaN. Trans. A. llA, 861 (1980). 20. H. H. Johnson and R. W. Lin, Hydrogen E&as in Metals (edited by I. M. Bernstein and A. W. Thompson), p. 3. T.M.S.-A.I.M.E., Warrendale, Pa (1980). 21. H. G. Nelson and J. E. Stein, NASA Rep. TND-7665 (1973). 22. T. P. Pemg and C. J. Altstetter, Scripta metait. 18, 67 (1984). 23. T. P. Pemg and C. J. Altstetter. In press. 24. N. R. Quick and H. H. Johnson, MetaN. Trans. A lOA, 67 (1979). 25. M. R. Louthan Jr and R. G. Derrick, Corrosion Sci. 15, I
565 (1975).
26. R. H. Collins and J. C. Turnbull, Vacuum 11, 114 (1961). 27. H. L. Eschbach, F. Gross and S. Schulien, Vacuum 13, 543 (1963). 28. J. R. Phillips and B. F. Dodge, A.Z.Ch.E.JI 14, 392 (1968).
29. K. F. Chaney and G. W. Powell, Metall. Trans. 1,2356 (1970). 30. W. J. Kass and W. J. Andrzejewski, AEC Rep. Sc-Dr72-0136 (1972).
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
31. J. H. Austin, T. S. Elleman and K. Verghese, J. nuci. Muter. 48, 307 (1973). 32. W. A. Swansinger, R. G. Musket, L. J. Wierick and W. Bauer, J. nuct. Mater. 53, 307 (1974). 33. M. R. Louthan Jr, J. A. Donovan and G. R. Caskey Jr, Nucl. Teclr. 26, 192 (1975). 34. V. A. Maroni, E. H. van Deventer, 1. A. Renner, R. H. Pelto and C. J. Wierdak, CONF-750989, IV-329 (1975). 35. J. T. Bell. J. D. Redman and F. J. Smith. ORNL 5297. 28 (1977): 36. E. H. Van Deventer and V. A. Maroni, ANL/FPP-77-7, 17 (1977). 37. H. K. Perkins and T. Noda, J. nucl. Mater. 71,349 (1978).
PERMEATION
IN STAINLESS STEELS
1781
38. W. Conley, P. Studebaker and C. Altstetter, Ref. [20], p. 169. 39. M. Matsuyama and J. D. Redman, MetaN. Trans. A 14A, 498 (1983). 40. R. A. Outlaw and D. T. Peterson, Metal!. Trans. A 14A, 1869 (1983). 41. G. J. Thomas, Ref. [20], p. 77. 42. H. K. Bimbaum and C. A. Wert, Ber. Btmse. Phyz. Chem. 76, 806 (1972). 43. R. A. Oriani, Acra meratt. 18, 147 (1970). 44. J. F. Breedis and W. D. Robertson, Acta metatt. 10, 1077 (1962). 45. D. S. McLachlan, J. Whys. C. To be published.
EFFECTS
oool-6160/86 $3.00+ 0.00 Copyright 0 1986 Pergamon Journals Ltd
OF DEFORMATION ON HYDROGEN PERMEATION IN AUSTENITIC STAINLESS STEELS
TSONGPYNG
PERNG and C, J. ALTSTE-M’ER
Department of Metallurgy and Mining Engineering, University of Illinois at Urbana-Champaign, Urbana, IL 61801, U.S.A. (Received 20 June 1985; in reoised form 3 January 1986) A~-Transient and steady state guxes of hydrogen were measured for annealed and deformed AISI 301, 304 and 310 austenitic and annealed AL 2942 ferritic stainless steel membranes using a gas phase permeation technique at T = lOO-350°C. Permeability and effective diffusivity and solubility constants were calculated from these data. Up to 80% deformation of the stable AISI 310 alloy made only a relatively small change in the transport parameters. Deformation of AISI 301 and 304 resulted in various amounts of stress-induced a’ martensite, which greatly enhanced the effective hydrogen diffusivity and permeability. The relationship between phase changes and hydrogen transport parameters was modeled using various assumptions about the microstructure. Effective solubility and diffusivity values am discussed in terms of dislocation trapping and transport. R6a11&--Nous avons mesure les flux transitoires et stationnaires d’hydrogene dans des membranes d’aciers inoxydables austenitiques recuits et deform&s AISI 301, 304 et 310 et d’acier inoxydable ferritique recuit Al 29-4-2, I l’aide dune technique de permeation en phase gazeuse entre 100 et 350°C. A partir de ces resultats, nous avons calculi: les constantes de permeabilite, de diffusivid effective et de solubilitt effective. Jusqu’ B 80%, la deformation de l’alliage stable AISI 310 ne prod& qu’un changement relativ~ent petit des parametres de transport. La d~fo~ation des alliages AISI 301 et 304 produisait diverses quantitCs de martensite a‘ induite par la contrainte, ce qui augmentait beaucoup la diffusivite et la permeabilite de l’hydrogene. Nous avons mod&id la relation entre les changements de phases et les parametres de transport de l’hydrogene en faisant different= hypotheses sur la microstructure. Nous discutons les valeurs de la solubilite et de la diffusivite effectives en considerant le piegeage et le transport par les dislocations.
ZItppmmeaiassung-Der
Wasserstoff-FluB im Ubergang und im station&en Zustand wurde an metallischen Membranen mit einer G~pha~n-Du~hd~n~n~t~hnik im Te~~ratur~~ich zwischen 100 und 350°C gemessen. Als Membranmaterial dienten ausgeheilte und verformte austenitische rostfreie Stiihle AISI 301, 304, 310 und der ausgeheilte ferritische rostfreie Stahl AL 29-4-2. Aus diesen Megdaten wurden die Permeabilitiit und die effectiven Diffusivittits- und Liislichkeitskonstanten ermittelt. Bis zu einer Verformung von 80% fand sich in dem stabilen Stahl AISI 310 nur eine relativ geringe binderung der Transportkonstanten. Die Verformung von AISI 301 und 304 filhrte zu verschiedenen Mengen an spannungs-induxiertem a’-Martensit, der die effektive Wasserstoffdiffusividit und permeabilitat stark erhohte. Der Zusa~~hang xwischen P~n~nde~ngen und Parametem des Was~~to~rans~rtes wurden untet verschiedenen Annahmen iiber die Mikrostruktur modelliert. Die Werte fir egektive Liislichkeit und Dilfusivitiit werden anhand des Einfangs an Versetxungen und des Transports entlang von Versetzungen diskutiert.
1. ~RODU~ON Knowledge of hydrogen transport behavior in a metal is important to the understanding of the kinetics of hydrogen embrittlement of that metal. Austenitic stainless steels are generally considered less susceptible to hydrogen embrittlement than ferritic alloys due, in part, to the lower diffusivity of hydrogen in the austenite than in the ferrite phase. Troiano and his coworkers, however, found that the embrittlement of austenitic stainless steel had essentially the same dependence on strain rate and test temperature as bee metals and the embrittlement phenomenon was probably ~ffusion~ontroll~ [l]. To embrittle f.c.c. metals, one to two orders of magnitude
greater hydrogen content was required than in b,c.c. metals. Generally speaking, metastable austenitic stainless steel alloys such as AI.91 30 1 or 304 are more susceptible to hydrogen embrittlement than stable alloys such as AISI 310. The metastable alloys are those which are relatively easy to transform to martensite when stressed. The pro~nsity for this transfo~ation depends primarily on the alloy composition and temperature. Alloys with more austenite stabilizers or higher stability factor [2] are found to be more resistant to martensitic transformation as a result of deformation or formation of a as a result of casting or annealing. It is generatly acknowledged that tr~sfo~ed CC’(b.c.c.) martensite is a sufficient, but not necessary condition to allow hydrogen 1771
1772
TSONG-PYNG
PERNG and ALTSTETTER:
embrittlement. Alloys which do not form a’ also suffer hydrogen embrittlement, but with less severity. Previous studies have shown that when notched unstable austenitic stainless steel alloys were stressed in gaseous hydrogen or after cathodic charging, stable (su~~tical) slow crack growth (SCG) under constant load was observed, and a‘ martensite was associated with the fracture surface [3-71. Various theories have been proposed to explain the kinetics of hydrogen-induced slow crack growth behavior in steels and other structural materials. It is frequently postulated that in order for a crack to initiate and propagate, hydrogen has to be transported to and accumulated in an “embrittlement region” at or near the crack tip. The rate of crack advance is controlled by the rate of supply and accumulation of hydrogen in that region [g-14]. Knowledge of hydrogen entry and transport behavior in plastically deformed material is, therefore, important to understand the kinetics of SCG of that material. Since the SCG of austenitic stainless steel may also involve phase transformation and a’ martensite may act as a suitable medium for entry into and transport of hydrogen within the matrix, info~ation on the effects of a’ phase on hydrogen transport in austenitic stainless steel is needed to understand the SCG behavior. In this paper we use annealed and homogeneously deformed specimens to simulate the state of material at a crack tip. Hydrogen transport parameters (diffusivity and permeability) and solubility were ~lculat~ from hydro~n flux measurements in a number of stainless steels in both annealed and deformed conditions using a gas-phase permeation technique. The effects of deformation and a’ martensite on hydrogen transport are discussed.
A schematic diagram of the permeation apparatus is shown in Fig. 1. The ion pumped ultrahigh vacuum (UHV) system and operation procedures have been described elsewhere [15]. Briefly, the technique employed a circular membrane specimen (net surface area 13.85cm2) clamped between two UHV chambers, both evacuated to less than 10G6Pa. Hydrogen gas was purified by diffusion through palladium and then introduced into one UHV chamber. The flux of hydrogen passing through the specimen was measured as an ion current in an ion pump in the other UHV chamber, The detection ion pump, Fig. 1, was calibrated by periodically opening the valve to a known capillary leak. The sensitivity of the detection ion pump was 2.0 x lo-*cm3 H,(STP)/s. The inlet hydrogen pressure ranged from 0.1 to 30 kPa. The furnace was capable of heating to 6OO”C, and the specimen tem~rature was constant to within +0.2”C. Most experiments were carried out at IOOtAL 29-4-2 is a product of the Allegheny Ludlum Steel Corporation.
HYDROGEN PERMEATION IN STAINLESS STEELS IDN FLANGE
CONYAINING
GUN
SPECIMEN
YITRANCE !ON PWP-
CALIBW\lED CAPILLARY
HYDROGEN LEAK
Fig. 1. Permeation apparatus schematic.
350°C. Under these conditions, the pressure build-up in the detection chamber was always below 1.33 mPa and was thus negligible with respect to the input pressure. The steady state flux, J,, and lag time, tL, were measured to calculate the ~~eability constant, 4, and an effective diffusivity, I), using the following equations, respectively
(1)
where h is the thickness of specimen and PH, is the inlet pressure of hydrogen gas. Flux transients were continuously recorded after increments in hydrogen pressure in the inlet chamber. Lag times, t,, were taken as the time to a~omplish 0.630 of the change to a new steady state after an inlet pressure increment. Lag time was independent of inlet hydrogen pressure. The ratio of permeability constant to diffusivity is the effective solubility constant, S. Specimen materials were commercial AISI 301, 304, 310 austenitic and AL 29”4-2t ferritic stainless steel alloys. The compositions are listed in Table 1. Some specimens of AISI 301 and 304 which had undergone various degrees of deformation and therefore contained various amounts of a’ (b.c.c.) martensite phase were prepared by a schedule of rolling, annealing, and electro-polishing to the desired thickness. Deformed specimens of AISI 310 were also prepared. The content of x‘ phase was measured magnetically with a commercial ferrite meter using stacks of sheet to simulate a thicker specimen. A sheet specimen of AL 29-4-2 was used in
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
PERMEATION
IN STAINLESS
1773
STEELS
Table 1. Chemical comwsitions fwt%f of the alloys used for permeation tests Material
C
AISI 301 AN 304 AISI 310
0.052 0.062 0.054
AL 29-4-2
0.0029 0.10
P
S
Si
0.023 0.022
0.009 0.020 0.004
0.48 0.57 0.65
17.1 18.35 24.76
0.01
0.009
0.10
29.5
Mn 1.28 1.31 1.90
Cr
the annealed condition. All specimens were electropolished in a chilled (- 2°C) bath of perchloric and glacial acetic acids (1:4), ion sputtered in Ar at 0.3 mA/cm2 for 120min, and then immediately coated with Pd by vapor deposition. Both entry and exit surfaces were sputtered and Pd-coated. In a bulk diffusion-controlled process, the permeated flux of hydrogen is expected to be not only proportional to the square root of input hydrogen pressure, but also inversely proportional to the thickness of specimens. To check this, three or four measurements of diffusivity and permeability under various pressures of hydrogen at each temperature were made for each specimen, In addition, three annealed 301 specimens with different thicknesses were also studied. The reproducibility was excellent for either case. The average values of permeability and diffusivity were used. The results for deformed 301 and 304 alloys which contained various amounts of a’ phase were compared with those. for annealed 301 and 304 austenitic and AL 29-4-2 ferritic alloys. The effects of Tmz
lo+ +
Ni
MO
N
7.25 9.19 19.55
0.24 0.14 0.11
0.038
2.23
3.93
0.012
0.049
Ti
Cu
Co
Sn
0.27 0.17
0.16
0.016 0.012
a’ phase on hydrogen diffusion and permeation were also studied by comparing the results of deformed 301 and 310 alloys. 3. RESULTS
3.1. AL 29-4-2 ferritic stainfess steel The results for a 195 pm thick specimen of AL 29-4-2 ferritic stainless steel are shown in Figs 2 and 3. A distinct break in the diffusivity curve at - 170°C was observed. The diffusivity in the high temperature region (T > 170°C) was approximately 20-30 times less than that of pure a-Fe, but had an identical activation energy of 7 kJ/mol [16]. No break was observed in the permeability curve in the same temperature range. It has been suggested that trapping effects may account for the downturn of diffusivity in the low t~~rature range [l&-20]. It is generally acknowledged that alloying elements in a-Fe may reduce the diffusivity by various degrees, depending upon the types and amounts of alloying elements [ 171. Nelson et al. reported that, depending on the microstructure, the medium alloy steel AISI 4130 had almost the same as or somewhat smaller diffusivity than pure a-Fe [Zl]. It is expected that AL 29-4-2,
AlSi 301 : A
," 24.0
,"
134
*,
38.3
183
10-g
u
t: -
AL 29 -4
1o-‘2[ 1.5
2.0
1000/r
2.5 (K-l)
Fig. 2. Hydrogen diffuaivity in annealed and deformed AISI 301 and in annealed AL 29-4-2.
1.5
I
I
I
1
I
,
I
‘r 8 i
2.0
1000/T
-2
,
I
>
2.5 (K-l)
Fig. 3. Hydrogen permeability in annealed and deformed AISI 301 and in annealed AL 29-4-2.
PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS
with 35% of alloying element in a-Fe matrix, would exhibit much lower diffusivity. The permeability in AL 29-4-2 was a factor of about 10-l 5 less than that in pure a-Fe. Since the effective solubility is the quotient of permeability over diffusivity, this means that the solubility of hydrogen in AL 29-4-2 was higher than in a-Fe. This, again, might be due to the high amount of alloying. The diffusivity and ~~eability constant of this specimen are described by the following equations
of a’ martensite phase (38.5, 63.5, 65.0 and 86.5%, respectively) were then tested. The content of a’ for each specimen did not change as a result of the test. A separate experiment showed that the stress-induced a’ in both AISI 301 and 304 was stable during annealing at 435°C for a period of 4 days. The results for these specimens along with the annealed alloy are summarized in Table 2 and piotted in Figs 2 and 3. With increasing defo~ation the diffusi~ty increased sharply from that of annealed AISI 301 toward that of AL 29-4-2. Furthermore, a distinct break in the diffusivity curve became increasingly apparent at the higher a’ contents. The peak was similar to that in AL 29-4-2, but at a somewhat higher temperature. The activation energy for diffusion in the high temperature region decreased with increasing a’ content but not as markedly as one might expect from the presence of so much b.c.c. phase (Table 2). The permeability of deformed specimens also increased with the degree of deformation. In fact, the specimen with 86.5% 01’was found to have higher permeability than AL 29-4-2. This has also been observed in a cryoformed 301 alloy with approximately 90% a’, which had higher permeability than both AL 29-4-2 and 301 with 86.5% a’ [22]. The f.c.c. phase dissolved more hydrogen than the b.c.c. phase, but hydrogen diffused more slowly in f.c.c. than in b.c.c. These opposing effects cause the permeability of severely deformed AISI 301 to exceed that of annealed AL 29-4-2. However, comparisons between 301 with a large amount of a’ and the 100% a in AL 29-4-2 must be made with caution because of the different solute contents (N 26% vs N 36%), dislocation densities and grain sizes. A comparison between severely deformed unstable alloy and a ferritic steel with the same amount of solute might be more appropriate, though the nature of the solute elements must also be considered, not just the total solute content. The effects of a’ phase on hydrogen effusion and permeation can be seen better in Figs 4 and 5, where the diffusivity and permeability are plotted as functions of a’ content.
1774
TSONG-PYNG
D = 6.40 x 10-sexp
- 7.0 (kJ/mol) RT m*/s, (T > 170°C)
= 8.45 x lo-* exp
(3)
- 33.7 (kJ/mol) RT Cm2Js, (T < 170°C)
Cp= 2.20 x 10m6exp
- 384(kJ/mol) RT
(4)
cm3 (SIP) cm-s-Pa”2’ (5)
3.2. AISI 301 austenitic stainless steel Type 301 alloy is the most unstable of the standard austenitic stainless steels. Even in the fully annealed condition, a small amount of residual ferritic phase might be present in the f.c.c. matrix, depending on the exact composition and annealing temperature. Both annealed and deformed specimens were studied. Three annealed 301 specimens with different thicknesses (81, 109, and 134 pm) were tested first. They all yielded the same diffusivity and ~~eability constants. The diffusivity and permeability constant of hydrogen in the annealed 301 stainless steel follow the Arrhenius relations
D = 1.30 x 10e3exp
-49.2 (kJ/mol) cm2/s RT
4 = 4.96 x 10e6 exp
- 55.0 (kJ/mol) cm3 (STP) RT cm.s*Pa’/*’
(6)
(7) Specimens with various degrees of deformation (24.0, 37.0, 37.5 and 67.5%) and therefore various amounts
3.3. AISI 304 austenitic stainless steel To examine further the effects of defo~ation and the attendant phase transfo~ation two specimens of
Table 2. Pm-exponential factors and activation energies for effective diffusivity and permeability of hydrogen in anneakd and defamed AK1 301 stainless steel % Deformation (% a’) 0 (0) 24 (38.5) 37 163.5) 37.j (65j
(T > 2Oo?z) (T < 200°C)
67.5 (86.5) (T (T < > 2OO*c) 2OWC)
4 (cm’is) 1.30 x 10-S 8.67 x 1O-4 2.70 x IO-’ 3.10 x 10-3 1.25 x IO-’ 4.60 x 10-S 1.64 x 10-J
(kg& 49.2 42.5 39.3 40.1 45.6
(cm3 (STPgn.s.
Pal@) (k.$L)
4.96 x 1O-6 2.01 x 10-b 6.00 x lo-”
55.0 47.4 45.4
4.46 x 1O-6
44.6
5.50 x 10-G
40.2
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
PERMEATION
IN STAINLESS STEELS
1775
r (“C) lo+
350
300
250
I
I
200 I def. WI
AL 29-4-2. AISI
304
-*-
o
-
o’l% 0
50.0
9.5
W7 10-g
-
10-a
2.0
1.6
1.6
1000/T
Fig. 6. Hydrogen diffusivity in annealed and deformed AISI 304.
lo+ 0
20
40
60
(K-‘1
60
a’%
Fig. 4. Effect of a’ martensite on enhancement of hydrogen diffusivity in deformed AISI 301.
a more stable, type 304, stainless steel were tested. One specimen with a thickness of 97 pm was in the annealed condition. Another one was cold rolled to a 50% reduction of thickness, resulting in 9.5% a’
to a thickness and permeation
martensite, and then electropolished of 112 pm. The hydrogen diffusion
results at T = 200-350°C for the two specimens are plotted in Figs 6 and 7. The data for the annealed one can be described best by the relations D = 7.69 x lo-’ exp
- 53.3 (kJ/mol) ~*/s RT
(8)
10-e
T (‘C) 10-g
i
350
300
250 I
I
200 I def.(%I
a’(%
lo- ‘2
1 lo- l2 1 0
I 20
I 40
I 60
I 60
100
1o-‘31 16
a’%
Fig. 5. Effect of a’ martensite on enhancement of hydrogen Permeability in deformed AISI 301.
1.6
2.0 lOOO/
Fig.
7.
T (K-l)
Hydrogen permeability in annealed and deformed AISI 304.
1776
TSONG-PYNG
PERNG and ALTSTE’ITER:
HYDROGEN PERMEATION
IN STAINLESS STEELS
Table 3. Effective hydrogen transport parameters in annealed and deformed AISI 310 stainless stee1
Deformation cm
4 m is)
38
7.45 x 10-J 7.16
52.8 53.0
3.30 x 10-s 3.50 10-J
58.5 57.3
2.15 3.04
4.52 5.47
60 80
3.70 x 10-j 3.48 x IO-’
52.6 52.8
3.09 x 10-S 3.30 x 10-s
58.7 59.1
5.19 5.89
6.15 6.33
4 = 2.60 x lO-‘exp
(cm3(~T~~~~.s.~a’/‘~
- 57.8 (kJ/mol) cm3 (STP) RT cm*s*Pa1’2 (9)
and the following equations are used to describe the deformed one -40.0 (kJ/mol) cm2/s D = 4.94 x 10m4exp RT
(~,~,~
(at. ~~,~a’l’~
observed after up to 80% deformation. The results for these specimens are shown in Table 3 and plotted in Figs 8 and 9. The effective diffusivity and permeability constant of hydrogen in the annealed 310 alloy can be described by the following equations D = 7.16 x 10e3exp
- 53.0 (kJ/moi) RT cm*/s
Cp= 3.50 x low5 exp
- 58.5 (kJ/mol) cm3 (STP) RT cm.s*Pa”2’
(10)
-49.1 (kJ/mol) cm3 (SIP) 4 = 7.27 x lOa exp RT cm*s.Pa”2’
~~,~
(13)
(11) It was found that both hydrogen diffusivity and permeability in the deformed specimen were higher than those for the annealed one, probably due to the presence of a’ phase. The deformed specimen also had lower activation energies of diffusivity and permeability.
3.4. AISI 310
austenitic stainless Steel
Finally, a completely stable austenite, AISI 310, was used to assess the effects of deformation alone on hydrogen transport. Four specimens of this alloy which had undergone 0,30,60 and 80% deformation (with thicknesses of 175, 188, 187 and 82 pm, respectively) were tested. No phase transformation was
(12)
Both quantities were slightIy lower (about a factor of two at 300°C) than those reported by Quick and Johnson [24]). The permeabilities for the deformed specimens were found to differ by less than 25% from those for the annealed specimen. The ~ffusi~ties in the 60 and 80% deformed specimens were about a factor of 2 lower than those in the annealed and 30% deformed alloys. No break in the diffusivity curve in the low temperature region was observed.
4. DISCUSSION
Three annealed austenitic stainless steel alloys have been studied. It has been reported that hydrogen transport parameters in austenitic stainless steels are Tf%l
T PC) ,o_9y0
3yo
150 ,
200 ,
270
,
o Annealed
AlSl 310
1.6
18
2.0
1000/T
0
30’hCW 6O%CW
.
eo%CW
l
2.2
AISI 310
.
\
2.4
(K-l)
Fig. 8. Hydrogen diffusivity in anneaied and deformed AISl 310.
1.6
1.8
. 30% 13 60Y.
cw CW
aOY.CW
2.0
2.2
24
1000/T (K-‘1
Fig. 9. Hydrogen liability in annealed and deformed AISi 310.
TSONG-PYNG PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS 4 = 1.20 x IOU5exp
1771
56.1 (kJjmo1) cm3 (STP) RT
cm*s.Pa”*’
(15) Type 301 0
l
Equations (14) and (15) are shown as straight lines in Figs 10 and 11, respectively, and all data points fall within a factor of two of these lines. A general expression for effective hydrogen solubility constant in austenitic stainless steels can be obtained as the quotient of equations (14) and (15).
0
302 A&
S = 5.97 x 10s3 exp -6.8~~‘mo1
lo-lot/ , , , , i.6
2.0
1.8
1000/T
I
1
I
I
2.4
2.2
(K-‘I
Fig. 10. Comparison of hydrogen diffusivity in four types of
austenitic stainless steel.
relatively independent of the austenite composition [25]. In this experiment we have also found that the differences of diffusivity and permeability among these alloys were relatively small. The hydrogen diffusivity and permeability in the annealed AISI 301, 304 and 310 alloys are summary in Figs 10 and Il. The data for AISI 302 from a separate study [23] using the same apparatus are also included. A least-squares fit of the diffusivity and permeability constant data for the four alloys yielded I) = 2.01 x low3 exp
-49.3 (kJ/mol) cm2js (14) RT
TIT) 350 10-91
300
250
200
150
I
I
I
I
I
Aurtenitic stainless steels
1000/T (K-‘I
Fig. 11. Comparison of hydrogen permeability in four types of austenitic stainless steel.
~~~~~.
(16)
The results of this work are in good agreement with previously reported values using similar experimental conditions [21,24-40]. The effects of defo~ation on hydrogen permeation in austenitic stainless steels were studied using AISI 301,304 and 3 10 alloys. Hydrogen permeability and diffusivity for each temperature were repeatedly measured and found to be the same each time, even when using different inlet hydrogen pressures. This implies that the filling of pre-existing deep traps was not an important factor under our test conditions. Deformation of type 301 alloy yielded up to 86% of a’ martensite and enhanced both hydrogen diffusivity and permeability by up to two orders of magnitude. In type 310 alloy, no a’ was observed after up to 80% deformation. The permeability was found to be almost the same for the annealed and deformed specimens, whereas the diffusivity was slightly smaller in the severely deformed ones. The difference in behavior between type 301 and 310 alloys is almost certainly due to the different phases present after deformation and not due to different dislocation confi~rations and densities. The effect of dislocations on hydrogen diffusion is somewhat controversial. In annealed f.c.c. metals, it has been concluded that dislocation trapping is frequently not important [41,42]. In deformed f.c.c. alloys, Louthan et al. reported some data which indicated that hydrogen diffusi~ty was not affected by cofd work 1251.In the present work, however, a slight decrease of hydrogen diffusivity was observed in the severely deformed AISI 310 specimens. It is possible that the lower hydrogen diffusivities were caused by dislocation trapping, and this hypothesis will now be examined. When hydrogen enters a specimen with traps the time for the transient to appear on the exit surface is increased, partly because the hydrogen must first fill the deep traps. This is reported, then, as a lower effective diffusivity. However, once the trap sites are saturated the steady state hydrogen flux is essentially unchanged, and thus the permeation is unchanged. This is the effect observed in the case of deformed AISI 310 steel. In order to minimize systematic errors in thickness measurement, thickness uniformity, surface condition, etc. we have taken the ratio of permeability to effective diffusivity, i.e. the
1778
PERNG and ALTSTETTER:
TSONG-PYNG
effective solubility, to characterize the trap population introduced by cold work. From equations (12) and (I 3), the effective hydrogen solubility constant in the annealed 310 alloy can be obtained S =4.89 x lo-‘exp
= 2.63 x 10”exp
= 3.04exp
-5.47 (kJ/mol) RT
- 5.47 (~/mol) RT
- 5.47 (kJ/mol) R1”
cm3H,(STP) cm3metal. Pa”2
atom ~ cm3 *Pall*
at ppm/Pa”*.
(17)
The formulae for the deformed specimens can also be constructed from the constants included in Table 3. The total concentration of hydrogen in the annealed alloy in PH2= 1 atm and at 573 and 423 K (within the temperature range of this experiment) are calculated to be 307 and 205 at. ppm, respectively. These values along with those for the deformed specimens are shown in Table 4. The increments of hydrogen concentration in the deformed alloys can be seen. We assume that in the annealed specimen dislocation trapping is negligible, so that the effective (or total) ~n~ntration, C,, is equal to the lattice concentration, C,. If we express concentrations in units of H atoms/cm3 and use the symbols C; and C;, then C; is calculated to be 2.66 x lOI and 1.77 x lOl9 H/cm3 at 573 and 423 K, respectively. The increments in hydrogen concentration, C; - C;, due to deformation are taken to be the amounts of hydrogen in trap sites. These are tabulated in Table 4, along with the fraction of the hydrogen in trap sites, (C+-C;)/Ci. This fraction increases rapidly at first and then less rapidly, which is not unexpected. It might be commented here that one would expect this behavior to be rather different in ferritic materials, where the lattice solubility is several orders of mag~tude lower. In this case one would expect trapping effects to dominate. Experiments are underway to examine this in a ferritic stainless steel. Thomas [41] reported that the hydrogen binding to dislocations in f.c.c. metals was essentially due to
Table 4. Cakulated
HYDROGEN
PERMEATION
IN STAINLESS STEELS
elastic interactions, with a binding energy of 9.6-13.5 kJ/moI. This is different from that in iron or ferritic steels, where core interaction constitutes a significant contribution to hydrogen trapping [19]. From the binding energy the fraction of trap sites occupied by hydrogen atoms, C,, can be calculated at any temperature. Assuming that all trap sites have the same energy of 12 kJ/mol and that this energy is independent of dislocation density or hydrogen content, a very great simplification for the case of a dislocation stress field, we estimate that the density of trap sites, N, (= C;CC;/C,), ranges from 0.5 to 5.0 x 102’ cm-j. This is the same order of magnitude as that estimated by Oriani for deformed iron and ferritic steels (431. The estimated increase in N, was not large in going from 30 to 80% deformation, considering the orders of magnitude increase in dislocation density. At a relatively high average dislocation density of lO’*cn-* there would be several H atoms per lattice plane along a dislocation line. We note that taking 12 kJ/mol as the binding energy, C, at 423 K should have been higher by 60% than that at 573 K. The fraction of hydrogen in trap sites, (C; - Ci)/C;, seemed to be insensitive to this difference. Nevertheless, this simple calculation did illustrate that in the deformed 310 stainless steel, it is reasonable to assume that increased dislocation density could be effective in trapping a certain amount of hydrogen, increasing the effective hydrogen concentration, and slightly reducing the effective hydrogen diffusitivity. In the deformed type 301 alloy, both hydrogen diffusivity and permeability were increased. Dislocation trapping certainly could not explain this result. On the other hand, if disl~ations provided preferential short-circuit paths for fast diffusion, one might expect to see this behavior. However, the experiments with two alloys of different austenite stability showed that the extent of enhancement was a function of u’ content rather than the amount of deformation. Type 304 was more stable than type 301 alloy, since only 9.5% a‘ phase was induced after 50% deformation, whereas 65.0% LX’phase was formed in a 37.5% deformed 301 alloy. Had dislocations enhanced hydrogen diffusion, the enhancement in the 50% deformed 304 alloy should have been higher than in
effective solubility’ and trapping parameters for hydrogen in AISI 310 stainless steel Total concentration
Trapped concentration c;-c; (lOI H/cm’)
Fraction of H in traps c; - c;jc;
Deformation (W
(at.zptn)
0 30 60 80
307 340 456 498
0 :
205 293 244
2.53 2.11
0.76 0.34
0.30 0.16
80
311
2.69
0.92
0.34
(10’e&m3)
At T = 573 K (300°C) 2.66 2.95 0.3 3.95 1.3 4.31 1.6 At T-423K
‘1 atm pressure of H,.
1.77
0.10 0.33 0.38
(150°C)
TSONG-PYNG PERNG and ALTSTETTER:
HYDROGEN PERMEATION IN STAINLESS STEELS
Table 5. Effective hydrogen solubility” in AISI 301 alloy at 200 and 300°C Deformation (X) 0 24.0 37.0 37.5 67.5
& 0 38.5 63.5 65.0 86.5
Concentration of H (at. ppm) 200°C 300°C 173 132 93 91 60
223 I64 122 111 76
‘1 atm pressure of H,.
the 37.5% deformed type 301. This was not the case (see Figs 2 and 6). Other evidence for this argument is the decrease of effective hydrogen solubilities resulting from the deformation in 301. The effective hydrogen solubilities at 200 and 300°C in the annealed and deformed 301 specimens calculated from the ~~eability and ~ffusivity are shown in Table 5. Unlike the deformed AISI 310 alloy, the effective hydrogen solubility decreased with increasing amount of deformation. This was ascribed to the presence of tl’ phase which accommodated much less hydrogen in solution than the y phase. The results of this series of experiments led us to conclude that the enhan~ment of hydrogen diffusivity and permeability in deformed 301 was not due to the enhanced dislocation density, but the presence of a’ martensite. Only in the deformed 310 alloy, where no tl’ was formed was there an effect which could be ascribed to dislocation trapping. It is noted that the activation energy for hydrogen diffusion at low tem~ratures in the 67.5% deformed 301 steel (48.8 Wfmol) was much larger than that in annealed AL 29-4-2 (33.7 kJ/mol). This is not unexpected if we consider that the highly dislocated Co martensite may yield significant trapping. Deformed AL 29-4-2 may well exhibit higher activation energies than the annealed alloy. Experiments are underway to examine this possibility. It is also interesting to note that the hydrogen transport behavior in deformed 304 was different from both 301 and 310 alloys. For deformed 304, the hydrogen ~~eability increased more than the diffusivity, resulting in a higher apparent solubility of hydrogen. Since Type 304 is more stable than type 301, only 9.5% tl’ was induced after 50% deformation. Interpolation from Table 3 shows that with this amount of deformation, there was -35% increase of hydrogen solubility in AISI 310 at T = 200-3~°C. On the other hand, inte~olation from Table 5 shows that the presence of 9.5% a’ in AISI 301 may result in only -5% decrease of hydrogen solubility in the same temperature range, neglecting the effect of dislocation trapping of hydrogen. These two opposite effects may lead to a higher effective hydrogen solubility in deformed 304. From Figs 6 and 7, the ratios of hydrogen solubility in deformed and annealed 304 were found to be 1.34 and 1.66 at 200 and 3OO”C, respectively. The motivation for studying hydrogen transport
1779
behavior in the deformed austenitic stainless steel originated from the fact that ~1’ martensite was associated with the fracture surface and the region near the crack tip when unstable alloys are stressed in hydrogen. The much higher hydrogen diffusivity and ~~eability in martensite than in the f.c.c. matrix rationalizes the hydrogen embrittlement susceptibility and hydrogen-induced SCG behavior of AISI 301. While it is very difficult to directly measure hydrogen transport at a crack tip, it is desirable to obtain these parameters for various martensite contents in a unifo~ly deformed alloy, so that fluxes into and out of the crack tip region may be estimated. Several models were used to quantitatively analyze the enhanced diffusivity and permeability in alloys containing various amounts of a’ in a y matrix. The quantitative predictions for the phase mixtures depend very greatly on the ~~rostructure. First, upper and lower bound estimates of the diffusivity arose from assuming that the a’ and y were present as continuous lamellae and the hydrogen flux passed either parallel to the lamellae (upper bound) or through the lamellae in series (lower bound). The diffusivity in the parallel case can be written as a linear combination of diffusivities of the a’ and y phases. The series case uses a Iinear combination of resistivities (reciprocal diffusivity) of the two phases. These two estimates are compared in Fig. 12 to the values of diffusivity measured to 200°C for cold rolled AISI 301 (shown as points in Fig. 12). For this figure the computed curves used the measured value of diffusivity (4.79 x 10-9cm2/s) for austenite and an assumed value of 1.5 x 10m6cm’/s for the martensite. Figure 12 shows that this assumed lamellar micro-
I
’
’
0
20
40
60
80
I
100
Volume % 0’
Fig. 12. Estimation of the enhanced hydrogen diffusivity in deformed AISI 301 . at 2OOT.. Points are experimentally .
determined values.
1780
TSONG-PYNG
PERNG and ALTSTE’ITER:
structure gives values which can differ by more than an order of magnitude from the measured diffusivities. A model used by Louthan et al. [33] gave somewhat better results than the series curve in Fig. 12, but values were still almost an order of magnitude too low. The interpolation formula, equation (18), was derived by McLachlan [451 for the conductivity of a mixture containing a volume
HYDROGEN
f
(Pin- PJ
For a thr~~imensional body Pi= (l)Jvbi3 where i stands for ~1,y or the mixture and A is a fitting parameter, which can be thought of as a fractal dimension. The plot of equation (18) shown in Fig. 12 for A = 1.487 and D,. = 1.5 x 10V6cm*/s, fits the experimental points remarkably well. The values of A and DES were optimized to give this close fit. It is seen that the optimized value of I), for a 100% a’ AISI 301 alloy is about an order of magnitude less than the measured value for AL 29-4-2. This difference is probably due to differences between, (a) the model micro-structure and the true microstructure of AISI 301, (b) the two alloy compositions and (c) the defect contents of 100% (Y’and annealed ferrite. The application of the above models to prediction of permeation constants gave results similar to Fig. 12. The good fit to the experimental data was obtained for A = 1.487. and & = 5 x lo-” cm3/ cm*s.Pa’lz. This is a factor of 4 greater than that for AL 2942. Since saturated deep traps are not expected to affect the permeability, we are led to the conclusion that the alloy content is one of the principal reasons for this difference. The estimated values of D and # for 301 containing 100% tl’ are consistent with a linear dependence of hydrogen solubility on a’ content. 5. CONCLUSIONS 1. A gas-phase permeation technique fully used to study hydrogen transport
was success-
in a number of stainless steels. The results for the annealed austenitic stainless steels were in good agreement with reported values determined by similar techniques. 2. The highly alloyed AL 29-4-2 ferritic stainless steel had lower diffusivity and permeability than lower alloy ferritic steels by a factor of 10-20. 3. Deformation of the metastable austenitic AISI 301 and 304 alloys resulted in various amounts of a’ martensite and enhanced both hydrogen diffusivity and permeability. 4. No phase transformation was observed in AISI 310 after up to 80% deformation. The effective hydrogen diffusivities in the deformed 310 were slightly decreased, probably due to dislocation trapping.
IN STAINLESS STEELS
Acknowledgements-This research was supported by the National Science Foundation under Contract NSF DMR 83-03421. The AL 2942 ferritic stainless steel was supplied by the Allegheny Ludlum Steel Corporation. We wish to thank Professor D. S. McLachlan of the University of the Witwatersrand, Johannesburg, South Africa for his help in fitting the hydrogen diffusivity data.
fraction, f, of spheres of highly conducting material (cc’) immersed in a lower conductivity medium (y).
P,,, + (A - l)P,
PERMEATION
REFERENCE 1. M. B. Whiteman and A. R. Troiano, Corrosion 21, 53 (1965). 2. C. J. Novak, Handbook of Stainless Steels (edited by D. Peckner and I. M. Bernstein). DD. 4-29. McGraw-Hill. New York (1977). ‘. __ 3. D. Eliezer, D. G..Chakrapani, C. J. Altstetter and E. N. Push. Metall. Trans. A 1OA. 935 (1979). 4. R.iitu, N. Narita, C. Altstetter, H.‘Bimbaum and E. N. Pugh, Merall. Truns. A llA, 1563 (1980). 5. N. Narita and H. K. Bimbaum, Scrip& metall. 14.1355 (1980). 6. S. Singh and C. Altstetter, Melall. Trans. A 13A, 1799 (1982). 7. G. Schuster and C. Altstetter. Metal% Trans. A I4A. 2085 (1983). Acta metall. 22. 1965 8. R. A. Oriani and P. H. Joseohic. . (1974). 9. R. A. Oriani and P. H. Josephic, Acta metall. 25, 979 (1977). 10. W. W. Gerberich, Y. T. Chen and C. St. John, Metall. Trans. A 6A, 1485 (1975). 11. W. W. Gerberich and Y. T. Chen, Metall. Trans. A 6& 271 (1975). 12. W. W. Gerberich, J. Garry and J. F. Lessar, E#ecfs of Hydrogen on Behavior of Materials (edited by A. W. Thompson and I. M. Bernstein), p. 70. T.M.S.-A.I.M.E, Moran, Wyoming (1976). H. P. van Leeuwen, Corrosion 29, 197 (1973). ::: H. P. van Leeuwen, Corrosion 31, 154 (1975). 15. P. H. Studebaker, M.S. thesis, Univ. of Illinois at Urbana-Champaign, Ill. (1981). 16. N. R. Quick, Ph.D. thesis, Cornell Univ., N.Y. (1975). 17. J. VBlkl and G. Alefeld, Diffusion in Solids (edited by A. S. Nowick and J. J. Burton), p. 252. Academic Press, New York (1975). 18. C. A. Wert, Hydrogen in iwetals, Topics in Appiied Physics (edited by G. Alefeld and J. Vitlkl), Vol. 29, p. 305 (1978). 19. J. P. Hirth, MetaN. Trans. A. llA, 861 (1980). 20. H. H. Johnson and R. W. Lin, Hydrogen E&as in Metals (edited by I. M. Bernstein and A. W. Thompson), p. 3. T.M.S.-A.I.M.E., Warrendale, Pa (1980). 21. H. G. Nelson and J. E. Stein, NASA Rep. TND-7665 (1973). 22. T. P. Pemg and C. J. Altstetter, Scripta metait. 18, 67 (1984). 23. T. P. Pemg and C. J. Altstetter. In press. 24. N. R. Quick and H. H. Johnson, MetaN. Trans. A lOA, 67 (1979). 25. M. R. Louthan Jr and R. G. Derrick, Corrosion Sci. 15, I
565 (1975).
26. R. H. Collins and J. C. Turnbull, Vacuum 11, 114 (1961). 27. H. L. Eschbach, F. Gross and S. Schulien, Vacuum 13, 543 (1963). 28. J. R. Phillips and B. F. Dodge, A.Z.Ch.E.JI 14, 392 (1968).
29. K. F. Chaney and G. W. Powell, Metall. Trans. 1,2356 (1970). 30. W. J. Kass and W. J. Andrzejewski, AEC Rep. Sc-Dr72-0136 (1972).
TSONG-PYNG
PERNG and ALTSTETTER:
HYDROGEN
31. J. H. Austin, T. S. Elleman and K. Verghese, J. nuci. Muter. 48, 307 (1973). 32. W. A. Swansinger, R. G. Musket, L. J. Wierick and W. Bauer, J. nuct. Mater. 53, 307 (1974). 33. M. R. Louthan Jr, J. A. Donovan and G. R. Caskey Jr, Nucl. Teclr. 26, 192 (1975). 34. V. A. Maroni, E. H. van Deventer, 1. A. Renner, R. H. Pelto and C. J. Wierdak, CONF-750989, IV-329 (1975). 35. J. T. Bell. J. D. Redman and F. J. Smith. ORNL 5297. 28 (1977): 36. E. H. Van Deventer and V. A. Maroni, ANL/FPP-77-7, 17 (1977). 37. H. K. Perkins and T. Noda, J. nucl. Mater. 71,349 (1978).
PERMEATION
IN STAINLESS STEELS
1781
38. W. Conley, P. Studebaker and C. Altstetter, Ref. [20], p. 169. 39. M. Matsuyama and J. D. Redman, MetaN. Trans. A 14A, 498 (1983). 40. R. A. Outlaw and D. T. Peterson, Metal!. Trans. A 14A, 1869 (1983). 41. G. J. Thomas, Ref. [20], p. 77. 42. H. K. Bimbaum and C. A. Wert, Ber. Btmse. Phyz. Chem. 76, 806 (1972). 43. R. A. Oriani, Acra meratt. 18, 147 (1970). 44. J. F. Breedis and W. D. Robertson, Acta metatt. 10, 1077 (1962). 45. D. S. McLachlan, J. Whys. C. To be published.