Measurement of Heavy Ion Irradiation Induced In-Plane Strain in

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Measurement of heavy ion irradiation induced in-plane strain in patterned face-centered-cubic metal films: an in situ study Kaiyuan Yu, Youxing Chen, Jin Li, Yue Liu, Haiyan Wang, Marquis A Kirk, Meimei Li, and Xinghang Zhang Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.6b03195 • Publication Date (Web): 28 Nov 2016 Downloaded from http://pubs.acs.org on November 30, 2016

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Measurement of heavy ion irradiation induced in-plane strain in patterned face-centered-cubic metal films: an in situ study K. Y. Yua*, Y. Chenb,c, J. Lib, Y. Liub, H. Wangd,e, M. A. Kirkf, M. Lif, X. Zhangd* a

State Key Laboratory of Heavy Oil Processing and Department of Materials Science and

Engineering, China University of Petroleum-Beijing, Beijing 102249, China b

Department of Materials Science and Engineering, Texas A&M University, College Station, TX

77843-3123, United States c

d

e

f

MPA-CINT, Los Alamos National Laboratory, Los Alamos, NM 87545, USA School of Materials Engineering, Purdue University, West Lafayette, IN, 47907 United States Department of Electrical Engineering, Purdue University, West Lafayette, IN, 47907 United States Nuclear Engineering Division, Argonne National Laboratory, Argonne, IL 60439, United States

*Corresponding authors: K. Y. Yu, [email protected]; X. Zhang, [email protected]

Abstract Nanocrystalline Ag, Cu and Ni thin films and their coarse grained counterparts are patterned using focused ion beam and then irradiated by Kr ions within an electron microscope at room temperature. Irradiation induced in-plane strain of the films is measured by tracking the location of nanosized holes. The magnitude of the strain in all specimens is linearly dose-dependent and the strain rates of nanocrystalline metals are significantly greater as compared to that of the coarse grained metals. Real-time microscopic observation suggests that substantial grain boundary migration and grain rotation are responsible for the significant in-plane strain. Key words: in situ irradiation, patterned films, in-plane strain, grain activities

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Irradiation produces abundant point defects and defect clusters in metallic materials and causes significant dimensional variations such as void swelling [1-3]. More than 5% swelling in materials may result in catastrophic failure of materials in the form of embrittlement [1]. Thus the control and quantification of irradiation-induced swelling is essential for the safety evaluation of materials in nuclear power plants. Swelling can be minimized by controlling irradiation induced defects. Singh and Foreman suggested decades ago that metals with smaller grains exhibit lower defect density and radiation induced hardening [4, 5]. Moreover, grain boundary denuded zone has later been observed in numerous metals [5-8] where high-angle grain boundaries (HAGBs) act as defect sinks to capture and eliminate radiation induced defect clusters. Reducing grain size is thereby an effective strategy for the enhancement of void swelling resistance. Recent studies show that nanocrystalline (NC) and ultra-fine grained (UFG) materials have significantly less swelling than their coarse-grained counterparts [9-11]. Meanwhile, other types of defect sinks have also been investigated, including dissimilar interfaces, twin boundaries and free surfaces [12-19]. These interfaces not only store inert gas, like helium (He), but also enable frequent recombination of interstitials and vacancies. For instance outstanding radiation tolerance (in terms of substantial reduction of void swelling) has been demonstrated in oxide dispersion strengthened (ODS) alloys [20] and nanostructured ferritic alloys (NFA) [21] with abundant metal/ceramic heterophase interfaces. He ion irradiation induced hardening and helium bubble density in nanoscale multilayers with metal/metal heterophase interfaces decreases significantly with the reduction of individual layer thickness [12, 22-25]. Such a strategy has also been proved valid in bulk nanolayered metals prepared by using accumulative roll 2


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bonding (ARB) method [24]. Moreover, the ‘sink’ effect of boundaries and interfaces has been captured using in situ irradiation technique. Sun et al. has employed in situ irradiation experiments to observe rapid migration of dislocation loops toward HAGBs in NC Ni [26], and similar phenomena have been observed in NC Fe [27, 28] and multilayered Ag/Ni [29]. More recently, in situ radiation studies have shown that twin boundaries can significantly reduce the density of stacking fault tetrahedra (SFTs) in nanotwinned metals [15], and irradiation induced fast twin boundary migration has been observed [30, 31]. In addition, free surfaces have been proposed to be another effective type of defect sinks as shown by ex situ and in situ irradiation studies on nanoporous Au and Ag [18, 19]. Despite all these achievements in controlling swelling by using high density engineered interfaces, the quantification of swelling of irradiated metallic materials is primarily performed ex situ. Conventionally, the swelling is calculated by measuring the areal ratio of voids and bubbles under microscopes [32, 33], while alternative methods include probing the strain using the glancing or incidence X-ray diffraction [34, 35] as well as a few in situ techniques [36]. However, resolvable voids and bubbles require high dose irradiation, which is costly and time-consuming. Therefore, it is necessary to develop new methods to estimate the irradiation induced dimensional changes of irradiated materials in a fast and convenient manner. Here we report a simple new method to quantify the in-plane strain induced by in situ heavy ion irradiation in face-centered-cubic (FCC) metals. A series of metal films (Ag, Cu and Ni) are synthesized by sputtering followed by patterning with an array of nanosized holes by using focused ion beam (FIB) milling technique. The holes serve as markers (reference) to track the strain during in situ irradiation. Our measurement demonstrates 3


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apparent in-plane strain and NC specimens expand faster than their coarse-grained (CG) counterparts. This study provides a simple yet significant methodology that links irradiation induced microstructural evolution to macroscopic strain and may shed light on the development of a universal technique for the determination of radiation induced strain in materials in general. Fig.1 shows the microstructure of a patterned NC Ag film. Fig.1A and B are the scanning electron microscopy (SEM) images showing the thin film deposited on Cu TEM mesh and the morphology of the FIB-milled arrays of holes. The bright-field and dark-field TEM micrographs of the film in Fig.1C show the microstructure of equiaxed nanograins and arrays of FIB patterned holes. These holes have similar geometry and pre-designed inter-separation distance. The inserted selected area diffraction (SAD) pattern shows continuous rings arising from the polycrystalline nature in NC Ag. A magnified TEM micrograph of a hole outlined in Fig.1D and the inset HRTEM image assure that FIB milling did not cause apparent damage to the TEM specimen. The amorphous rim near the inner surface of the hole arises from the amorphous carbon film on the TEM grid. The microstructure and statistics of grain size for the six TEM specimens are displayed in Fig.2. The average grain size of NC Ag and CG Ag as shown in Fig.2A1 and A2 is ~ 100 nm and 1300 nm, respectively. NC Cu and CG Cu shown in Fig.2B1 and B2 have an average grain sizes of 70 and 900 nm, respectively, while NC Ni and CG Ni (Fig.2C1 and C2) have a respective average grain size of 23 and 700 nm. Note that NC Cu has a bi-model grain size distribution: a majority of the grains had an average diameter of ~ 70 nm while a handful of larger grains were also present. 4


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These NC and CG metal films were then irradiated by Kr ions using the in situ ion irradiation

facility

in

a

TEM

microscope

up

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dose

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displacements-per-atom (dpa) at room temperature and at a constant dose rate. The morphological evolution of the irradiated films was recorded with a digital video camera, which was later used for extensive post irradiation strain analyses. An example of irradiation-induced strain measurement for NC Cu at different doses is shown in Fig.3. The strain measurement procedures are established as follows. Numerous patterned holes in the irradiated specimen were outlined as references for the strain measurement. A central point, which was usually located around the center of the patterned area, was then chosen as a set-point and marked by a red cross, as shown in Fig.3A. Then multiple points of interest surrounding the set-point (as shown by the blue crosses in Fig.3A) were selected to track strain evolution during irradiation. Thus the separation distances between the red and blue crosses at 0 dpa are considered as initial gauge length Li (Supplementary Fig.S1). During in situ irradiation experiments, the separation distances along different directions increased and the variation is denoted by ∆Li. The strain induced during radiation is calculated by taking average of in situ measured values of ∆Li /Li along various directions as shown in Fig.3B and 3C. The area of the holes barely changed during the irradiation experiments. More details on the measurement of strain in other irradiated FCC metal films are provided in supplementary materials (Supplementary Fig. S2). Fig.4 compares the irradiation induced in-plane strain for all TEM specimens. As shown in Fig.4A1-A3, the strain for all irradiated metal films increased linearly with increasing irradiation dose and did not exhibit saturation up to 2 dpa. For all the three irradiated FCC 5


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metals, the NC specimens showed evidently greater in-plane strain than their CG counterparts at all doses. Meanwhile, the strain rates (defined as strain per dpa) of NC metals were 2 times greater than those of the CG counterparts as shown in Fig.4B. Among all NC specimens, the overall in-plane strain and the strain rate of NC Ag was the greatest (4%, 2.9%/dpa), followed by Cu (2.0%, 1.1%/dpa), and NC Ni (1.4%, 0.6%/dpa). Origin of the in-plane strain in irradiated CG and NC metals. It is known that significant swelling of most metals occurs at elevated temperatures (ex. 200-500℃ for pure Cu, fission neutron irradiation) [7]. However, we have observed substantial void/bubble-free in-plane strain in Kr ion irradiated CG and NC metals at room temperature and the strain rates are essentially higher than other bulk or thin foil FCC metals irradiated by neutrons, heavy ions and in situ high energy electrons [32, 33, 37-44]. The following factors may contribute to the observation of significant strain in the current in situ studies and help to explain the discrepancy. First, the fact that the measured strain in CG specimen is greater than the void swelling measured from irradiated bulk specimens could be a consequence of TEM thin foil effect. The free surface of TEM thin foils may accelerate the annihilation of interstitial loops and leave more invisible vacancy clusters behind. Hence the void swelling or radiation damage in TEM thin foils (due to vacancy clusters) could be greater than that in bulk specimens. Second, such a discrepancy is probably due to the difference of measurement methodology. The traditional methods for swelling measurement often count on measurement of visible voids (void-based measurement) by TEM techniques. For thin foil metals irradiated by high energy electrons within HVEM, most swelling data were typically acquired above 5-10 dpa when the nanovoids were visible (sufficiently large) for effective TEM 6


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measurement, while swelling data below 2 dpa was lacking due to the difficulty in detecting voids[32, 37-40, 42-45]. However, there are studies using positron annihilation technique showing that the density of invisible vacancy clusters in irradiated metals could be two orders of magnitude greater than that of visible ones [46-49]. Hence it is likely that the current measurement technique may be more sensitive to gauge radiation induced invisible defect clusters in CG metals than classical TEM based void measurement technique. Third, for CG metals, the in-plane strain might come from the high density irradiation induced defect clusters. As shown in Supplementary Fig. S3, the defect density in irradiated CG Cu at 2 dpa was measured to be 8.1×1023m-3. It is known that the major type of defects in heavy ion irradiated FCC metals is SFT. A high density of SFTs could contribute to the increase of in-plane strain in the CG metals as they are most of vacancy-type in nature. In comparison, the greater magnitude of strain observed in NC metals is somewhat counterintuitive. This is because a classical rationale often predicts that the magnitude of swelling is NC metals is less than in CG metals as the abundant defect sinks in NC metals could absorb and eliminate much more defects than in CG counterparts. Hence the discrepancy between the current experiments and general anticipation could indicate that other mechanisms are playing a major role in the determination of the strain, which will be discussed in the following section. Controversy may also rise on the uniformity of the strain. During irradiation, the stress was not concentrated at the edges of the holes (no sign of contrast change or crack initiation), indicating measured strain was likely to be uniform. Moreover, there was little variation in the dimension of the holes (Supplementary Fig. S4), suggesting the measured strain stemmed primarily from the materials between the holes. 7


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Grain size dependent strain in irradiated NC and CG metals. It appears surprising that the accumulation of irradiation induced strain in NC metals is much greater and faster than their CG counterparts. This is somewhat counterintuitive as one would anticipate that HAGBs are effective defect sinks for defect clusters, and consequently NC metals should have much less strain and strain rates than CG counterparts. As defect clusters were barely observed during in situ radiation of NC metals, we hypothesize that the irradiation induced large strain in NC metals arised from other mechanisms. Detailed analyses on microstructural evolution during in situ irradiation reveal two possible mechanisms. First, we observed substantial irradiation induced grain boundary migration (Fig.5). It is well known that during deformation, grain boundary migration can accommodate plastic strain, and the magnitude of grain boundary migration is typically much greater in metallic materials with smaller grain sizes. Such grain boundary assisted plastic deformation has been widely observed during in situ deformation of NC metals [50, 51]. In the current study, our NC metals are free-standing thin films and hence it is less likely that there will be large stress to trigger substantial strain during irradiation. However, as shown in Fig.5, the outlined grain boundary in NC Ag migrated over a distance of 32 nm during radiation from 0.375 to 1 dpa, at a speed of ~ 0.05nm/s (See Supplementary Fig.S5 and video for details). In mechanically deformed NC metals, it has been argued that during deformation, enhanced diffusion along grain boundaries may be responsible for the stress induced grain boundary migration [52] and at higher temperatures grain boundary sliding and diffusion tend to dominate the plastic deformation mechanism of NC metals [53-55]. The mechanisms of irradiation induced grain boundary migration have also been investigated by others. Simulation results of NC Ni have 8


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shown that ion beam induced grain growth is a direct result of recrystallizaton of the thermal spike whose volume overlaps with the grain boundaries, and hence the growth mechanism is inherently different from that during annealing [56]. In this study, the grain boundary migration was likely to arise from the large thermal spike volume at the grain boundaries since the ion energy (1MeV) and flux (10-3 dpa/s) were high and the temperature increase during irradiation (4K) was negligible. Second, in situ studies show frequent radiation induced grain rotation or recrystallization. As shown in Fig.6, the two small equiaxed grains in NC Ag (both with diameters of ~ 20 nm) indicated by arrows rotated prominently (as indicated by gradual loss of contrast) from 1.19 to 1.56 dpa. Grain rotation has been suggested as an important plastic deformation mechanism and widely observed in various NC materials during plastic flow, such as Cu [57], Ni [58, 59], and TiN [60]. Ion irradiation induced grain recrystallization and rotation has also been reported. Nita and coworkers irradiated NC Ni using 590 MeV protons to 0.56 dpa and observed grain refinement attributed to irradiation induced recrystallization [61]. One possible mechanism is that defect clusters might migrate to subgrain boundaries and form a cell structure, which later evolves into new smaller grains [61]. Another possible mechanism is that if the cascade region is larger than the grain diameter, a stacking fault can form across the grain, breaking the grain into two smaller ones [62]. Similar recrystallization processes were also explored in NC Ag [63], Cu alloys [64] and simulated in NC Ni [56]. More recently, Zizak and coworkers reported grain rotation in NC Ti irradiated by 350 MeV Au ions [65, 66]. It has been proposed that the thin layers of grain boundaries were sheared during the ion bombardment and led to the slide between adjacent grains. Besides NC Ag, we have also 9


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observed irradiation induced grain rotation in NC Ni during in situ radiation studies (Supplementary Fig.S6). Such grain activities have not been observed in the CG metals during irradiation, and thus CG metals exhibit smaller strain and strain accumulation rate. From a statistical viewpoint, we have compared the intensity of (111) and (200) rings in the SAD patterns of NC Ag and NC Ni to quantify the grain rotation. The SAD patterns of NC Ag and NC Ni were respectively taken at the same temperature, electron beam current, ion beam condition and focus condition at different doses. We integrated the intensity of Ag (111) and (200) rings from 0 to 360 degree (azimuthal angle) to measure the overall intensity of the rings, generating a plot of intensity as a function of d-spacing (Supplementary Fig. S7). The measured data are indicated by the red dots and fitted using Gauss function (blue solid curves). As seen in Supplementary Fig. S7a and b, after radiation of NC Ag to 1.56 dpa, the intensity difference between Ag (111) and Ag (200) peak becomes much greater. The intensity ratio between the (200) and (111) peaks (I200/I111) decreases from 0.4 (which is very close to the random polycrystalline Ag without texture) to 0.2. As the SAD pattern is recorded for the plan-view TEM micrograph, the diffraction ring of any index derives from atomic planes oriented nearly parallel to the electron beam. Hence this large variation of peak intensity ratio suggests that a significant amount of (200) planes in nanograins have rotated (or replaced) by (111) oriented planes in the grains during irradiation as shown schematically in Supplementary Fig. S8. In short, the film became more textured and the average in-plane d-spacing increased, leading to positive in-plane strain. Since the atoms in (111) planes are more close-packed than those in the (200) planes, such rotations reduced the grain boundary energy of the nanomaterials and were thereby energetically favorable [67, 68]. 10


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Strain accumulation rate difference among NC metals. In FCC metals, a large number of defects induced by ion irradiation are SFTs, which might affect the local stress within the films due to their interactions with other defects. It has been shown that the density of SFTs is closely correlated to SFE [69]. Lower SFE typically leads to higher density of SFTs in irradiated FCC metals. Since the SFEs of the metals in this study are 19 mJ/m2 for Ag, 41 mJ/m2 for Cu and 125 mJ/m2 for Ni, it is likely that the greater strain rate of Ag is due to its much greater SFT density. Another possibility is that the minimum displacement energy is different for the three metals (14.5 eV for Ag [70, 71], 19 eV for Cu and 23 eV for Ni [72]), hence more displacement damage could occur in Ag during irradiation than in Cu and followed by Ni, resulting in larger defect density. Furthermore, the difference in the amount of rotated grains in different materials might also lead to distinct strain rates. Comparisons among the measured SAD patterns show that the intensity ratio change of Ag(200)/Ag(111) is greater than that of Ni(200)/Ni(111) after irradiation, as shown in Supplementary Table S1, indicating that more Ag grains rotated during irradiation. Therefore irradiation induced in-plane strain of NC Ag was higher than that of NC Ni, which is consistent with our strain measurements. In summary, we have reported a new method to measure the heavy ion irradiation induced in-plane strain of metal films by patterning the specimens with arrays of holes. The magnitude of strain in all specimens increases linearly with increasing irradiation dose. NC metals exhibit drastically greater strain and strain accumulation rate as compared to their CG counterparts, primarily due to substantial GB migration and grain rotations.

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METHODS NC Ag, Cu and Ni films of ~ 50 nm in thickness were deposited on carbon-coated TEM mesh grids using the magnetron sputtering technique at room temperature. The vacuum chamber was evacuated to a base pressure of 6 × 10-8 Torr prior to deposition. The thicknesses of the foils were precisely controlled by the deposition time during synthesis. The deposition rate was measured to be 1.9±0.1 nm/s for Ag, 1.1±0.1 nm/s for Cu, and 0.6±0.05 nm/s for Ni. Therefore, the calculated error of the thickness for Ag, Cu and Ni foils is 2.4 nm, 4.5 nm and 4.2 nm, respectively. CG Ag, Cu and Ni films with large grain sizes were obtained by annealing in vacuum at 600℃ for 30 min for comparison. All TEM specimens were patterned by introducing arrays of holes with 5 KeV Ga+ ions at a beam current of ~ 0.05 nA, using an FEI X835 Focused Ion Beam (FIB) system. Each film was patterned with multiple holes in form of N by N (N =3-7) matrices. The time for creating one hole was about 2 s, thus the ion beam induced damage in the TEM specimen was minimized. The size and shape of the holes was expected to vary to some extent due to the slight change of beam profile during FIB experiments. Film morphology was characterized by using an FEI Quanta 600 scanning electron microscope (SEM). The TEM specimens were then irradiated by 1MeV Kr++ ions at the Intermediate Voltage Electron Microscope (IVEM)-Tandem facility at Argonne National Laboratory. Calculation shows that 99.995% percent of Kr ion penetrated through the specimens using Stopping Range and Ion for Materials (SRIM) using the Kinch-Pease method (Supplementary Fig. S9) [73]. The displacement energies were assumed to be 25eV for all three materials. 1 dpa in this study corresponds to a fluence of 2×1014 ions/cm2 and the damage rate was about 10-3 dpa/s. The typical duration for one irradiation 12


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experiment up to a dose of 2 dpa was about 30 min, excluding the time for imaging and data collection when the irradiation was paused. In situ TEM experiments were performed in a Hitachi-9000NAR microscope, which was attached to an ion accelerator, at Argonne National Laboratory in United States. Post-irradiation analyses were performed using a Tecnai F20 electron microscope. The films were sufficiently thin to be electron-transparent and thus no TEM specimen preparation was necessary.

SUPPORTING INFORMAITON AVAILABLE Strain measurement methodology, strain measurements of all six specimens, defect distribution, hole size statistics, additional snapshots of grain boundary migration and grain rotation, schematics and quantification of grain rotation, SRIM simulation, and in situ movies. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected], [email protected]. Notes The authors declare no competing financial interest.

ACKNOWLEDGEMENTS KY acknowledges financial supports from National Science Foundation of China (51501225) and Start-up Program of China University of Petroleum-Beijing (2462014YJRC019 and 2462015YQ0602). XZ et.al at Purdue University acknowledge financial support by NSF-DMR-Metallic Materials and Nanostructures Program under grant no. 1304101. HW acknowledges financial support by ONR. We also acknowledge the use of microscopes at the DOE Center for Integrated Nanotechnologies managed by Los Alamos National Laboratory. The IVEM facility at Argonne National Laboratory is supported by DOE-Office of Nuclear Energy.

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Manuscript Figures

Figure 1 | Microstructure of patterned nanocrystalline (NC) Ag. (A) Scanning electron microscopy (SEM) micrograph of NC Ag film deposited on carbon-coated mesh grid. (B) Magnified SEM micrograph of the area outlined in (A), showing patterned holes with an average diameter of ~ 150 nm. (C) Bright field and dark field transmission electron microscopy (TEM) micrographs showing morphology of the holes and the equiaxed nanograins. Inset is the selected area diffraction (SAD) rings confirming the polycrystalline structure of FCC Ag film. (D) Magnified TEM micrograph of the area outlined in (C) showing the morphology of the hole. Inset is a high resolution TEM image of the hole surface, verifying that limited damage was introduced to the specimen during Focused Ion Beam (FIB) milling.

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Figure 2 | Microstructure and grain size distributions of patterned metals. NC Ag (A1), CG Ag (A2), NC Cu (B1), CG Cu (B2), NC Ni (C1) and CG Ni (C2). The average grain size of each specimen was labeled.

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Figure 3 | Measurements of in-plane strain at different dose in NC Cu. (A) The outlined holes at 0 dpa (as-deposited specimen). The red cross in the center was chosen as set point and the blue crosses are markers of interest. The distance labeled is the gauge length. (B) The average strain was calculated to be 1.1 % at 1 dpa. A portion of tracked points were not shown intentionally. Actual measurements tracked 6-10 points. (C) The average strain at 2 dpa was measured to be 2.0 %. See supplementary Figure S1 and S2 for more details.

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A1

B NC Ag

NC metals CG metals

CG Ag

A2 NC Cu CG Cu

A3

Strain rate (%/dpa)

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Strain (%)

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2 Ag Cu

1 Ni

NC Ni CG Ni

0.0

0.5

1.0 Dose (dpa)

1.5

2.0

0 Specimen

Figure 4 | Strain and strain rate of all irradiated specimens. (A) Strain as a function of irradiation dose. Strain was found to increase with irradiation dose in all the specimens. (B) The strain rates (%/dpa) exhibited significant difference between the CG and NC specimens.

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Figure 5 | Grain boundary migration in nc Ag captured during the in situ irradiation experiment. (A1) Bright field image before GB migration, (A2) Dark field image before GB migration, (B1) Bright field image after GB migration, (B2) Dark field image after GB migration. As indicated by the dotted line, the GB migrated over a distance of 32 nm during irradiation from 0.375 dpa to 1 dpa. The specimen was not tilted during irradiation and the electron and ion beams were kept constant. See supplementary video for more.

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Figure 6 | Grain rotation induced by irradiation in NC Ag. The two grains indicated by arrows rotated over a dose range from 1.19 to 1.56 dpa. See supplementary figures and videos for more.

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Measurement of Heavy Ion Irradiation Induced In-Plane Strain in

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