Determination of the dominant factors of hydrogen embrittlementin austenitic stainless steel
オーステナイト系ステンレス鋼における水素脆化支配因子の決定
Principal Investigator: Chiba University, Graduate School of Engineering,
Assistant Professor Luca Chiari
研究代表者 千葉大学大学院工学研究院 助教 キアリ ルカ
1. Introduction
Although the utilization of hydrogen as a clean energy vector to meet the needs of amore environmentally friendly society is getting full-fledged, hydrogen embrittlement inmetals poses a significant problem for its widespread application. Hydrogen
embrittlement is a reduction in the mechanical properties of a large class of susceptible
metals, such as their decrease in ductility and early fracture, when they are used in ahydrogen environment. Several theories of hydrogen embrittlement have been developed
and consistent supporting experimental evidence has been accrued over the years [1].
Notwithstanding this large research effort in attempting to understand theembrittlement process, the underlying mechanisms are still not completely understood.
Amongst the proposed hydrogen-embrittlement mechanisms, the hydrogen-enhanced
strain-induced vacancy formation (HESIV) model argues that the atomic vacancies
formed by dislocation cutting upon tensile testing are stabilized by bonding withhydrogen, which in turn promotes their aggregation by forming vacancy clusters andinduces the destabilization of the plastic deformation [2]. In fact, theoretical calculations
show that hydrogen reduces the vacancy formation energy in γ-iron by 0.23 eV [3]. Hence,the experimental elucidation at the atomic level of the interaction between hydrogen,vacancies and dislocation motion is very important to detect the dominant defects inhydrogen embrittlement. For this purpose, a non-invasive and highly sensitive analysis oflattice defects is required.
Positron annihilation lifetime spectroscopy (PALS) is a non-destructive technique used
to detect vacancy-type defects, such as dislocations, atomic vacancies, and vacancy
clusters, with a high sensitivity of up to 10-7 in atomic ratio, from measurements of thepositron lifetime in the material [4]. The application of PALS to hydrogen embrittlement
studies revealed so far that stretching susceptible metals in the presence of hydrogenpromoted the formation of vacancy clusters. However, the formation of vacancy clusters is
promoted in α-iron (bcc phase) even in the presence of hydrogen under stretchingconditions that do not lead to hydrogen embrittlement (e.g. fast strain rate). Therefore,－ 115 －

vacancy clusters were ruled out as the primary defects responsible for hydrogen
embrittlement. In α-iron, the hydrogen diffusion coefficient at room temperature is verylarge and, therefore, the hydrogen-charged samples were quenched in liquid nitrogen
immediately after stretching, and PALS measurements were carried out at variabletemperatures. Those results have shown that, under the conditions where hydrogen
embrittlement occurred, mobile monovacancies were detected in the range between 250 K
to 420 K, whereas already grown vacancy clusters had formed when hydrogen
embrittlement did not occur. Therefore, it was concluded that hydrogen binds to andstabilizes monovacancies by forming vacancy-hydrogen complexes, which are the primarycontrolling factor of hydrogen embrittlement in α-iron. Those defects accumulate locallyon the slip surface due to the dislocation motion and become the starting point of cracks.
In austenitic stainless steel SUS 304, which is a γ-phase (fcc phase), it is believed thathydrogen embrittlement depends on the γ-phase stability, as the phase transformation to
the processing-induced α' martensite phase and stacking faults are closely related tohydrogen [5]. However, the formation of vacancy clusters was observed at a strain value
at which no processing-induced α' martensitic transformation occurred. Therefore,hydrogen embrittlement in the austenitic stainless steel phase cannot be explained onlyby the formation of the processing-induced martensite phase, and it is necessary toconsider the hydrogen embrittlement process of the γ-phase itself. One of the possible
causes of hydrogen embrittlement may be that hydrogen-induced vacancies aggregate ina high-strain region such as the boundary between the γ-phase and the ε phase (hcp) toform vacancy clusters and serve as a starting point for fracture [6]. Preliminary PALS
measurements in SUS 304 pointed out that the vacancy-hydrogen complexes diffuse into
the steel already at the initial stage of elongation.
In order to elucidate the hydrogen embrittlement behaviour in austenitic stainlesssteels, it is necessary to examine the vacancy generation process, and a non-destructive
and highly sensitive analysis is required. In this research, the formation process ofhydrogen-induced vacancies in austenitic stainless steel SUS 304 after exposure to ahydrogen environment was investigated. In-situ analysis of the defect formation
behaviour at the boundary between the elastic region and the plastic deformation region
during tensile testing was carried out by PALS measurements.
2. Experimental methods
2.1 Sample preparation
All the samples were manufactured from austenitic stainless steel SUS 304 thin plates.
The chemical composition of the present steel SUS 304 is shown in Table 1 and matches－ 116 －

that specified by the Japanese Industrial Standards.
Table 1 Chemical composition of the present SUS 304 samples
C
<0.03

Si
<1.00

Mn
<2.00

P
<0.045S
Ni<0.03 8.00~10.50

Cr
18.00~20.00

Febalance

The SUS 304 thin plates were processed into dumbbell-shaped specimens for use intensile testing by electric discharge machining. Fig. 1 shows a schematic diagram of thesample shape and size. The specimen width is 10 mm, gauge length is 20 mm, andthickness is 0.2 mm. A notch is provided at the center of the gauge section to concentrate
the tensile stress in that area.

Fig. 1 Schematic diagram of the sample shape and size Units: mm
The samples underwent solution treatment to remove the defects generated during therolling and processing of the material. The solution treatment was performed in an Ar
atmosphere using an infrared gold image furnace. The solution treatment temperature
was 900°C, the heating time was 4 h (heating rate 4°C/min), and the holding time was 5
min. Thereafter, the sample was left to cool down to room temperature. As the samples
showed a black oxide film of several μm in thickness after the solution treatment, theirsurface layer was removed by mechanical and chemical polishing.
2.2 Hydrogen charging
Hydrogen was introduced into the samples by the electrolytic hydrogen charging
method. A 3% NaCl + 3 g/l NH4SCN aqueous solution was used as the electrolyte. The
charging time was 48 hours at a current density of 50 A/m2.

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2.3 Tensile tests
After hydrogen charging, the samples were subject to tensile testing. The strain rate
was 2.75 × 10-5/s for the samples strained up to fracture, whereas the strain rate was 4.17
× 10-3/s for the samples strained up to strain values of 5%. For the latter samples, themeasurements were carried out in situ on the sample maintained stretched at the samestrain value for the whole duration of the measurement.
2.4 PALS measurements
Positron annihilation lifetime spectroscopy is a non-destructive technique used todetect open-volume defects in a variety of materials with high sensitivity. PALS candetermine the type, size and relative concentration of vacancy-type defects based on thepositron lifetime in the material, i.e. the time from the positron injection into the sample
to the annihilation with an electron of the material. The positron source is made of 22Na
sealed within Kapton films from which high-energy positrons are generated by β+ decay
and can penetrate into metals up to about 100 µm in depth before annihilation. The γrays produced by the positron generation (start signal) and the positron annihilation (stopsignal) are detected by two photomultiplier tubes coupled to BaF2 scintillators. The output
signals are recorded and processed by a digital storage oscilloscope. Each positron lifetime
spectrum was obtained by integrating at least one million counts and at least three
spectra were collected for each sample to achieve enough statistical accuracy. The present
PALS measurements were carried out in situ while keeping the samples stretched at agiven strain value.
Fig. 2 shows a typical lifetime spectrum of stainless steel SUS 304 and its
decomposition into various components. The spectrum is analysed by fitting the data withthe sum of exponential decay functions using least squares regression. The number offitting components corresponds to the number of defect types in which positrons
annihilated. Typically, the lifetime components correspond to positron annihilation in thebulk (lifetime value of around 110 ps), at dislocations and/or monovacancies (~160-180 ps),and in vacancy clusters (~200-400 ps).

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Fig. 2 Typical positron lifetime spectrum for hydrogen-charged and fractured stainlesssteel SUS 304 and spectrum decomposition into two lifetime components
The vacancy size can be quantitatively obtained from measurements of the positron
lifetime value by comparing with first-principle calculations which correlate defect sizeand positron lifetime (Fig. 3). The relative concentration of each defect type in the sample
is proportional to the relative intensity of the respective lifetime component.

Fig. 3 First-principle calculation correlating the positron lifetime value to the number of
atomic vacancies in a vacancy cluster in α-iron [7].

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3. Results and discussion
First of all, the effect of the hydrogenation on the mechanical properties of the SUS
304 samples was investigated by tensile testing. Fig. 4 shows the results of the tensile
tests on the hydrogen-free and hydrogen-charged samples. The sample subject tohydrogen charging is clearly hydrogen embrittled, as its strain value at fracture and its
maximum stress value are significantly reduced with respect to the non-hydrogenatedsample. However, the stress-strain curves of the two samples are not significantly

800

400

700

350

600

300

Stress(MPa)Stress(MPa)different to each other in the elastic region.

500
400
300
200
100
0

250
200
150
100
50

0

5

10

15

20

25

Strain(%)

30

35

40

0

0

0.5

Strain(%)

1

1.5

Fig. 4 Left: stress-strain curves of the hydrogen-free (red) and hydrogen-charged (blue)
samples. Right: expanded view of the elastic region.
Fig. 5 shows the positron lifetimes of the hydrogenfree and hydrogen-charged samples before the tensile
test. Both samples show a single lifetime of about 110
ps, which corresponds to the known lifetime value forpositrons annihilating in the bulk. Since only the bulk
component appeared in the hydrogen-charged sample,
it is clear that no defects responsible for hydrogen
embrittlement had formed during the hydrogencharging stage. Therefore, the hydrogen addition itself
is not the cause of the defect generation.

Fig. 5 Positron lifetimesof the hydrogen-free andhydrogen-charged samples

The lifetimes of the hydrogen-free and hydrogencharged samples after tensile testing until fracture are shown in Fig. 6. In the hydrogenfree fractured sample, a single lifetime of around 165 ps was observed. This value is－ 120 －

consistent with the known lifetimes for positrons annihilating at dislocations andmonovacancies. This confirms that dislocations and monovacancies are generated when
the sample is subject to tensile deformation. On the other hand, in the hydrogen-charged
fractured sample, two lifetime components were observed. The first component was
similar in value to that observed in the hydrogen-free fractured sample, thus indicating
the formation of dislocations and monovacancies. The second component was significantly
longer (285 ps) and can be attributed to positron annihilation at vacancy clusters. Basedon the results of the first-principle calculation shown in Fig. 3, we expect these vacancy
clusters to be of around 10 atomic vacancies in size. This result shows that vacancy
clusters are formed when the hydrogenated sample is stretched and fractured comparedto the non-hydrogenated sample. In the tensile tests, a premature fracture was observedin the hydrogenated sample due to the hydrogen embrittlement. Hence, it is concluded
that there must be a correlation between the formation of vacancy clusters and thehydrogen embrittlement-related premature fracture.

Fig. 6

Positron lifetimes of the hydrogen-free and hydrogen-charged samples after

tensile straining until fracture. The blue data correspond to the first component(dislocations and monovacancies), while the green data show the second component(vacancy clusters).
The origin of the relation between the vacancy clusters and the premature fracture canbe clarified by studying the defect generation process at the early stage of the tensile
deformation. For this purpose, the lifetime was measured at various tensile strain values
at the boundary between the elastic region and the plastic deformation region, andcompared between the hydrogen-charged and uncharged samples.

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First of all, we determined the extent of the elastic region and ensured that the sample
condition does not change before and after tensile stretching in the elastic region. For thispurpose, PALS measurements were carried out in three stages: before tensile strain isapplied (0% strain), while the sample is stretched at a strain value of 0.5%, and after
releasing the sample back to its original shape (0% strain because of the elasticity of thesample in this strain region). The measurements were carried out in situ while keeping
the samples in each condition. Fig. 7 shows the lifetimes measured in these three stages
for both the hydrogen-free and hydrogen-charged samples. In both the hydrogen-free andhydrogen-charged samples, only the bulk component of around 110 ps was observed in all
three stages. This confirms the elasticity of the sample and that no defects wereintroduced in the sample upon tensile straining up to this strain value and byhydrogenation. Such measurements were also repeated at a strain value of 1% and theresults confirmed the elasticity of the samples up to this strain value and the absence of
defects.

Fig. 7

Positron lifetimes of the hydrogen-free (left) and hydrogen-charged (right)

samples before elongation (0%), at a tensile strain of 0.5%, and after release
Next, PALS measurements were carried out at various tensile strain values at theboundary between the elastic region and the initial plastic deformation region to study
the defect formation process. The measurements were carried out in situ while
maintaining the samples in the stretched condition at each strain value. Fig. 8 shows the
lifetimes and the intensity of the second component (where present) measured at thestrain values of 0%, 0.5% and 1% in the elastic region, and at 3% and 5% in the initialplastic deformation region. The results of the hydrogen-free and hydrogen-charged
samples at each strain value are compared. In both the hydrogen-free and hydrogencharged samples, only one lifetime component corresponding to the bulk was detected up－ 122 －

to the strain value of 1%. This confirms the elastic region and the absence of defects inthis range upon tensile testing and hydrogenation. In the hydrogen-free sample, at strain
values beyond 1% a second component of 170-180 ps was observed. This defect size isconsistent with the generation of dislocations and monovacancies upon the initial plastic
deformation of the sample. Although the lifetime of this component did not change byincreasing the strain value, its relative intensity doubled when the strain was increased
from 3% to 5%. This reflects an increased concentration of dislocations andmonovacancies as expected due to the larger tensile strain. On the other hand, in thehydrogen-charged sample a second component of around 250 ps was observed at strain
values beyond 1%. This defect size corresponds to that of vacancy clusters composed of ~5
atomic vacancies. As the strain value was increased from 3% to 5%, the relative intensityof this component increased slightly, indicating that their relative concentration alsoslightly increased. The small change might be due to the fact that the vacancy clusters
are not very mobile during room temperature annealing.
The results shown in Fig. 8 resemble those observed in the fractured samples (Fig. 6),in that only dislocations and monovacancies were found in the hydrogen-free sample,
whereas vacancy clusters were only detected in the hydrogen-charged sample. Therefore,
this corroborates our hypothesis that it is important to study the defect generation
process at the boundary between the elastic region and the initial plastic deformation
region in order to understand the relation between the formation of vacancy clusters andthe hydrogen embrittlement.

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Fig. 8 PALS results for the hydrogen-free (left) and hydrogen-charged (right) samples atvarious tensile strain values between 0% and 1% in the elastic region, and 3% and 5% in
the initial plastic deformation region. Lower panels: positron lifetime of each component;
upper panels: relative intensity of the second component. The red data correspond to thebulk component, the blue data to dislocations and monovacancies, and the green data tovacancy clusters.
In terms of the defect formation process and their relevance in the hydrogen
embrittlement mechanism, the PALS results of strain-dependence measurements can beinterpreted in the following way (see Fig. 9). Upon application of tensile strain, no defects
are generated within the elastic region, whereas dislocations and monovacancies are
generated as soon as the tensile strain value reaches into the plastic deformation region.
In the presence of hydrogen, hydrogen becomes trapped into the monovacancies andstabilizes them by forming vacancy-hydrogen complexes, as suggested by our previous
measurements and first-principle calculations. As hydrogen is desorbed by room
temperature annealing, the leftover vacancies become mobile and, upon tensile strain,
they agglomerate and form vacancy clusters. Through this clustering process, thevacancies become bigger in size but smaller in number, as suggested by the lower
intensity of the vacancy cluster component in the hydrogen-charged sample compared tothat of the dislocations and monovacancies in the hydrogen-free sample. These vacancy
clusters then progressively become bigger in size as the tensile strain increases andeventually become the trigger of the premature fracture.

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Fig. 9 Schematic diagram of the strain-dependent defect dynamics in the hydrogen-free(left) and hydrogen-charged (right) samples based on the results of Fig. 8: ① dislocations
and monovacancies are generated upon the application of tensile strain, ② in thepresence of hydrogen, monovacancies and hydrogen form complexes, ③ by room
temperature annealing, hydrogen is desorbed and the leftover monovacancies fromvacancy clusters upon tensile straining.
4. Conclusions
In the present study, the generation process of hydrogen-related defects in austenitic
stainless steel SUS 304 was investigated to shed light on the mechanism of the stillunresolved important issue of hydrogen embrittlement. The non-destructive and highly
sensitive positron annihilation lifetime spectroscopy was used to determine the defect sizeand concentration. Using this technique, in-situ analysis of the defect formation
behaviour at the boundary between the elastic region and the plastic deformation region
during tensile testing was carried out. The measurements revealed the formation ofvacancy clusters already at the early stages of the plastic deformation in the hydrogencharged samples. These defects are believed to originate from the agglomeration of thevacancies leftover by the vacancy-hydrogen complexes after hydrogen desorption at room
temperature and application of tensile strain. Upon application of further tensile
straining, the vacancy clusters then become the trigger of the brittle fracture.

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Acknowledgement
This research was supported by a research grant from the JFE 21st Century
Foundation.
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