Deuterium and helium ion irradiation of nanograined tungsten and tungsten–titanium alloys

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Deuterium and helium ion irradiation of nanograined tungsten and tungsten-titanium alloys L. Buzi , M. Yeh , Y-W. Yeh , O.K. Donaldson , M.I. Patino , J.R. Trelewicz , N. Yao , R. Doerner , B.E. Koel PII: DOI: Reference:

S2352-1791(19)30058-4 https://doi.org/10.1016/j.nme.2019.100713 NME 100713

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Nuclear Materials and Energy

Received date: Revised date: Accepted date:

16 June 2019 14 September 2019 1 November 2019

Please cite this article as: L. Buzi , M. Yeh , Y-W. Yeh , O.K. Donaldson , M.I. Patino , J.R. Trelewicz , N. Yao , R. Doerner , B.E. Koel , Deuterium and helium ion irradiation of nanograined tungsten and tungsten-titanium alloys, Nuclear Materials and Energy (2019), doi: https://doi.org/10.1016/j.nme.2019.100713

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Highlights: • Engineered W and W-20Ti alloy (W alloy with 20 at.% Ti) were exposed to D and He plasmas and gas retention as well as surface modifications were evaluated. • Nanograined W was observed to have superior performance (smaller D-induced blisters and less developed He-induced fuzz) compared to larger grain W.

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Deuterium and helium ion irradiation of nanograined tungsten and tungsten-titanium alloys L. Buzia1, M. Yeha, Y-W. Yehb, O.K. Donaldsonc, M.I. Patinod, J.R. Trelewiczc,e, N. Yaob, R. Doerner d, B.E. Koela a

Department of Chemical and Biological Engineering, Princeton University, Princeton NJ 08544 Princeton Institute for the Science and Technology of Materials (PRISM), Princeton University, Princeton NJ 08544 c Department of Materials Science and Chemical Engineering, Stony Brook University, Stony Brook, NY 11794 d Center for Energy Research, UCSD, La Jolla, CA. 92093-0417 e Institute for Advanced Computational Science, Stony Brook University, Stony Brook, NY 11794 1 Present address: IBM T.J. Watson Research Center, Yorktown Heights, NY 10598 b

Abstract: Tungsten (W), a primary candidate for the plasma-facing components of nuclear fusion reactors (e.g. the divertor region in ITER) is susceptible to cracks, blisters, bubbles, and other morphological changes when irradiated with energetic particles. This work investigated two new materials, nanograined W and a nanograined W-Ti alloy, for potential use as plasma-facing materials. Their retention properties and morphological changes after exposure to deuterium (D) and helium (He) plasma at 50 eV and surface temperatures of 500 and 1000 K were analyzed. Nanograined W was found to have smaller blisters and be less prone to fuzz formation than commonly-utilized micro-grain polycrystalline W. Additionally, the nanograined W-Ti alloy exhibited a lower concentration of blisters on its surface than pure W, including nanograined W. Keywords: Tungsten, Plasma-facing Materials, Deuterium, Helium

1. Introduction Tungsten (W) is one of the most promising candidates for use as a plasma-facing component (PFC) in nuclear fusion reactors due to its favorable thermal and mechanical properties, including high melting temperature, low erosion yield, and low fuel retention [1]. However, studies have shown that irradiation of W with high fluences of energetic deuterium (D) and helium (He) ions causes nanostructure [2] and blister [3] formation (among other morphological changes [4]). These changes may lead to degradation of material properties, impurity release, and plasma contamination [5]. Previous work has investigated the effect of D and He irradiation on W properties for several types of W and at various ion fluences, fluxes, energies, and surface temperatures [6,7]. In many cases of pure D irradiation, morphological changes of W occurred, the most common of which was the formation of blisters [6]. Blister formation was found to depend on the surface temperature (i.e. micrometer size blisters formed at Ts < 700 K) and grain orientation [8,9]. The size and shape of these blisters were seen to vary with the microstructure and dopant materials of the W [7]. Additionally, D retention was observed to have a direct correlation with blister size and density. For implantation of W with low-energy He (i.e. > 5-9 eV, fluence > 2×1024 m-2), nanometer-sized pores (bubbles) just below the sample surface (i.e. topmost 30 nm) were observed. At sample temperatures above 1000 K, bubbles migrate to the surface and result in surface morphological changes described as “fuzz” [10]. The grain orientation of the W was shown to influence

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the severity of the morphological changes presented [11]. He retention was demonstrated to be much lower than D retention at similar sample temperatures and ion fluences [12]. Studies have also shown that He retention has a direct correlation with grain size [13,14]. The increased grain boundary density of smaller-grain materials is thought to act as a sink for He ions, trapping He in the grain boundaries and then allowing it to diffuse back out of the sample. High temperature He ion irradiation showed both bubble and nanostructure formation in nanograined metals [13], whose severity depended on the grain size. Studies on the effects of D ion irradiation of nanograined metals showed mixed results. Two studies found that three times more D was trapped in nanograined W than commonly utilized large-grain polycrystalline W [15,16]. However, other studies that utilized the same equipment and similar operating conditions found the opposite trend [17]. Surface blistering was observed less frequently in nanograined W [15]. Similarly, no obvious pore or hole formations were observed in these cross-sections [15]. W alloys have also shown promise as potential plasma-facing components. Alloyed W was formed to create more grain boundaries, which was thought to reduce D(H) retention. For example, studies showed that a tungsten-tantalum (W-Ta) alloy retained less D [13] and experienced smaller and less frequent blistering from D irradiation than pure W [18]. Additionally, alloying can improve the thermo-mechanical stability of W [9][19], including for nanocrystalline materials [20] that may be prone to thermodynamic stability challenges [21–23]. The work reported herein investigated the potential of engineered, nanograined W and W alloys for use as plasma-facing components in fusion devices. In particular, a W-20Ti alloy (W alloy with 20 at.% Ti) has been identified as a promising material for the stabilization of nanocrystalline grains [24]. Given the ability to retain its nanostructure at elevated temperature rather than undergoing rapid microstructural coarsening common in W [25], W-20Ti has the potential for improving the resistance of W to particleinduced surface and microstructural modifications. In this study, we tested this prediction by exposing nanograined W-20Ti samples at 500 and 1000 K to D and He plasma. Morphological changes and retention properties were measured. In addition, measurements were made for nanograined W samples exposed at identical conditions (i.e. 500 and 1000 K to D and He plasma) for comparison. Section II details the experimental approach, and results from experiments using both materials are presented in section III.

2. Experiments Thin films ~100-nm thick of W and W20Ti were sputter deposited onto bulk molybdenum (Mo) substrates in an AJA ATC-1500 stainless-steel chamber at the University of Alabama, Tuscaloosa using >99.95% pure elemental targets obtained from the Kurt J. Lesker company. The 1×10×10 mm Mo substrates were outgassed at 1270 K for 30 min before the films were deposited and outgassed again at 870 K for 30 min after film deposition. W20Ti films were annealed ex-situ using a ramp rate of approximately 100°C/5min, isothermally held at 1000 °C for 1 hour and then quenched using a Lindberg furnace (Model #54434) equipped with an Omega controller. The unalloyed tungsten films were instead annealed in situ to match the annealing temperature of the alloyed films but avoid drastic coarsening of the nanocrystalline grain structure. Using a Gatan heating holder (Model #628) in a Philips CM-30 300kV TEM, the temperature was manually ramped up to 1000 °C where the films were held for 20 minutes and followed by open-loop cooling to room temperature. Microstructure and phase analysis was accomplished using bright-field imaging and indexing of the selected area diffraction (SAD) patterns coupled with precession electron diffraction (PED) mapping using the NanoMEGAS ASTAR system. Phase identification was completed using the method described by Edington [25] to compare the ring spacing ratios with the reciprocal of the known d-spacings. PED scans were performed at 0° tilt and a reduced magnification to capture full-field grain structures for statistical analysis as a function of total dose. Scans

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were cleaned through a removal of ambiguities using the procedure described in Ref. [25] with a zeroorder extrapolation and noise reduction limited to removal of anomalous spikes. Additional characterization of the samples before plasma exposure was accomplished at Princeton University using TEM, scanning TEM (STEM), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDX). The samples were also weighed using a microbalance and their mass was recorded for plasma erosion calculations. Plasma exposures of the samples were performed on the PISCES-A linear plasma device at the University of California in San Diego [26]. A D or He plasma column was generated from an arc discharge source in a 150 cm long chamber, confined by magnetic field coils (0.07-0.14 T), and directed at the sample holder. A downstream Langmuir probe, 126 cm away from the cathode, was used to characterize the plasma in the vicinity of the sample. Electron temperature and density for this plasma system are in the range ≤20 eV and 5×1010 ∼ 1013cm3, respectively. The plasma exposure conditions for the current experiments are outlined in Table 1 and are similar to those found at the divertor region of fusion reactors. For the data in Table 1, the flux was varied to obtain similar fluence values for the high and low temperature exposures. The samples are identified in Table 1 as WBx-y or W20TiBx-y, where Bx-y identify the batch number and sample id. The sample holder held two samples for each plasma exposure - one W sample and one W20Ti sample - such that each pair of samples was exposed to the exact same plasma conditions. The holder was made from Ta with a Mo cap. Grafoil was used between the sample and the air-cooled holder to ensure good thermal contact and allow for variability in the exposure temperature. A thermocouple mounted on the sample holder was used to measure the temperature at the back of the sample. The temperature uncertainty was ±45 K. The energy of the incident ions (~50 eV) was set by biasing the sample negative with respect to the plasma (-60 V) and calculated from the empirical formula Eion  =  Ubias  −  2kTe [27]. Following plasma exposure, the samples were re-characterized using Auger electron spectroscopy (AES), imaging techniques, and thermal desorption spectroscopy (TDS). AES was used to determine the elemental composition of the surface (topmost 5 nm) of the sample and revealed carbon, oxygen and nitrogen concentrations of 13at.%, 18at% and 5at%, respectively. Surface and cross-sectional images of the samples were taken using SEM, TEM, and STEM. TDS was used to monitor species desorbed from the samples while the samples were heated linearly from room temperature to 1273 K at 0.5 K/s. During TDS, the sample was placed adjacent to a thermocouple inside a quartz tube and heated by flash lamps. MKS Microvision II (MKS) and Stanford Research System (SRS) mass spectrometers recorded signals from species with masses between 1.97 - 2.0288, 2.99 - 3.0288, and 3.97 - 4.0488 amu corresponding to H2, HD, and D2/He, respectively.

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Table 1: Plasma exposure conditions Samples

Plasma D D D D He He He He D D

Temperature (K) 1000 1000 500 500 1000 1000 500 500 1000 1000

Flux (1022 ions / m2 s) 0.90 1.0 0.15 0.20 1.0 0.72 0.19 0.20 0.90 1.0

Fluence (1025 ions / m2) 10 10 1.6 2.2 11 7.8 2.1 2.2 10 10

WB1-1 WB1-2 WB1-3 WB2-2 WB2-3 WB2-4 WB2-5 WB2-1 W20TiB1-1 W20TiB1-2 W20TiB2-1

D

500

0.15

1.6

W20TiB2-2

D

500

0.20

2.2

W20TiB2-3

He

1000

1.0

11

W20TiB2-4

He

1000

0.72

7.8

W20TiB2-5

He

500

0.19

2.1

W20TiB2-6

He

500

0.20

2.2

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3. Results and Discussion

The W and W20Ti samples were engineered to contain grains with a size of less than 100 nm through magnetron sputtering and post-deposition annealing at 1000 °C. In the pristine condition, the W films exhibited the nanocrystalline grain structure shown in the bright-field image in Figure 1a with a random polycrystalline texture from the PED map in Figure 1b. Indexing of the SAD pattern revealed the films were composed entirely of body-centered cubic (BCC) α-tungsten with no preferential texturing, consistent with the orientation analysis from the PED results. The grain structure of the W20Ti films shown in the bright-field image in Figure 1c was generally finer than the W films but retained the random polycrystalline texture from the orientation map in Figure 1d. Reflections specific to the hexagonal closepacked (HCP) Ti phase were evident in the SAD pattern, indicating the presence of a Ti-rich second phase in the alloy film and consistent with prior results on annealed W-Ti thin films [25]. Grain size distributions were determined directly from the PED maps using crystallographic orientation to delineate the various grain geometries. Plotted as cumulative area fractions in Figure 1e, the mean grain size, , for W20Ti was = 41 nm, which was marginally smaller than the grain size for W of = 71 nm. Despite this subtle difference in grain size, both structures were well within the nanocrystalline regime with only a small fraction of grains in the range of 125 – 150 nm, and thus provide a good basis for comparison on W and W20Ti with comparable microstructures.

Figure 1. Bright-field image and selected area diffraction pattern for nanograined (a) W and (c) W20Ti with corresponding PED grain orientation maps in (b) and (d), respectively. Grain size distributions from the PED maps with mean grain size, , included in the legend for the the W and W20Ti films. Both distributions were within the nanocrystalline regime with only a small fraction of larger grains between 125 – 150 nm.

For each material and plasma exposure condition (i.e., D or He, high or low temperature) two sets of samples were developed. One set of samples was analyzed utilizing AES, SEM, and TEM to observe composition and morphological changes. The other set of samples was analyzed immediately after plasma exposure by TDS to obtain D and He retention data. The mass loss for a few of the samples after plasma exposure (and before TDS) was also recorded.

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AES spectra probing the surface elemental composition of two W samples exposed to He plasma were taken ex-situ. The low and high temperature exposures resulted in minimal differences in the AES spectra: 64% W, 18% O, 13% C, and 5% N. The presence of W peaks (and lack of Mo peaks) proved that the W film was not completely sputtered or delaminated during plasma exposure. The O, C, and N impurities are likely due to the interaction of the sample with air during post-exposure handling. Surface morphology changes were observed by using SEM (Figure 2). At the low sample temperature, D plasma exposure caused blistering in both W (Figure 2c) and W20Ti (Figure 2d). The blister number density is higher in W than in W20Ti. However, the blisters present on the W20Ti sample were slightly larger than those on the pure W sample. The pure W sample had blisters of 0.1 - 0.4 m in diameter, while the blisters on the W20Ti sample ranged from 0.2 - 2 m in diameter. For polycrystalline samples with larger grains, low-energy, high-flux D exposures induced surface blisters 0.1 - 2 m in diameter and near-surface blisters 30 nm in diameter [28]. Therefore, in comparison to large-grain W, nanograined W had smaller blisters, while nanograined W20Ti had similar sized blisters to the nanograined W.

Figure 2. Top-down SEM images of the samples after plasma exposure (SE detector). On the images are identified the plasma ions (D or He) and the exposure temperature (500 K or 1000 K).

Additionally, the images in Figure 2 show that both W and W20Ti samples manifest step-like features when exposed to D plasma at high temperatures and He plasma at both low and high temperatures. In all cases, the erosion and its severity varied significantly throughout the sample surface. Each change in erosion pattern likely corresponds to a new grain in the sample, with the direction and severity of the erosion corresponding to the direction of the grain orientation. Tilted SEM images at 45 and TEM analysis indicated that the observed structures were a few nanometers deep. For W and W20Ti samples exposed at high temperature to He plasma, small bubbles were observed throughout the new nanostructures (Figure 2e and 2f), as shown in Figures 2e and 2f, but with no significant differences observed between the W and the W20Ti samples. The bubbles are the beginning of “fuzz,” defined as fibers 20-50 nm in diameter [22]. These plasma-induced features are consistent with previous studies on the initial phase of fuzz growth [29]. At 1175 K and 1×1026 ions/m2, El-Atwani et al. found that while large-grain W samples showed full fuzz formation, nanograined W only displayed prefuzz growth similar to that shown in Figure 2e and 1f [22]. Additionally, large-grain W at 1175 K first exhibited fuzz at a fluence of 1×1025 ions/m2 [22]. This is several orders of magnitude below the fluences used here that only show pre-fuzz formation. Therefore, nanograined W (pure and alloyed) exhibits a higher resistance to surface morphology changes than large-grain polycrystalline W.

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Figure 3 Cross-sectional TEM images of the samples after plasma exposure. The scale bar indicated in the (b) image applies to all images.

Cross-sectional samples prepared by using a focused ion beam (FIB) were used to investigate morphological changes to the bulk of the samples utilizing TEM/STEM. Some results are illustrated in Figure 3. After D plasma exposure, Figure 3(a-d) show a similar homogenous internal W structure. However, the greyscale stratification was no longer evident, and the surface was no longer smooth. Instead, a sawtooth-like surface structure was present for pure W samples, with variations in the surface height ranging up to 12.5 nm. SEM scanning showed good center to edge consistency of these features so this is most likely a result of preferential sputtering along different grain orientations [30]. Additionally, Figure 3b shows small cracks beneath the erosion structures for the W20Ti sample exposed to D plasma at 1000 K. Cracks are also evident on W and W20Ti samples exposed to He plasma at low temperature. In Figure 3(e-f), the lighter grey layer appearing above the nanostructure formation are nanotendrils indicative of pre-fuzz formation. Additionally, small bubbles were seen in the W layer, consistent with previous studies [31]. No significant differences were found between W and W20Ti samples exposed to He at high or low temperatures.

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Figure 4 TDS spectra of (a) D and (b) He desorbed from nanograined W and nanograined W-20Ti.

Table 2: Total D and He retention from TDS analysis. D/m2 [x1019]

He/m2 [x1019]

WB2-2 (D/500K) W20TiB2-2 (D/500K)

128 18.4

WB2-1 (He/500K) W20TiB2-6 (He/500K)

6.48 8.05

WB1-2 (D/1000K) W20TiB1-2 (D/1000K)

2.99 6.15

WB2-4 (He/1000K) W20TiB2-4 (He/1000K)

11.9 13.9

The total amount of D and He retention was studied using TDS (Figure 4). Samples exposed to D at low temperature had peaks at 630 K (WB2-2/D/500K) and 610 K (W20TiB2-2/D/500K), consistent with previous findings that showed low temperature peaks for exposure temperatures below 700 K. Samples exposed at high temperature had D release peaks at 760 and 860 K (W20TiB1-2/D/1000K) and one broad D peak at 1000 K (WB1-2/D/1000K), with D release peaks at higher exposure temperature being one to two orders of magnitude lower [32]. Nanograined W20Ti showed lower D retention than nanograined W at the low exposure temperature, but slightly higher D retention than nanograined W at the higher exposure temperature (see

Table 2: Total D and He retention from TDS analysis. ). The total D retention in W and W20Ti at 500 K was within the spread of previous results for larger grain W at 500 K (30-200 eV, fluence >10 25 m-2), while total D retention values at the higher temperature of 1000 K were 10-20 times higher than larger grain samples [33], which is probably due to increased diffusion of implanted D through the grain boundaries. Kolasinski et al. performed similar experiments at PISCES-A with ultra-fine grain W bulk samples at temperatures 470 - 770 K and found the total D retention to be 3 times higher than conventional polycrystalline W [15]. These differences are likely due to the finer grain size and columnar nanostructure of the films providing rapid diffusivity paths through the grain boundaries for desorption of D at the film surface. TDS results for the He plasma-exposed samples show scattered He peaks from internal bubbles bursting and He accumulation at grain boundaries [34]. General trends in Figure 4 include characteristic low and high temperature He desorption peaks, with no distinct differences between W and W20Ti samples. However, there is a desorption peak coinciding at ~650 K similarly for W and W20Ti samples (WB2-1 (He/500K) and W20TiB2-4 (He/1000K)). While the exact mechanism responsible for this peak is

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not entirely understood, we note that exposure temperature plays a critical role in the size and morphology of the He defects formed in nanograined tungsten [35], which will inevitably influence the desorption behavior and likely responsible for the differences in the TDS results between the 500 and 1000 K exposures. Additional work is needed to correlate anomalous TDS peaks with the underlying defect microstructure. TDS of He desorbed from powder-metallurgical W samples irradiated at 1.11×1026 m-2 and 933 K showed both high and low temperature peaks near the same temperatures shown in Figure 4 [36]. The total He retention in both W and W20Ti samples increased slightly at the higher temperature, similar to He retention in large-grain W. However, the retention values shown here are 2-4 times larger than in large-grain W [26], which may be attributed to the additional grain boundaries in the nano-grain samples. SEM analysis was also performed after TDS measurements. These results showed a lack of blisters on the lower temperature D plasma-exposed samples where blisters had previously been observed [37]. Similarly, the nanostructure and pre-fuzz formations that were formed during He plasma exposures on samples at the higher temperature were no longer present. These changes were likely the result of blisters bursting and nanostructure formation annealing at the high temperatures reached during TDS analysis. The masses of four samples were measured before and after He plasma exposures and used to calculate the mass loss due to He ion sputtering. Since the ion energy during plasma exposure was ~50 eV, below the threshold for W sputtering by He+, the sputtering measured was most likely a result of a small fraction of O, C, or N impurity ions in the plasma. The sputtering yields for the four samples are shown in Table 3. From Table 3, the sputtering yields of nanograined W20Ti is lower than that of nanograined W at 500 K but higher than that of nanograined W at 1000 K. Table 3: Sputtering yield of W and W20Ti nanograined samples exposed to He plasma with E(He+)~50 eV WB2-1 (500 K) Sputtering Yield (W atoms / He ion)

3.09×10-4

W20TiB2-6 (500 K) 1.96×10-4

WB2-3 (1000 K) 1.80×10-5

W20TiB2-3 (1000 K) 4.75×10-5

4. Conclusions Samples of W and a W-Ti alloy (20 at.% Ti) with nano-sized grains were exposed at 500 and 1000 K to He and D plasmas (~50 eV, 1.6-11×1025 m-2). Nanograined W was observed to have superior performance (smaller D-induced blisters and less developed He-induced fuzz) compared to larger grain W. This is attributed to the increased density of grain boundaries in nanograined W that form a pathway for gas diffusion out of the sample. In addition, it was found that nanograined W-Ti had fewer D-induced blisters than nanograined W. Further research involving bulk samples rather than thin films and a broader range of plasma irradiation conditions is needed to more fully assess the capabilities of nanograined W and nanograined W-Ti alloys as plasma-facing materials.

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5. Acknowledgements Support for this work was provided for ODK and JRT at Stony Brook University through Department of Energy Grant DE-SC0017899. The authors gratefully acknowledge Dr. Tyler Kaub and Prof. Gregory Thompson at the University of Alabama Tuscaloosa for supplying the magnetron sputtered thin films. OKD and JRT also thank their collaborators at Sandia National Laboratories, Dr. Khalid Hattar and Dr. Christopher Barr, for assistance with the PED measurements as well as valuable insights and feedback. This work was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science by Los Alamos National Laboratory (Contract 89233218CNA000001) and Sandia National Laboratories (Contract DENA-0003525).

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