Response of 14YWT alloys under neutron irradiation: A complementary study on microstructure and mechanical properties

Acta Materialia 167 (2019) 181e196

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Response of 14YWT alloys under neutron irradiation: A complementary study on microstructure and mechanical properties E. Aydogan a, b, * , J.S. Weaver a, c, U. Carvajal-Nunez a, M.M. Schneider a, J.G. Gigax a, D.L. Krumwiede d, P. Hosemann d, T.A. Saleh a, N.A. Mara e, D.T. Hoelzer f, B. Hilton g, S.A. Maloy a a

Los Alamos National Laboratory, Los Alamos, NM, 87544, USA Integrated Manufacturing Center, Sabanci University, Istanbul, 34906, Turkey National Institute of Standards and Technology, Gaithersburg, MD, 20899, USA d University of California Berkeley, Berkeley, CA, 94720, USA e University of Minnesota, Minneapolis, MN, 55455, USA f Oak Ridge National Laboratory, Oak Ridge, TN, 37830, USA g TerraPower LLC, Bellevue, WA, 98005, USA b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 31 August 2018 Received in revised form 25 January 2019 Accepted 27 January 2019 Available online 29 January 2019

Nanostructured ferritic alloys (NFAs) having sub-micron grain size with a high density of nano-oxides (NOs) (size of ~2e3 nm) are one of the best candidates for structural components in Generation IV nuclear systems. In this study, 14YWT NFA cladding tubes were irradiated in BOR60 reactor up to 7 dpa at 360e370  C. Detailed microstructural analysis has been conducted using bright field transmission electron microscopy, bright field scanning transmission electron microscopy, energy filtered transmission electron microscopy, energy dispersive X-ray spectroscopy, electron energy loss spectroscopy and transmission Kikuchi diffraction techniques. This revealed cavities, <100> and <111> type dislocation loops, and a0 precipitates forming after irradiation with relationships between cavities and NOs, and a0 precipitates and NOs. Cavities mostly form on the NOs; whereas, a0 precipitates form between the NOs where the point defect concentration is high. Moreover, a0 precipitates are distributed homogenously on and around the dislocation loops which is consistent with the observation that there is no Cr segregation on dislocation loops. Grain boundaries were found to be mostly depleted in Cr; however, the characteristics of each grain boundary determines the Cr behavior and the a0 denuded zone around the grain boundaries. Mechanical properties of the irradiated tubes have been determined by using both low force and high force nanoindentation techniques, resulting in 1.03 ± 0.33 GPa and 0.82 ± 0.20 GPa hardening, respectively. Dispersed barrier hardening calculations and nanoindentation measurements are in good agreement. In this study, 14YWT NFA has been systematically studied after neutron irradiation to better understand its superior performance: low a0 concentration, low swelling and low radiation-induced hardening. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Nanostructured ferritic alloys (NFAs) Neutron irradiation Nano-oxides (NOs) Alpha prime (a0 ) Radiation-induced hardening

1. Introduction Nanostructured ferritic alloys (NFAs) are attractive materials for Generation IV reactors due to their high temperature strength, thermal stability, creep resistance and radiation tolerance, provided by a high density of Y-Ti-O nano-oxides (NOs) having the size less

* Corresponding author. Integrated Manufacturing Center, Sabanci University, Istanbul, 34906, Turkey. E-mail address: [email protected] (E. Aydogan). https://doi.org/10.1016/j.actamat.2019.01.041 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

than 5 nm [1e3]. NOs were determined to be mostly Y2Ti2O7 particles having a pyrochlore structure [4,5]. They can pin grain boundaries and dislocations, leading to high strength and superior structural stability besides acting as recombination sites for both point defects created from neutron damage and trapping sites for helium atoms created by transmutation reactions [1,6]. It has been found that Y-Ti-O NOs are extremely stable under both neutron and ion irradiations in 14Cr NFAs (MA957 and 14YWT) at the temperatures between 300  C and 450  C. For instance, Mathon et al. [7] reported that NOs are stable at 325  C up

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to 5.5 dpa in neutron irradiated MA957. Similarly, Ribis et al. [8,9] found no significant changes in the NO size and distribution in MA957 neutron irradiated at 412  C and 430  C to 50 and 75 dpa, respectively. Similar to the neutron irradiations on MA957 alloys, Certain et al. [10] and He et al. [11] reported that NOs are stable at 300  C and 450  C under heavy ion irradiations in 14YWT alloys. Moreover, we have reported the same trend in 14YWT alloys at 450  C up to 1100 peak dpa damage using Fe2þ ion irradiations [12]. Even though there are a few ion irradiation studies on 14YWT, as mentioned above, there are no neutron irradiation studies investigating the microstructural changes on these alloys to the best of the authors’ knowledge. In ferritic alloys having a high Cr content, alpha prime (a0 ) formation occurs at certain temperatures which can cause embrittlement of the materials [13,14]. These a0 precipitates are a Cr-rich body centered cubic (bcc) phase coherent to the Fe-rich bcc matrix [15]. According to a recent work by Bonny et al. [16], the Cr threshold concentration to form a0 precipitates is less than 10 at.% at temperatures below 400  C. a0 formation can occur under thermal annealing, however, the formation temperature is reduced under irradiation due to local enrichment of Cr atoms as a result of displacement damage and kinetically accelerated radiation enhanced diffusion [17,18]. a0 formation has been reported to occur in Fe-Cr binary and Fe-Cr-Al ternary systems [17e22] together with more complex engineering alloys such as high Cr ferritic/ martensitic and oxide dispersion strengthened (ODS) alloys under neutron irradiation [19,23,24]. Radiation induced segregation (RIS) of Cr in high Cr alloys affects the performance of the alloys [25]. Depletion of Cr at grain boundaries is well known in face centered cubic (fcc) iron-based stainless steels (austenitic stainless steels); however, the behavior of Cr is complex in ferritic steels [26e29]. In the literature, most of the RIS studies for ferritic alloys were performed on ion irradiated alloys [26,30,31], and there are only a few studies on neutron irradiated high Cr ferritic alloys [24,25]. Among both ion and neutron irradiation studies, there is somewhat of a contradiction in RIS of Cr. Contradictions in the ion irradiation experiments might have resulted from the experimental variations such as ion type, dose and temperature. For instance, Was et al. [32] and Wharry et al. [33] have concluded that RIS behavior in ferritic/martensitic

(F/M) alloys depends closely on the Cr concentration and irradiation temperatures. They found that the grain boundaries are prone to get enriched at 400  C in T91 alloys having 9 at.% Cr while they are prone to be depleted at 500  C in HT9 and HCM12A alloys, both having 12 at.% Cr, after 2.0 MeV proton irradiations. This has been attributed to the dominance of a vacancy mechanism at high temperatures, and an interstitial mechanism at low temperatures [34]. In the case of neutron irradiations, even on the same MA957 alloy irradiated to ~110 dpa at 412  C, energy dispersive X-ray spectroscopy work by Toloczko et al. [35] has found Cr depletion at the grain boundaries while atom probe tomography work by Bailey et al. found Cr enrichment [24]. This contradiction has been attributed to the existence of Cr-rich particles (a0 ) in the microstructure [24]. However, recent studies have reported grain boundary characteristic dependent segregation behavior of Cr in a neutron irradiated 9Cr Model F/M alloy [25]. Therefore, it is crucial to understand the relationship between RIS of Cr at grain boundaries, grain boundary character, and a0 formation in high Cr ferritic alloys. In this study, we have investigated the radiation response of a 14YWT (14 wt.% Cr) NFA material, which was irradiated in the BOR60 reactor to 7 dpa at 360e370  C. This study investigates radiation effects on the microstructure and mechanical properties in a complementary way. It has been found that there are certain correlations/anti-correlations between nano-oxides, cavities, dislocation loops and a0 precipitates. Moreover, detailed analysis on the grain boundary characteristics and RIS has been performed. It has been demonstrated that based on the grain boundary characteristics, there is either a depletion of Cr or no change in the composition at the grain boundaries. Finally, mechanical properties were correlated with the microstructure by using both nanoindentation tests and hardening models. 2. Experimental procedure The nominal composition of the 14YWT NFA is 14Cr-3W-0.4Ti0.21Y-Fe (in wt.%). A description of the fabrication of the plate can be found elsewhere [36]. Cladding tubes having ~0.5 mm wall thickness were cut by electrical discharge machining (EDM) through the extrusion direction of the plate. Fig. 1 shows the deformation

Fig. 1. Schematic showing the plate deformation directions and sampling directions of the EDMed tubes. Corresponding bright field transmission electron microscopy images of the 14YWT tube neutron irradiated at 360e370  C up to 7 dpa are shown along azimuthal and radial directions. (ED: extrusion direction, RD: rolling direction, ND: normal direction).

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directions of the plate together with the EDMed cladding tube and corresponding azimuthal and radial microstructures in correlation with the initial plate deformation directions. The tubes were loaded into capsules for insertion into the BOR60 reactor (RIAR, Dimitrovgrad, Russia, 2014) through a collaborative irradiation campaign with TerraPower, LLC. Fast flux neutron irradiations were conducted at a temperature of 360e370  C up to 7 dpa. Temperature was determined by measurement-based analysis. In other words, it was first measured by thermocouples prior to commissioning and then determined by linear regression analysis based on the reactor parameters. The total time for the irradiation, and hence the duration at which the sample was at temperature, was ~1 year. Neutron flux was determined to be ~1.7  1015 n/cm2-s for energies greater than 0.1 MeV. Moreover, the dose and He production rates are estimated as 6.5  107 dpa/sec and ~0.1 He appm/dpa, respectively. Table 1 shows the dimensions and irradiation conditions of the 14YWT tube. It should be noted that this temperature range corresponds to the typical inlet temperatures in fast reactors. Moreover, radiation induced hardening occurs at these temperatures in ferritic alloys due to the increased density of vacancy and interstitial clusters resulting in the formation of dislocation loops and a0 particles which limits the lifetime of high Cr ferritic alloys [37]. Therefore, this temperature range has been selected to investigate the behavior of 14YWT alloys under the most life-limiting conditions. After neutron irradiation, coupons having the size of 1.0 mm  0.5 mm x 0.5 mm were cut from both radial and azimuthal directions of the tube (Fig. 1) via EDM to reduce the radioactivity [24]. The specimens were then mechanically polished (up to 1200 grit) at the University of California, Berkeley's nuclear materials laboratory. To reduce the activity to undetectable levels, samples were cut into a 15 mm  10 mm x 2 mm foil and attached to the Cu grids specimens using a standard lift-out technique with a FEI Quanta 3D FEG Focused Ion Beam (FIB) tool, instrumented with a Omniprobe G200 and a platinum gas injection system. Final thinning and cleaning of the foils were performed at Los Alamos National Laboratory using an FEI Helios Nanolab 600 dual beam FIB. An FEI Tecnai F30 TEM operating at 300 kV was used for bright field transmission electron microscopy (BFTEM) to investigate the swelling properties of the tube and bright field scanning TEM (BFSTEM) analysis to conduct dislocation studies. It should be noted that BFTEM was performed with 1e1.5 mm under-focus to enhance the visibility of the cavities. To further characterize these tubes, an FEI Titan 80 operating at 300 kV equipped with Gatan Tridiem 863 ER/S electron energy loss spectrometer was used to collect energy filtered TEM (EFTEM) jump ratio images at characteristic energy losses specific to the Fe-L2,3, Ti-L2,3 and Cr-M4,5 transitions in order to spatially map the distribution of the NOs and the a0 features. Also, the foil thickness was determined using a well-accepted methodology that compares the ratio of elastic to the inelastically scattered electrons that are transmitted through the sample [38]. To investigate the RIS along the grain boundaries and dislocation loops, transmission Kikuchi diffraction (TKD), energy dispersive Xray spectroscopy (EDX) and electron energy loss spectrometry (EELS) techniques were utilized. Grain boundary characteristics in

Table 1 Tube dimensions and neutron irradiation conditions. Length (mm) Outer diameter (mm) Inner diameter (mm) Wall thickness (mm) Dose (dpa) Irradiation temperature ( C) Dose rate (dpa/s) He production rate (appm/dpa)

30.99 4.57 3.51 0.53 7 360e370 6.5  107 ~0.1

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the TEM foil were determined by TKD using an FEI Inspect FEG SEM equipped with an EDAX Hikari Ultra EBSD system at an accelerating voltage of 20 kV and an aperture size of 50 mm. Line scans across the grain boundaries were performed along 40e60 nm line lengths with ~1.5 s dwell times using the FEI Titan 80 microscope at 300 kV also equipped with an EDAX EDX detector. It should be noted that the correlative TKD/EDX analysis was performed on 30 different grain boundaries. STEM-EEL spectral images were collected on dislocation loops at a camera length of 100 mm, with a 2.5 mm GIF entrance aperture, and a STEM probe current of ~1 nA. Due to the fact that spectra were collected at two-beam conditions to reveal the dislocation loops, diffraction contrast varied on the images. These differences in the total diffracted intensities resulted in variations in the total amount of signal entering the spectrometer. While these differences in the total amount of signal existed, the ratio of the inelastically scattered electrons of interest, namely the Fe-L2,3 and Cr-L2,3 energy losses, was only affected by compositional differences. As such, in order to accurately characterize the composition on dislocation loops, the ratios of the integrated intensities of the Cr-L2,3 edge to the Fe-L2,3 were computed. The sizes of dislocation loops, NOs and a0 were measured manually using the ImageJ v1.49 digital processing software including 200 to 700 counts for each foil. Statistics on dislocation loop, NO and a0 size and number density were obtained from at least 3 different regions on both the specimens from the radial and the azimuthal foil orientations. For dislocation loop analysis, the foils were first tilted to a <100> zone axis, then tilted to a twobeam condition with g<110> type direction. At this kinematic condition of g<110>, types of the dislocation loops were determined by their shapes and orientations as described in detail by Yao et al. [39]. According to this study, a<100> type loops have much larger aspect ratio compared to the a/2<111> loops along g<110> of <100> zone axis [39]. Since g<110> reveals 1/2 of the a/2<111> loops and 2/3 of the a<100> loops, correction factors of 2 and 3/2 were applied to a/2<111> and <100> loops, respectively. It should be noted that despite attempts to orient the sample into a perfect two-beam condition, at 300 kV the flattening of the Ewald Sphere results in the weaker excitation of other nearby diffraction vectors; as a result, some secondary g·b conditions were weakly excited and some loops having weaker contrast were present in the recorded images. These weaker loops were taken into account and error is represented as standard deviation in measurements to satisfy the possible error coming from the measurements. Error in the case of NO and a0 measurements were also represented as the standard deviation in the measurements. Nanoindentation tests were performed using a Hysitron Triboindenter 950 with a Berkovich, diamond indenter tip. A combination of low force and high force transducer tests were used to make reliable measurements at shallow (100s of nm) and deep (1000s of nm) indentation depths. The deeper depths using the high force transducer should be more comparable to microindentation testing since indentation size effects are significantly reduced and the maximum force is approaching the microindentation regime [40e42]. For low and high force tests, the continuous stiffness (CMX) module was used which allows for the continuous measurement of the unloading stiffness by oscillating the tip to generate many small elastic unloads. The CMX displacement amplitude was ~2 nm in all cases, and the frequency, which depends in part on the transducer, was 80 Hz for the high force transducer and 100 Hz for the low force transducer. The OliverPharr method was used to calculate the modulus and hardness using a tip area function calibrated on a quartz standard [43]. It should be noted that a total of two different Tribo-950 nanoindenters and four different Berkovich tips were used: one system and set of low force and high force transducer tips for the

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unirradiated material and one system and set of tips for the irradiated material in a radiological area. 3. Results 3.1. Microstructure 14YWT alloys are reported to have tri-modal grain size [36,44], and the NO distribution varies with respect to the grain size. The NO size and density are found to be 2.0 ± 0.9 nm and 6.3 ± 1.2  1023 m3, respectively. The distribution of NOs is not perfectly homogenous and NOs were observed to precipitate along some of the grain boundaries [12]. Fig. 2 shows the under-focused BFTEM micrographs of the asreceived plate before irradiation and the EDMed tube after irradiation taken to investigate the cavity distribution. Fig. 2a and b show cavity distribution at different regions in the as-received plate before irradiation. There are clearly both large cavities and small cavities in the as-received plate. The large cavities are typically larger than 100 nm in diameter and their volume fraction is 0.03 ± 0.02%. Moreover, the size of the small cavities is less than 10 nm in diameter and their volume fraction is 0.005 ± 0.003%. Fig. 2c and d show microstructures of the tube along the radial and azimuthal directions after 7 dpa irradiation at 360e370  C. As seen from the micrographs, in both directions, cavity sizes are less than 10 nm. Excluding the large cavities existing before irradiation, the volume fraction of cavities after irradiation is measured to be 0.015 ± 0.01%. If the fraction of pre-existing small cavities are subtracted, it can be

concluded that 0.01 ± 0.01% of the cavities form as a result of neutron irradiation. Considering the variations in the grain size and NO distribution in the initial microstructure [12], cavity size and density measurements were performed in 5e7 different grains along both radial and azimuthal directions of the tube. It has been found that the dislocation loop types, sizes, densities and distributions are similar along both the radial and the azimuthal directions. Preliminary examination of an area and roughly 60 loops from both the radial and azimuthal directions showed only a 6% variance in the values of loop size, fraction of types, and densities. Based on these preliminary results, more detailed studies were only continued to pursue quantification along the radial direction of the 14YWT tube. Fig. 3a shows BF-STEM micrographs of dislocation loops on different grains and different kinematic conditions in 14YWT tube along radial direction after irradiation to 7 dpa at 360e370  C. g<110> conditions were obtained for all images by tilting from a <100> zone axis. It should be noted that the dislocation loops are a/2<111> and a<100> type loops as their shapes are similar to the shapes of the theoretically calculated <111> and <100> type loops illustrated in Fig. 3a [39]. Fig. 3b shows the size distribution of a/2<111> and a<100> type loops. Both of the loop distributions are skewed with values up to 50e60 nm. Their sizes and number densities were determined based on their shapes as 12.9 ± 2.6 nm and 1.9 ± 0.5  1022 m3 for <111> type loops and 16.9 ± 3.4 nm and 6.2 ± 2.6  1021 m3 for <100> type loops, respectively. The size of <111> loops is slightly smaller while their number density is ~3 times larger compared to <100> loops. It should be noted that the correction explained in the

Fig. 2. BFTEM micrographs showing the cavity distributions in (a) and (b) as-received plate at different grains; in EDMed tube after neutron irradiation at 360e370  C to 7 dpa at (c) radial and (d) azimuthal directions. Images were taken at 1e1.5 mm under-focus. Red arrows highlight some of the cavities. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

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Fig. 3. (a) BF-STEM two-beam images of neutron irradiated 14YWT tube along radial direction at different grains and kinematic conditions (b) the size distribution of <111> and <100> type loops showing skewness. Images were taken by tilting from <100> zone axis. Scale bar in (a) is the same for all four images.

experimental part has been applied to the loop densities due to the invisibility criteria along g<110>. Fig. 4 shows the EFTEM Fe and Ti jump ratio maps of the NOs and their size distributions in irradiated tubes at different regions along both the radial and the azimuthal directions. The average NO size is measured as 2.04 ± 1.04 nm with a distribution skewed up to ~11 nm, with the highest number fraction of oxides having the size between 1 and 2 nm. Moreover, number density of the NOs was found to be 10.5 ± 5.9  1023 m3. It should be noted that size and number density of the NOs are similar to the results before irradiation, which were 2.0 ± 0.9 nm and 6.3 ± 1.2  1023 m3, respectively. Size and number density of NOs before and after irradiation indicate that NOs are stable under ~7 dpa neutron irradiation at 360e370  C. In other words, there is no dissolution or growth meaning that dissolution due to radiation cascade and growth due to thermal annealing at 360e370  C have been balanced as discussed in detail elsewhere [12]. The precipitation of a0 has been shown to be closely related with irradiation temperature, time and alloy composition [23,45,46]. Such precipitation has been observed in many other high Cr alloys (>9 wt.%) under neutron irradiation between 300 and 440  C [17e24]. Similarly, the 14YWT alloy neutron irradiated to 7 dpa at

360e370  C in this study shows a0 formation. EFTEM Cr jump ratio maps to visualize the a0 distribution in the neutron irradiated 14YWT tube are shown in Fig. 5. As seen in Fig. 5, the a0 distribution is quite homogenous and it shows almost a Gaussian distribution. The a0 size and density are measured as 3.37 ± 0.85 nm and 4.97 ± 0.99  1023 m3, respectively. It should be noted that either NOs (Y-Ti-O particles) or their outer rims are reported to be rich in Cr [12,47]. Therefore, EFTEM Ti jump ratio maps have been shown in Fig. 5 to distinguish between NOs and a0 . In the present study, we have observed that depending on the specific grain boundary type, boundaries might be enriched, depleted or unaffected. However, the terms of ‘depletion’ and ‘enrichment’ should be defined first. For the case of qualitative assessments (mapping as in Fig. 6), these terms will be used relative to the Cr concentration in the matrix. In other words, enrichment and depletion are determined by looking at the contrast differences present between the matrix and the grain boundaries. For instance, Fig. 6a shows two grain boundaries; one showing Cr depletion with an a0 free zone, a denuded zone, Fe enrichment and a homogenous Ti (and Y-Ti-O particle) distribution, second showing Cr and Ti enrichment and Fe depletion. Fig. 6b shows a grain boundary having Cr depletion (or no change), Fe depletion and Ti enrichment. It is not

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Fig. 4. EFTEM Fe and Ti jump ratio maps showing NO distribution, obtained from neutron irradiated tube along (a) radial and (b) azimuthal directions, and (c) NO distribution showing skewed distribution up to ~11 nm. Dark and bright contrasts in Fe and Ti maps correspond to depletion and enrichment, respectively.

Fig. 5. EFTEM Cr and Ti jump ratio maps in neutron irradiated 14YWT tube to show a0 distribution along (a) radial and (b) azimuthal directions, and (c) the a0 size distribution showing almost a Gaussian distribution.

clear if there is an a0 denuded zone or not. Fig. 6c shows enrichment of Cr and Ti with an a0 denuded zone coupled with the depletion of Fe at the grain boundary while Fig. 6d shows Cr and Ti enrichments without an a0 denuded zone besides Fe depletion at the grain boundary. Finally, Fig. 6e demonstrates a grain boundary showing no enrichment or depletion of Cr, Ti and Fe. In this figure, Y-Ti-O particles are positioned on and around the grain boundary. However, the terms of ‘enrichment’ and ‘depletion’ can also be used based on the pre-irradiation condition as will be discussed in the discussion part. 3.2. Mechanical properties Nanoindentation tests were performed to measure the modulus

and hardness as a function of indentation depth and quantify the radiation hardening (change in hardness after irradiation) of the 14YWT tube. In contrast to the ion irradiations, neutron irradiation indentation studies can be done to larger displacements since the radiation damage is not limited to a shallow region at the surface. Thus, low and high force transducers were used to measure the hardness at displacement ranges of 200e300 nm and 3800e4000 nm. The measurements at 3800e4000 nm displacement mitigate indentation size effects and the influence it has on the radiation hardening measurement. The reduced modulus and hardness versus displacement curves are shown in Fig. 7 using the low and high force transducers on unirradiated and irradiated samples. The average properties are given in Table 2. The average

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Fig. 6. EFTEM jump ratio maps of Cr, Ti and Fe at grain boundaries having (a) Cr depletion with denuded zone on one; Cr enrichment and denuded zone on another (b) Cr depletion (or no change) without denuded zone (c) Cr enrichment with denuded zone (d) Cr enrichment without denuded zone and (e) no enrichment/depletion and no denuded zone. White dashed lines show a0 denuded zones. Arrows point the grain boundaries. þ, -, 0 on maps represent enrichment, depletion and no change, respectively. dz: denuded zone.

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Fig. 7. Nanoindentation hardness and modulus plots as a function of depth obtained using a (a) low load and (b) high load transducer. Areas enclosed by solid and dashed lines represent one standard deviation of the hardness and modulus, respectively.

reduced (effective) modulus ranges from 198 to 210 GPa. This closely matches the estimated reduced modulus of 195e206 GPa which depends on the crystal grain orientation. The estimated reduced modulus was calculated using a diamond tip, Young's modulus and Poisson's ratio of 1140 GPa and 0.07, respectively, elastic constants for 14YWT measured by Stoica et al. [48] and theory for anisotropic cubic materials from Vlassak and Nix [49,50]. There does not appear to be a trend in the modulus between the unirradiated and irradiated sample. The modulus curves on the high force tests do show a slightly decreasing trend with increasing displacement compared to the low force tests which show a constant modulus with displacement. On the other hand, high force tests on the quartz standard showed a constant modulus with displacement meaning the machine compliance and area function are calibrated correctly. There can be some additional compliance in the system based on how the sample is mounted which is likely the cause for the slightly decreasing modulus on the high force tests on 14YWT. The modulus measurements in Table 2 show that the high force tests are in good agreement with the low force tests where this additional compliance was not an issue; therefore, the slightly decreasing modulus on the high force tests has a negligible effect on the measured properties.

There is an overall increase in hardness after irradiation which reflects a net increase in the defect density thereby increasing the force required to plastically deform the material under the tip. The average hardness measured for unirradiated and irradiated samples decreases as one changes from the low to the high force transducer. This is primarily due to the decreasing hardness with increasing indentation depth combined with the fact that the hardness values are determined at displacements of 200e300 nm for low force and 3800e4000 nm for the high force. The change in hardness values, DH, are nearly the same for the low and high force measurements, D Hðlow forceÞ ¼ 1:03±0:33 GPa and DHðhigh forceÞ ¼ 0:82±0:20 GPa; with a slightly lower average for the high force tests. Typically the change in hardness will increase with increasing indentation depth due to differences in the indentation size effect between the two materials [42]. Namely, the indentation size effect is more pronounced for materials with a low defect density (i.e., unirradiated material) compared to the materials with a high defect density (i.e., irradiated material). This results in a smaller change in hardness at small displacements compared to larger displacements [42]. The absence of the expected trend in the measurements in this study can be explained through consideration of the microstructure under the indenter tip. The very small grain size, 100s of nm, means that the material probed even at 200e300 nm displacement contains many grains, grain boundaries, and nano-oxide particles (i.e., there are many defects present). This effectively reduces the difference in the indentation size effect before and after introducing radiation damage, thus reducing the sensitivity of the change in hardness to differences in the indentation size effect before and after irradiation. Alam et al. [51] measured the micro Vickers hardness on the same unirradiated 14YWT as 352e376 kgf/mm2. In order to compare our results with these values, the high force transducer hardness (H) measurements were converted to microhardness (HV)
 kgf using the relationship of HV mm ¼ H½GPa*94:5. This conversion 2 accounts for the difference in the area which is the projected contact area for nanohardness and the surface area for Vickers microhardness. This results in an average hardness of 445±16 kgf/ mm2 of the unirradiated material. This is slightly higher than the values measured by Alam et al. [51]. It should be noted that a force of 0.50 kgf was used in the case of micro Vickers hardness and a force of only ~0.15 kgf is achieved during the high force nanoindentation tests. Since the nanohardness is still dropping even at 4000 nm displacement, it is possible that testing to higher displacements/loads might bring these values in better agreement. Alam et al. [51] also noted the presence of micro-cracks in their sample while we did not directly observe any. The absence of micro-cracks in the volume of material tested in this study could also explain why we estimate a higher value for Vickers microhardness. 4. Discussion 4.1. Radiation induced microstructural changes in 14YWT alloys Nanostructured ferritic alloys have been reported to be radiation resistant due to the existence of high number density of NOs. These

Table 2 Summary of the nanoindentation measurements on irradiated and unirradiated tubes. Condition

Transducer/Displacement [nm]

Reduced Modulus [GPa]

Hardness [GPa]

Unirradiated Irradiated Unirradiated Irradiated

Low/200-300 Low/200-300 High/3800-4000 High/3800-4000

198.2 ± 3.0 209.7 ± 3.2 203.1 ± 6.0 193.1 ± 4.5

5.29 ± 0.22 6.32 ± 0.25 4.71 ± 0.17 5.54 ± 0.11

DH [GPa] 1.03 ± 0.33 0.82 ± 0.20

No. Tests 25 20 45 8

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NOs may act as recombination centers for both point defects created by neutron damage and trapping sites for helium atoms created by transmutation reactions [1e3]. Under very high dose self-ion irradiations (up to 1100 peak dpa) at 450  C, NOs were found to be extremely stable. Their size decreased slightly while their density increased just slightly [12]. Moreover, NFAs were found to exhibit almost zero swelling with ~0.02% swelling at the depth of 400e600 nm [52]. Similar to the ion irradiations, after 7 dpa of neutron dose at 360e370  C, swelling is determined to be ~0.01%. Considering the error in the measurements of cavities before and after irradiation and small value of cavity volume fraction, swelling can be counted almost zero proving the fact that 14YWT alloys are radiation resistant materials. <111> and <100> type loops have been reported in Fe-Cr and Fe-Cr-Al alloys under both ion and neutron irradiation at the irradiation temperatures of 300e500  C [39,45,46,53,54]. It has been reported that there is a transition temperature (500e550  C) above which only a<100> type loops are seen while there is a mixture of a/2<111> and a<100> type loops existing together below this temperature [55e57]. In this study, the number density of <111> type loops is found to be ~3 times larger than that of <100> loops. Existence of a high density of <111> loops after neutron irradiation at 360e370  C infers that <111> type glissile loops are not completely lost to the defect sinks during irradiations, as also observed by Field et al. [46] in FeCrAl alloys neutron irradiated to 1.8 dpa at 382  C. In this study, the loop types were determined based on their apparent morphology. Due to the resolution limits of the recorded BF-STEM images, the morphology of the black dots cannot be determined, and this uncertainty may affect the proportions of <111> and <100> type loops counted. Moreover, the

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size of the <100> loops is found to be larger than that of the <111> loops. Experiments and molecular dynamics simulations have shown that a<100> type loops are formed through the interaction of similar size a/2<111> loops [58,59], and they grow with point defect and <111> type loop absorption [60]. It is known that in ferrous alloys having a high Cr amount (>9 wt.% Cr), a0 formation occurs at certain temperatures which can cause embrittlement of the materials [13,14]. It has been found that Fe-Cr-Al alloys having 10e18 wt.% Cr form a0 under neutron irradiation at 320e382  C [20e22,46,61]. Moreover, Bachhav et al. [19] reported a0 precipitation at 290  C in binary alloys having 9e18 wt.% of Cr; Kuksenko et al. [62,63], Chen et al. [64] and Pareige et al. [65] observed a0 precipitation at 300  C in Fe-(10e16) wt.% Cr model alloys under neutron irradiation. Similarly, a0 precipitation has been observed in high Cr (>9 wt.% Cr) engineering alloys. For instance, Anderoglu et al. [23] reported the formation of a0 in 12Cr F/M HT9 alloy neutron irradiated at 380  C and 440  C while Bailey et al. [24] observed a0 precipitation in 14Cr MA957 alloy neutron irradiated at 412  C. In this study, a0 precipitation has been observed also in 14YWT alloys after 7 dpa of neutron irradiation at 360e370  C. 4.1.1. Correlation between cavities and NOs Cavities have been found to form along dislocation lines within the grains as seen in Fig. 8a. Similarly, in Fig. 8b, NOs are shown to exist on dislocation lines. This may be the indication of either dislocation pinning of NOs or the precipitation of NOs on the dislocations during high temperature processing. Yamashita et al. [66] have reported dislocation pinning by the NOs in MA957 alloys after irradiation to ~100 dpa at ~700  C. On the other hand, Fig. 8b shows

Fig. 8. (a) BFTEM image taken under focus of ~1.5 mm to visualize the cavity distribution (cavities are aligned, probably on dislocation lines) (b) BFTEM image showing the distribution of NOs on dislocation lines (c) High magnification BFTEM image of NOs and cavities in under focus and over focus conditions, showing that cavities are correlated with NOs. Red and white arrows indicate cavities and NOs, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

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that the dislocations are rather decorated with NOs. We have shown that during processing of the alloys, deformation results in dissolution of some of the NOs by dislocation shearing, and then, during high temperature annealing, NOs re-precipitate mostly on dislocations [44]. This indicates that there exists a certain correlation between NOs and cavities. Fig. 8c shows both under- and overfocused images; the cavities are more readily visible in the underfocused condition and NOs in the over-focused condition. It is shown that the cavities form on the NOs. Odette et al. [67] and Yang et al. [68] reported that most of the NOs are associated with helium bubbles, and their existence results in a reduced size of helium bubbles in the ferrite matrix as well as a lower number of helium bubbles along grain boundaries. On the other hand, in the BOR60 reactor, helium production rate is very low and the helium effect within the microstructure is negligible. Therefore, the formation of cavities as a result of irradiation is assumed to be due to either migration of vacancy clusters initially formed inside the matrix to NO/matrix interface or direct formation of vacancy clusters at the NO/matrix interface. The formation energies of vacancies in the iron matrix, grain boundary and NO/matrix interface have been calculated as 2.14, 0.91 and 0.62 eV, respectively [68,69]. This indicates that the vacancies preferentially form at the NO/matrix interface rather than at the grain boundaries or within the matrix. Moreover, the migration energies for the vacancy clusters forming at the matrix have been reported to be less than 0.69 eV [70]. On the other hand, as shown in the previous section, some of the cavities were determined to come from processing of the alloy due to mechanical alloying and heat treatment processes. For those cavities, it is not possible to speculate whether cavities are precipitating on the NOs or vice versa. 4.1.2. Correlation between a0 precipitates, dislocation loops and NO-cavity complexes It is known that the presence of Cr reduces radiation induced swelling [71e75]. Initial theories proposed a mechanism based on the dislocation bias towards interstitials which leads to the clustering of vacancies to form cavities. It has been proposed that the Cr atoms trap vacancy clusters and enhance the recombination rate, resulting in reduced swelling [71]. However, Cr and vacancies form a weak bond and it is difficult to justify the enhanced recombination in the presence of weak bonding, especially at high temperatures [72]. On the other hand, recently, Terentyev et al. [73,74] have shown that the Cr atoms interact with interstitial loops and restrict their movement. Therefore, vacancies can find the interstitial clusters and recombination occurs [75]. In ODS alloys, the existence of NOs is presumed to lead to the formation of vacancies mostly at the NO/matrix interface (due to their lowest formation energies [68,69]) and restrict the swelling. However, due to the reduced movement of interstitial complexes as a result of the Cr trapping, interstitial-vacancy recombination might be limited at the NO/ matrix interface. In other words, the existence of a high Cr content may result in higher swelling. Under ion irradiation, it has been reported that the Cr can segregate at the dislocation loops restricting their motion [26,76]. For instance, Jiao et al. [26] has reported the Cr segregation at dislocation loops in ferritic/martensitic alloys HCM12A, T91 and HT-9 following proton irradiation to 7 dpa at 400  C and in HCM12A following Fe2þ ion irradiation to 100 dpa at 500  C, using atom probe tomography technique. Similarly, Heintze et al. [76] presumed that Cr segregation occurs at dislocation loops in 9 at.% Cr alloy after neutron irradiation at 300  C up to 0.6 and 1.5 dpa, indirectly by using small angle neutron scattering. Contrary to these studies, Fig. 9 shows no segregation on the dislocation loops, however, it does show the heterogeneity of the Cr distribution caused by the presence of the a0 precipitates. Even though this can

be due to the resolution limits of different techniques, resolution limit of the EELS system is as low as ~0.05 at.%, which is comparable with the other techniques. Therefore, the contradiction can be explained in two ways: First, it does not necessarily mean that Cr is not coupling with interstitial clusters; the interstitial clusters might not be visible under the current imaging conditions as also suggested by Hernandez-Mayoral et al. [77]. Second, even though the Cr-interstitial complexes are initially stable under irradiation, with Cr diffusion from the grain boundaries, in time, the driving force for the formation of a0 precipitates might become larger than the Crinterstitial complex binding energy. This may explain why segregation of Cr is not seen and a homogenous distribution of a0 is seen (Fig. 9), which is consistent with the findings in Ref. [78]. Little and Stow [71] initially proposed that the a0 formation promotes the swelling since Cr atoms will be used by the precipitates and the Fe-rich matrix will be depleted in Cr. Contrary to this, high Cr alloys (>10 at.% Cr) have been shown to have better swelling resistance [79]. Therefore, it has been suggested that repulsive force between the a0 precipitates and the interstitial clusters traps the interstitial clusters between the precipitates which in turn leads to enhanced recombination resulting in reduced swelling [75,80]. Fig. 9 demonstrates that a0 precipitate distribution is quite homogenous and there is no evidence of depletion in the vicinity of the loops. It is not possible to claim this result is contradictory to the above stated studies. Again, there may be some interstitial clusters away from the a0 precipitates that cannot be resolved in the current study. Moreover, the existence of NOs in the high Cr alloys adds more complexity to the evolution of defects and precipitates under irradiation. As discussed above, due to the preferential formation of the vacancies at the NO/matrix interface, the recombination might be limited compared to other high Cr ferritic or F/M alloys even though small (unresolved) interstitial clusters are trapped between a0 precipitates. It should be also noted that there is a certain anti-correlation between a0 precipitates and NOs. In other words, a0 formation occurs between the NOs which is also reported before by Bailey et al. [24]. This is attributed to the existence of high density point defects away from the NOs. Therefore, a0 formation is the easiest away from NOs. 4.1.3. Radiation induced grain boundary segregation/depletion Radiation induced segregation or depletion occurs in Fe-Cr alloys due to preferential solute atom - point defect coupling migrating to defect sinks, such as grain boundaries (GBs) [25]. In both austenitic stainless steels and F/M alloys, it has been shown that the RIS response of the materials is related with the grain boundary characteristics [25,81,82]. Similarly, in this study, grain boundaries show either enrichment or depletion or no change in Cr concentration, shown in Fig. 6. The relationship between grain boundary characteristics and segregation/depletion behavior of Cr has been investigated by correlative TKD/EDX along 30 grain boundaries. Fig. 10 shows the measured Cr concentrations from the EDX line profiles through the grain boundaries having different misorientation angles. At least three line scans were performed through each grain boundary to confirm the Cr concentration trend. As mentioned above, ‘enrichment’ and ‘depletion’ terms are used relative to the pre-irradiation conditions where Cr concentration of these alloys was reported as ~14 wt.% Cr in the matrix before irradiation [12]. It should be noted that the grain boundaries might be enriched or depleted in unirradiated condition. Regardless of that, comparison is made considering the matrix composition of unirradiated condition, and red dashed line in Fig. 10 corresponds to the pre-irradiation matrix composition. As seen in Fig. 10, the Cr concentration shows both enrichment and depletion within the grain due to the existence of a0 precipitates. In other words, a0 precipitates show up as Cr enrichment and then between the a0

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Fig. 9. EELS maps of Cr, Fe and Cr/Fe on dislocation loops along g<110> (as indicated by white arrows on the HAADF STEM images). White dashed ellipsoids on Cr/Fe maps mark the dislocation loops seen on the HAADF STEM images. Red arrows showing enrichment in Cr/Fe maps point some of the a0 precipitates. Note that the apparent deficiencies in the Cr and Fe maps are not caused by a compositional change, but rather the strong diffraction condition (as evident in the HAADF image) causing more signal to be deflected out of the energy loss spectrometer's entrance aperture. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

particles, the matrix is depleted in Cr. Low angle grain boundaries having misorientation angles of less than 5 do not show any enrichment or depletion of Cr. Also, there is no a0 denuded zone around the grain boundaries. Similarly, low angle grain boundaries having a misorientation angle between 5 and 15 show no change in the Cr concentration. In contrast, random high angle grain boundaries show depletion at the grain boundaries and there is a low Cr concentration region around the grain boundary between the two Cr peaks. Therefore, the denuded zone around the random high angle grain boundary in Fig. 10 is determined to be ~23 nm. In addition to the low- and high-angle boundaries, coincidence site lattice (CSL) boundaries of Ʃ5 show a local enrichment at the grain boundaries compared to its immediate environment. The Cr concentration at the grain boundary is ~14 wt.% and there is a low Cr concentration region around these types of boundaries. The width of the denuded zone is found to be ~13 nm in Fig. 10. In the case of Ʃ3 CSL boundaries, no enrichment or depletion is observed, nor is there a depleted zone, as observed in Fig. 10. It should be noted that the deviation from CSL misorientation angle is determined by Palumbo-Aust criterion, defined in Ref. [83]. Field et al. [25] reported that there is always segregation of Cr in

9Cr Model alloys under neutron irradiation to 3 dpa at 500  C and that the degree of segregation decreases considerably in the case of CSL boundaries. Moreover, while the random high angle boundaries have the highest Cr concentration, the low angle boundaries have a lower amount of segregation, though still higher than the CSL boundaries. However, they did not report the formation of Crrich particles. In this study, even though the trend is similar to the above stated study, we have found that there is either a depletion or no change of the Cr concentration at the grain boundaries compared to the Cr concentration of the matrix in unirradiated condition. The existence of a0 precipitates in the matrix confirms this finding in such a way that Cr atoms move from grain boundaries as a result of point defect migration under irradiation and form a0 precipitates. On the other hand, one should be careful when concluding either enrichment or depletion due to the possibility of hitting a0 precipitates on grain boundaries during scanning. For instance, on MA957, Bailey et al. [24] performed atom probe tomography and found that grain boundaries are rich in Cr; yet, on the same sample, Toloczko et al. [35] performed EDX line scans across the grain boundaries and reported a Cr depletion. In both cases, the boundary characteristics were not reported. The

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considerable hardening (DH ¼ 1.03 ± 0.33 GPa for low force, DH ¼ 0.82 ± 0.20 GPa for high force) due to neutron irradiation in 14YWT alloys at 360e370  C. The change in tensile yield strength ðDsy Þ can be estimated using the Busby model [42]: Dsy ðMPaÞ ¼ 289:17DHðGPaÞ: The Busby model was developed for microindentation data; therefore, it requires the nanohardness to be converted to microhardness following the conversion mentioned previously. Krumewiede et al. [42] critically evaluated the use of nanoindentation data with the Busby model. It was recommended to use the characteristic hardness to reduce the influence of indentation size effects. Here we note that there is less than a 2% change between the characteristic hardness and hardness at 3800e4000 nm displacement. In addition, the change in hardness remains relatively unchanged between the low and high force nanoindentation tests. Using the Busby model results in a change in tensile yield stresses, Dsy , for low and high force of 297.8 ± 95.4 MPa and 237.1 ± 57.8 MPa, respectively. In the current study, the radiation-induced hardening was calculated considering dislocation networks, dislocation loops, cavities and a0 precipitates. It should be noted that the black dots were not taken into account as it is not possible to determine their types (<100> or <111>) by looking at their shapes due the resolution limit of the BF-STEM. As has been explained above, there is no considerable change on the size and number density of the NOs after irradiation. The dispersed barrier hardening model described in Ref. [85] can be used to calculate the strengthening caused by the radiation induced discrete defects as in equation (1).

pffiffiffiffiffiffiffi

DsSR ¼ MaGb dN

Fig. 10. Cr concentrations from EDX line scans across the grain boundaries having different misorientation angles (GB1 to 5), as determined by TKD. Red dashed lines (14 wt.% Cr) correspond to the matrix composition in unirradiated condition. (dz ¼ depleted or denuded zone). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

difference in the reported results can likely be attributed to a variation in the grain boundary types on which each measurement was conducted. If the grain boundary is low angle boundary or CSL type, as shown in both Fig. 6b and e, the a0 distribution will be random. If the line scans cross across a particle, the grain boundary will likely be labeled as enriched rather than being depleted. It has been reported that the precipitate free zone occurs as a result of small point defect concentrations around the grain boundaries which slows down the diffusion of Cr and thereby restricts the Cr-rich particle precipitation [84]. In other words, a precipitate free zone is an indication of the effectiveness of a grain boundary to annihilate the point defects in its immediate vicinity. As seen in Fig. 10, GB 3, a high angle random grain boundary, has the largest depleted zone compared to the other low angle and CSL boundaries, meaning that high angle random boundaries are the best at the annihilation of nearby point defects. While proving efficacious as defect-sinks, these boundaries exhibit the lowest Cr concentration. From these observations, we conclude that highangle grain boundaries are necessary for the effective annihilation of the point defects and the restriction of a0 formation; however, extensive Cr depletion on these boundaries could prove harmful for the structural integrity and corrosion resistance of these materials.

4.2. Radiation induced changes in mechanical properties Both low force and high force indentation tests indicated a

(1)

where M is Taylor factor which is determined to be 3.06 for most bcc polycrystalline materials [86]; a is the dispersed barrier constant; G is the shear modulus (81.07 GPa), b is the Burgers vector, and N and d are the number density and size of the defects, respectively. Since a<100> type dislocations are sessile at the current irradiation temperatures, slip occurs mostly by a/2<111> type dislocations; therefore, b ¼ 2.48 Å, for the 〈111〉{110} slip system [87]. Furthermore, when the barriers are impenetrable, as in the case of the Orowan by-pass mechanism, then a ¼ 1. When a is in the range of 0.05 and 0.3, the particles are soft and susceptible to shearing [1]. Field et al. [46] have calculated the dispersed barrier coefficients for linear dislocations, dislocation loops and a0 precipitates by fitting the experimental data to the calculated data using both linear summation and root-sum-square summation methods. a values for both methods were determined to be relatively close. Therefore, in this study, the values obtained by these two methods were averaged. A barrier strength coefficient of 0.05, 0.11, 0.32 and 0.62 were used for a0 precipitates, <111> loops, <100> loops and dislocation lines, respectively. The high barrier strength coefficient of <100> type loops compared to <111> type loops is attributed to the sessile characteristics of <100> loops [46,55]. Moreover, a value for the cavities is determined to be 0.45 from the barrier strength vs. size plots reported by Tan et al. [88]. Indeed, they have shown that the a values for all discrete obstacles are size dependent in austenitic alloys neutron irradiated at 275e375  C. However, the size dependency of the a values for precipitates and loops has been shown to be weaker compared to the cavities. Therefore, it is ignored for a0 precipitates and loops in this study. Consequently, strengthening due to a0 precipitates and cavities has been calculated to be 126 ± 25 MPa and 89 ± 18 MPa, respectively while it is calculated as 106 ± 21 MPa and 201 ± 40 MPa for <111> and <100> type loops, respectively. On the other hand, dislocation network density decreased from ~8  1014 m2 to ~6  1014 m2 (measured by the method described in Ref. [83]). Therefore, softening due to the reduction in dislocation density is calculated to be 144 ± 29 MPa.

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The change in total strength can be estimated using a non-linear superposition model developed by Dunn et al. [89] using equation (2):



Dstot ¼ Dsna0 þ Dsn< 111 > þ Dsn< 100 > þ Dsncav  Dsndisl

1=n

(2)

The value of n has been found to be 2.22 using dislocation dynamics simulations [90]. Total strengthening using Eq. (2) can be found as 225.6 ± 45.1 MPa while the measured strengthening values from low force and high force tests are 297.8 ± 95.4 MPa and 237.1 ± 57.8 MPa, respectively. It is obvious that measured and calculated strengthening values are similar. Fig. 11a presents the contribution of a0 precipitates, dislocation loops, cavities and dislocation networks to the irradiation hardening, and Fig. 11b compares the calculated total hardening and measured hardening by both low force and high force tests. In Fig. 11, the strengthening has been converted to hardening using the correlation of Dsy ðMPaÞ ¼ 289:17DHðGPaÞ. Consequently, low force and high force tests result in 1.03 ± 0.33 GPa and 0.82 ± 0.20 GPa hardening, respectively, while total calculated hardening is found as 0.79 ± 0.16 GPa. It should be noted that the hardening coming from black dots has been disregarded in this study. Even if it is considered, the total calculated hardening would still be between the values obtained by low force and high force tests. However, there are some uncertainties in the calculations. First of all, there is an uncertainty on the barrier strength coefficients of the obstacles. Second, even though the dispersed barrier hardening model is widely used, Sobie et al. [90] have shown that the Bacon Kocks Scattergood (BKS) model was found to accurately predict hardening due to cavities, whereas the Friedel Kroupa Hirsch (FKH) model was superior for SIA loops. Third, error may be coming from the hardening calculations due to cavities. In Fig. 8, we have shown that most of the cavities are associated with NOs. Therefore, cavities cannot be taken as the only medium individually interacting with the dislocations. Barrier strength coefficients of NOs in 14YWT NFAs were reported as ~0.23 [52]. Thus, the barrier strength coefficient of NO/cavity complex might be somewhere between 0.23 and 0.45. In this case, in Eq. (2), there should be another term for NO/cavity complexes rather than individual cavity strengthening term. Also, since there still is a high density of NOs without cavities, a NO strengthening term should be included as well. Consequently, the calculated hardening will be lower. Rather than hardening, NO/cavity complex may even result in softening. Wang et al. [91] have reported that if the cavity sizes become comparable with the size of the nanostructure (in their case nano-layers), cavities might facilitate the dislocation emission through the interface of the nanostructure. Wei et al. [92] also found that in the presence of cavities on the grain boundaries, grain boundary sliding occurs easily resulting in softening after irradiation.

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4.3. Overall radiation resistance of 14YWT alloys compared to other materials In the literature, NFAs have been reported to have good high temperature microstructural stability, creep and radiation resistance [1e3]. However, mechanical properties, nano-oxide stability, a0 formation and dislocation loop evolution have been investigated separately. This study reports the microstructural evolution which is correlated with mechanical properties in a complementary way. In this study, even though there is a0 , dislocation loops and small amount of void formation (almost zero though), their number densities are lower compared to the Fe-Cr and Fe-Cr-Al model alloys and high Cr structural materials. For instance, Bachhav et al. [19] performed neutron irradiations on different compositions of Fe-Cr alloys at 390  C up to 1.8 dpa. They reported that the number density of a0 is 9.5 ± 0.2  1023 m3 in Fe-12Cr alloy while it is 32±3  1023 m3 in Fe-15Cr alloy. At slightly lower temperature, 320  C, Reese et al. [93] reported the number density of a0 particles as 3.1 ± 0.6  1023 m3 for Fe-12Cr alloys while it is 11±3  1023 m3 for Fe-15Cr alloys after 7 dpa neutron irradiation. Similarly, Field et al. [46] reported a0 formation in Fe-12Cr-4.4Al and Fe-15Cr-3.9Al alloys with number densities of ~20  1023 m3 and ~48  1023 m3, respectively after neutron irradiation at 382  C to 1.8 dpa. Moreover, Edmondson et al. [20] found that the number densities of a0 are ~7  1023 m3 for Fe-12Cr-4.4Al and ~23  1023 m3 for Fe-15Cr3.9Al alloys after neutron irradiation at 320  C to 7 dpa. All these studies reported a0 size between 2 nm and 4 nm. In this study, we report a0 size of ~3.4 nm and density of ~5  1023 m3. It should be noted that number density of the a0 precipitates vary considerably with dose and dose rate while size changes slightly [18,64,94]. In other words, while the volume fraction of a0 decreases with increasing dose rate, it increases with increasing dose [64]. Therefore, considering that the above stated studies have been performed at the dose rates of 3.4-8.1  107 dpa/sec, similar to the current study, but some being irradiated to lower doses, 14YWT alloy having 14Cr shows much lower density of a0 . In other words, even though a0 formation occurs in 14YWT alloys, they are confined in between NOs and their number density is much smaller compared to the other alloys. Moreover, since they are confined to the space between NOs, their size and number density is expected to saturate in time. However, this hypothesis requires further evidence at higher dose irradiations. Lower hardening compared to the other alloys confirms the efficiency of these alloys. Krumwiede et al. [42] analyzed the hardening behavior of various alloys irradiated at 320  C up to 6.5 dpa. They found that while HT9 alloys show ~1.2 GPa of hardening, 14YWT alloys show almost no change in their hardness. Moreover, FeCrAl alloys irradiated to 1.8 dpa at 382  C show ~0.6 GPa of hardening [46]. However, Anderoglu et al. [95] have compiled the mechanical data for Fe-Cr and ferritic/martensitic alloys and found that hardness is saturated after 5e10 dpa. Therefore, the hardening

Fig. 11. Plots showing (a) the contributions of a0 precipitates, <111> and <100> loops, cavities and dislocation lines to the hardening, and (b) comparison between total calculated, low force and high force measured hardening.

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amount in those FeCrAl alloys is expected to increase at higher doses. In fact, irradiation hardening of ~0.8 GPa in 14YWT alloys after 7 dpa irradiation at 360e370  C seems to be quite reasonable and lower compared to the other high Cr alloys. 5. Conclusions In this study, radiation induced microstructural and mechanical properties in 14YWT NFA material have been investigated after neutron irradiation to 7 dpa at 360e370  C. After irradiation, cavity, dislocation loop and a0 formation occur in addition to radiation induced segregation/depletion of Cr at the grain boundaries. The complex microstructural evolution of 14YWT due to neutron irradiation has been systematically quantified. The consortium of defects existing before and generated during irradiation result in low a0 precipitation, swelling, and hardening. The conclusions are as follows:  NOs mostly decorate the dislocation lines and most of the cavities are associated with NOs. This has been proposed to occur due to the low formation energies of vacancies at the NO/matrix interface compared to the matrix and the grain boundaries.  <100> and <111> type loops have been observed after irradiation. Even at the temperatures of 360e370  C, the proportion of <111> type loops is ~3 times larger, which can possibly be attributed to the analysis ignoring the black dots.  Cr has previously been reported to segregate at the dislocation loops. However, in this study, there is a homogenous distribution of a0 precipitates around the loops after neutron irradiation. Moreover, Cr enrichment or depletion at the grain boundaries is strongly dependent on grain boundary characteristics.  We have demonstrated that low angle grain boundaries investigated in this study do not show any change in Cr concentration while Cr depletion is observed at random high angle grain boundaries with a large a0 denuded zone. Furthermore, Ʃ3 boundaries show neither a change in the Cr concentration nor an a0 denuded zone. Even though Ʃ5 boundaries show no change in the Cr concentration at the boundaries, a denuded zone does exist around these boundaries, though the size of the denuded zone is smaller than those observed at the random high angle boundaries.  Hardening as a result of neutron irradiation has been determined to be 1.03 ± 0.33 GPa for low force tests and 0.82 ± 0.20 GPa for high force tests. Combining the effect of all the defect types observed using the dispersed barrier hardening model and a nonlinear summation method to predict the change in hardness matches well with the measurements.  Compared to the other Fe-Cr and Fe-Cr-Al alloys, 14YWT alloy exhibits low concentrations of a0 formation and lower amount of radiation induced hardening which makes it one of the best candidates for the Generation IV reactor applications. Acknowledgements This research was partially supported by the DOE-NE Fuel Cycle Research and Development Program under the Contract number DEAC52-06NA25396. Moreover, it 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. The authors would like to give special thanks to Prof. G.R. Odette for his contributions during production of the alloys and Dr. R. McCabe for his help on preliminary TKD experiments. Moreover, the authors would like to give their appreciation to Dr. A. Nelson for allowing them to use his facility for radioactive materials and Drs. W. Chen and E. Martinez for helpful discussions on hardening models.

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