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Surface damage in ultrafine and multimodal
grained tungsten materials induced by low energy
helium irradiation
Osman El-Atwani
Birck Nanotechnology Center, Purdue University, [email protected]
Mert Efe
Purdue University, [email protected]
Bryan Heim
Purdue University
Jean Paul Allain
Birck Nanotechnology Center, Purdue University, [email protected]
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El-Atwani, Osman; Efe, Mert; Heim, Bryan; and Allain, Jean Paul, "Surface damage in ultrafine and multimodal grained tungsten
materials induced by low energy helium irradiation" (2013). Birck and NCN Publications. Paper 1358.
http://dx.doi.org/10.1016/j.jnucmat.2012.11.012
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Journal of Nuclear Materials 434 (2013) 170–177
Contents lists available at SciVerse ScienceDirect
Journal of Nuclear Materials
journal homepage: www.elsevier.com/locate/jnucmat
Surface damage in ultrafine and multimodal grained tungsten materials induced
by low energy helium irradiation
Osman El-Atwani a,b,⇑, Mert Efe a,c, Bryan Heim d, Jean Paul Allain a,b,d
a
School of Materials Engineering, Purdue University, West Lafayette, IN 47907, United States
Birck Nanotechnology Center, Purdue University, West Lafayette, IN 47907, United States
c
Center for Materials Processing and Tribology, Purdue University, West Lafayette, IN 47907, United States
d
School of Nuclear Engineering, Purdue University, West Lafayette, IN 47907, United States
b
g r a p h i c a l a b s t r a c t
SEM micrograph of a multimodal grained tungsten material after being irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2 showing the
detachment and the nanostructuring of the small grains as well as pore formation on top of the nanostructures. The sample was etched before irradiation
for better imaging of the grain boundaries.
a r t i c l e
i n f o
Article history:
Received 3 August 2012
Accepted 8 November 2012
Available online 20 November 2012
a b s t r a c t
Although tungsten is considered the best candidate as a plasma facing component (PFC) in the divertor
region in the International Thermonuclear Experimental Reactor (ITER), severe morphology changes such
as cavities, blisters, bubbles and nanostructure formation are expected. Increasing defect sinks in the
tungsten microstructure is one of the possible solutions to mitigate the irradiation damage. In this work,
helium irradiation at low energy (50 and 200 eV) and temperatures of 600 (threshold of vacancy migration) and 950 °C were performed on multimodal and ultrafine grained tungsten prepared by spark plasma
sintering and severe plastic deformation (SPD), respectively. The multimodal samples consisted of small
grains (300–700 nm size) juxtaposed to larger grains (1–3 lm size). Detachment of the small grains was
observed in the multimodal grained tungsten irradiated at 600 °C and a fluence of 1 1022 m2 due to
grain boundary grooving. On the same sample but at 950 °C, detachment and nanostructuring of the
small grains were observed together with recrystallization of the large grains. Irradiation of the SPD
⇑ Corresponding author at: School of Materials Engineering, Purdue University,
West Lafayette, IN 47907, United States.
E-mail address: [email protected] (O. El-Atwani).
0022-3115/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.jnucmat.2012.11.012
O. El-Atwani et al. / Journal of Nuclear Materials 434 (2013) 170–177
171
samples at 200 eV and 950 °C to a fluence of about 2 1022 m2, resulted in nanostructuring of the ultrafine grained shear bands in the microstructure.
Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction
Several advantages make tungsten the primary material choice
for the divertor region in ITER [1,2]. These advantages range from
high melting point, high thermal conductivity, high sputtering
threshold to low sputter yield and low tritium retention. Furthermore, tungsten is also considered in all-metal concepts for ITER
[3]. However, one drawback of tungsten is its irradiation intolerance in the extreme environment of fusion reactors [4,5] where helium is one of the main outputs of the nuclear reaction.
Both helium and neutron irradiation introduce damage to a
tungsten material facing the plasma. Helium irradiation on tungsten was studied under systematic conditions, ranging from low
[6] to high [7] temperatures, and 1–10 eV irradiation [8,9] up to
tens of keV [10] with different fluences that reached in some cases
over 1027 m2 [7,9].
Depending on the irradiation conditions, different changes in
the microstructure of tungsten occur. Formation of He bubbles is
the most known and is believed, in some studies, [11,12] to be
the cause of the other microstructural changes. Bubble formation
is related to the defect dynamics in irradiated tungsten. If we assume that the binary collision approximation holds, then the energy needed for a helium atom to displace a tungsten atom upon
head to head collision can be calculated by the following equation:
Ed ¼
4M 1 M 2
ðM1 þ M2 Þ2
Ei
ð1Þ
where Ed = displacement energy, Ei = incident energy of helium
atom, and M1 and M2 are the mass number of tungsten and helium,
respectively.
For 40 eV displacement energy of tungsten atoms, helium energy of 480 eV is needed (Eq. (1)). It was reported [9] that helium
energy of 500 eV can lead to tungsten atoms displacement and that
if tungsten is irradiated with energies higher than this energy,
vacancies and interstitials are created.
Interstitials in tungsten can migrate at very low temperatures,
including RT, due to its low migration energy (0.08 eV) [6]. Migration of interstitials will leave vacancies behind that can coalesce
and form small voids, and which in turn can permit some He gas
to go into and form a helium bubble. Even if tungsten is irradiated
with energies below the displacement energy, bubbles can grow
due to loop punching by ejecting tungsten atoms from their lattice
sites to form interstitial loops [9,13]. Minyou [9] found that 5 eV
was the bubble formation threshold energy of helium for samples
irradiated at temperatures between 2400 and 3200 °C and fluence
of 1027 m2. Bubble size was demonstrated to increase with fluence and irradiation energy.
Increasing the temperature to above 600 °C allows vacancies to
migrate in tungsten in addition to interstitials, and the formation
of voids and bubbles is enhanced. Bubbles, also can migrate and
coalesce, [11] and this is the reason for larger bubbles being observed at higher temperatures. Bubbles diffusion to the surface
can result in surface pores, which are another microstructural
damage due to helium irradiation of tungsten. Moreover, internal
stresses are believed to play a role [14] in the observed blistering
of ion irradiated tungsten surfaces [15,16].
Bubbles, pores and blister formation are not the only microstructural modification in tungsten under helium irradiation. It
was shown [12] that nanostructure formation can also occur on
the surface of irradiated tungsten. Small and dense nanostructures
on irradiated tungsten surfaces are known as ‘‘fuzz’’ [17]. Kajita
et al. [12] developed a map of energy, temperature and fluence
for which bubbles and nanostructure formation can occur. According to that study, nanostructures can appear in a temperature range
on which an upper and lower bound represent the fast surface diffusion and the slow bubble growth and migration, respectively.
2. Mitigating the irradiation damage: increasing defect sinks
Since defect migration and annihilation affect the bubble and
the nanostructure formation in tungsten under helium irradiation,
suppression of point defects by annihilating the freely migrating
defects (vacancies and interstitials) to defects sinks, such as grain
boundaries can lead to higher irradiation tolerance of tungsten
materials. In this case, higher fluence threshold of helium particles
[18] will be necessary to introduce the above mentioned microstructural changes to tungsten materials. To increase the sink volume in tungsten materials, refining the microstructure to the
nanoscale is essential. Ultrafine and nanocrystalline tungsten
materials [19,20] have higher grain boundary area to hold higher
helium fluences. Moreover, in small grain size samples, the diffusion path of helium particles to sinks is shorter.
Another important aspect to discuss is the mechanical properties of tungsten materials. Large grain tungsten materials are
known for their poor ductility [21,22]. In addition, the ductile to
brittle transition temperature of tungsten is high (600 K or
900 °C) [14]. Thus, working with tungsten as a plasma facing component means that the operation temperature has to exceed
900 °C, but not the temperature limit of creep [14]. Thus, it is
important to begin with a tungsten material of higher ductility.
Ductility in tungsten is postulated to be enhanced by the removal
of impurities, particularly oxygen, from the grain boundaries [23].
Ultrafine (<500 nm) and nanocrystalline tungsten (<100 nm) were
proven to have increased plastic flow, [20] higher yield strength,
and improved ductility. It is also postulated that improvement in
ductility and strength can be achieved through the formation of bimodal grain size distribution materials, [24] where the small grains
give the high strength and the large grains enable ductility. ElAtwani et al. [25] observed an increase in the material flow during
the microindentation measurement of multimodal grain size distribution tungsten, an indication of increased plastic flow.
Formation of ultrafine and nanocrystalline tungsten can be performed using severe plastic deformation (SPD) methods such as
equal-channel angular pressing (ECAP) and high-pressure torsion
(HPT). Wei et al. [20] used HPT to prepare superstrong and less
brittle tungsten materials of hardness approximately 11 GPa. SPD
methods, however, have some limitation on the sample size and
geometry. For preparation of samples scalable to industrial application, powder consolidation processes can be used. Spark plasma
sintering (SPS) may be the best sintering option to prepare tungsten materials of small grain sizes due to its shorter sintering time
and lower sintering temperature. These advantages are crucial in
sintering tungsten, due to its low secondary recrystallization temperature (1400 °C), [26] where rapid grain growth occurs. In addition, some sintering studies on tungsten have shown that SPS leads
to impurity removal during sintering [27,28]. However, formation
of ultrafine and nanocrystalline tungsten with high relative density
(near the theoretical density) requires advances in tungsten powder synthesis.
In this work, we present an irradiation study on two types of
tungsten materials. The first type is multimodal grain size distribution tungsten, with ultrafine grains juxtaposed to fine grain
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Table 1
Average grain sizes, densities and Vickers hardness of the samples. The average grain size of each mode in the multimodal size distribution samples (W3 and W4) was found from
the Gaussian fit of the grain size distribution guide-the-eye plot.
Vickers hardness (kg/mm2)
± standard deviation
Sample
Peak #
Mode average
grain size (lm)
Standard deviation
of peaks (lm)
W3
Mode
Mode
Mode
Mode
1
2
3
4
0.330
0.670
1.27
2.18
0.067
0.147
0.372
0.132
93.2
523 ± 23
W4
Mode 1
Mode 2
Mode 3
0.737
1.79
3.38
0.363
0.818
0.002
94.0
518 ± 14
Commercial W
SPD
17
0.8
Relative density (%)
Average grain size (lm)
tungsten, prepared using spark plasma sintering [25]. The other
type of the tungsten material in ribbon forms is prepared by
orthogonal machining of bulk tungsten (a severe plastic deformation method, SPD). Irradiations were performed using He at different energies and temperatures on these types of tungsten samples
described above.
3. Experimental
Powder metallurgy tungsten samples (W3 and W4) of multimodal grain size distribution were prepared with powder consolidation using the spark plasma sintering apparatus (Sumitomo Coal
Mining Co., Japan, Model 825S). The sintering temperature was
1300 °C and 1400 °C for W3 and W4 respectively, and the pressure
was 266 MPa. Consolidation time was 5 min for both samples. Details about the consolidation conditions are mentioned in our previous work [25].
For production of the SPD samples, commercial purity (99.95%)
tungsten in disk form (American Elements, Los Angeles, CA) was
subjected to plane strain orthogonal machining. The same machining method has been investigated before as a severe plastic deformation technique to produce nanocrystalline or ultrafine grained
bulk materials [29]. Large strains can be imposed in a narrow
and confined shear deformation zone in a single pass in order to refine microstructures to the ultrafine level. The other thermomechanical conditions such as strain rate and temperature can also
be modeled and controlled in the deformation zone as accurate as
strain. Since machining is a high strain rate process by its nature,
adiabatic heating due to the large plastic deformation can cause
a local temperature rise in the deformation zone. The temperature
rise in the deformation zone was calculated and reported, even
though all the machining was done at room temperature without
the pre-heating of the disk. The process parameters and equations
governing the strain, strain rate and temperature in the deformation zone are detailed in the previous work [29].
Semi-continuous and curled platelet chips (ribbons) were produced in 6.5–7.5 mm in length, 5 mm in width and 0.35–
0.45 mm in thickness with rectangular cross-sections. For the
tungsten chips used in this study, the effective strain, e ¼ 1:9, the
effective strain rate, de=dt ¼ 3:8 104 and the deformation temperature, T 900 °C was obtained in the deformation zone. The
time spent at the deformation temperature was only 10 ms, which
corresponds to the time taken for a material element to traverse
the deformation zone (100 lm). The chip rapidly quenches back
to the room temperature in few seconds. The chips were ground
and polished from the surfaces and irradiation experiments were
performed on the middle cross-section (i.e. mid-thickness) in order
to avoid the oxide layer that was formed on the surfaces.
98
100
448 ± 21
556 ± 31
The commercial tungsten sample, used in this work, was supplied by Materials Modification Inc., Fairfax, VA (MMI). The sample
was prepared using the plasma pressure compaction (PPC) technique, and then were hot isostatically pressed. Table 1 summarizes
the samples used in the irradiation experiments.
Irradiation of the samples was performed in the Particle and
Irradiation Interaction of Hard and Soft Matter (PRIHSM) facility
at Purdue University. The irradiation in this case, was performed
at 50 and 200 eV He ions, which are relevant energies to the conditions expected in the divertor region of ITER [30]. The base pressure, during irradiation, was 2 108 torr. The samples were
heated to the intended temperature using an in situ e-beam heater.
For microstructure characterization, scanning electron microscopy (SEM), and electron backscattering diffraction (EBSD) imaging
were performed on the samples using a Philips XL 40 SEM (with
Schottky FEG) and a Hitachi S4800 FESEM. Cross section imaging
was performed using FEI xT Nova NanoLab Dual Beam focused
ion beam/scanning electron microscope (FIB/SEM), and secondary
electron images were taken using the electron beam. Transmission
electron microscopy (TEM) characterization of the SPD samples
was performed (after mechanical thinning) using FEI Titan 80/
300 field emission TEM, operating at 300 kV, equipped with a Tridiem GIF and were zero-loss filtered. The samples were prepared
using the twin-jet electropolishing technique with a 0.5% NaOH –
water solution at RT.
4. Results and discussion
Fig. 1a and b shows the scanning electron microscopy images of
multimodal grain size distribution samples before irradiation. The
number of modes in the distribution with the average grain size in
each mode is given in Table 1. Microstructure characterization
showed a higher relative density of the samples than those measured by the Archimedes approach (Table 1). To better illustrate
the densification characteristics, cross section imaging was performed. Fig. 1c and d shows the cross section FIB and TEM images,
respectively, of W3 sample before irradiation. Limited porosity and
a high relative density are observed.
4.1. Irradiation of multimodal (powder metallurgy) samples at 600 °C
As discussed before, vacancy migration in tungsten has a
threshold temperature of 600 °C [6]. Interstitial migration, on the
other hand, occurs even at RT. To examine the effect of this fact,
the samples were irradiated at 600 °C (only interstitial migration
occur) and 950 °C (both vacancy and interstitial migration occur).
Irradiation results of 200 eV He ions at 600 °C and 1 1022 m2
fluence are shown in Fig. 2a and b. Small grains detach from the
O. El-Atwani et al. / Journal of Nuclear Materials 434 (2013) 170–177
173
Fig. 1. (a and b) SEM micrographs of W3 and W4, respectively (c) cross-section micrograph of W3 using the electron beam in a focused ion beam (FIB) system (d) Bright field
transmission electron microscopy cross section micrograph of W3.
Fig. 2. (a and b) SEM micrograph of W3 and W4, respectively, showing the regions of fine and ultrafine grains as irradiated with 200 eV He ions at a temperature of 600 °C and
a fluence of 1 1022 m2. (c and d) higher magnification regions showing the small (ultrafine grains) of the irradiated W3 and W4, respectively.
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O. El-Atwani et al. / Journal of Nuclear Materials 434 (2013) 170–177
large grains after the irradiation. Using the Stopping and Range of
Ions in Matter (SRIM) software, [31] the penetration depth of He in
tungsten at 200 eV is simulated to be around 5 nm. The density of
He atoms (=fluence/depth) on the surface is then determined to be
2 1029 ions m3. While 200 eV He atoms do not cause a displacement of W atoms, the high concentration of He atoms can generate
a high density of defects. This occurs due to loop punching, where
helium atoms are trapped by vacancies and grow as bubbles by
ejecting tungsten atoms. Migration of defects, in this case interstitials, in the 5 nm layer, initiates the necessary diffusion to start
grain boundary grooving. The grain boundary grooving process,
modeled by Mullins (1957), [32] is based on the conservation of
the total mass. Grain boundary grooving occurs when atoms near
the triple junction diffuse or evaporate. Since the vapor pressure
of tungsten is very low at the irradiation temperature, [33] grain
boundary grooving can occur through diffusion. Diffusion, in turn,
can be volume or surface diffusion. With a penetration depth of
5 nm in tungsten, surface diffusion dominates [34].
In addition to its defect generation, we believe that ion beam
irradiation on the surface results in momentum transfer of the ions
Fig. 3. SEM micrographs of: (a) W3 irradiated with 200 eV He ions at 950 °C and a fluence of 1 1021 m2 (b) W4 irradiated with 200 eV He ions at 950 °C and a fluence of
1 1021 m2 (c) W3 irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2 (d) W4 irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2 (e
and f) higher magnification of the large grains of W3 and W4 respectively after being irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2 (g) cross-section
using the electron beam in a focused ion beam system of W3 after being irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2 (h) commercial tungsten after
being irradiated with 200 eV He ions at 950 °C and a fluence of 1 1022 m2.
O. El-Atwani et al. / Journal of Nuclear Materials 434 (2013) 170–177
to the diffusing atoms which enhances the grain boundary grooving by promoting higher surface diffusion (i.e. radiation enhanced
diffusion). It is expected that the small grains will be more affected
by grooving, due to the smaller grain boundary curvature than the
large grains.
Surface diffusion being the main mechanism of grain boundary
grooving, helium bubbles on grain boundaries exacerbate the
grooving process. Helium bubbles form first on the grain boundaries and the dislocations [35]. Helium ions trapped in grain
boundaries (2D trapping) and in dislocation cores (1D trapping),
have limited diffusion than in bulk (3D trapping). Helium bubble
formation on the surface of grain boundaries and triple junctions,
and the subsequent bubble growth and burst, intensify the grain
boundary grooving process described above.
In addition to the detachment of the small grains, pore (10–
25 nm) formation on the small grains is observed (Fig. 2c and d),
where the pore size increases with the grain size. On the large
grains, however, very small bumps (20 nm) on the surfaces are
observed, which will be discussed in the next section. Large pores,
as the ones observed by Cipiti and Kulcinski [10] were not formed.
At 200 eV energy, tungsten displacement does not occur, and
the thermal defects on the surface are of low density due to the
low irradiation temperature (600 °C). At this temperature, vacancy
migration and vacancy cluster formation is very low. Moreover, the
mobility and coalescence of any small bubbles formed at this temperature is also low, [9] which can explain the low damage density
(small pores) on the surface. Pore formation on the small grains is
believed to occur after the detachment from the large grains, when
no helium sinks (grain boundaries) are available to trap He ions.
4.2. Irradiation of multimodal (powder metallurgy) samples at 950 °C
As the temperature increases from 600 to 950 °C, vacancy
migration begins to occur, in addition to higher thermal defect
generation. Fig. 3a–f shows the microstructures after He irradiation
at 200 eV and 950 °C for two different fluences: 1 1021 m2 and
1 1022 m2. At 1 1021 m2, not only the small grains are detached from the large grains, but structuring of the small grains occurred. Structures of preferred orientation are observed on W4. The
small grains in W3 is more severely damaged than in W4. This may
be due to more rapid grain boundary grooving of the small grains
in W3 and the subsequent detachment from the large grains. As
the fluence increased one order of magnitude to 1 1022 m2,
the damage to the small grains increased. Cross section imaging
of W3 at 1 1022 m2 (Fig. 3g) shows nanostructure formation
on the small grains.
Small bumps on the surface are now clearer on the large grains in
case of W3 and W4 samples due to the higher irradiation temperature. We consider these bumps to be small grains (triple junctions
are obvious), formed due to the recrystallization of the samples.
While recrystallization of tungsten occurs at 1700 °C, irradiation-
175
Fig. 5. SEM micrograph of W3 after being irradiated with 200 eV He ions at 950 °C
and a fluence of 1 1022 m2 showing the detachment and the nanostructuring of
the small grains as well as pore formation on top of the nanostructures. The sample
was etched before the irradiation for better imaging of the grain boundaries.
enhanced recrystallization [36,37] can result in recrystallization of
the grains at temperatures lower than 1700 °C. Irradiation enhanced
recrystallization happens due to the following mechanisms: the creation of defect clusters that increase the system free energy (more
nucleation driving force), [36] the increase in defect and grain
boundary mobility [38,39] (radiation-enhanced diffusion), [40]
and the increase in dislocation density (loop punching) [41,42].
Radiation enhanced diffusion is a result of high defect concentration
which increases the diffusion rate [43]. As mentioned earlier, defect
generation in our case, occurs due to loop punching. Due to low
implementation depth of He (5 nm) at the irradiation conditions
used in this study (200 eV), defect concentration is expected to be
high, and the transfer of momentum to the diffusing atoms by collision is expected to exacerbate their mobility; therefore, recrystallization will be enhanced. Irradiating commercial tungsten samples
showed similar recrystallization behavior (Fig. 3h). It should be
noted that damage to a small grain at a depth of 700 nm was observed in the W3 sample irradiated at 200 eV and 950 °C for a fluence of 1 1022 m2 (Fig. 3g). However, damage at similar depths
was not observed in other parts of the sample, and this study has
no strong evidence of irradiation damage at higher depths than
the ions implementation depth (5 nm at the irradiation conditions
as determined by SRIM). Thus, the damage is considered to be consistent with the implementation depth of He.
Decreasing the irradiation energy from 200 eV to 50 eV did not
have much effect on the results (Fig. 4a and b). This is due to the
fact that neither energy cause displacement in tungsten, and the
damage observed is due to the defect generation on a shallow
depth in tungsten.
Fig. 4. SEM micrograph of: (a) W3 and (b) W4 as being irradiated with 50 eV He ions at 950 °C and a fluence of 1 1022 m2.
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Fig. 6. (a) TEM bright field image of the SPD sample showing the shear band region. The region consists of elongated ultrafine and nanocrystalline grains. (b) EBSD
micrograph of the SPD sample showing the ultrafine grained shear bands and large grained matrix. (c) Number fraction vs. grain orientation angle of the SPD sample with an
inset for the inverse pole figure (d) SEM micrographs of the SPD samples after being irradiated with 200 eV He ions at 950 °C and a fluence of 2 1022 m2 (e and f) higher
magnifications of the bright contrast region in (d).
In addition to the enhanced recrystallization behavior of the
large grains, nanostructuring of the small grains (at 950 °C) was
observed on W3 and W4 samples at fluences starting from
1 1021 m2. Nanostructuring of tungsten surfaces is believed to
be caused by bubble coalescence and migration [12]. According
to that theory, pinholes are formed on the surface when migrating
bubbles bisect each other. At later irradiation stages, the pinholes
grow in height as the bubbles continue to bisect each other at
the bottom of the pinholes. Swallowing then follows inside the
nanostructures themselves through the coalescence of the bubbles.
We believe that this process is analogous to the grain boundary
grooving observed for the small grains, which is enhanced also
by bubble coalescence and bisection. Fig. 5 shows an SEM image
of a small grain (W3 sample) nanostructured after helium ion irradiation at 200 eV and 950 °C. Nanostructures and pores on top of
the structures are observed suggesting bubble formation inside
the nanostructures. Bubble formation on the sides of the detached
grains at depths of 500 nm confirms the detachment of the grains
before the nanostructure formation.
4.3. Irradiation of ultrafine (SPD) samples at 950 °C
The microstructure of the ultrafine grained tungsten sample
(formed by severe plastic deformation) is shown in Fig. 6a. Microstructure consists of elongated ultrafine grains (width < 0.4 lm) as
the result of the severe plastic deformation, much like the
ECAP + rolling [19] and HPT [20] processed tungsten. EBSD micrograph from a larger area (Fig. 6b) shows a non-uniform microstructure which is a mixture of ultrafine grained shear bands and large
grained matrix. While the shear bands consist mostly of relatively
lower-angle ultrafine grains due to the localized deformation, as
shown in number fraction vs. grain boundary misorientation angle
plot (Fig. 6c); the rest of the sample is formed of larger grain size
(>5 lm). The original grains in the bulk sample are thought to be
preserved after the deformation due to the less strain in the matrix.
It is a known fact that, the shear bands form due to the adiabatic
heating during the high strain rate deformation [19,44]. The temperature rise in the deformation zone, in our case T 900 °C, re-
sults in flow softening followed by plastic instability and shear
localization [19,44]. Meyers and co-workers [44] explain the
microstructure development inside the shear bands through the
dynamic recovery mechanism. Dynamic recovery (i.e. continuous
rotational dynamic recrystallization) occurs inside the shear bands
via dislocation rearrangement into elongated dislocation cells. The
dislocation cells subsequently increase their misorientation angle
through dislocation accumulation and become subgrains as strain
continues to localize in the shear bands. Finally, the elongated subgrains fragment into recrystallized high-angle ultrafine grains
leaving the microstructure almost free of dislocations. We believe
that our microstructure inside the shear band did not reach to
the final stage of dynamic recrystallization and still consists of
elongated subgrains and dislocations due to the relatively lower
effective strain, e ¼ 1:9.
Irrespective of the deformation mechanisms, the microstructure
after SPD was highly non-uniform. The sample was irradiated at
950 °C and 200 eV He up to a fluence of 2 1022 m2. As seen in
Fig. 6d, the irradiation damage was also non-uniform. The shear
banded regions were nanostructured unlike the other regions
resulting in a contrast difference in SEM micrograph (Fig. 6d).
Figures 6e and f show the higher magnification images of the
damaged region having the white contrast.
We attribute this observation to the high density of dislocations
in the subgrains and their boundaries inside the shear bands,
which would trap more He ions. Bubble formation then should occur faster than the large grains, and the coalescence of bubbles initiates nanostructure formation as illustrated before in Section 4.2.
The behavior of the grain size and orientation during the irradiation process is best analyzed using an in situ TEM work, the focus
of ongoing work in our group.
5. Conclusions
Tungsten is arguably the best material choice for the divertor region in the future nuclear fusion reactors. However, its drawbacks,
such as its tenuous brittle behavior and irradiation-induced damage
morphology, require innovative studies to modify tungsten
O. El-Atwani et al. / Journal of Nuclear Materials 434 (2013) 170–177
materials that can sustain the severe plasma conditions in fusion
reactors. In this work, irradiation of multimodal grained tungsten
(of fine and ultrafine grains) produced by powder consolidation
and ultrafine grained tungsten processed by severe plastic deformation (orthogonal machining) are irradiated with He ions at low
energy and high temperatures for fluences up to (1–2) 1022 m2.
Irradiation of these materials revealed the following:
– Irradiation of multimodal grain tungsten at 600 °C (below the
threshold of vacancy migration) at 200 eV, up to fluences of
1 1022 m2, showed detachment of small grains (ultrafine
grains) from the large grains (fine grains), and pore formation
on the small grains. Detachment of small grains is a result of
irradiation-enhanced diffusion of defects and grain boundary
grooving.
– Irradiation of the same materials at 950 °C (where both vacancy
and interstitial migration occur) to fluences of 1 1022 m2
showed complete damage (nanostructuring) of the small grain
regions and recrystallization of the large grains. Detachment
of the small grains was clear at fluence of 1 1021 m2 (concentration of He = 2 1029 m3). Preferred orientation of the nucleated nanostructures was observed. Nanostructure formation
was attributed to large bubble formation in the small grains,
which become single crystalline tungsten regions after the
detachment from the large grains. Pores on top of the structures
were observed due to bubbles bursting on the structure surface.
Similar observations were noticed as the irradiation energy was
decreased from 200 to 50 eV, since neither energy induce tungsten displacement.
– Irradiation of ultrafine grain tungsten materials (SPD) at 950 °C
and 200 eV He showed nanostructure formation on the shear
bands (ultrafine and nanocrystalline grains), unlike the large
grained matrix.
It is concluded that the grain boundaries and the dislocations
can play a crucial role in the microstructural damage of tungsten
materials exposed to high helium irradiation doses. The refined
microstructures of the tungsten materials used in this study, however, did not improve their resistance to He modification. Additional studies at higher fluxes and fluences are necessary to
judge about their resistance to fuzz formation. In addition to high
irradiation fluences, early stage irradiation work is also essential to
study the nanostructure formation mechanism, and to understand
the behavior of grain boundaries in small grained tungsten materials. Ongoing work in our group is focusing on morphology characterization at early, moderate and high irradiation fluences to study
the structure formation on these materials and their effect on the
performance of tungsten as a plasma facing component. In situ
irradiation-TEM studies are being performed on ultrafine grained
tungsten for that purpose as well as to examine the behavior of
grain boundaries and defects during the irradiation process. The
in situ TEM results will be presented in another paper.
Acknowledgements
This work was supported in part by the US Department of Energy’s 2010 Early Career Award DE-SC0004032. We would like to
thank Prof. Anter El-Azab for the valuable discussion regarding
the irradiation damage of tungsten materials, Prof. Joanna Groza
177
and Dr. Dat Quach for the spark plasma sintering of the tungsten
powders, and Bradley Schultz for the assistance with the
metallography.
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