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: iuliia.ipatova@postgrad.manchester.ac.uk . 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

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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.

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* *

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,

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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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