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Atomistic simulations of hydrogen embrittlement
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i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 4
Avai lab le at www.sc iencedi rect .com
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Atomistic simulations of hydrogen embrittlement
Ryosuke Matsumoto a,b, Shinya Taketomi a,b, Sohei Matsumoto a, Noriyuki Miyazaki a,b,*a Department of Mechanical Engineering and Science, Kyoto University, Yoshida-Honmachi, Sakyo-ku, Kyoto 606-8501, Japanb National Institute of Advaned Industrial Science and Technology (AIST), 744 Motooka, Nishi-ku, Fukuoka 819-0395, Japan
a r t i c l e i n f o
Article history:
Received 29 May 2009
Received in revised form
12 September 2009
Accepted 19 September 2009
Available online 16 October 2009
Keywords:
Hydrogen embrittlement
Atomistic simulation
Molecular dynamics method
Molecular statics method
Crack propagation
Dislocation
* Corresponding author at: Department of Me606-8501, Japan. Tel.: þ81 75 753 5213; fax: þ
E-mail address: [email protected]/$ – see front matter ª 2009 Profesdoi:10.1016/j.ijhydene.2009.09.052
a b s t r a c t
It is well known that hydrogen weakens strengths of metals, and this phenomenon is
called hydrogen embrittlement. Despite the extensive investigation concerning hydrogen
related fractures, the mechanism has not been enough clarified yet. In this study, we
applied the molecular dynamics method to the mode I crack growth in a-Fe single crystals
with and without hydrogen, and analyzed the hydrogen effects from atomistic viewpoints.
We estimated the hydrogen trap energy in the vicinity of an edge dislocation in order to
clarify the distribution of hydrogen atoms, using the molecular statics method. We also
evaluated the energy barrier for dislocation motion under a low hydrogen concentration.
Based on these results, we propose a mechanism for hydrogen embrittlement of a-Fe under
monotonic loading.
ª 2009 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved.
1. Introduction low to weaken the bonds. In the HELP mechanism, a plastic
Metals absorbing hydrogen show the reduction of ductility
due to hydrogen [1] and the acceleration of fatigue crack
growth [2]. This phenomenon is known as hydrogen embrit-
tlement. Much attention has been paid to hydrogen as a clean
energy source to solve environmental problems and to cope
with the global warming problem. Increase in hydrogen use
would result in increase of failure accidents related with
hydrogen embrittlement.
Various mechanisms for hydrogen embrittlement have been
proposed so far. Among them, the hydrogen enhanced deco-
hesion (HEDE) [3,4] and the hydrogen enhanced localized plas-
ticity (HELP) [5] are typical ones. In the HEDE mechanism, the
bonds of metal atoms are weakened by hydrogen atoms. It has
been, however, supposed that a hydrogen concentration is too
chanical Engineering an81 75 753 5719.ac.jp (N. Miyazaki).sor T. Nejat Veziroglu. Pu
behavior of a material is affected by hydrogen atoms. This
mechanism is supported by experimental results. For example,
increase in dislocation mobility is observed in in situ TEM
observations [6–8]. It is also observed that slip bands are local-
ized in the vicinity of a crack tip in fatigue tests using hydrogen-
chargedtest specimens [2].Thefracture phenomenon causedby
hydrogen embrittlement cannot be explained only by the HELP
mechanism.
Despite a lot of experimental works concerning hydrogen
embrittlement, for example Refs. [9] and [10], the mechanism of
hydrogen embrittlement is not fully understood. Hydrogen has
a high diffusivity and a low concentration in metal. It is, there-
fore, difficult to perform direct observation of hydrogen and to
answer the question from experimental results what is a correct
mechanism for hydrogen embrittlement. Atomistic simulations
d Science, Kyoto University, Yoshida-Honmachi, Sakyo-ku, Kyoto
blished by Elsevier Ltd. All rights reserved.
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Nomenclature
Roman symbols
b Magnitude of Burger’s vector, m
KI Mode I stress intensity factor, MPaffiffiffiffiffimp
R Gas constant, J/mol K
t Time, s
T Absolute temperature, K
xeq Hydrogen concentration expressed by
the atomic ratio -
x0 Hydrogen concentration expressed by the
atomic ratio without hydrostatic stress, -
Greek symbols
dUH Partial molar volume of hydrogen, m3/mol
3 Ratio of number of hydrogen atoms to the
number of iron atoms, -
shyd Hydrostatic stress, MPa
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 4 9577
such as the molecular dynamics method, the molecular statics
method and soonare powerful tools tostudythe mechanism for
hydrogen embrittlement.
In the present study, after choosing an adequate inter-
atomic potential for a-Fe and hydrogen system (hereafter
abbreviated as aFe-H system), we perform several kinds of
atomistic simulations for such a system, that is, the molecular
dynamics analyses of crack propagation and the molecular
statics analyses of the interaction between dislocation and
hydrogen atoms, and we propose a mechanism for hydrogen
embrittlement of a-Fe under monotonic loading, based on the
results of the atomistic simulations.
2. Interatomic potential for aFe–H system
Adequate selection of an interatomic potential is of crucial
importance for molecular dynamics and molecular statics
calculations. Only three kinds of interatomic potential have
been proposed so far for the aFe–H system. They are the
embedded-atom-method (EAM) potential by Ruda et al.
(abbreviated as EAM-R) [11], the Morse type potential by Hu
et al. (abbreviated as Morse) [12] and the EAM potential by Wen
et al. (abbreviated as EAM-W) [13]. EAM-W was formulated by
improving EAM-R so as to reproduce more properties of the
aFe–H system accurately. Thus, EAM-W is superior to EAM-R.
In comparison with Morse, EAM-W provides accurate results
for elastic constants of a-Fe and the properties of hydrogen in
a-Fe, as shown in Tables 1 and 2. The elastic constants of a-Fe
calculated from EAM-W and Morse are shown in Table 1,
compared with the experimental results [14]. The
Table 1 – Elastic constants of a-iron.
Exp. [14] EAM-W Morse
C11 (GPa) 243.1 230.2 244.3
C12 (GPa) 138.3 135.8 80.7
C44 (GPa) 138.1 116.7 80.7
experimental results agree well with the calculated results
using EAM-W. The heat of solution and migration energy of
hydrogen in a-Fe are shown in Table 2. They are also
compared with experimental results [15,16]. Again EAM-W
provides better results than Morse in comparison with the
experimental results. Morse neglects H–H interactions.
Moreover it takes account of the long-range (about 1 nm)
interaction between iron and hydrogen atoms, so that it
requires a large amount of computational time for a molecular
dynamics calculation despite a pair potential. Because of the
above reasons, we selected EAM-W as the best interatomic
potential for the aFe–H system and employed it in the subse-
quent atomistic simulations.
3. Molecular dynamics analyses of crackpropagation
We applied molecular dynamics simulations to the mode I
crack propagation in a-Fe single crystal with and without
hydrogen for several analysis conditions, i.e. two kinds of
crystal orientation and two levels of temperature.
3.1. Analysis model
Fig. 1 shows an analysis model whose shape is a circular disk
with 9.7 nm in radius, 2.8 nm in thickness and 0.7 nm in
thickness of the boundary region. The displacements corre-
sponding to KI of 0.9 MPaffiffiffiffiffimp
are prescribed to all atoms to
introduce an initial crack. The origin is set at the crack tip and
the x-axis and z-axis correspond to the forward direction of
the initial crack and the thickness direction, respectively. This
analysis model is a quasi three-dimensional model consisting
of about 71,000 atoms, on which a periodic boundary condi-
tion in the z-direction is imposed.
According to Ref. [17], stable locations of hydrogen atoms
are tetrahedral sites (T-sites), and a hydrogen distribution
depends on the hydrostatic stress. Thus, hydrogen atoms are
introduced at the T-sites in iron atoms in accordance with the
following equations [18], using random numbers:
xeq ¼(
x0exp�
shydðXÞdUHRT
�ðjXj � jr0jÞ
x0exp�
shydðr0ÞdUHRT
�ðjXj < jr0jÞ
(1)
where xeq is the hydrogen concentration expressed by the
atomic ratio at a location X, x0 the hydrogen concentration
without hydrostatic stress shyd, dUH the partial molar volume
of hydrogen in a-Fe, R the gas constant and T is the absolute
Table 2 – Heat of solution and migration energy ofhydrogen atom.
Exp. EAM-W Morse
heat of solution (eV) 0.30 [15] 0.28 –
migration energy (eV) 0.035 [16] 0.037a 0.05a
a potential energy difference between hydrogen atoms at a tetra-
hedral site and an octahedral site.
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y [110]
{112} slip plain
a
x
y
z
9.7nm
2.8nm
0.7nm
boundary region
Fig. 1 – Analysis model.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 49578
temperature. We used dUH of 1.2�10�6m3/mol [18]. We
selected jr0j ¼ 5:0� 10�10 m to avoid an infinite quantity of the
hydrogen concentration at a crack tip. In the present study, we
changed the hydrogen concentration x0 from 0 (no hydrogen
atom) to 5.0�10�4 as an analysis parameter. The initial
distribution of hydrogen atoms is shown in Fig. 2 for
x0¼3.0�10�4 (z 5.4 mass ppm) and 400 K, where red points
and blue points denote iron atoms and hydrogen atoms,
respectively. It is found from the figure that the hydrogen
concentration is higher around the crack tip than in the
periphery of the analysis model because of stress concentra-
tion around the crack tip.
x [110]
crack surface
54.7°
54.7°
crack plane
3.2. Analysis conditions
Crack propagation analyses were performed for two kinds of
crystal orientation, crystal orientations (A) and (B), as shown in
Fig. 2 – Initial distribution of hydrogen atoms
(x0 [ 3.0 3 10L4).
Fig. 3, in which the gray planes indicate the {112} slip planes of
a-Fe. In the crystal orientation (A), the crack plane is the (112)
plane, and the forward direction of the initial crack is the ½110�direction. In this case, there is no slip plane in the xy-plane, and
we can expect no dislocation emission from the crack tip. In the
crystal orientation (B), the crack plane is the (110) plane, and
the forward direction of the initial crack is the [001] direction. In
this case, there exists {112} slip planes in the xy-plane, and we
can expect the emissions of dislocations from the crack tip.
Although the ductile-brittle transition temperature (DBTT)
for a-Fe is less than 100 K, the DBTT calculated using EAM-W is
between 200 K and 300 K. In the crystal orientation (B) where
dislocation emissions are expected, we performed the crack
propagation analyses at the initial temperature of 400 K, which
is above the DBTT, and that of 100 K, which is below the DBTT.
Dislocation emissions are expected at 400 K, and no dislocation
emission is expected at 100 K. In the crystal orientation (A), no
dislocation emission is expected at any temperature, so that
we performed the crack propagation analyses only at the initial
temperatures of 400 K. We can discuss the effect of dislocations
on the crack growth behavior by comparing the result of the
crystal orientation (A) at 400 K with that of the crystal orien-
tation (B) at 400 K. In the present analyses, we did not control
the temperature of the analysis system to keep the tempera-
ture at a constant value during crack propagation.
The crack propagation analyses were performed by
imposing the displacement rate corresponding to the rate of
z [001]
crack tip
x [001]
y [110]
z [110]
{112} slip plane
crack plane
54.7°
54.7°
crack tipcrack surface
b
Fig. 3 – Arrangement of [112] slip planes. (a) Crystal
orientation (A) (b) Crystal orientation (B).
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Fig. 4 – Snapshots of crack propagation behavior for the crystal orientation (A) at 400 K (x0 [ 0).
Fig. 5 – Time evolution of crack growth length in the crystal
orientation (A) at 400k.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 4 9579
stress intensity factor dKI=dt ¼ 5:0� 109MPaffiffiffiffiffimp
=s on the
atoms in the boundary region of the analysis model. This
analysis condition is equivalent to the crack opening velocity
of 0.4 m/s, which is very slow for a molecular dynamics
analysis. In the present analyses, we performed several
numbers of analyses with different initial locations of
hydrogen atoms and different initial velocities of iron and
hydrogen atoms for the respective analysis conditions by
changing random numbers in order to avoid the effect of
initial conditions.
3.3. Results and discussion
3.3.1. The cases without dislocation emissionIn the crystal orientation (A), crack propagation analyses were
performed for four cases of the initial hydrogen concentra-
tion, x0¼ 0 (no hydrogen atom), 1.0�10�4, 3.0�10�4 and
5.0�10�4 at the initial temperature of 400 K. A crack propa-
gation behavior is shown for the case without hydrogen
(x0¼ 0) in Fig. 4, where a green part indicates the bcc crystal
structure and a black one other crystal structures. In this case,
a crack propagates straight without dislocation emission.
Although the figures are omitted here, similar crack propa-
gation behaviors were also obtained for the cases with
hydrogen atoms (x0¼ 1.0� 10�4, 3.0�10�4 and 5.0�10�4). Fig. 5
shows the time evolution of crack growth length in the crystal
orientation (A) at 400 K. Significant differences in the initiation
time for crack growth and the crack growth velocity are not
observed among the respective hydrogen concentrations.
In the crystal orientation (B), crack propagation analyses
were performed for x0 ¼ 0 (no hydrogen atom), 1.0� 10�4 and
2.0� 10�4 at the initial temperature of 100 K. Fig. 6 shows
a crack propagation behavior for x0¼ 1.0� 10�4. As expected,
no dislocation emission is observed because of a lower
temperature than the DBTT, and a crack propagates nearly
straight. No meaningful difference is observed between the
cases with and without hydrogen atom.
3.3.2. The cases with dislocation emissionsIn the crystal orientation (B), crack propagation analyses were
performed for five cases of the initial hydrogen concentration,
x0¼ 0 (no hydrogen atom), 0.5� 10�4, 1.0� 10�4, 3.0� 10�4 and
5.0� 10�4 at the initial temperature of 400 K. The distributions
of hydrogen atoms around dislocation cores are shown for
x0 ¼ 3:0� 10�4 in Fig. 7, where a grey circle and a black one
denote an iron atom and a hydrogen atom, respectively. It is
found from the figure that hydrogen atoms are trapped
around dislocation cores within 100ps. Detailed discussion on
the interaction between dislocations and hydrogen atoms will
be given in the following section.
Figs. 8(a), (b) and (c) respectively represent the crack prop-
agation behaviors for three cases of hydrogen concentration,
that is, (a) x0 ¼ 1:0� 10�4, (b) x0 ¼ 3:0� 10�4 and (c) x0 ¼ 0 (no
hydrogen atom). As shown in Figs. 8(a) and (b), not only dislo-
cation emissions from the crack tip but also crack propagation
along the {112} slip planes are observed in high frequency in the
cases with hydrogen atoms. On the other hand, as shown in
Fig. 8(c), only crack tip blunting caused by dislocation emis-
sions tends to be observed without crack propagation in the
case where no hydrogen atom is included. The fracture in the
slip plane is observed in only one case in the case without
hydrogen atom and in three cases in the cases with hydrogen
atoms out of four cases with different initial conditions.
Figs. 9(a) and (b) show the enlarged figures near the crack
tip during slip plane fracture. It is found that hydrogen atoms
gather on the {112} slip planes, from which slip plane fracture
occurs. The hydrogen atoms trapped near dislocation cores
seem to promote the crack propagation along the slip plane.
In the molecular dynamic analysis, we can deal with
atomic motion for only a very short period of pico-second
order. In actual, a lot of hydrogen atoms would be trapped for
a long period of time, and hydrogen atoms trapped with high
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Fig. 6 – Snapshots of crack propagation behavior for the crystal orientation (A) at 100 K (x0 [ 1.0 3 10L4).
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 49580
density at the dislocations emitted from a crack tip would
induce the fracture in the slip plane easier than expected in
the molecular dynamics simulations.
4. Molecular statics analyses of interactionbetween dislocation and hydrogen
In section 3, we presented the results of the molecular
dynamics analyses showing that hydrogen atoms gather
around the cores of edge dislocations on the {112} slip planes.
Here we will deal with this phenomenon in detail using the
molecular statics method. We will also show how hydrogen
atom affects dislocation motion. We have already published
the paper concerning these phenomena [19]. So we will show
the results briefly.
4.1. Hydrogen occupation sites around dislocation core
We consider the interaction between an edge dislocation on
the {112} slip planes and hydrogen atoms. Fig. 10 shows an
analysis model, in which the (112)[111] edge dislocation is
introduced on the xz-plane by removing an atomic plane and
relaxing the atomic structure using the conjugate gradient
(CG) method. The analysis model contains 8054 iron atoms,
and the dimensions of the unit cell are 11.05 nm in the
x-direction, 4.91 nm in the y-direction and 2.02 nm in the
z-direction. A periodic boundary condition was imposed on
Fig. 7 – Hydrogen distributions during cleavage in the slip plan
the x- and z-directions. The dislocation density in this system
is approximately 0.018nm�2. We used EAM-W as the inter-
atomic potential for the aFe–H system. A hydrogen atom was
allocated either at the T-site or the O-site near the dislocation
core, and then the positions of iron and hydrogen atoms were
relaxed to minimize the total potential energy using the CG
method. The hydrogen trap energy at each occupation site is
shown in Fig. 11. We can observe three regions with strong
hydrogen trap energy. The hydrogen trap energy is strongest
around the dislocation core. It is also relatively strong around
a high hydrostatic stress region (region A in Fig. 11) and along
a slip plane (region B in Fig. 11). It should be noted that the
region B has strong hydrogen trap energy. This could not be
predicted by the theory of elasticity that hydrogen and dislo-
cation interact mechanically as a result of lattice dilatation
caused by hydrostatic stress. Furthermore, the result suggests
that a lot of hydrogen atoms accumulate on the slip plane
around the dislocation core.
4.2. Effect of hydrogen on dislocation mobility
In situ TEM observation performed by Robertson et al. [6–8]
revealed that the distance between dislocations decreases
when hydrogen gas is introduced during TEM observation. This
fact indicates increase in dislocation mobility by hydrogen
atoms. The elastic analysis performed by Sofronis and Birn-
baum [20] showed that the shear stress acting on dislocation
decreases with increase in hydrogen concentration. This
e for the crystal orientation (B) at 400 K. (x0 [ 3.0 3 10L4).
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Fig. 8 – Snapshots of crack propagation behavior for the crystal orientation (B) at 400 K. (a) x0 [ 1.0 3 10L4 (b) x0 [ 3.0 3 10L4
(c) x0 [ 0.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 4 9581
phenomenon is called the hydrogen-induced shielding effect.
They concluded that the hydrogen-induced shielding effect
causes increase in dislocation mobility. Their conclusion needs
to be examined, because they did not consider the effect of
Fig. 9 – Enlarged views near the crack tip during cleavage in th
(a) x0 [ 1.0 3 10L4 (b) x0 [ 3.0 3 10L4.
hydrogen atoms at a dislocation core and dealt with extremely
high hydrogen concentrations such as 3 (the ratio of the
number of hydrogen atoms to the number of iron atoms) of 0.1
and 0.01. It is not confirmed whether the hydrogen-induced
e slip plane for the crystal orientation (B) at 400 K.
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Fig. 10 – Analysis model for molecular statics analyses of
interaction between dislocation and hydrogen.
Fig. 12 – Variations of the energy barrier for dislocation
motion.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 49582
shielding effect holds for a low hydrogen concentration.
Therefore we study the effect of hydrogen from the viewpoint
of energy barrier for dislocation motion.
According to the hydrogen trap energy obtained in 4.1, it is
the highest at a dislocation core. Thus the probability of
hydrogen occupation is the highest at the dislocation core.
From this reason, we placed hydrogen atoms at the disloca-
tion core. We used the same analysis model shown in Fig. 10.
We evaluated the energy barrier for the edge dislocation
motion of 1b (b: the magnitude of Burger’s vector) with and
without hydrogen, using the nudged elastic band (NEB)
method [21]. We obtained the energy barrier for the following
three cases; (a) without hydrogen atom, (b) with a hydrogen
atom at the dislocation core in the initial state and the dislo-
cation moving forward by 1b, and (c) with a hydrogen atom 1b
ahead of the initial dislocation and the dislocation moving to
the hydrogen atom. The hydrogen concentration of this
system is 2.24 mass ppm, and the number of hydrogen atoms
per unit length of a dislocation line is 0.49nm�1. The variations
of energy barrier with dislocation motion are shown in Fig. 12
for the case without a hydrogen atom and two cases with
a hydrogen atom. As shown in Fig. 12, the energy barrier for
dislocation motion is 2:65� 10�20J for the case without
a hydrogen atom, while it decreases to 2:35� 10�20J for the
Fig. 11 – Distribution of hydrogen trap energy at each site of
hydrogen.
case (b) with a hydrogen atom and 1:18� 10�20J for the case (c)
with a hydrogen atom. It is concluded that the energy barrier
for dislocation motion decreases due to hydrogen atoms.
Next we performed atomistic analyses in order to examine
whether the hydrogen-induced shielding effect holds under
a low hydrogen concentration. We obtained the stress field
around an edge dislocation based on atomistic model shown
in Fig. 10. The stress field around the dislocation is calculated
using the molecular statics method both for the case without
a hydrogen atom and for the case with hydrogen atoms. Fig. 13
shows the shear stress distributions along the slip plane near
a dislocation core. It is found from the figure that the shear
stress distribution is not affected by hydrogen atoms. Even if
the number of hydrogen atoms per unit length of a dislocation
line is increased up to 7.35nm�1, no significant difference in
the stress distribution is observed. So the hydrogen-induced
shielding effect is not observed under low hydrogen concen-
tration conditions.
Fig. 13 – Shear stress distributions along the slip plane near
a dislocation core.
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Fig. 14 – Effect of hydrogen atoms on the surface energy.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 4 ( 2 0 0 9 ) 9 5 7 6 – 9 5 8 4 9583
It is shown that hydrogen at a dislocation core reduces the
energy barrier for dislocation motion. It is also shown that the
hydrogen-induced shielding effect is very small. It is therefore
suggested that one reason for increase in dislocation mobility
under low hydrogen concentration conditions is not the
hydrogen shielding effect but the reduction of the energy
barrier for dislocation motion due to hydrogen.
5. Mechanism for hydrogen embrittlement
The molecular statics analysis using EAM-W provides the
result that hydrogen atoms existing on a slip plane promote
the separation of the slip plane because of decrease in its
surface energy caused by hydrogen atoms. Fig. 14 shows the
effect of hydrogen atoms on the surface energy of a-Fe for
{100}, {110} and {112} surfaces. Considering this fact and the
results shown in sections 3 and 4, we can propose the
following mechanism for hydrogen embrittlement of a-Fe
under monotonic loading, as follows:
(1) Dislocations are emitted from a crack tip and they exist
along a slip plane.
(2) A lot of hydrogen atoms are trapped at dislocation cores
and along a slip plane in the vicinity of a dislocation core.
(3) The hydrogen atoms at a dislocation core reduce the
energy barrier for dislocation motion and increases dislo-
cation mobility. Thus the distance between dislocations is
reduced.
(4) Separation of a slip plane is caused due to the hydrogen
atoms trapped by a dislocation, and such separation is
connected among pile-up dislocations.
Our proposed mechanism is a hybrid of the HELP and the
HEDE. The fracture is associated with the HELP mechanism in
that plastic deformation with dislocations is needed prior to
the fracture. On the other hand, the fracture is associated with
the HEDE mechanism in that the fracture results from the
separation of a slip plane.
Our proposed mechanism for the hydrogen embrittlement
agrees well with several experimental observations [22,23]
showing that the fracture of a hydrogen-charged test spec-
imen occurs at {110} or {112} slip planes.
6. Concluding remarks
We chose EAM-W as the best interatomic potential for the
aFe-H system. We performed the molecular dynamics anal-
yses of crack propagation in a-Fe including hydrogen atoms
under monotonic loading. We estimated the hydrogen trap
energy in the vicinity of a (112)[111] edge dislocation in order
to clarify the distribution of hydrogen atoms. We also evalu-
ated the energy barrier for dislocation motion under a low
hydrogen concentration. Based on the above results, we have
proposed a mechanism for hydrogen embrittlement of a-Fe
under monotonic loading. Our proposed mechanism agrees
well with several experimental observations.
Acknowledgements
This research was performed as a part of the Fundamental
Research Project on Advanced Hydrogen Science funded by
the New Energy and Industrial Technology Development
Organization (NEDO).
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