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TRANSCRIPT
Characterization of lattice damage formation in tantalum irradiated at
variable temperatures
I. Ipatovaa,b,*, P.T. Wadyb, S.M. Shubeitab, C. Barcellinia, A. Impagnatielloa, E.
Jimenez-Meleroa
a School of Materials, The University of Manchester, Manchester M13 9PL, UK
bDalton Cumbrian Facility, The University of Manchester, Moor Row CA24 3HA, UK
*email: [email protected] . Phone: +44 78 49290480
Summary
The formation of radiation-induced dislocation loops and voids in tantalum at 180(2), 345(3)
and 590(5)C was assessed by 3MeV proton irradiation experiments and subsequent damage
characterization using transmission electron microscopy. Voids formed at 345(3)C and were
arranged into a body centred cubic lattice at a damage level of 0.55dpa. The low vacancy
mobility at 180(2)C impedes enough vacancy clustering and therefore the formation of voids
visible by TEM. At 590(5)°C the Burgers vector of the interstitial-type dislocation loops is
a<100>, instead of the a/2 <111> Burgers vector characteristic of the loops at 180(2) and
345(3)C. The lower mobility of a<100> loops hinders the formation of voids at 590(5)C up
to a damage level of 0.55dpa.
Lay description
High-temperature metallic materials for demanding technological applications in radiation
environments, such as future nuclear reactor systems and enhanced-output targets for
spallation sources, will be subject to the continuous bombardment of energetic particles at
elevated temperatures. Tantalum constitutes an advanced candidate material for those
applications due to its high radiation tolerance, ductility and water corrosion resistance. We
have characterized the damage caused by proton bombardment to tantalum at variable
temperatures using transmission electron microscopy. The results revealed significant
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differences in structural damage as a function of temperature and damage level, and therefore
help to understand the formation, or otherwise, of voids induced by radiation. These results
constitute unique evidence of the occurrence of structural damage and voids in tantalum,
which will pave the way to reliable predictions of radiation-induced swelling of tantalum-
based components in future applications.
Keywords: Nuclear materials, Tantalum, Electron microscopy, Radiation damage,
Dislocation analysis, Void formation.
Introduction
The continuous exposure of structural materials to intense radiation fields causes atomic
displacement cascades that generate a high density of vacancies and self-interstitial atoms
(SIAs) in equal proportions. Those lattice point defects can evolve into a range of defect
structures such as dislocation loops, nano-clusters, stacking fault tetrahedra or voids
(Was, 2007; Odette et al., 2008). The long-term lattice damage brings along phenomena such
as radiation-induced hardening or void swelling that compromises the integrity of key
structural components in applications such as nuclear reactor cores or targets in neutron
spallation sources (Yvon et al., 2009; Azevedo, 2011; Zinkle et al., 2013). Radiation-induced
void swelling was initially reported in face-centred cubic metals and alloys, such as austenitic
stainless steel or nickel (Cawthrone et al, 1967; Norris, 1971). Material candidates for future
nuclear reactor systems and higher-output targets for spallation sources are mainly based on
body-centred cubic (bcc) metals due to their enhanced radiation tolerance and, in particular,
to their higher void swelling resistance (Singh et al., 1995; Stork et al., 2014). Tantalum
currently stands out as an advanced high-temperature material candidate due to its additional
benefits of a high fluence threshold for He+ ion-induced surface nano-structuring, high water
corrosion resistance, good workability, and ductility retention at relatively high radiation
damage levels (Chen et al., 2003; Nelson et al., 2012; Novakowski et al., 2016).
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Void formation and swelling in bcc metals is interpreted in terms of the biased absorption of
SIAs at defect sinks such as dislocation loops, and the evolution of the resultant vacancy
excess in the matrix into vacancy clusters and eventually voids (Evan et al., 1985; Trinkaus et
al., 1993). This overall process depends on the mobility of SIAs, vacancies and defect sinks,
and therefore on temperature. However, the existing experimental evidence of the formation
and evolution of SIA and vacancy arrangements at variable temperatures in tantalum is still
fragmented. Earlier radiation damage studies in tantalum bombarded with neutrons up to a
fluence of 2.5×1022 neutron/cm2 at a temperature of 585C reported the appearance of voids
ordered on a bcc superlattice. Those voids randomise at higher irradiation temperatures up to
1050C and are absent at 425C (Wiffen, 1977). Heavy ion bombardment of tantalum induces
the formation of vacancy clusters that evolve into randomly distributed voids at temperatures
higher than 400C (Yasunaga et al., 2000). In this work we aimed to assess the formation of
radiation-induced dislocation structures in tantalum at variable temperatures and their impact
on the occurrence, or otherwise, of void formation. We used a 3MeV proton beam produced
by an electrostatic particle accelerator as a surrogate of neutron bombardment in void
formation studies, since it generates a uniform radiation damage profile through a sample
thickness of 30 μm. Ion irradiation experiments have been performed successfully since
1970s on metallic materials to mimic void swelling induced by neutron damage (Nelson et
al., 1970; Mazey, 1990).
Experimental
Proton irradiation
The starting Ta material was provided by Goodfellow Cambridge Ltd. in the form 1mm-thick
sheet. The as-received material was annealed at 1400°C during 2 hours for recrystallization.
Irradiation experiments were performed using the 3 MeV proton beam generated by a 5 MV
tandem ion accelerator and a high-current TORVIS source installed at the Dalton Cumbrian
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Facility of the University of Manchester (Wady et al., 2016). Equivalent samples were
irradiated at either 180(2), 345(3) or 590(5)C. During irradiation experiment the temperature
was monitored by a non-contact IR camera placed at an angle of 30° with respect to the
sample surface. Additionally, thermocouples were spot welded onto the samples close to the
irradiated area. Liquid indium was used in order to improve the thermal contact between the
samples and the NIMONIC75® alloy block of the sample stage, the latter containing the
heaters and the water cooling loop. The average temperature values correspond to the
irradiation period when the temperature was held relatively constant, i.e. excluding the
heating to reach the target temperature and the final cooling down to room temperature, and
the error in the irradiation temperature corresponds to the standard deviation of the
temperature distribution during that period. The charge accumulated on the sample was
measured during the irradiation experiment using a picoammeter. The accumulated charge
was used to calculate the proton fluence. The irradiation conditions for the studied samples
are summarised in Table 1. The damage rate was 1×10-5dpa/s in all irradiations. The
thickness of the irradiated layer of the samples was determined by nano-indentation
measurements using a NanoIndenter-XP (MTS Systems Corp.) after sample cross sectioning,
and supported by simulations using the SRIM software with the quick Kinchin–Pease
approach and default values for other software settings (Ziegler et al., 2010; ASTM E521-96,
2009), see Fig. 1.
Characterization of radiation damage
3mm-diameter discs for transmission electron microscopy (TEM) were prepared by electro-
polishing using the Struers TenuPol-5 and an electrolyte containing 15vol.% sulphuric acid
(95%) – 85vol.% methanol at a temperature of 5C. Regions of the irradiated layer at two
damage levels, namely 0.15 and 0.55 dpa, were selected for further analysis. In order to
achieve the targeted level of damage, prior to electropolishing the non-irradiated side of the
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sample was covered with Elektron Technology’s acid-resistant Lacomit varnish which was
simply removed thereafter by acetone. The sample was electro-polished from the irradiated
side to a preselected depth below the surface. We assessed the amount of material removed
by electro-polishing using a Keyence VK-X200K 3D Laser Scanning Microscope. A sample
thickness of 20μm and 28μm was removed in order to obtain a TEM thin foil at the
damage level of 0.15 dpa and 0.55 dpa respectively. Afterwards, the Lacomit varnish was
removed and applied to the irradiated side of the sample, and finally the sample was back
thinned for electron transparency by electropolishing (Z. Yao et al., 2008). The structural
analysis was performed using the 200kV FEI Tecnai T20 Transmission Electron microscope
equipped with a double-tilt specimen holder. The determination of the Burgers vector b of the
dislocations was based on the g.b=0 invisibility criterion, where g denotes the scattering
vector. This determination was based on estimating all g.b possible values for the relevant
reflections from the TEM analysis made by tilting the sample to specific two-beam
conditions. Prismatic dislocations in bcc metals are characterised by the Burgers vector
a/2<111>, a<001> or a/2<110> (Okamoto et al., 1966). The predicted invisibility, or
otherwise, of the loops was confronted with the experimental evidence from the acquired
TEM images. The foil thickness was derived from the fringes spacing of the convergent beam
electron diffraction (CBED) pattern. The error in thickness measurements using CBED
patterns from the exact Bragg condition is assumed to be ±10% (Harte et al., 2017). Bright-
field imaging of small radiation-induced voids was based on the “out-of-focus” imaging
technique (Jenkins et al., 2001).
Results and discussion
Dislocation structure
The lattice damage in tantalum is evidenced by the presence of a relatively high density of
ellipsoidal dislocation loops at the three studied irradiation temperatures and a damage level
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of 0.15dpa (Fig. 2). The TEM images of dislocation loops presented in Figure 2 have been
taken along the [111], [113] or [012] zone axis. These loops are formed in the debris of the
damage cascade by the migration and clustering of SIAs. At that damage level, the number
density of dislocation loops observed at 180(2)°C is the highest (~110×1021m-3), and
corresponds to approximately three times the loop density at 345(3)°C (~40×1021m-3) and
eight times the value at 590(5)°C (~13×1021m-3). The reduction in loop density with
increasing temperature occurs simultaneously with an increase in the average loop
dimensions (Fig. 3). These results manifest the predominance of loop growth over nucleation
as the irradiation temperature increases. The dislocation loops are determined to be of
interstitial nature based on the inside-outside technique (Fig. 4). Interstitial loop nucleation is
easier than vacancy loop nucleation, since interstitial loop nucleation is much less sensitive to
vacancy involvement than vacancy loop nucleation is to interstitial involvement
(Russell et al., 1973). The loop growth rate increases with temperature in irradiated, annealed
metals, i.e. at a higher temperature there are fewer loops but they are larger (Norris, 1972).
The escape probability for an interstitial close to a vacancy may increase with increasing
temperature (Urban et al., 1971), and small loops would therefore be unstable at higher
temperatures. Moreover, the attractive elastic interaction with dislocations is larger for
interstitials than for vacancies (Urban, 1970), so that larger loops would be able to further
increase in size by absorbing more interstitials at higher temperatures. Besides that, recent in
situ TEM studies revealed the small loop-loop interaction and coalescence in bcc W-based
materials at higher temperatures (Yi et al., 2015). In both cases, small loops disappear at high
temperatures in bcc metals in favour of larger loops.
The loops are characterised by the a/2 <111> Burgers vector at 180(2) and 345(3)C.
However, the dislocation analysis revealed that the loops present at 590(5)C correspond to
the a<100> Burgers vector (Fig. 5). Since the elastic energy associated with the dislocations
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is proportional to b2, where b denotes the length of the Burgers vector, a/2<111> type loops
would be favoured (Hull et al., 2011). Traditionally the formation of both types of perfect
dislocations is described by the shearing of a/2 <110> partial dislocations at an early stage of
growth, according to the classical dislocation reaction theory (Eyre et al., 1965):
a2
⟨110 ⟩+ a2
⟨001 ⟩ → a2
⟨111 ⟩
a2
⟨110 ⟩+ a2
⟨110 ⟩ → a2
⟨100 ⟩
The latter higher-energy shear only occurs at elevated temperatures. In this work we have not
observed the formation of a/2 <110> partial dislocations at an early stage of radiation-
induced growth at any of the studied irradiation temperatures. An in-situ TEM study from
literature on the structure evolution of bcc Fe during electron irradiation and heating revealed
the direct transformation of a/2<111> loops into a<100> loops, without the coalescence of
the a/2<111> loop with an external loop. The process was proposed to involve the nucleation
and propagation of a proper shear loop, triggered either by thermal fluctuations at high
temperature or by the presence in the vicinity of an external loop acting as a source of a
significant shear stress (Arawaka et al., 2006). Recently molecular dynamics simulations
pointed at the formation of a dumbbell configuration of SIAs, that evolves into a
configuration of parallel crowdions on nearest neighbour <111> lines when the number of
SIAs is 7, and subsequently into a/2<111> loops (Chen et al., 2013). Those a/2<111> loops
can directly evolve into <100> loops at high temperatures by the proposed mechanism of
gliding of <111> crowdions and jumping between different <111> directions, rotating into
<100> orientation, and gliding of segments of {100} loops along <100>.The estimated
energy barrier of a four SIA cluster jump from a/2[111](211) to a <100>{100}loop
configuration was estimated to be 1.2 eV (Chen et al., 2013). Furthermore, in the less likely
event that mobile a/2<111> loops with similar size and shape encounter themselves at high
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temperature, their interaction can lead to the formation of a<100> dislocation loops, as
predicted by kinetic Monte Carlo simulations for bcc metals under irradiation (Xu et al.,
2013). The resultant a<100> type dislocations would be glissile or sessile depending on the
slip planes of the reacting dislocations (Spitzig et al., 1966), but the shear stress for gliding
a<100> type dislocations is larger than for a/2<111> type dislocations (Hull et al., 2011).
Recent molecular dynamic simulations on bcc metals yielded a value of 2eV for the one-
dimensional migration energy of a<100> loops when the number of SIAs is 24 (Chen et al.,
2013). Furthermore, an increase in damage level from 0.15 to 0.55dpa causes a reduction in
the loop number density at the three irradiation temperatures, coupled with an increase in the
average loop dimensions; see Fig. 3 and Table 2. At 345C and 0.55dpa there is a relatively
high density of dislocation loops presented simultaneously with the developed dislocation
tangles. (Fig. 2). As the proton irradiation proceeds, the dislocation loops are in continuous
movement (Norris, 1972) scavenging additional SIAs from the tantalum matrix, and some of
those growing loops also coalesce.
Void formation
Fig. 6 shows the bright-field imaging of radiation-induced voids. The microstructure at
345(3)C and a damage level of 0.15dpa contains a high density of radiation-induced voids
that are randomly distributed in the matrix. However, at 0.55dpa voids are arranged in a bcc
lattice, oriented parallel to the underlying (111) bcc lattice plane of the Ta matrix and with an
average void distance of 7.3(2)nm. The average void diameter increases from 1.2(1)nm
(0.15dpa) to 2.2(1)nm (0.55dpa), and the number density from 2.3×1023m-3 (0.15dpa) to
3.1×1023m-3 (0.55dpa). The SIA trapping at mobile a/2<111> dislocation loops leaves a
vacancy excess in the matrix that evolves into voids. A void disorderorder transition occurs
with increasing damage level at 345(3)C. Existing models propose that non-aligned voids
shrink and collapse, and consequently void ordering develops. This process is based on the
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anisotropic diffusion of SIAs or small interstitial dislocation loops along closed packed
directions or planes (Jäger et al., 1993; Semenov et al., 2006; Semenov et al., 2008;
Barashev et al., 2010) or on the anisotropic energy transfer provided by long propagating
discrete breeders (Dubinko et al., 2009; Dubinko et al., 2011; Murzaev et al., 2015).
In contrast, the TEM data for tantalum irradiated at 180(2)C reveal the absence of voids in
the microstructure. At this lower temperature, interstitial dislocation loops are still formed
and generate a vacancy excess in the matrix. However, the vacancy mobility is significantly
reduced (Johnson , 1960; Satta et al., 1999) and, together with a number of vacancy-SIA
recombination events, does not lead to enough vacancy clustering so that voids, if present,
cannot be observed by TEM. Voids are also not observed in the sample irradiated at
590(5)C. At this temperature a<100> dislocation loops with a lower mobility predominate in
the microstructure. Dislocation mobility is needed to induce a local vacancy excess, and void
nucleation would take place in the vicinity of dislocations (Kitajima et al., 1979). The upper
temperature limit for void formation, which depends on particle fluence, is related to the
reduction in the supersaturation of vacancies due to an increased thermal-equilibrium
concentration of vacancies (Norris, 1972) and also, as observed in this study, to a change in
the type of predominant dislocation loops that act as scavengers for self-interstitial atoms in
the matrix.
The change in yield stress (∆ σ y ¿of the material caused by irradiation-induced defects, such
as dislocation loops and voids, can be derived using the dispersed barrier model
(Taylor, 1934; Seeger, 1958) :
∆ σ y=αMμb√ Nd (1)
where is the barrier strength coefficient, M the Taylor factor (2.7), the shear modulus of
tantalum (69GPa), b the Burgers vector of the dislocations (2.76Å), N the density and d the
average size of the obstacles (Seeger, 1958; Rosenberg (1971)). The value of depends on
9
the defect type and size for a given temperature (Hu et al., 2016). A value of = 0.2 has been
reported for dislocations in tantalum (Yasunaga et al., 2000), whereas it takes a value of
= 0.25 in the case of voids with a diameter of 1-2nm (Hu et al., 2016). The change in
hardness (∆ H ¿corresponds to:
∆ H=K ∆ σ y (2)
where the correlation factor K 3 and both ∆ H and ∆ σ y are given in Pa (Cahoon et al.,
1971; Busby et al., 2005). Using the values for N and d for the a/2 ⟨111⟩ interstitial
dislocation loops and for the voids observed at damage level of 0.15 dpa and at 350°C, we
obtain a value of ∆ H 0.6 GPa (loops) and ∆ H 0.7 GPa (voids) respectively.
Conclusions
Proton irradiation of tantalum causes the formation of interstitial-type dislocation loops
whose size increases with temperature and damage level, and concomitantly the loop number
density reduces. At 345(3)C the presence of mobile a/2<111> dislocation loops induced by
radiation results in a vacancy excess in the matrix that evolves into randomly distributed
voids at 0.15dpa. At the higher damage level of 0.55dpa, those voids are ordered into a bcc
lattice with an average void distance of 7.3(2)nm along the (111) plane. Voids are not
observed at either 180(2) or 590(5)C due to the low mobility of vacancies or a<100>
dislocation loops, respectively.
Acknowledgments
The work described was supported by the Dalton Cumbrian Facility Project, a joint facility of
The University of Manchester and the Nuclear Decommissioning Authority.
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Figure captions
Fig. 1. Simulated damage profile using the SRIM software, together with the nano-hardness
values measured along the cross section of the tantalum sample irradiated at 180(2)°C. The
asterisks indicate the regions of the sample from where TEM foils were extracted for
analysis.
Fig. 2. Evolution of dislocation structure in tantalum with damage level and temperature.
TEM data taken along the [111], [113], [012] zone axis.
Fig. 3. Average dislocation loop size and number density in tantalum as a function of
temperature at 0.15dpa and 0.55dpa.
Fig. 4. Determination of the nature of the dislocation loops by the inside-outside technique in
proton-irradiated tantalum at 590(5)°C. The loop show inside contrast in the g=121 reflection
and outside technique in the g=121 which belong to the [012] zone axis. The Burgers vector
of the loops was determined to be a<001> (see text).
Fig. 5. Bright-field TEM imaging of interstitial-type dislocation loops in tantalum proton
irradiated at: top: 180(2)°C; bottom – 590(5)°C, shown for different two-beam conditions.
These TEM data was used for the dislocation analysis using the g.b=0 invisibility criterion,
where g denotes the scattering vector and b the Burgers vector, and revealed the a/2<111>
(top – 180(2)C) and a<001> (bottom – 590(5)C) Burgers vector.
Fig. 6. TEM BF image showing the presence of radiation-induced voids in tantalum at
345(3)°C, and their absence at 180(2)°C and 590(5)°C.
16
Fig. 1. Simulated damage profile using the SRIM software, together with the nano-hardness values measured along the cross section of the tantalum sample irradiated at 180(2)°C. The asterisks indicate the regions of the sample from where TEM foils were extracted for analysis.
17
* *
Fig. 2. Evolution of dislocation structure in tantalum with damage level and temperature.
TEM data taken along the [111], [113], [012] zone axis.
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g110
g110
g110 g110 g200
g110 g110
g110 g110
Fig. 3. Average dislocation loop size and number density in tantalum as a function of temperature at 0.15dpa and 0.55dpa.
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Fig. 4. Determination of the nature of the dislocation loops by the inside-outside technique in proton-irradiated tantalum at 590(5)°C. The loop show inside contrast in the g=121 reflection and outside technique in the g=121 which belong to the [012] zone axis. The Burgers vector of the loops was determined to be a<001> (see text).
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Fig. 5. Bright-field TEM imaging of interstitial-type dislocation loops in tantalum proton irradiated at: top: 180(2)°C; bottom – 590(5)°C, shown for different two-beam conditions. These TEM data was used for the dislocation analysis using the g.b=0 invisibility criterion,
21
Temperature = 590(5)°C
Temperature = 180(2)°C
where g denotes the scattering vector and b the Burgers vector, and revealed the a/2<111> (top – 180(2)C) and a<100> (bottom – 590(5)C) Burgers vector.
Fig. 6. TEM BF image showing the presence of radiation-induced voids in tantalum at 345(3)°C, and their absence at 180(2)°C and 590(5)°C.
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Table 1. Summary of irradiation conditions used for tantalum at selected irradiation temperatures and radiation damage levels.
Irradiation temperature
(oC)
Fluence(1018 protons/cm2)
Flux(1018protons/cm2/h)
Damage level (dpa)
Damage rate(10-6 dpa/s)
Sampling depth(µm)
180(2) 4.4 0.20.15 2.2 20
0.55 9.3 29
345(3)0.9 0.9 0.15 8.9 28
3.7 0.8 0.55 5.9 29
590(5) 4.4 0.20.15 2.2 20
0.55 9.3 29
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