# Xe Ion Irradiation-Induced Microstructural Evolution and Hardening Effect of Nickel-Base Superalloy

Authors: Yu Hou; DeHui Li; YanLing Lu; HeFei Huang; WeiGuo Yang; RenDuo Liu
Journal: Science and Technology of Nuclear Installations
 (2021)

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## Abstract

The nickel-base superalloy Hastelloy N was irradiated using 1 MeV Xe20+ and 7 MeV Xe26+ ions with displacement damage ranging from 0.5 dPa to 10 dPa at room temperature (RT). The irradiated Ni-based superalloy was characterized with transmission electron microscopy (TEM), XRD, and nanoindenter to determine the changes in microstructural evolution and nanoindentation hardness. The TEM results showed that ion irradiation induced a large number of defects such as black spots and corrugated structures and the second phase was rapidly amorphized after being irradiated to a fluence of 0.5 dPa. The XRD results showed that the Hastelloy N alloy sample did not undergo lattice distortion after ion irradiation. An obvious irradiation hardening phenomenon was observed in this study, and the hardness increased with Xe ion fluence. The pinning effect in which the defects can become obstacles to the free movement of dislocation may be responsible for the irradiation-induced hardening.


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## Body

## 1. Introduction

Nuclear energy will play a more important role in meeting future energy needs for the fast-growing energy demands and concerns about climate changes [1]. The molten salt reactor (MSR), one of the six most promising generation IV nuclear reactor systems, has attracted increasing interest for its incomparable advantages: high inherent safety, reliability, higher power generation efficiency, etc. [2, 3]. However, the development and deployment of MSR systems are hindered by the structural materials. The structural alloys which must be subjected to extremely harsh environments involving high temperature, high neutron doses, and strong molten salt corrosion [4–7]. Thanks to its excellent corrosion resistance and adequate high-temperature strength, the nickel-based Hastelloy N alloy, which was developed by Oak Ridge National Laboratory (ORNL) in the 1960s [8], is still considered as the most promising candidate for the structural materials of MSR [9].As we know, neutron bombardment can induce a microstructural evolution of structural materials during the operation of the reactor, which can degrade their mechanical performances [10–12]. Hence, the evaluation of irradiation damage in structural materials is particularly important. The ion irradiation has advantages such as the easier control of irradiation parameters, reduction of cost and rapid damage production to neutron irradiation, and the high dose and high flux ion irradiation [13–19].In this study, xenon ions were selected to irradiate the Hastelloy N alloy for their significant displacement damage within the material, which is similar to the neutron irradiation damage behavior. Some thin-foil and bulk specimens of Hastelloy N alloy were irradiated, respectively, with 1 MeV Xe20+ ions and 7 MeV Xe26+ ions, at different fluences. The bulk and thin-foil specimens were characterized by nanoindentation and transmission electron microscopy (TEM) to study the changes. However, the ion irradiation-induced mechanical properties are difficult to be tested owing to the limitation of ion penetration depth, normally from nanometers to micrometers. In recent years, the nanoindentation test has been widely used to study the irradiation hardening of the reactor structural materials [20, 21], thanks to its powerful capability on thin-layer nanohardness measurement [22]. The correlation between the irradiation-induced formation of defects and irradiation-induced hardening was quantitatively analyzed.


## 2. Experimental

The Hastelloy N alloy used was subjected to cold-rolling and annealing treatment (for 6 min at 1177°C, water quenching). The overall chemical composition of Hastelloy N alloy is provided in Table1. The bulk materials were cut into several plates with a dimension of 10 × 10 × 1 mm3 using a wire electrode and polished with silicon carbide paper from grade 800 to 2000, by means of vibration polishing (including nanoalumina powder ∼20 nm), to obtain a mirror-like surface. These mechanically polished plates were subsequently prepared by electropolishing at 36 V for about 10 s in an aqueous solution of 50% H2SO4 and 40% glycerin under the temperature of 0°C, followed by ultrasonic cleaning using alcohol and deionized water sequentially. Several specimens were prepared as disks of 3 mm in diameter for TEM studies. Thin foils were obtained by mechanical thinning of the disks down to 100 μm, followed by twin-jet electropolishing, with an alcoholic solution of 5% HClO4.Table 1
Chemical composition of the Hastelloy N superalloy (wt. %).
ElementsNiMoCrFeMnSiAlTi/Cu/CoCHastelloy N alloy71.017.16.93.90.510.270.190.010.057The as-prepared samples were irradiated at room temperature (RT), on the 320 KV Highly Charged Ions Research Platform, at the Institute of Modern Physics, Chinese Academy of Sciences (HCIRP-IMP-CAS) [23]. The bulk samples were irradiated with 7 MeV Xe26+ ions to fluences of 0, 1.25 × 1014, 5.0 × 1014, 1.25 × 1015, and 2.5 × 1015 ions/cm2 equivalent to a maximum damage of 0, 0.5, 2, 5, and 10 dPa. For the thin-foil specimens, the irradiation of 1 MeV Xe20+ ions was chosen to make sure that the irradiation damage depth is suitable to the thickness of the thin area (less than 150 nm).The ion irradiation-induced microstructural evolution of the Hastelloy N alloy samples was characterized by a FEI Tecnai G2 F20 transmission electron microscope with an accelerating voltage of 200 kV. In order to reveal the correlation between the changes in mechanical properties and the microstructural evolution, the nanoindentation hardness of the un-irradiated and irradiated samples was measured on a G200 nanoindenter instrument. The experimental hardness was determined from the analysis of load-displacement (P–h) curves using the Oliver and Pharr method [24, 25]. In this study, five single indents were performed for each sample to ensure the reliability of the measurements.


## 3. Results and Discussion

### 3.1. TEM Characterization

Bright-field TEM micrographs of the un-irradiated Hastelloy N sample are shown in Figure1. Except for some dislocation lines, almost no other defects were observed in the un-irradiated sample. Figure 2 shows the TEM micrographs of Hastelloy N after ion irradiation to 0.5, 2, 5, and 10 dPa. Irradiating to a fluence equivalent to 0.5 dPa caused the disappearance of the original dislocations in the alloy and appearance of obvious black spots of about 5∼10 nm in the alloy. These black spots may be dislocation loops [26, 27], stacking fault tetrahedra (SFTs) [28], or defect clusters [29, 30]. When the irradiation fluence reached 2 dPa, corrugated-like structures began to gradually form, as shown in Figure 2(b), and this stage was the initial stage of the corrugated structure. At 5 dPa, black spots started to mix with the corrugated-like structures. Figures 2(c) and 2(d) show that the morphology of the corrugated-like structure became more and more obvious with the increase in ion fluence, which means that the corrugated structure is the final formation stage of the irradiated superalloy.Figure 1
TEM micrographs of un-irradiated Hastelloy N.Figure 2
TEM micrographs of Hastelloy N after ion irradiation to a fluence equivalent to (a) 0.5 dPa, (b) 2 dPa, (c) 5 dPa, and (d) 10 dPa.
(a)(b)(c)(d)To study the characteristics of the irradiated Hastelloy N alloy, the element mapping was performed to investigate the atomic migration and element distribution of the alloy after irradiation (Figure3). Figure 3(a) shows the TEM morphology results of the alloy viewed along the <110> axis, and Figures 3(b)–3(g) show the element mapping results of the corrugated-like structure. It can be seen that the ion irradiation has caused a huge change in the alloy and some elements are unevenly distributed in the alloy (e.g., Figure 3(b)), but the other elements are still uniformly distributed in the alloy (e.g., Figures 3(c)–3(g)), which indicate that Xe ion irradiation can make some specific elements to segregate. Considering that irradiation will only make the sample surface smoother, but not make the flat surface more uneven, in this respect, the change of elements may be due to the element segregation and may not be caused by the uneven surface.Figure 3
Morphology and element mapping results of Hastelloy N along the <110> belt axis after ion irradiation (10 dPa).
(a)(b)(c)(d)(e)(f)(g)Ion irradiation affects the alloy structure in many aspects. Here, the second phase of the alloy after ion irradiation was also studied, as shown in Figure4. Figure 4(a) shows that the second phase in the un-irradiated alloy has a face-centered cubic structure. Its lattice constant a was found to be 1.09 nm, which is consistent with the M6C, reported in Hastelloy N alloy by Oak Ridge National Laboratory (ORNL) [31]. However, the M6C can be rapidly amorphized with a lower fluence of 0.5 dPa (Figure 4(b)) and the amorphous state is still stable with a higher fluence (Figures 4(c)–4(e)). This is different from the MSRE results that the neutron irradiation would accelerate the formation of small M6C along the grain boundaries and would not change the structure of M6C.Figure 4
Phases in Hastelloy N alloy before and after ion irradiation: (a) 0 dPa, (b) 0.5 dPa, (c) 2 dPa, (d) 5 dPa, and (e) 10 dPa.
(a)(b)(c)(d)(e)It was reported in the literature that irradiation will cause the physical, chemical, and mechanical properties of the alloy grain boundaries to be reduced [32]. In this study, we analyzed the element distribution of the grain boundaries before and after ion irradiation (Figures 5 and 6). However, the results showed that there was no obvious phenomenon of element segregation along the grain boundaries.Figure 5
Distribution of elements near the grain boundary of the un-irradiated Hastelloy N.Figure 6
Distribution of elements near the grain boundary of Hastelloy N after irradiation (10 dPa).To investigate the elemental difference between the second phase and the matrix before and after irradiation, a line scan analysis was performed between the second phase and the matrix using transmission electron microscopy (Figures7 and 8). The results of the line scan showed that the second phase was enriched with Mo and Si, while the Ni, Fe, and Cr contents in the matrix were relatively high; as for Mn, it was relatively uniformly distributed in the second phase and the matrix.Figure 7
Distribution of elements in the second phase and Hastelloy N matrix.Figure 8
Distribution of elements in the second phase and Hastelloy N matrix after irradiation (10 dPa).


### 3.2. XRD

The alloy samples were analyzed by the grazing incidence X-ray diffraction (GIXRD) test method. Figure9 compares the strongest diffraction peaks of un-irradiated and irradiated samples. It can be seen that the positions of the diffraction peak of the sample have inconspicuous deflection at the experimental condition, which reveals that no obvious lattice distortions occur. This phenomenon may be true for the relatively lower irradiation fluence or less irradiation area; undoubtedly, the excellent irradiation-resistant property of the alloy may make the alloy sustain less irradiation damage at the experimental condition.Figure 9
Diffraction pattern of Hastelloy N alloy before and after Xe ion irradiation (enlarged view of the strongest matrix diffraction peak).


### 3.3. Nanoindentation

In this study, nanoindentation was performed on pristine and irradiated samples using a G200 nanoindenter system with a Berkovich-type indentation tip. 5 single indents were performed for each sample to ensure the reliability of the measurements. T hardness of the samples gradually decreased with increasing indentation depth as clearly observed in Figure10. In light of the presence of strain rate effects around the tip, the uncertainty in the indenter tip geometry, and other surface artifacts such as surface contamination films [32–34], the data in the depth lower than 200 nm were ignored. In the depth larger than 200 nm, the nanohardness decreased slightly with penetration depth. This can be explained by the Nix–Gao model [35] which predicts the hardness-depth profile to obey the following equation:(1)H=H01+h∗h,where H0 is the hardness at infinite depth and h∗ is a characteristic length which depends on the material and the shape of the indenter tip. h∗ can be considered as a constant for a given material and indenter. From this equation, it is quite clear that as the indentation depth (h) increases, the measured hardness value decreases gradually and finally approaches H0, which is consistent with results in Figure 10.Figure 10
Variation curve of the hardness of Hastelloy N alloy specimen before and after ion irradiation.The hardness of Hastelloy N alloy gradually increased with the increase in ion irradiation fluence. Compared to the un-irradiated sample, the irradiated Hastelloy N alloy samples showed an evident hardening phenomenon. In order to get more insight into the changing trend of the hardness with the irradiation fluence, the average hardness of the five groups of different samples before and after irradiation in the range of 200 to 1600 nm was calculated. The results are shown in Figure11. The hardness of the Hastelloy N alloy sample increased rapidly after being irradiated to 0.5 dPa, then increased slowly after irradiating from above 0.5 dPa to 5 dPa, and remained constant in the samples irradiated to an irradiation fluence equivalent to above 5 to 10 dPa. When the irradiation damage fluence was in the range of 5–10 dPa, the relation curve between the hardness and the depth basically coincided, indicating that there was a hardening saturation phenomenon. These results are similar to the results reported by Chen et al. [35].Figure 11
Trend of the average hardness value of Hastelloy N alloy samples in the depth range of 200–1600 nm before and after ion irradiation to different fluences.During the ion irradiation process, the bombardment of the surface of the material by high-energy particles forms a cascading collision near the surface of the material, which promotes the formation of a large number of supersaturated point defects (vacancies and interstitial atoms) inside the material [36]. It should be emphasized that most of the point defects disappear in a few picoseconds due to recombination and the remaining point defects will exist in the material in the form of monomers or fine clusters, respectively. The presence of these clusters has a pinning effect on the movement of dislocation lines inside the material resulting in an increase in the hardness of the material [37, 38]. Theoretically, as the irradiation fluence increases, the supersaturated point defects generated by the cascade of collisions inside the material gradually increase together with the density of the clusters resulting in the gradual increase in the hardness of the material. In this study, the hardness of the Hastelloy N alloy sample did not change after the ion irradiation from 5 dPa to 10 dPa. The reason may be that the Hastelloy N alloy did indeed produce more irradiation damage at 10 dPa than 5 dPa. However, as the concentration of point defects increases, the diffusion rate also increases accordingly, so that the probability of annihilation of vacancies and interstitial atoms increases resulting in a low number of final point defects remaining in the material irradiated to a fluence of 10 dPa. Hence, the defects are at the same level as in samples irradiated to 5 dPa so that the hardness of the material remains unchanged under these two irradiation fluences.It is worth noting that there was no formation of voids observed in the samples irradiated to different fluences at RT. In light of the statistical thermodynamic analysis in the literature [39], it is necessary to reach a critical vacancy concentration depending on the nucleation size of a void and the temperature for void nucleation. The no-void formation observed in this study may be attributed to the irradiation temperature, which is room temperature. Generally, the temperature of irradiation-induced void formation is about 0.3–0.55 Tm (Tm is the melting point) [40–44], and due to the fact that the Tm of Hastelloy N alloy is about 1400°C, the void is formed at about 400–700°C.


## 3.1. TEM Characterization

Bright-field TEM micrographs of the un-irradiated Hastelloy N sample are shown in Figure1. Except for some dislocation lines, almost no other defects were observed in the un-irradiated sample. Figure 2 shows the TEM micrographs of Hastelloy N after ion irradiation to 0.5, 2, 5, and 10 dPa. Irradiating to a fluence equivalent to 0.5 dPa caused the disappearance of the original dislocations in the alloy and appearance of obvious black spots of about 5∼10 nm in the alloy. These black spots may be dislocation loops [26, 27], stacking fault tetrahedra (SFTs) [28], or defect clusters [29, 30]. When the irradiation fluence reached 2 dPa, corrugated-like structures began to gradually form, as shown in Figure 2(b), and this stage was the initial stage of the corrugated structure. At 5 dPa, black spots started to mix with the corrugated-like structures. Figures 2(c) and 2(d) show that the morphology of the corrugated-like structure became more and more obvious with the increase in ion fluence, which means that the corrugated structure is the final formation stage of the irradiated superalloy.Figure 1
TEM micrographs of un-irradiated Hastelloy N.Figure 2
TEM micrographs of Hastelloy N after ion irradiation to a fluence equivalent to (a) 0.5 dPa, (b) 2 dPa, (c) 5 dPa, and (d) 10 dPa.
(a)(b)(c)(d)To study the characteristics of the irradiated Hastelloy N alloy, the element mapping was performed to investigate the atomic migration and element distribution of the alloy after irradiation (Figure3). Figure 3(a) shows the TEM morphology results of the alloy viewed along the <110> axis, and Figures 3(b)–3(g) show the element mapping results of the corrugated-like structure. It can be seen that the ion irradiation has caused a huge change in the alloy and some elements are unevenly distributed in the alloy (e.g., Figure 3(b)), but the other elements are still uniformly distributed in the alloy (e.g., Figures 3(c)–3(g)), which indicate that Xe ion irradiation can make some specific elements to segregate. Considering that irradiation will only make the sample surface smoother, but not make the flat surface more uneven, in this respect, the change of elements may be due to the element segregation and may not be caused by the uneven surface.Figure 3
Morphology and element mapping results of Hastelloy N along the <110> belt axis after ion irradiation (10 dPa).
(a)(b)(c)(d)(e)(f)(g)Ion irradiation affects the alloy structure in many aspects. Here, the second phase of the alloy after ion irradiation was also studied, as shown in Figure4. Figure 4(a) shows that the second phase in the un-irradiated alloy has a face-centered cubic structure. Its lattice constant a was found to be 1.09 nm, which is consistent with the M6C, reported in Hastelloy N alloy by Oak Ridge National Laboratory (ORNL) [31]. However, the M6C can be rapidly amorphized with a lower fluence of 0.5 dPa (Figure 4(b)) and the amorphous state is still stable with a higher fluence (Figures 4(c)–4(e)). This is different from the MSRE results that the neutron irradiation would accelerate the formation of small M6C along the grain boundaries and would not change the structure of M6C.Figure 4
Phases in Hastelloy N alloy before and after ion irradiation: (a) 0 dPa, (b) 0.5 dPa, (c) 2 dPa, (d) 5 dPa, and (e) 10 dPa.
(a)(b)(c)(d)(e)It was reported in the literature that irradiation will cause the physical, chemical, and mechanical properties of the alloy grain boundaries to be reduced [32]. In this study, we analyzed the element distribution of the grain boundaries before and after ion irradiation (Figures 5 and 6). However, the results showed that there was no obvious phenomenon of element segregation along the grain boundaries.Figure 5
Distribution of elements near the grain boundary of the un-irradiated Hastelloy N.Figure 6
Distribution of elements near the grain boundary of Hastelloy N after irradiation (10 dPa).To investigate the elemental difference between the second phase and the matrix before and after irradiation, a line scan analysis was performed between the second phase and the matrix using transmission electron microscopy (Figures7 and 8). The results of the line scan showed that the second phase was enriched with Mo and Si, while the Ni, Fe, and Cr contents in the matrix were relatively high; as for Mn, it was relatively uniformly distributed in the second phase and the matrix.Figure 7
Distribution of elements in the second phase and Hastelloy N matrix.Figure 8
Distribution of elements in the second phase and Hastelloy N matrix after irradiation (10 dPa).


## 3.2. XRD

The alloy samples were analyzed by the grazing incidence X-ray diffraction (GIXRD) test method. Figure9 compares the strongest diffraction peaks of un-irradiated and irradiated samples. It can be seen that the positions of the diffraction peak of the sample have inconspicuous deflection at the experimental condition, which reveals that no obvious lattice distortions occur. This phenomenon may be true for the relatively lower irradiation fluence or less irradiation area; undoubtedly, the excellent irradiation-resistant property of the alloy may make the alloy sustain less irradiation damage at the experimental condition.Figure 9
Diffraction pattern of Hastelloy N alloy before and after Xe ion irradiation (enlarged view of the strongest matrix diffraction peak).


## 3.3. Nanoindentation

In this study, nanoindentation was performed on pristine and irradiated samples using a G200 nanoindenter system with a Berkovich-type indentation tip. 5 single indents were performed for each sample to ensure the reliability of the measurements. T hardness of the samples gradually decreased with increasing indentation depth as clearly observed in Figure10. In light of the presence of strain rate effects around the tip, the uncertainty in the indenter tip geometry, and other surface artifacts such as surface contamination films [32–34], the data in the depth lower than 200 nm were ignored. In the depth larger than 200 nm, the nanohardness decreased slightly with penetration depth. This can be explained by the Nix–Gao model [35] which predicts the hardness-depth profile to obey the following equation:(1)H=H01+h∗h,where H0 is the hardness at infinite depth and h∗ is a characteristic length which depends on the material and the shape of the indenter tip. h∗ can be considered as a constant for a given material and indenter. From this equation, it is quite clear that as the indentation depth (h) increases, the measured hardness value decreases gradually and finally approaches H0, which is consistent with results in Figure 10.Figure 10
Variation curve of the hardness of Hastelloy N alloy specimen before and after ion irradiation.The hardness of Hastelloy N alloy gradually increased with the increase in ion irradiation fluence. Compared to the un-irradiated sample, the irradiated Hastelloy N alloy samples showed an evident hardening phenomenon. In order to get more insight into the changing trend of the hardness with the irradiation fluence, the average hardness of the five groups of different samples before and after irradiation in the range of 200 to 1600 nm was calculated. The results are shown in Figure11. The hardness of the Hastelloy N alloy sample increased rapidly after being irradiated to 0.5 dPa, then increased slowly after irradiating from above 0.5 dPa to 5 dPa, and remained constant in the samples irradiated to an irradiation fluence equivalent to above 5 to 10 dPa. When the irradiation damage fluence was in the range of 5–10 dPa, the relation curve between the hardness and the depth basically coincided, indicating that there was a hardening saturation phenomenon. These results are similar to the results reported by Chen et al. [35].Figure 11
Trend of the average hardness value of Hastelloy N alloy samples in the depth range of 200–1600 nm before and after ion irradiation to different fluences.During the ion irradiation process, the bombardment of the surface of the material by high-energy particles forms a cascading collision near the surface of the material, which promotes the formation of a large number of supersaturated point defects (vacancies and interstitial atoms) inside the material [36]. It should be emphasized that most of the point defects disappear in a few picoseconds due to recombination and the remaining point defects will exist in the material in the form of monomers or fine clusters, respectively. The presence of these clusters has a pinning effect on the movement of dislocation lines inside the material resulting in an increase in the hardness of the material [37, 38]. Theoretically, as the irradiation fluence increases, the supersaturated point defects generated by the cascade of collisions inside the material gradually increase together with the density of the clusters resulting in the gradual increase in the hardness of the material. In this study, the hardness of the Hastelloy N alloy sample did not change after the ion irradiation from 5 dPa to 10 dPa. The reason may be that the Hastelloy N alloy did indeed produce more irradiation damage at 10 dPa than 5 dPa. However, as the concentration of point defects increases, the diffusion rate also increases accordingly, so that the probability of annihilation of vacancies and interstitial atoms increases resulting in a low number of final point defects remaining in the material irradiated to a fluence of 10 dPa. Hence, the defects are at the same level as in samples irradiated to 5 dPa so that the hardness of the material remains unchanged under these two irradiation fluences.It is worth noting that there was no formation of voids observed in the samples irradiated to different fluences at RT. In light of the statistical thermodynamic analysis in the literature [39], it is necessary to reach a critical vacancy concentration depending on the nucleation size of a void and the temperature for void nucleation. The no-void formation observed in this study may be attributed to the irradiation temperature, which is room temperature. Generally, the temperature of irradiation-induced void formation is about 0.3–0.55 Tm (Tm is the melting point) [40–44], and due to the fact that the Tm of Hastelloy N alloy is about 1400°C, the void is formed at about 400–700°C.


## 4. Conclusion

This study was set out to investigate the Xe ion irradiation-induced microstructural evolution and hardening in Hastelloy N alloy:(1)
A large number of defects were caused by Xe ion irradiation inside the Hastelloy N alloy, such as black spots and corrugated structures. The corrugated structure formation was the final stage, and the element segregation was found in the corrugated structures.(2)
The second phase in Hastelloy N alloy is a face-centered cubic structure, with the lattice constanta = 1.09 nm, and it can be rapidly amorphized after being irradiated to 0.5 dPa. The second phase was enriched with Mo, Si, Ni, Fe, and Cr.(3)
Segregation of grain boundary elements was not observed in Hastelloy N alloy before and after irradiation in this study.(4)
Irradiation hardness was found in the Hastelloy N alloy, which increased with increase in irradiation fluence. The irradiation-induced hardening initiated at low ion fluence (0.5 dPa) and then saturated at a higher ion fluence in the range of 5 to 10 dPa.



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