Introduction

Next-generation nuclear reactors will need materials that can withstand very harsh operating conditions such as higher temperatures, high irradiation doses, and chemical interactions with a variety of coolants.1 A variety of different materials are being explored to meet these needs. One class of materials that is showing promise for next-generation nuclear reactor applications is nanostructured ferritic alloys (NFA) including oxide dispersion-strengthened (ODS) steels.2,3,4 Due to the initial favorable indications there is extensive ongoing worldwide research effort to develop ODS steels suitable for nuclear reactors, specifically 14YWT.5,6,7,8,9,10,11,12,13 14YWT is a NFA ODS alloy with 14% chromium in the material, and YWT stands for yttrium, tungsten, and titanium which contribute to the nanoscale oxide particles in the steel.

These NFA ODS steels have a microstructure with nano-sized grains and a high density of nano-oxide particles, which improve mechanical properties such as creep.14 Literature has shown that high densities of oxide nano-particles in the metal matrix increase irradiation tolerance by creating a critical density of interfaces, which enables effective defect annihilations and recombination.14,15 These interfaces also assist with the material’s ability to withstand helium production and accumulation in these steels.16,17 The creep of structural components during operation can lead to failure. In nuclear reactors, this is a particular concern because radiation damage can enhance creep from the production of additional point defects in the materials.18 Measuring this enhanced creep is challenging for both neutron and ion beam irradiations. In neutron irradiations, it is difficult to fit the appropriate testing apparatus in intense neutron flux locations to enable measuring creep under irradiation. However, several pressurized tube studies on different materials have been performed.19,20,21 In addition, neutron irradiations are time consuming and costly, limiting the number of data points that can be collected. The challenge for utilizing ion beams for irradiation-enhanced creep measurements is the limited penetration depth of ions into the material. This limited penetration depth restricts the testing to small-scale mechanical testing (SSMT) or meso-scale techniques.22 Measuring the irradiation-enhanced creep of structural materials is important to understanding their behavior during the operation of a nuclear reactor. In this study, SSMT techniques were used to compare the creep behavior of unirradiated and ion irradiated 14YWT.

The purpose of this study is to evaluate the effectiveness of small-scale creep techniques as such as in-situ transmission electron microscopy (TEM) nanopillars and Berkovich nanoindentation techniques. The choice of in-situ TEM nanopillars with ion beam irradiation would enable the measurement of irradiation-enhanced creep properties with the ions completely passing through the pillar to avoid the stopping peak. The ability to observe changes in the creep rate between the control and ion beam irradiated nanopillars would enable high throughput testing. In addition, exploring the possibility for nanoindentation methods to capture changes in deformation mechanism. The literature on 14YWT indicates a change in deformation mechanism using macro-scale mechanical testing > 873 K (600°C).5,13 The ability to also observe a similar change in the deformation mechanism with nanoindentation in the same temperature range would validate utilizing small-scale mechanical testing for irradiated materials. There have been a few studies on the nanoindentation of 14YWT but all have been performed at room temperature.11,34 There have also been a few studies on the irradiation-enhanced creep of ODS alloys.14,20 In Ref. 14, helium ions were used on 100-µm-thick samples of PM2000 and K1 to examine the irradiation-enhanced creep response, and it was observed that the size and distribution of dispersoids did not show a significant influence on the irradiation creep behavior. In Ref. 23, pressurized cylinders of MA 957 were used to examine neutron irradiation-enhanced creep. At Sandia National Laboratories, a TEM has been modified to enable a wide range of in-situ experimental capabilities.23 This instrument is called the In-situ Ion Irradiation TEM (I3TEM). Several of these capabilities can be combined enabling the evaluation of irradiation-enhanced creep of materials in-situ in the TEM.24,25

In this study, a Hysitron PI-95 in-situ mechanical testing holder along with an incoming ion beam was used to evaluate irradiation-enhanced creep of 14YWT utilizing nanopillars at room temperature. The use of the I3TEM with SSMT enables the otherwise difficult measurement of the creep properties during irradiation. However, there are still challenges with small sample size effects. In addition to these in-situ studies, nanoindentation creep measurements were performed on control samples of 14YWT using a Hysitron indenter from room temperature to 1073 K to evaluate any changes in the deformation mechanism of the material as a function of temperature. Observations indicated that there was an increase in the creep rate from the in-situ TEM studies during irradiation at room temperature. Ex-situ nanoindentation creep studies showed increased creep with temperature and a change in the deformation mechanism between 873 K and 1073 K, which agrees with literature on macro-scale testing of 14YWT.

Experimental Methods

In-situ TEM Experiments

A sample of 14YWT was initially obtained from a research study, which produced a large heat of 14YWT through gas atomization of the powders, attritor milling for 40 h, and consolidation through extrusion at 850°C.15 The composition and additional microstructural investigations of the material can be found in Refs. 5, 6, 9, 12, and 15. The composition of the material is Fe-14Cr-3W-0.4Ti-0.2Y (wt.%). The sample of 14YWT for the nanopillar TEM experiments was first mechanically polished to approximately 50 μm thickness. The grinding/polishing was performed with SiC papers and water as a lubricant. The samples were then sequentially polished with 3 and 1 μm diamond solution to a mirror finish. The sample thickness was periodically measured during the grinding/polishing with a micrometer. To ensure proper alignment for later testing in the I3TEM, the sample of 14YWT (1 mm × 1 mm × 50 μm-thick) was mounted on the JEOL PI-95 sample holder using silver paste prior to the fabrication of the pillars. FEI Helios dual-beam focused ion beam Ga + /scanning electron microscopy (FIB/SEM) was used to manufacture the nano pillars for the testing. The rough trenching at 65 nA left a section of the 14YWT approximately 5–10 μm thick and 100 μm-long. The nanopillars for the creep compression experiments were then milled from this section of the sample using subsequently lower currents in multiple shaping and cleaning steps to produce the final nanopillars. The nanopillars were approximately 280 nm-thick × 280 nm-wide × 650–700 nm-high and were spaced 5 µm apart as shown in Fig. 1B. The area for the nanopillars was measured top down for the width and thickness and from the side for height. The pillars were overtilted 2° to minimize the taper on the nanopillars. The area measurement for the cross-sectional area was performed at the top of the nanopillars. The 14YWT was a sub-micron grain size in the material that would indicate that each pillar had 2–3 grains in it. In the analysis the orientation of the grains in the pillars was not accounted for.

Fig. 1
figure 1

(A) Showing the geometry of the ion beam, electron beam, and pillar compression direction (not to scale) reprinted with permission from Ref. 23. (B) A schematic of the pillar size, spacing, and shape along with that of the pillar.

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The in-situ TEM experiments were performed on the I3TEM at Sandia National Laboratories. The I3TEM allows for imaging of the sample, while being irradiated with ions.23 In addition, the Hysitron PI-95 picoindenter enables instrumented indentation in the TEM to allow mechanical testing. Room temperature (RT) compressive creep experiments were performed with and without ion beam irradiation. The nanopillars tested as the control specimens (no ion beam irradiation) were tested perpendicular to the electron beam in the TEM. The nanopillars tested during the ion beam irradiation were tilted to 25° to allow simultaneous viewing of the pillars and ion beam irradiation. The nanopillars were irradiated with 2.8 MeV Au4+ ions. The flux of ion was known by measuring the area of an irradiation burn spot in an optical microscope and knowing the current (4.5 nA). Using SRIM26 with the composition of 14YWT and a displacement energy of 40 eV, with thickness of the pillars at 25° and 2.8 MeV Au4+ ions, the peak displacement damage rate in the material was calculated at 0.028 dpa/s as illustrated in Fig. 1A. It was verified that the ions would pass predominately through the pillar for a more constant damage profile throughout the pillar. The creep tests were performed while simultaneously irradiating the nanopillars. For the compression creep experiments, the PI-95 was used in load control mode. The pillars were loaded to approximately ~ 450 MPa (~ 50 μN load) and held for 300 s to measure the change in displacement with time, which is approximately half of the room temperature yield stress for most heats of 14YWT. For these experiments, three nanopillars were tested in each condition.

An example of a nanopillar in the testing configuration and the strain versus time for the in-situ TEM nanopillar compression creep experiments is displayed in Fig. 2A. The stress in the pillars was calculated using the cross-sectional area of the pillars and the load applied to the pillars from the PI-95 indenter. Digital image correlation (DIC) analysis of the tests in the TEM was used to produce accurate displacement values for the pillars using the base of the pillars and interface between the pillar and tip as the reference points. The strain rate for the pillars was calculated by first fitting a linear trend to the corrected displacement data from the DIC. Next, the slope of the linear fit was measured with the pillar height to calculate a strain rate for the test. The stress for the tests was calculated using the area of the pillars and load from the PI-95 instrument. The equation for the stress is given in Eq. 1. The strain was calculated by taking the displacement form DIC, which was measured using SHIFT DIC software at the top of the pillar and the bottom with example of the points (red crosses) chosen seen in Fig. 2A. The equation for the strain rate can be seen in Eq. 2, and Fig. 2B shows that slope of the line is also the strain rate. The engineering stress and strain were used to calculate the true stress and strain using the standard equations.

Fig. 2
figure 2

(A) TEM image of the nanopillar ready for testing in the I3TEM in the PI-95. A small amount of dust on the edge of the flat punch is obscuring the view. (B) Representative strain curves for the nanopillars tested in the irradiation and control configurations.

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$$\mathrm{Stress}(\mathrm{Pa})= \frac{\mathrm{Force }(\mathrm{N})}{\mathrm{area} ({\mathrm{m}}^{2})}$$
(1)
$$\mathrm{Strain}\; \mathrm{rate}\left(\frac{1}{s}\right)= \frac{\mathrm{Test}\; \mathrm{displacement} \left(\mathrm{m}\right)/\mathrm{Pillar}\, \mathrm{height} \left(\mathrm{m}\right)}{\mathrm{Time}\; \mathrm{of} \; \mathrm{test} (s)}$$
(2)

Elevated Temperature Nanoindentation

The sample for the elevated temperature nanoindentation experiments was from the same bulk piece of 14YWT as the in-situ TEM nano-compression experiments. The sample surface finish was also mechanically polished to a mirror finish using the same procedure described earlier. The nanoindentation and nanoindentation creep experiments were performed on the non-irradiated 14YWT at RT, 573 K (300°C), 873 K (600°C) and 1073 K (800°C) in a Hysitron Triboindenter with a XSOL 800 hot stage. The Hysitron Triboindenter at Los Alamos National Laboratory (LANL) has been modified to allow for flooding the whole chamber with an inert or reducing gas to inhibit the oxidation of the samples.27 In these experiments ultra-high-purity (UHP) argon was used as the cover gas and the oxygen was monitored using a GasAlert Extreme Single Gas Detector (BW technologies) with detection limit of 0.1% (1000 ppm) oxygen. The indenter chamber was purged overnight with UHP argon until an undetectable quantity of oxygen was measured with the oxygen sensor and was further flushed 2–3 h to ensure that sample was in a low oxygen environment.

The nanoindentation experiments were performed with a sapphire Berkovich indenter that was calibrated on fused silica before the nanoindentation experiments. In addition, the tip was evaluated on fused silica after the indentation experiments to assess the wear of the tip during the experiments, which was not significant.

The argon atmosphere in the nanoindenter chamber limits the maximum load available for testing to ~ 3–4 mN because higher loads would cause arcing in the transducer leading to anomalous displacement events in the load versus displacement curves. The nanoindentation and nanoindentation creep measurements were performed in load control mode with a maximum load of 3 mN. For the RT indentation the loading scheme was a 10 s load and dwell for 10 s followed by a 10 s unload. For the elevated temperature nanoindentation, the dwell period was adjusted so the unloading was elastic in nature and could be as high as 60 s. The nanoindentation creep experiments had a similar load-unload scheme except the hold period was 300 s. The 300 s hold was to enable the sample to reach a second stage steady state creep during the test.

The tip is not actively heated in the XSOL 800 setup, which is equilibrated using the image feature of the Hysitron indenter. This feature holds the tip of the indenter in contact with the sample surface to equilibrate the temperature difference between the tip and the sample surface. The tip was in contact with the sample surface 5–10 min prior to performing the nanoindentation creep experiments to minimize thermal drift during the indent. The sample was heated between the different testing temperatures at 20°C a minute and was stabilized at each testing temperature for 5–10 min prior to the tip being in contact with sample surface. At least ten indents, enough to ensure a random sampling of grains, for the nanoindentation or the nanoindentation creep experiments were performed at each temperature.

The reduced modulus and hardness from the nanoindentation creep experiments were calculated using the Oliver and Pharr method.28 The reduced modulus with equations in Ref. 28 was used to calculate the elastic modulus of the 14YWT regarding temperature with corrections for temperature-dependent elastic modulus of the sapphire tip.29 The hardness is calculated from the forced measured from the indenter and the contact area of the tip, which is calibrated on fused silica prior to the indentation experiments as seen in Eq. 3. The reduced modulus is calculated for the slope of the initial unloading portion of the curve as seen in Eq. 4. A schematic of a nanoindentation curve can be seen in Fig. 3 showing these values. If the tip material and Poisson’s ratio of the sample are known, Eq. 5 can be used to calculate Young’s modulus of the sample from the measured reduced modulus.

$$\mathrm{Hardness}= \frac{\mathrm{Pmax}}{A}$$
(3)

where Pmax is the maximum load measured with the indenter and A is the area in contact which is calibrated prior to the indents on fused silica.

$$S= \frac{\mathrm{d}P}{\mathrm{d}h}= \frac{2}{\sqrt{\pi }}{E}_{\mathrm{r}}\sqrt{A}$$
(4)

where Er is the reduce modulus, A is again the area, and S is the slope as shown in Fig. 4.

$$\frac{1}{{E}_{\mathrm{r}}}= \frac{\left(1- {\nu }_{s}^{2}\right)}{{E}_{\mathrm{s}}}+ \frac{\left(1- {\nu }_{i}^{2}\right)}{{E}_{i}}$$
(5)

where νs is Poisson’s ratio of the sample, Es is Young’s modulus of the sample, νi is Poisson’s ratio of the indenter, and Ei is the temperature appropriate Young’s modulus of the indenter. The indenter values were from Ref. 29.

Fig. 3
figure 3

Schematic of a nanoindentation curve showing the values used to calculate the hardness and reduced modulus.

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Fig. 4
figure 4

A plot showing the strain rate of the nano-compression specimens in the control and irradiated conditions.

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The stress exponent of a nanoindentation creep experiment is calculated from the slope of ln(\(\dot{\varepsilon }\)) and the ln(σ) or ln(H) curves, where \(\dot{\varepsilon }\), σ, and H are the steady state strain rate, uniaxial stress, and hardness, respectively.30,31,32 This originates from the power-law creep equation for conventional steady-state creep, which is shown in Eq. 6.

$$\dot{\varepsilon }=A{(\frac{\sigma }{G})}^{n}{e}^{(-\frac{Q}{\mathrm{R}T})}$$
(6)

where n is the uniaxial stress exponent, A is the uniaxial pre-exponential term, Q is the activation energy, R is the universal gas constant, and T is the temperature.

The strain rate for nanoindentation creep experiments can be calculated with Eq. 7:30,31,32

$${\dot{\varepsilon }}_{i}= \left(\frac{1}{h}\right)\left(\frac{\mathrm{d}h}{\mathrm{d}t}\right)$$
(7)

where h is the depth of the indent and dh/dt is the penetration rate in the steady-state section or linear portion of the nanoindentation creep curve, which is assumed to correspond to secondary or steady-state creep. The stress during nanoindentation creep was calculated from the area of the tip and the force from the indenter.

Results

In-situ TEM

The measured stresses and strain rate for the nanopillars tested in both the non-irradiated and ion beam irradiated conditions can be seen in Table I and Fig. 4. For similar stresses the strain rate increased for the nanopillar tests under ion beam irradiation. The strain rates for the different conditions are 2.7 × 10–5 s−1 for the control and 5.7 × 10–4 s−1 for the irradiated nanopillars calculated using engineering stress and strain. If the true stress and strain curves are used the strain rates are 2.69 × 10–5 s−1 for control and 5.24 × 10–4 s−1 for the irradiated case. If there was a jump in the displacement (one pillar in the irradiated set), the jump was not included in the calculation of the strain rate and the slope of the continuous curve was used. In addition, being able to watch the tests in situ ensures that pillars were loaded in a uniaxial fashion and there was no bending. The slope on either side of the jump was calculated and found to be within 5% of each other. The jumps were anomalous artifacts from the PI-95 during testing.

Table I Strain rates and stress for the nanopillars tested in both the control and irradiated conditions in situ in the TEM
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Elevated Temperature Nanoindentation

The hardness and Young’s modulus values measured with nanoindentation over temperature for 14YWT can be seen in Table II where the Young’s modulus and hardness of the 14YWT sample are decreasing with increasing temperature. The stress exponents for the 14YWT nanoindentation creep experiments at the tested temperatures are also provided in Table II. Stress exponent value at 873 K (600°C) is 3.0 ± 0.4 while at 1073 K (800°C) is 1.6 ± 0.3. These values of the stress exponent would indicate that dislocation motion is the deformation mechanism at 873 K and that grain boundary sliding is the deformation mechanism at 1073 K. The change in the deformation mechanism will be expanded on in the discussion section of the manuscript. In Fig. 5 is a plot of the hardness and Young’s modulus over temperature showing a linear decrease for Young’s modulus from RT to 873 K and exponential decrease in the hardness from RT to 1073 K.

Table II Hardness, Young’s modulus and stress exponent values of 14YWT over temperature measured with nanoindentation
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Fig. 5
figure 5

Young’s modulus and hardness of the 14YWT sample measured over temperature.

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Discussion

In-situ TEM

The true stress and strain calculations show that the strain rate for control pillars was 2.69E-5 s−1 and for the irradiated pillars 5.24E-4 s−1. It is believed that the increase in the strain rate was because of the increase in the defect density (vacancies, interstitials, and subsequent dislocations) caused by the irradiation. This has also been observed in previous studies using the same experimental setup in which the creep strain rate of a pure material increases with ion beam irradiation. In the case of Jawaharram et al.’s work on silver (Ag), the strain rate can be seen going from negligible at room temperature with no ion beam irradiation to substantially increasing to 0.0013 s−1 at temperature and irradiation.24 As discussed here, it is stated in Ref. 24 that the increase of point defects from irradiation and temperature is causing the increase in the creep rate. Comparing a structural alloy like 14YWT with a variety of alloying elements, nanocrystalline structure and nano-oxide particles would not be straightforward. However, it is useful to see that both materials had an increase in the creep rate with irradiation. Due to the limited number of samples and stress levels tested, it would be difficult to calculate stress exponent values from these results. It is promising though that a change in the strain rate is observed in these experiments on structural materials. More systematic studies which incorporate compensating for the size effects and large surface-to-volume ratio of these nanopillar geometries would be needed to fully understand the phenomena and apply them to bulk-scale phenomena.

In addition, these experiments show that it is possible to use SSMT techniques with I3TEM to rapidly evaluate the response of new and innovated materials to irradiation damage. Innovative sample mounting designs could introduce the possibility to measure these properties of magnetic structural materials at elevated temperature in the TEM.

Nanoindentation

As expected, Young’s modulus and hardness decreased with temperature. Young’s modulus values go from 221 ± 16 GPa at room temperature to 123 ± 19 GPa at 1073 K (800°C). This is in good agreement with literature values that report 230 GPa at room temperature, 175 GPa at 873 K, and 132 GPa at 1073 K,33 lending credence to the nanoindentation measurements. Unfortunately,33 does not provide values for 573 K; however, it reports a value of 205 GPa at 673 K, which agrees well with our results. Other nanoindentation data at room temperature show Young’s modulus values of 220–235 GPa.34 E. Aydogen et al. measured similar values of 211–234 GPa during room temperature nanoindentation for Young’s modulus.11 The nanoindentation studies also report hardness values of 5.5 GPa34 at the same depth as these indents at room temperature and 5.29 ± 0.22 GPa in11 at 200–300 nm depth.

Literature provides very few hardness data for 14YWT at elevated temperatures; however, several tensile studies have been completed.5,8,13,15,35 Many empirical relationships have also been developed in other studies to relate the hardness of material with the tensile properties. A generally accepted conversion from the yield stress of the material to the hardness is to divide by 3.36 This conversion is performed on the hardness data collected in this study and compared with macro-scale tensile testing of 14YWT at temperature shown in Table III and Fig. 6. When using this general conversion there is good agreement between the hardness values measured at 873 K and 1073 K and the macro-scale tensile tests. The hardness values at RT are much higher than the values of the macro-scale results, which is believed to be due to the size effects in the material. Indentation size effects have been seen in a variety of materials and are the phenomenon of smaller indents measuring higher hardness values. Due to the restrictions because of the Ar gas in the sample chamber (arcing issue described earlier), the RT indents only had a depth of 100 nm while the 873 K and 1073 K indents were > 200 nm. In the case of the RT indents, the 100 nm depth could still be in the indentation size effect zone and would influence the results. It has been shown in literature that as the temperature of nanoindentation increases the size effect becomes less pronounced.37

Table III Comparison of converted hardness values from this study with yield stress values of a tensile tests from literature
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Fig. 6
figure 6

A plot of the yield stress of the control 14YWT over temperature. The hardness values are converted to yield stress by dividing by 3.

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The stress exponents for the nanoindentation creep studies at 873 K and 1073 K are displayed in Table II; the stress exponent at 873 K is 3.0 ± 0.4 and at 1073 K is 1.6 ± 0.3. Current literature reports show that nanoindentation creep tests give similar results to macro-scale uniaxial creep tests for the creep exponent.38,39,40 However, it should be noted for the following discussion that nanoindentation creep tests have a complicated stress state under the tip due to the pyramidal nature versus the uniaxial stress state in most macro-scale tests. A stress exponent of 3 at 873 K would indicate that the dislocation motion is the primary deformation mechanism in the material, which agrees well with other studies of the creep and tensile testing of 14YWT and NFAs.41,42 The value of 1.6 for the stress exponent at 1073 K corresponds more to grain boundary sliding as the primary deformation mechanism in the material. The literature shows that at around 1073 K there is change in the deformation mechanism of 14YWT. Creep specimens tested at 1073 K still had a cup and cone fracture surface while those at 1173 K had an intergranular fracture surface with grains being separated by decohesion.15 In Ref. 43 it is suspected that there was grain boundary migration before this decohesion. In Ref. 5 they observed a decrease in the elongation of tensile tests performed at 1073 K and above. This was believed to be due to a reduction in the low-energy decohesion at grain boundaries. In Ref. 44 fracture toughness measurements at 973 K saw a change in the fracture surface that is believed to be related to grain boundary decohesion. These studies list several reasons for the change in the behavior such as the powder-forming oxides along grain boundaries prior to processing. As this change in the deformation behavior of the material is occurring around 1073 K, it would be expected that the same phenomena are occurring here. It is believed that even with smaller volume samples in these creep experiments there is still the possibility for grain boundary effects due to the nano-grain size of the material. In nanoindentation, it is approximated that the volume probed with an indent is half sphere shaped with a radius 20 times the depth of the indent.45,46 In the 1073 K (800°C) case the indents had depths of 418 ± 39 nm. This would be a half sphere with a radius of 8.36 µm of material probed including a sufficient number of grain boundaries for evaluation.

To evaluate whether there is a change in the deformation mechanism, the activation energy of the plastic deformation should be examined. The activation energy can be calculated from Eq. 8 where Qc is the activation energy, E is the Young’s modulus, a constant pre-factor G, and R is the gas constant [8.314 J/(mol K)]. The activation energy for plastic deformation can be determined by a graphical analysis, where the natural logarithm of the hardness divided by Young’s modulus is plotted against the reciprocal of absolute temperature. The slope of the fitted straight lines multiplied by R is Qc. The literature values for the yield stress in Table III were converted to hardness values for this plot. All the hardness values were divided by Young’s modulus measured in this study at the appropriate temperature. The plot with the fitted lines can be seen in Fig. 7. It shows that there is a substantial change in the activation energy of plastic deformation between RT-873 K and 873–1073 K when the value for the average of nanoindentation and tensile literature values for RT-873 K is 2.0 kJ/mol and the value for 873–1073 K is 15.5 kJ/mol. The value for just the indentation performed in this study for RT-873 K is 3.8 kJ/mol and the value for 873–1073 K is 16.8 kJ/mol, which agree well with the overall average values. The literature on mechanical testing of 14YWT would suggest that the deformation is similar from room temperature up to 873 K.5,13,33,35,43 This would also suggest a change in the deformation mechanism occurred from 873 K to 1073 K. The literature data from multiple sources are plotted on the graph in Fig. 7, where the data from RT until 873 K exhibit a linear decrease as seen with the fit to the average values at the temperatures tested in this study. One data set also has overall higher values but still is following the same trends in the data with a change of slope around 873 K.

Fig. 7
figure 7

Activation energy plot of the natural logarithm of the measured mechanical ratio of the properties as a function of inverse temperature.

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$$\frac{H}{E}=G{e}^{-\frac{{Q}_{c}}{RT}}$$
(8)

Conclusion

The in-situ TEM nanopillar experiments showed a significant increase in the creep rate of 14YWT when performed under ion beam irradiation. It is believed that an increased density of point defects resulting from the irradiation caused the increase in the creep rate. The nanoindentation creep tests showed a change in the deformation behavior at 873 K that agrees well with macro-scale results. It is believed that creep deformation went from dislocation-controlled creep to grain boundary sliding. While there are challenges with these small-scale creep tests, further refinement could demonstrate the efficacy of performing creep measurements under ion beam irradiations utilizing one of the several in-situ SEM indenters currently available.47 This could provide an avenue for exploring the irradiation-enhanced creep properties of materials at elevated temperature, especially for materials that are not suitable for TEM applications.