Hydrogen Absorption into Austenitic Stainless Steels Under High-Pressure Gaseous Hydrogen and Cathodic Charge in Aqueous Solution
Hydrogen Absorption into Austenitic Stainless Steels Under High-Pressure Gaseous Hydrogen and Cathodic Charge in Aqueous Solution
MASATO ENOMOTO 0
LIN CHENG 0
HIROYUKI MIZUNO 0
YOSHINORI WATANABE 0
TOMOHIKO OMURA 0
JUN'ICHI SAKAI 0
KEN'ICHI YOKOYAMA 0
HIROSHI SUZUKI 0
RYUJI OKUMA 0
0 MASATO ENOMOTO, Emeritus Professor, is with Ibaraki University , Nakanarusawa, Hitachi 316-8511 , Japan. Contact
Type 316L and Type 304 austenitic stainless steels, both deformed and non-deformed, were hydrogen charged cathodically in an aqueous solution as well as by exposure to high-pressure gaseous hydrogen in an attempt to identify suitable conditions of cathodic charge for simulating hydrogen absorption from gaseous hydrogen environments. Thermal desorption analysis (TDA) was conducted, and the amount of absorbed hydrogen and the spectrum shape were compared between the two charging methods. Simulations were performed by means of the McNabb-Foster model to analyze the spectrum shape and peak temperature, and understand the effects of deformation on the spectra. It was revealed that the spectrum shape and peak temperature were dependent directly upon the initial distribution of hydrogen within the specimen, which varied widely according to the hydrogen charge condition. Deformation also had a marked effect on the amount of absorbed hydrogen in Type 304 steel due to the straininduced martensitic transformation.
HYDROGEN is a clean energy source which has the
potential to be widely used in various energy systems in
the near future. Fuel cell vehicles equipped with
cylinders containing high-pressure hydrogen gas up to
70 MPa will be commercialized within a few years,
and hydrogen stations which supply fuel cell vehicles
with gaseous hydrogen are under construction.[
safe use and public acceptance of compressed hydrogen
gas are important issues for the developments of
However, it is well known that high-pressure hydrogen
gas degrades mechanical properties of many kinds of
metallic materials, which are used for the storage and
transportation of hydrogen.[
] The environmental
degradation in gaseous hydrogen is called hydrogen
environment embrittlement (HEE) or hydrogen gas embrittlement
(HGE), which occurs when a hydrogen-free material is
mechanically tested in gaseous hydrogen at ambient
temperature. HGE is recognized as one type of hydrogen
embrittlement (HE) caused by hydrogen entry into metallic
materials from gaseous hydrogen environments. Thus,
effects of gaseous hydrogen on tensile or fatigue properties
of several candidate materials have been the subject of keen
interest in the past.[
] Up to now, it was reported that
metals with a face-centered cubic (fcc) structure such as
stable austenitic stainless steels, e.g., Type 316L or
aluminum alloys, are not susceptible to HGE.[
] Because the
assessment of HGE needs special testing equipment
capable of controlling extremely high-pressure gaseous
hydrogen, the limited availability of such testing
equipment hampers is a serious obstacle to the collection of a
large amount of test data necessary for establishing
standards and regulations for hydrogen storage materials.
Cathodic hydrogen charge in an aqueous solution is a
convenient test method to assess HE. Indeed, a large
amount of data has been reported with respect to tensile
properties of metals under cathodic charge in acidic
solutions with catalysts which promote hydrogen entry
into the materials. Hence, cathodic charging tests can be
a convenient alternative to evaluate HGE. However,
there are a limited number of studies regarding the
relationship between HGE and embrittlement in an
aqueous solution. It is reported that the degrees of
Fig. 1—Configuration and dimensions (in mm) of specimen for
mechanical degradation of the tensile properties under
cathodic charge and high-pressure gaseous hydrogen are
similar for Type 316L and 304L austenitic stainless
In this study, to explore the suitable conditions of
cathodic charge for simulating hydrogen absorption
from high-pressure gaseous hydrogen, Type 316L and
Type 304 austenitic stainless steels, both deformed and
non-deformed, were hydrogen charged cathodically and
by exposure to high-pressure gaseous hydrogen. It is
well known that thermal desorption analysis (TDA) is a
useful tool to analyze the diffusion and trapping of
hydrogen in the material which is thought to be closely
related to HE. Hence, TDA was conducted to compare
the amount of absorbed hydrogen and TDA spectrum
shape between the two charging methods. Furthermore,
simulations were performed using the McNabb–Foster
model to analyze in detail the differences between the
steels in the TDA spectrum shape and peak temperature,
defined as the temperature at which the hydrogen
desorption rate passes through a maximum, and
understand the effects of deformation on hydrogen
absorption. It will be shown that TDA spectra in stainless steel
vary widely with the distribution of hydrogen prior to
the analysis, and this has to be taken into account in
establishing the relationship between different hydrogen
charging methods, which are usually performed under
different conditions with respect to hydrogen pressure,
temperature, and time.
II. EXPERIMENTAL PROCEDURE
Commercial Type 304 and 316L stainless steels, cold
rolled to 0.77-mm-thickness sheet and annealed above
1273 K (1000 C), were used in this study. Table I
shows the chemical composition of the alloys studied. In
order to investigate the effects of deformation on
hydrogen absorption, tensile specimens whose configuration
is given in Figure 1 were machined from the sheet. They
were deformed in tension to a strain of 10 and 40 pct.
These specimens were polished and etched in aqua regia
(a mixture of 3 hydrochloric acid and 1 nitric acid) for
microstructural observation. Figures 2(a) through (d)
show the optical micrographs of non-strained and
40 pct strained specimens of the two steels. The
microstructure of non-strained specimens consists of
polygonal austenite grains with annealing twins, while grains
are somewhat elongated along the tensile direction in the
40 pct strained specimens.
The microstructure of Type 304 stainless steel is
metastable austenite. The formation of martensite was
detected by both X-ray diffractometry using Co Ka
radiation and magnetic measurement: the relative
permeability, i.e., ratio of the permeability of specimen to
vacuum, was measured using the so-called l-meter
(LP141, Denshijiki Industry Co., Ltd) in which the contact
probe detects a shift in the magnetic field lines in the
applied magnetic field. The XRD spectra in Figure 3
indicate the formation of strain-induced martensite in
the 40 pct strained Type 304 steel. In the 10 pct strained
specimen, the strain-induced martensitic transformation
was not detected by the measurement of permeability,
either, as shown in Figure 4.
Hydrogen charge was conducted by two methods. One
was exposure to gaseous hydrogen, and the other was
cathodic electrolysis. The apparatus used for exposure to
high-pressure hydrogen is illustrated in Figure 5(a). It can
simulate the gaseous environment of high-pressure
hydrogen in the container of fuel cell vehicle and
hydrogen station. Table II shows the condition of
exposure to gaseous hydrogen. According to the preliminary
estimate, it is assumed in Condition 1, i.e., exposure to
hydrogen atmosphere of 98 MPa at 523 K (250 C) for
72 hours, that absorbed hydrogen is distributed
uniformly from the surface to the center of specimen. In
contrast, Condition 2, exposed at 358 K (85 C) for
1000 hours, which is supposed to represent the condition
of practical use, assumes that the absorbed hydrogen
concentration is non-uniform, high near the surface of
specimen and low in the central region. The electrolysis
cell used for cathodic charge is illustrated in Figure 5(b).
This charging method is widely used because the amount
of charged hydrogen can be controlled relatively easily.
Here, cathodic charge was conducted at a constant
electric current of 1 mA/cm2 for 96 hours in an aqueous
solution of 3 mass pct NaCl with the addition of 3 g/L
NH4SCN. The parallel part, 20 mm in length, 4 mm in
width, and 0.77 mm in thickness, cut from the
hydrogencharged specimen was subjected to TDA using a gas
chromatograph or a quadrupole mass spectrometer at a
heating rate of 100 K/h.
Figures 6(a) and (b) show TDA spectra of Type 304
and 316L steels, respectively, exposed to gaseous
hydrogen under Condition 1. It is seen that all spectra exhibit
Fig. 6—TDA spectra of (a) Type 304 and (b) Type 316L steels
exposed to gaseous hydrogen under Condition 1.
single desorption peaks no matter whether specimens
are deformed or not. The desorption peak of the 40 pct
strained Type 304 steel is considerably lower and is
located at a temperature 60 K to 70 K lower than the
specimens of no-strain and 10 pct strained. In contrast,
the desorption peak of Type 316L steel specimen is
higher, and the difference in temperature from the other
two specimens is only 20 K to 30 K.
In Figure 7, the desorption spectrum area, i.e., the
amount of evolved hydrogen expressed in the mass
fraction (ppm) of hydrogen in the specimen, is plotted
against tensile strain. Although the increase in the
amount of absorbed hydrogen in the 40 pct strained
Type 316L specimen is not significant, it is clear that
deformation decreased the peak area in Type 304 steel.
It is tentatively proposed that the decrease in the peak
area of 40 pct strained Type 304 steel is caused by the
formation of martensite upon deformation, as will be
TDA spectra of Type 304 and 316L specimens exposed
to gaseous hydrogen under Condition 2 are shown in
Figures 8(a) and (b), respectively. In Figure 8(a), only the
spectrum of 40 pct strained Type 304 specimen has a
single peak. The spectra of the other specimens have a
broad peak which could be divided into two flat peaks. On
the other hand, in Figure 8(b), the spectra of all three
specimens have broad peaks in Type 316L steel. Figure 9
shows the variation of peak area with tensile strain. It is
seen that the peak area increased with tensile strain in
Type 304 steel, while no appreciable change is observed in
the other steel. These are contrary to the results obtained
under charging Condition 1.
C. Cathodic Electrolysis
Figures 10(a) and (b) show TDA spectra of Type 304 and
316L specimens, respectively, which were cathodically
charged. The peak desorption temperature is around
373 K (100 C) in all specimens. However, there is a sharp
increase in the peak height with tensile strain in the former
steel, while the increase in peak area is small in Type 316L
steel. The TDA spectrum peak area is plotted against tensile
strain in Figure 11. A marked increase in the peak area is
observed in the 40 pct strained Type 304 steel, while only a
very moderate increase is observed in Type 316L steel. These
changes in peak area with strain are similar in sign, but
occurred to a greater extent than those in Condition 2. This
is tentatively ascribed to the inhomogeneous initial
distribution of hydrogen caused by insufficient charging time.
IV. SIMULATION OF TDA SPECTRA BY
A. Non-deformed Specimen Charged Under High
According to McNabb and Foster,[
] the diffusion
and trapping of hydrogen in metallic materials can be
described by the following equations:
where c is the hydrogen concentration in the lattice, Nt
is the density of trap sites, ht is the occupancy fraction
at the trap site, D is the diffusivity of hydrogen in the
k ¼ k0 exp
p ¼ p0 exp
¼ p0 exp
EB þ ED
are the kinetic coefficients of trapping and detrapping,
respectively. Here, ED is the activation energy of lattice
diffusion and EB is the binding energy of the trap site. Ed
(=EB + ED) is the activation energy of detrapping. k0
and p0 are related to each other as k0 = p0/NV,[
NV (=8.77 9 1028 m 3) is the number of available sites
of hydrogen in the fcc lattice per unit volume. Table III
shows the activation energy and the pre-exponential
factor of diffusivity of hydrogen in Type 304 and 316L
steels reported in the literature.
Since the fraction of hydrogen in the trap sites is
expected to be small in austenitic stainless steel due to
the small binding energy of hydrogen to dislocations,
the above equations were solved assuming Nt = 0 for
specimens without tensile strain. Figures 12(a) and (b)
show the TDA spectra of hydrogen in Type 304 and
316L specimens exposed to a hydrogen atmosphere of
98 MPa (Condition 1), calculated assuming the surface
concentration of hydrogen equal to 100 ppm. The
hydrogen concentration at the specimen surface was
determined by trial and error comparing the whole area
between simulated and experimental TDA spectra.
Diffusivities reported by Sakamoto and Katayama
] and Mine and Kimoto (M–K)[
] were used
for the Type 304 steel McNabb–Foster model
calculations. For the Type 316L calculations, diffusivities
published by Brass and Chene (B–C)[
] and Tanabe
and Imoto (T–I)[
] were employed. If one employs these
diffusivities, the hydrogen distribution is uniform all
over both specimen types under Condition 1. It is seen
in Figure 12 that the S–K and B–C diffusivities can
reproduce TDA peaks quite well in Type 304 and 316L
specimens, respectively, while the other diffusivities (S–K
and T–I) both predict higher peaks at lower temperatures.
Figure 13 compares calculated and experimental
TDA spectra in Type 304 steel charged under
Condition 2. The surface concentration of hydrogen was
adjusted to 71 ppm so that the peak area matches that
of experimental peak. As mentioned above, the experimental
Fig. 11—Variation with tensile strain of peak area of cathodically
charged Type 304 and 316L steels.
spectrum appears as if two peaks were overlapping, the
left peak lying at 523 K (250 C), considerably lower
than that found under Condition 1 (Figure 12(a)).
These features are reproduced well in the simulation in
which no trap sites were assumed. Figure 14(a) shows
the hydrogen profile prior to temperature ramp, i.e., in
the as-charged specimen. In contrast to Condition 1, the
hydrogen concentration is not uniform, only ~20 pct of
the surface concentration (denoted cs hereafter) in the
central region. In order to see the reason why the
peculiar TDA spectrum shape mentioned above was
formed, the variation of hydrogen concentration profile
with time during heating is calculated in Figure 14(b). It
is seen that hydrogen near the outside of the specimen
quickly desorbed from the surface, which may have
formed the first peak at a lower temperature. At the
same time, a substantial fraction of hydrogen diffused
inward and came back to the surface later to form the
second higher temperature peak. The TDA spectrum of
Type 316L steel (Figure 8(b)) can also be explained in
this way. The shape and relative heights of the first and
second peaks may vary with the diffusivity of hydrogen
and presumably with charging time.
B. Specimen Without Tensile Strain Charged Under
Figure 15 compares simulated TDA spectra for Type
304 steel specimens which were cathodically charged in
the unstrained condition. In order to match the peak
area of the experimental spectrum, the surface
concentration of hydrogen was determined to be very high
(510 ppm), as expected in cathodic charge.[
temperature at the peak hydrogen evolution rate was
relatively low [~373 K (~100 C)], and this peak
temperature is reproduced very well by calculations
using both the reported S–K and M–K diffusivities
(Figure 10(b)). Such a low peak temperature was also
obtained in cathodically charged Type 316L specimens.
Clearly, these low gas evolution peak temperatures are
the consequence of the initial hydrogen distribution
confined to the narrow peripheral region of the specimen
shown in Figure 14(a) (cathodic charge). Almost all
hydrogen in the peripheral region desorbed quickly from
the surface at an early stage, thus forming a single
desorption peak at a lower temperature.
Takai et al.[
] reported that the peak temperature and
shape of TDA spectra of Type 316L steel became very
similar irrespective of charging method if the
temperature and time of hydrogen charge were the same and
sufficiently long. The present results indicate that TDA
spectra in stainless steels, both peak temperature and
shape, depend very much on the condition of hydrogen
charge, or more specifically the initial distribution of
hydrogen in the specimen. When hydrogen is distributed
uniformly prior to temperature ramp, one broad peak
appears above 623 K (350 C). If the initial distribution
is confined to the narrow peripheral region of the
specimen, a peak appears at a temperature as low as
373 K (100 C). An initial distribution between these
two extreme cases can cause a broad peak with humps,
ranging from 473 K to 673 K (200 C to 400 C), and
Fig. 15—Comparison of calculated and experimental TDA spectra
of cathodically charged Type 304 without pre-strain (Condition 3).
The surface concentration of hydrogen was assumed to be 510 ppm.
the whole peak often appears as if more than one peak is
C. Effects of Deformation on TDA Spectrum
Maroef et al. compiled the binding energies hydrogen
at various trap sites, e.g., dislocation, grain boundary,
oxide and carbide interfaces, etc., in ferritic steel.[
Compared to ferritic steel, much less data are available
for austenitic steel. Atrens and Fiore[
] reported the
binding energy of hydrogen at dislocation in Type
310 stainless steel (EB = 13.5 kJ/mol) using a
dislocation-damping experiment. This binding energy is
considerably smaller than that in bcc iron (EB = 42 kJ/
]). Thus, in austenitic steel, the amount of
hydrogen trapped by dislocations can be less than that
dissolved in the lattice. Mine et al.[
] reported that the
amount of hydrogen trapped by grain boundary can be
neglected in Type 310S stainless steel in which the grain
size was refined to less than 100 nm by high-pressure
torsion. This also seems to be the case with the present
alloys because the grain size was much greater (13 to
15 lm). Furthermore, in the alloys studied, a low
concentration of oxide inclusions or carbides is likely
to be present because the specimen alloys were annealed
at a high temperature. Hence, dislocations are assumed
to be the primary hydrogen trap sites in the simulation
of TDA spectra of specimens with tensile strain.
Shintani and Murata[
] reported that the dislocation
density does not exceed 5.0 9 1014 m 2 with increasing
strain in Type 304 stainless steel. In fcc iron, hydrogen
occupies the octahedral interstitial sites. Since the
number of octahedral interstitial sites is equal to the
number of Fe atoms, the number of occupation sites is
evaluated to be 5 to 10 per Fe atom plane if hydrogen
occupies the octahedral sites within the distance of two
times the Burgers vector from the dislocation core. It is
also assumed that local equilibrium was achieved
between the lattice and the trap sites, which requires
that p0 is equal to 108 s 1 or greater.
Fig. 16—Comparison with experiment (dashed curve) of TDA
spectra of cathodically charged Type 316L with 40 pct strain calculated
assuming the dislocation density of qd = 5 9 1014 and
5 9 1015 m 2. The surface concentration of hydrogen was assumed
to be 550 ppm.
With these parameters, the TDA spectrum of
cathodically charged Type 316L specimen with 40 pct strain
was simulated as shown in Figure 16. No appreciable
influence of hydrogen trapping at dislocations was
observed in the spectrum until the dislocation density
was increased to 5 9 1015 m 2. As long as the
dislocation density reported by Shintani and Murata[
assumed, the trap site of the excess amount of hydrogen
is not clear. Dislocation kinks and/or vacancies formed
upon the intersection of dislocations could serve as
additional trap sites.
In Figure 6(b), the peak temperature decreased
slightly in the 40 pct strained specimen compared to
non-strained and 10 pct strained Type 316L specimens.
This was presumably caused by pipe diffusion of
hydrogen along dislocations at a high hydrogen
As shown earlier, the effects of pre-strain in Type 304
steel are very different from those in Type 316L steel; the
TDA spectrum area of 40 pct pre-strained specimen
decreased under Condition 1, while it increased
remarkably when specimens were exposed to hydrogen under
Condition 2 or through cathodic charging. The
solubility of hydrogen in ferrite is considerably smaller than
that in austenite.[
] Recently, it is observed that the
amount of absorbed hydrogen does not exceed a few
ppm in Type 420 martensitic stainless steel after
prolonged exposure at high pressure.[
] From the XRD
spectrum in Figure 3, the volume fraction of martensite
was evaluated to be in the range of 0.21 to 0.28.[
Indeed, it is reported that ca. 20 pct of the austenite
matrix transformed to martensite upon tensile
deformation in Type 304 steel.[
] Moreover, according to
Shintani and Murata,[
] the fraction of martensite was
50 to 60 pct in 40 pct cold rolled Type 304 steel. Thus,
the decrease in the TDA spectrum area of the 40 pct
strained Type 304 steel (Figure 7) is primarily due to the
formation of strain-induced martensite.
Fig. 17—(a) Comparison of calculated and experimental TDA
spectra of Type 304 with 40 pct strain exposed to high-pressure
atmosphere at 358 K (85 C) for 1000 h (Condition 2). Thin and thick
dashed curves were calculated using the diffusivity reported by
Sakamoto and Katayama for non-deformed and deformed specimens.
The surface concentration of hydrogen was assumed to be 81 ppm.
(b) Diffusion penetration curves in the as-charged condition
calculated in Type 304 with 40 pct strain and without strain, charged
under Condition 2.
Sakamoto and Katayama[
] reported that the
diffusivity of hydrogen in deformed Type 304 steel increased
by a factor of 3 to 5 from that without deformation.
They ascribed this increase in effective diffusivity to the
formation of strain-induced martensite, which can serve
as a high-diffusivity path for hydrogen at a high
hydrogen concentration. Figure 17(a) compares with
experiment the calculated and experimental TDA
spectra of Type 304 specimen with 40 pct strain charged
under Condition 2. In Figure 17(b), the diffusion
penetration curves are shown for the same specimen in the
as-charged specimens with and without deformation.
Thus, the marked increase in peak area can be attributed
to the increase in effective diffusivity due to the
formation of martensite. This can also account for the
large increase in the TDA spectrum area in the
cathodically charged specimen with 40 pct strain
(Figure 10(a)). The absence of a peak around 373 K
(100 C) reported in martensitic steel[
] is possibly
attributed to the fact that strain-induced martensite
crystals are embedded in austenite in which the
hydrogen diffusivity is much slower than in martensite.
Hydrogen absorption during cathodic charging and
exposure to high-pressure gaseous hydrogen were
studied in Type 316L and 304 stainless steels by means of
TDA in an attempt to establish the relationship between
the conditions of the two different hydrogen charging
methods. TDA spectra were simulated by the McNabb–
Foster model to analyze the difference in the spectrum
shape and peak gas evolution temperature between the
two steels and the effects of charging temperature, time,
and pre-deformation strain. The results are summarized
1. Both Type 316L and Type 304 steels without tensile
strain, which were exposed to high-pressure gaseous
hydrogen at 523 K (250 C) for 72 hours, exhibited
a large gas evolution peak above 573 K (300 C),
while those exposed to high-pressure gaseous
hydrogen at 358 K (85 C) for 1000 hours exhibited a flat
and broad peak ranging from 473 K to 673 K
(200 C to 400 C). The peak appeared as if it
consisted of closely lying two peaks. The simulation
indicated that hydrogen penetrated fully into the
specimen in the former condition, while it did only
partially in the latter condition. Simulation also
revealed that nearly a half of hydrogen in the
specimens exposed to hydrogen at the lower temperature
condition quickly desorbed at the surface, and the
other remaining portion diffused toward the interior
of the specimen and then returned to the surface to
form a second flat broad peak in the gas evolution
rate, which produced two peaks in the TDA
2. Both steels cathodically charged at ambient
temperature exhibited a peak in hydrogen gas evolution at
approximately 373 K (100 C). According to the
simulation, the initial hydrogen distribution was
confined to the narrow peripheral region of the
3. Tensile strain exhibited only a marginal increase in
the amount of hydrogen in Type 316L steel, both
cathodically charged and exposed under
high-pressure hydrogen atmosphere. The simulation using
the reported dislocation density and the binding
energy of hydrogen in fcc iron did not have an
appreciable effect on the amount of hydrogen, and thus
lend support to the observed absorption behavior in
4. Tensile strain increased the amount of absorbed
hydrogen from high-pressure hydrogen gas at
358 K (85 C) in Type 304 steel, while it decreased
significantly the amount of hydrogen absorbed from
high-pressure hydrogen gas at 523 K (250 C). The
tensile strain also markedly increased the amount of
hydrogen in the same steel cathodically charged at
ambient temperature. These are likely to be caused
by strain-induced martensitic transformation, which
has a lower hydrogen solubility and faster hydrogen
diffusivity than in austenite.
These results indicate that the initial distribution of
hydrogen prior to the TDA has a large influence on the
spectrum shape and peak hydrogen gas evolution
temperature. Moreover, one can predict the absorption
of hydrogen under high-pressure atmosphere at elevated
temperatures from that under cathodic charge as long as
the surface concentration of hydrogen or effective
hydrogen fugacity at the specimen surface is available.
The authors express their sincere thanks to the
members of forum of building fundamental platform
for research of HE, sponsored by The Iron and Steel
Institute of Japan, Tokyo.
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