Understanding the Impact of He Ion Irradiation on the Microstructure and Mechanical Properties of a Ni-Cr-Mo Alloy: A Comprehensive Study

doi:10.21203/rs.3.rs-4805274/v1

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Understanding the Impact of He Ion Irradiation on the Microstructure and Mechanical Properties of a Ni-Cr-Mo Alloy: A Comprehensive Study

https://doi.org/10.21203/rs.3.rs-4805274/v1

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The impact of helium ion irradiation on the evolution of microstructure and mechanical properties of a Ni-45%Cr-1.4%Mo (wt.%) alloy has been investigated. Analysis via transmission electron microscopy (TEM) of a pristine sample revealed an uneven distribution of alpha-Cr phase precipitates along with a few isolated planar arrangements of dislocations within the fcc γ-Ni matrix phase. A detailed study of irradiated samples established the presence of radiation-induced defect clusters, dislocation loops, network dislocations and He-stabilized voids as perceptible irradiation effects. Furthermore, examination of the irradiated sample revealed nearly complete dissolution of the native alpha-Cr precipitates, accompanied by a simultaneous increase in the Cr and Mo content of the matrix. X-ray diffraction (XRD) and positron annihilation spectroscopy (PAS) studies also confirmed lattice swelling as an effect of irradiation in this alloy, likely caused by helium-stabilized voids. A line profile analysis method which takes into account of planar defects, was employed on GIXRD (Grazing Incidence X-ray Diffraction) data to analyze changes in microstructural parameters such as domain size, microstrain, dislocation density, and lattice parameter as a function of irradiation dose. All of these irradiation-damage manifestations led to the increase in the hardness of the alloy in nano-indentation experiment. Overall, the findings of this work provide significant insights into the radiation-induced microstructural changes and their effect on the hardening behavior of the studied Ni-Cr-Mo alloy.

Nuclear materials

Ion-implantation

Radiation effects

Microstructure

Ni-based alloys and their modifications have emerged as promising structural materials, possessing excellent creep properties at elevated temperatures and adequate swelling resistance [1–5]. Their potential for use in nuclear reactors is undeniable; however, they are prone to substantial ductility loss under irradiation, which may severely limit their application. Despite this limitation, the invaluable characteristics of Ni-based alloys still position them as potential candidates for structural materials in heat exchangers or core components where the dose of irradiation is limited to a few dpa (displacement per atom) at most. The majority of published studies on Ni-based alloys concern Ni-Cr-Fe-Mo systems such as Inconel 625, Inconel 718, Hastelloy X, Nimonic PE16, etc. and Ni-Cr-Fe systems such as Alloy 600, Alloy 690, Inconel X-750, Alloy 903, Incoloy 800, Nimonic 80A, etc. [4, 6, 7]. Notably, a critical concern with Ni-based alloys in nuclear environments is the production of helium due to Ni (n, α) Fe transmutation reaction, which leads to He-embrittlement [8–10]. Wiffen et al. showed that voids can form in nickel-based alloys even at low-temperature irradiation up to 4–9 dpa [11].

Heavy ion irradiation can effectively simulate the effects of neutron irradiation on the microstructural integrity and mechanical properties of materials [12–18]. In order to analyze microstructural and mechanical property changes after irradiation, depth-sensitive techniques capable of probing a thin irradiated layer are required. Positron annihilation spectroscopy (PAS), grazing incidence X-ray diffraction (GIXRD), transmission electron microscopy (TEM), and nanoindentation are some of the techniques that can assess irradiation damage caused by heavy ions at shallow depths and provide valuable insights into the material's response to irradiation [19–21]. Nanoindentation has also proven to be one of the effective tools for evaluating the mechanical properties of irradiated materials with small sample volumes. In nano-indentation, the continuous stiffness measurement (CSM) mode gives the hardness of irradiated materials as a function of depth in the order of nm.

Ni-Cr-Mo alloy system is considered to be one of the most promising material for nuclear applications even at moderate damage levels. Gurovich et al. have investigated the structural evolution of the 42XNM (Ni-Cr-Mo) alloy during neutron irradiation up to 12 dpa at ~ 300°C and brought out the role of He bubble formation in microstructural changes due to irradiation [11]. This alloy system is being used in nuclear reactors as a control rod shell and is being considered for fuel element shell in Gen IV reactors [3, 5]. We have already reported irradiation response of a similar alloy (Ni-42 wt% Cr), upon proton irradiation to a very low dose of ~ 0.07 dpa [22]. To ensure the safe functioning of reactors, it is crucial to have a comprehensive understanding of the impact of He accumulation in structural materials. In the current study, we have selected a Ni-45Cr-1.4Mo (wt.%) alloy for investigation. This alloy has been exposed to He ion irradiation at room temperature to achieve a range of irradiation doses up to 33.9 dpa of peak damage. This study seeks to comprehensively elucidate the effects of very high doses of He ion on the microstructure and mechanical properties of Ni-45Cr-1.4Mo alloy employing a variety of depth-sensitive techniques, such as Grazing Incidence X-ray Diffraction (GIXRD), Transmission Electron Microscopy (TEM), Positron Annihilation Spectroscopy (PAS), and nanoindentation.

A sheet of Ni-45Cr-1.4Mo (wt.%) alloy was vacuum-sealed in a quartz tube and subjected to heat treatment at 1150°C for 10 minutes, followed by quenching in water. To prepare the samples for irradiation experiments, samples with dimensions of 10 mm × 10 mm × 1 mm were cut from the heat-treated sheet. The surfaces of these samples were prepared employing standard metallographic procedures followed by electropolishing with a solution of 10% perchloric acid and 90% methanol (v/v) at 20 V and − 40 ^oC.

The samples were irradiated using He ion beam from a 6.45 GHz electron cyclotron resonance (ECR) ion source available at Variable Energy Cyclotron Centre, Kolkata. Three samples were exposed to He ion beam of energy 65 keV under 10^− 6 mbar vacuum to achieve total doses of 3×10¹⁶, 2×10¹⁷, and 7×10¹⁷ ions/cm², respectively, at an implantation rate of 3.5×10¹² ions/cm²-sec at room temperature.

Since the damage due to ion irradiation is limited to few hundreds of nanometers from the exposed surface of the sample, these were characterized by Grazing Incidence X-ray Diffraction (GIXRD) technique with a PANanalytical system employing CuKα X-rays of 1.545 Å wavelength. The diffraction profiles of unirradiated and irradiated samples were obtained in the 2θ range of 40–95̊ with a 0.02̊ step size.

For further insight into microstructural changes, un-irradiated and irradiated samples was examined using a 300 kV JEOL 3010 TEM. TEM samples were prepared by mechanically grinding from the unirradiated side to get thin foil (80–100 µm) and subsequently 3 mm diameter discs were punched from these foils. The discs were then thinned to the point of perforation using the twin jet electropolishing process, in which Lacomit peelable varnish was applied to the irradiated side to prevent its polishing and ensure thinning only from the unirradiated side. This procedure ensured that the irradiated zone of the samples is transparent to electrons. Electropolishing was performed at 20 V and − 45 ºC using an electrolyte consisting of 10% perchloric acid in methanol. In addition to the imaging of microstructural features, EDS (energy dispersive spectroscopy) analysis of the composition of second-phase precipitates and that of the matrix phase was also performed.

The samples were also subjected to depth-dependent Doppler broadening measurements using a slow positron beam of variable energy. All the measurements were performed at room temperature using incident positron energies ranging from 0.2 to 17.2 keV with a high purity Ge detector having an energy resolution of 1.4 keV at 514 keV of ⁸⁵Sr. The Doppler broadening spectrum at different implantation energies or implantation depths (total counts ~ 5x10⁵) was characterized by two line-shape parameters, viz., S- and W-parameters. The S-parameter was evaluated by taking the ratio of counts in the low momentum region (511.000 ± 0.684 keV) to the total counts under the annihilation peak, whereas the W-parameter was evaluated by taking the ratio of counts under a 4.564 keV window centered at (511.000 ± 13.692 keV) in the high momentum region to the total area of the annihilation peak.

The hardness of unirradiated and irradiated samples was evaluated using the nanoindentation technique employing a Berkovich diamond indenter (UNHT, CSM, Switzerland). The continuous stiffness measurement method was used to measure the hardness of the irradiated samples as a function of depth. The indentation was carried out up to a depth of 450 nm from the irradiated surface. At least 10 indentations were performed at each depth and the average of these measurements was considered for interpretation. The spacing between two successive indents was kept at 50 µm to prevent interference with plastic zones formed by adjacent indents.

The nominal composition of the Ni-Cr-Mo alloy obtained by chemical analysis is given in Table 1. A widely known Monte-Carlo simulation-based software SRIM [23] has been used to estimate the depth profile of the distribution of implanted He ions and vacancies formed during irradiation and the same is shown in Fig. 1. The irradiation damage has been evaluated in terms of dpa (displacements per atom) using the following formula:

$$\:dpa=\left({{\nu\:}}_{\text{N}\text{R}\text{T}}\:\times\:{10}^{8}\times\:D\right)/{N}_{o}$$

where D signifies the irradiation fluence (ions/cm²), N_o denotes the atom density of the target material in atoms/cm³, and ν_NRT stands for the number of displaced atoms per Å per ion determined by the method outlined by Stoller et al. using Eq. 2 [24].

$$\:{{\nu\:}}_{\text{N}\text{R}\text{T}}=(0.8\times\:{T}_{dam})/(2\times\:{E}_{d})$$

where $\:{T}_{dam}$ represents the total damage energy calculated from the phonon contribution to the ions and recoils. The estimated peak damage for the samples irradiated with fluences 3×10¹⁶, 2×10¹⁷, and 7×10¹⁷ ions/cm² is 1.5, 9.6, and 33.9 dpa, respectively. The number density of He atoms implanted during the irradiation experiments for a fluence equivalent to a dose of 7×10¹⁷ ions/cm² is 2x10²² ions/cm³. Figure 1 shows that irradiation damage is extended up to a depth of ~ 350 nm with a peak at around 200 nm. The peak becomes more pronounced with an increase in irradiation dose.

In order to probe the damage region which is confined up to a depth of 350 nm from the surface using GIXRD, the effective penetration depth ($\:{x}_{\alpha\:}$) of an X-ray beam was calculated using Eq. (3).

$$\:{x}_{\alpha\:}=\frac{-\text{l}\text{n}(1-G)}{\mu\:(\frac{1}{Sin\alpha\:}+\frac{1}{Sin(2\theta\:-\alpha\:)})}$$

where G is the fraction of the incident beam that is absorbed up to a depth of $\:{x}_{\alpha\:}$, µ is the linear attenuation coefficient of the material,$\:\:\alpha\:$ is the grazing angle, and 2$\:\theta\:$ is the diffraction angle. In this study, the fraction of unabsorbed X-rays is taken as 1/e and the diffraction angle of the most intense peak corresponding to the (111) plane is taken as 2$\:\theta\:$. The mass attenuation coefficient of Ni-45Cr-1.4Mo (wt.%) alloy is 110 cm²/g for CuKα X-rays. Based on Eq. 3, the grazing angle of 1.5º was found to be most appropriate to investigate the irradiation-induced microstructural changes up to a depth of 350 nm and therefore GIXRD experiments were carried out at 1.5º angle of incidence.

3.1. Microstructure prior to irradiation

The microstructure of the pristine alloy is presented in Fig. 2. An overview of the gross microstructural features of the alloy can be obtained from Fig. 2a. The irregularly shaped precipitates seen in Fig. 2a correspond to alpha-Cr. A non-uniform distribution of such precipitates was noticed within the matrix phase with grain boundaries acting as the preferred site for nucleation. The volume fraction of the precipitates was found to be small with their sizes varying between 100–300 nm. The average composition of the precipitates was estimated to be Cr-10.3Ni-3.5Mo (wt.%) through EDS analysis. The matrix phase showed an average composition of Ni-43.4Cr-1.1 Mo (wt.%). The micrograph shown in Fig. 2b was obtained under two-beam diffracting condition. It depicts the planar arrangement of dislocations within the matrix phase without any sort of dislocation interaction or entanglement.

3.2. X-Ray Diffraction Line Profile analysis

Figure 3 shows a typical GIXRD profile for the un-irradiated and irradiated NiCrMo alloy samples. The peaks observed in the diffractogram before and after irradiation are indexed to the fcc γ-Ni phase. It can be concluded from these diffractograms that no new phases have formed after irradiation. Figure 4 shows the enlarged view of the strongest intensity (111) peak before and after irradiation. This figure suggests that this peak shifts gradually to a lower 2θ value with an increase of irradiation dose, indicating an increase in the lattice spacing with the increase in the irradiation dose. Indeed, other peaks were also observed to shift to lower diffraction angle with an increase in irradiation dose. Therefore, it could be inferred that the He ion irradiation caused lattice expansion in the irradiated samples.

In addition to the shift in Bragg’s peak, width of the peak also changes with irradiation. The broadening of Bragg peaks is mainly due to a change in domain size and microstrain after accounting for the broadening due to instrumental parameters. Line profile analysis of the diffractograms can bring out important microstructural parameters such as the root mean square microstrain and the average domain size, which help in understanding the irradiation effect on these samples. In this study, Williamson–Hall method based on whole powder pattern fitting technique has been employed.

3.2.1 Modified Williamson Hall Method

The method assumes that the contributions to the broadening of Bragg peaks due to the domain size and the microstrain effects are linearly superimposed and independent of each other. Under these assumptions, Williamson and Hall (WH) relationship between the broadening (β), the volume weighted average domain size (D), and the maximum microstrain ($\:\epsilon\:$) is given by Eq. 4.

$$\:\frac{\beta\:\:Cos\:\theta\:}{\lambda\:}=\:\frac{1}{D}+2\epsilon\:\:\left(\frac{2\:Sin\:\theta\:}{\lambda\:}\right)$$

Thus, D_v and ε can be determined from the linear plot of $\:\frac{\beta\:\:Cos\:\theta\:}{\lambda\:}\:\left({\Delta\:}\text{K}\right)$ vs$\:\frac{2\:Sin\:\theta\:}{\lambda\:}\left(\text{K}\right)$ as shown in Fig. 5 for both unirradiated and irradiated samples. It is apparent that the data points in the WH plots are scattered and therefore, not suitable for linear fitting with acceptable accuracy. In the WH method, it is assumed that the micro-strain is isotropic along the crystallographic axes. The highly scattered data points in the WH plots reveal the existence of anisotropic broadening in the samples. The scattering tendency of the data points is also observed to increase with the increase in dose, suggesting that the micro-strain anisotropy along the different crystallographic planes increases with dose. To take care of such anisotropies, Ungár et. al.[25–27] suggested a modified version of the W-H method that is described by the Eq. 5.

$$\:{\Delta\:}\text{K}=\:\frac{1}{D}+\:{\left(\frac{\pi\:{b}^{2}\rho\:}{2B}\right)}^{1/2}K{C}^{1/2}$$

here b is the magnitude of the Burgers vector, $\:\rho\:$ is the dislocation line density, B is a constant considered to be 10 over a wide range of dislocation distribution and C is the average contrast factor of the dislocations based on their positions relative to the diffraction vector [25–27]. The C factor for a cubic crystal is a linear function of the fourth-order invariant of the hkl indices of the various reflections and depends on the Burgers vector, line vectors characterizing the dislocations, the elastic anisotropy constant, A_i and the ratio of C₁₂/C₄₄, where C₁₁, C₁₂ and C₄₄ are the elastic constants. Using C₄₄ = 1.25 and A_i = 2.54, C is calculated employing the procedure given in ref [28]. However, the C factor in the modified Williamson-Hall method (Eq. 5) incorporates only the anisotropy component of strain. According to Warren [29], if a sample contains stacking faults or twin boundaries, or both, then apparent domain size is less than the actual domain size. Therefore, the presence of planar defects also influences the broadening of the diffraction peaks. Ungár et al.[30] included the effect of twinning and stacking faults in the modified W-H plot. With this change, Eq. (5) can be rewritten as

$$\:\:\:\:{\Delta\:}\text{K}-{\upvarsigma\:}\text{W}\left(\text{k}\right)=\:\frac{1}{D}+\:{\left(\frac{\pi\:{b}^{2}\rho\:}{2B}\right)}^{1/2}K{C}^{1/2}$$

where the parameter, $\:{\upvarsigma\:}$ represents the density of twin boundaries and deformation faults and it is set as a fitting parameter. The value of the parameter W(k) for fcc unit cell is determined by Ungár et al. [28] for different peaks. Figure 6 shows linear fitting of data points in the plot of$\:\:{\Delta\:}\text{K}-{\upvarsigma\:}\text{W}\left(\text{k}\right)$ vs. $\:K{C}^{1/2}$ for all the samples. The domain size and dislocation density were evaluated from these plots using Eq. (6) and their variation as a function of irradiation dose is given in Table 2.

The domain size decreases significantly from 356 Å to 136 Å when irradiated at highest dose of 33.9 dpa. The domain size is seen to decrease with increase in irradiation dose and the rate of its decrease is ~ 74 Å /dpa, 42.5 Å /dpa, and 5.7 Å /dpa at doses of 1.5 dpa, 9.6 dpa, and 33.9 dpa, respectively. This indicates that the maximum reduction in domain size occurred during the initial phase of irradiation and approaches saturation as the dose of irradiation increases.

3.3. Transmission electron microscopy of irradiated samples

The irradiated samples were investigated in TEM to reveal the geometric configurations attained by the clusters of point defects generated by irradiation. The bright-field TEM micrographs of Figs. 7(a-b) and 7(c-d) show the defect microstructures of irradiated samples recorded under two-beam diffraction imaging conditions at doses of 9.6 dpa and 33.9 dpa, respectively. Two different diffraction g-vectors were used to image the defect structure. A network of dislocation loops could be observed in both the irradiated samples very clearly besides small, black dot-like feature which represents defect clusters/small dislocation loops. No discernible difference could be noticed in terms of the density of the loops in the two samples (9.6 vis-à-vis 33.9 dpa) although more of dot-like feature was observed in the 9.6 dpa sample in contrast to 33.9 dpa when imaged with 200 diffraction vector (Fig. 7(a) vs. 7(c)). This suggested towards an increase in the size of the defect clusters/small dislocation loops with irradiation dose, leading to the attainment of a more distinct loop-like character. Although the nature of the loops has not been determined in the present work, it is expected to be similar to that exhibited by the microstructure of any irradiated fcc material with predominance of faulted (Frank) loops (1/3 < 111 > Burgers vector) over perfect loops (1/2 < 110 > Burgers vector) owing to lower stacking fault energy.

It is to be noted that all these defect features are an outcome of the irradiation process alone, as the unirradiated sample imaged under similar diffracting condition did not exhibit any of these features except for a few isolated planar dislocations (Fig. 2(b)). Another major irradiation effect on the microstructure of the alloy, which became apparent through TEM examination, was the dissolution of the native alpha-Cr precipitates. The alpha-Cr precipitates present in the unirradiated matrix (Fig. 2(a)) were hardly visible after irradiation; Only their leftover vestiges were noticed as shown encircled in Fig. 8. The dissolution of alpha-Cr precipitates got accompanied by a corresponding increase in the Cr- and Mo-content of the matrix phase to 44.7 wt.% and 1.8 wt.%, respectively from that of 43.4 wt.% and 1.1 wt.% in the unirradiated condition. Elemental analysis carried out on the precipitate remnants showed composition very close to that of the irradiated matrix namely Ni-47Cr-2.3Mo (wt.%), thereby proving their dissolution. However, the re-precipitation of alpha-Cr phase was not observed. It should be noted that pristine samples were heat treated in such a way so as to minimize the precipitate volume fraction during processing. Thus, the dissolution of small fraction of pre-existing precipitate may not be considered detrimental for application purpose.

Focus-series experiments were performed on the highest dose (33.9 dpa) sample to examine the generation of voids on account of irradiation. In such experiments, the presence of voids gets evinced as black dots surrounded by a white fringe in over focused image (under positive defocus of the objective lens) and as white dots surrounded by a dark fringe in under focused image (under negative defocus of the objective lens). The features encircled in Figs. 9(a) and (b), present near the sample’s edge, showed this sort of contrast reversal, which established unequivocally the presence of voids in the irradiated sample. The size of the voids was observed to be in the range of ~ 2 to 6 nm. He has a negligible solubility in metals and have the ability to stabilize vacancy clusters thereby leading to the formation of voids [10]. In Ni-based alloys, bubble formation due to He ion implantation has been studied by several authors [31–35]. Van et al. studied the growth of helium bubbles and the formation of blisters during low energy He ion irradiation in nickel and amorphous Metglass alloy (Fe₄₀Ni₃₈Mo₄B₈) in the temperature range of -73 °С to 327 °С and dose range of 5x10²⁰ He/m²-1x10²² He/m². In crystalline nickel even at 0 °С, the bubbles grew to a saturation value of mean diameter equal to 2.5 nm [36].

3.4. Positron annihilation spectroscopy study

The variation of S-parameter as a function of mean positron implantation depth (S(Z) profiles) from the sample surface is presented in Fig. 10 for both unirradiated and irradiated samples. The experimental S-Z profiles could be analyzed using the variable energy positron fit (VEPFIT) software [37–39]. The solid line shows the fitting of VEPFIT program. It can be seen that the S-parameter of the unirradiated sample increase sharply close to the surface, before mean implantation depth of 50 nm. This behavior is generally observed in metallic samples and is attributed to the back diffusion of positrons towards the surface, which results in the formation of a positronium-like state that has higher S-parameter values [40]. The S-parameter values of the irradiated samples are higher than that of unirradiated samples throughout the depth and the S-parameter value increases as the dose increases. The S-W plot is analyzed to identify the defects that have been produced by ion implantation. The typical S-W plots for the sample irradiated to the highest dose of 33.9 dpa is shown in Fig. 11.

3.5. Nano indentation

The hardness H of the samples before and after irradiation were evaluated using nano-indentation. The inset in Fig. 12(a) shows the average nanoindentation hardness as a function of depth. The hardness of all irradiated samples was observed to be higher than that of the unirradiated sample together with a progressive increase in hardness with irradiation dose. There is a decrease in the hardness of all the samples with increasing indentation depth in irradiated as well in unirradiated samples. This is attributed to the ‘indentation size effect’ (ISE) which can be described by Nix–Gao model based on geometrically necessary dislocations [41, 42] using Eq. 7.

$$\:\frac{H}{{H}_{0}}=\sqrt{1+\frac{{h}^{*}}{h}}$$

where H is the measured hardness corresponding to indentation depth h, H₀ is the hardness corresponding to infinite depth which is also known as the bulk equivalent hardness or ISE free hardness, $\:{h}^{*}$ is the characteristic constant of indention depth which depends on the indenter shape and shear modulus. In evaluating the hardness of shallow-depth low-energy ion-irradiated materials, the estimated hardness often turns out to be a convoluted hardness having contributions from both the hard-irradiated surface and the underneath soft unirradiated substrate. An approach developed by Kasada et al. [43] allows one to evaluate the hardness of irradiated region without interference from the unirradiated substrate.

A fitting of H² against 1/h of the hardness data was carried out for depths beyond 125 nm and the same is shown in the Fig. 12. For all irradiated samples, the plot of H² versus 1/h exhibits an inflection at a critical indentation depth (h_c) and this inflection point is more pronounced at higher dose of 33.9 dpa and can be seen at around 300 nm. The hardness of the irradiated samples, H₀, and h* evaluated by fitting the data in the range of 125 nm < h < hc are given in Table 3. It is observed that the value of H₀ increases and h* decreases with increase in irradiation dose.

In order to understand the effect of He ion irradiation on the mechanical properties, the change in the average nanohardness value (∆H), defined as ∆H = H_irradiated - H_unirradiated was evaluated and shown in Table 3. The hardening rates at 1.5 dpa, 9.6 dpa and 33.9 dpa were found to be 347 MPa/dpa, 195 MPa/dpa and 25 MPa/dpa, respectively. This shows that the hardening rate is higher at lower irradiation doses. The irradiation hardening can be fairly assumed to be mainly due to a change in dislocation density and therefore dislocation density can be estimated using Eq. 8.

$$\:\rho\:={\left(\frac{{H}_{0}}{MC\alpha\:Gb}\right)}^{2}$$

where M is the Taylor factor (3.06 for cubic lattice) and α (= 0.5) is defect barrier strength (dependent on dislocation structures) [44, 45]. G represents the shear modulus (87 GPa, [1]), b represents the magnitude of the Burgers vector, and C represents the Tabor's factor. For metals C = 3 can be considered [44, 45]. Using Eq. (8), the dislocation density was calculated for each sample and given in Table 3. It is observed that the dislocation density estimated from nano-hardness is higher than that estimated from XRD analysis.

Different characterization techniques have been used in this study to understand the role of irradiation induced defects on microstructural and mechanical property changes of Ni-Cr alloy owing to He ion irradiation. The SRIM profile (Fig. 1) shows that the damage region extends up to a depth of approximately 350 nm, whereas the S-Z profiles of irradiated samples suggested the damage to be extended beyond 350 nm. This is due to the thermal effects that SRIM ignores. The vacancies introduced by He ion irradiation can easily diffuse beyond the irradiation zone as the vacancies in Ni have been reported to be mobile even at room temperature [51]. It can be seen from Fig. 10 that the S parameter value is higher in the irradiated samples in the region with positron penetration depth of 100 to 300 nm indicating that the defect density in this region is higher than that in the unirradiated sample or in bulk region (beyond 350 nm) as expected. As per SRIM estimation, peaking in S-Z profile should occur around 200 to 300 nm, but S- parameter is observed to continuously decrease beyond 100 nm. This can be understood considering the following. As per SRIM implantation profile, nearly 87% of He ions stopped in the range of 100 nm to 350 nm. Authors and other researchers have established that He ions have a strong affinity for vacancies, vacancy clusters and can form ion-vacancy complexes which have lower S-parameter values than that of a vacancy or a vacancy cluster [46–49]. Hence, the damage region having higher deposition of He ions may show lower S-parameter value. This inference is complemented by the S-W plots of the irradiated samples, a typical representation of which is shown in Fig. 11 for the 33.9 dpa sample. It shows two different slopes corresponding to the depth below 100 nm (Layer I, higher S- values) and between 100 to 300 nm (Layer II, lower S- values). Different slopes suggest the presence of different types of defects in Layer I and Layer II. In addition, it can be noticed that the slope of the line in layer II is higher compared to the slope of the line in layer I, indicating that layer II might have a smaller size of defect compared to layer I, which can be attributed to the formation of He ion-vacancy complexes in layer II. The effect of the size of defects on the slope change in S-W parameter plot has also been brought to notice by other researchers [50, 51]. Authors have shown that these ion-complexes have lower S-parameter values [48].

The formation of defect clusters/dislocation loops and of voids in the irradiated microstructure has been brought to light by TEM investigation. The extent of the irradiated region which got probed in TEM by virtue of the sample preparation technique is ~ 50 to 100 nm from the irradiated surface. This is in view of the fact that the approximate film thickness required for electron transparency is also of the same order. TEM analysis could bring out extensive dislocation loop formation within this depth in comparison of the density of voids. The explanation for such observation can be readily obtained if the PAS results corresponding to the same depth (depicted as Layer I) are put in perspective. According to PAS analysis, Layer I comprised mostly of vacancy clusters and less of He-vacancy complexes whereas layer II has dominating He-vacancy complexes. As collapse of point defect clusters (be it interstitial or vacancy) leads to the formation of dislocation loops, hence predominance of vacancy clusters in this layer justified the observed higher dislocation loops within it. At the same time, it is also a known fact that He atoms help stabilize voids (a three-dimensional configuration of vacancy clusters) by preventing the collapse of the void nuclei into vacancy loops. Thus, the lower concentration of He-vacancy complexes in Layer I could clearly corroborate the observed lower density of voids in the 50–100 nm irradiated zone.

Analysis of XRD data of irradiated samples indicated an increase in the lattice parameter of these samples. The lattice parameter expansion (Δa) corresponding to the highest irradiation dose of 33.9 dpa calculated from the peak shift is about 0.14% with respect to the unirradiated sample. This lattice expansion may be attributed to the contribution from beam heating i.e. due to electronic energy loss (ionization, electronic excitation or phonon interaction of target atoms) as well as due to nuclear energy loss (displacement effects) with the later playing major role as the irradiation temperature did not exceed room temperature [52, 53]. TEM study clearly established formation of voids as an irradiation induced effect in these samples, which may contribute to lattice swelling. Hofmann et al. observed lattice swelling in tungsten alloy due to helium-filled vacancy defects after helium implantation [54]. Wang et al. also showed lattice parameter expansion of about 0.2% in high-entropy carbide after Zr ion irradiation [55].

Further, line-profile analysis of XRD data showed a decrease in domain size and an increase in microstrain and dislocation density. A complementary observation is seen in S-parameter values where the broad hump in the S-Z profile of irradiated samples became more prominent as the irradiation dose increased suggesting an increase in the inhomogeneous distribution of irradiation-induced open volume defects along the depth. At lower doses, a large change in these parameters could be noticed from Table 2, whereas the magnitude of the change became smaller with increase in dose suggesting the onset of a saturation-type behavior at higher irradiation doses. The decrease in domain size and increase in dislocation density, microstrain and S-parameter with increase in irradiation dose can be attributed to the formation of point defects, defect clusters, and dislocation loops due to irradiation. Because the damage caused by ion irradiation is primarily confined within 350 nm of the sample surface, a high concentration of defects remains localized within this region, promoting defect migration. Generally, the reactions that the point defects usually encounter while migration are: (1) annihilation by mutual recombination of self-interstitial atoms and vacancies; (2) elimination at pre-existing as well as newly created defect sinks, such as grain boundaries and dislocation loops; (3) accumulation (of similar type) into defect clusters; and (4) collapse to form dislocation loops. Thus, although the concentration of point defects keeps on increasing with increasing dose, these point defects are also eliminated during irradiation. Consequently, an equilibrium between formation and elimination of point defects sets in at a higher dose of irradiation. Thus, although lower domain size, higher microstrain, higher dislocation density, higher S-parameter and higher hardness values at higher doses indicated an increase in defect density, the decrease in the corresponding rate of change in these parameters with dose suggested the defect density to increase at a slower rate on account of their annihilation by cascade overlap, mutual recombination, and annihilation at sinks.

It is expected that the microstructural changes induced by He ion irradiation would also affect the mechanical properties. A progressive increase could be noticed in the hardness of the alloy with irradiation dose as seen in the inset of Fig. 12. Such hardening effect of irradiation can be attributed to the production of point defects, point defect clusters, dislocation loops, He-vacancy complexes and voids, which hinder dislocation glide. The hardness values were used to estimate dislocation density and the results are shown in Table 3. The hardness values were found to increase with irradiation, a trend similar to that observed with dislocation density estimated from the XRD data. However, the estimated dislocation density from the hardness data were observed to be higher in comparison to that estimated from XRD data. This difference in the estimated values of dislocation density from XRD and nanohardness measurements is expected since microstructural features apart from network dislocations and dislocation loops, as mentioned above, also contribute in increasing the hardness. In addition, the contribution of solid solution hardening in this hardness increment due to the rise in Cr and Mo content of the matrix phase on account of irradiation induced ‘recoil dissolution’ of native alpha-Cr precipitates cannot also be ignored as well.

The impact of He ion irradiation on the evolution of microstructure and mechanical properties of a Ni-45%Cr-1.4%Mo (wt.%) alloy was studied employing depth sensitive defect characterization techniques. TEM analysis of unirradiated sample showed a non-uniform distribution of irregularly shaped precipitates of alpha-Cr phase (sizes varying between 100–300 nm) in fcc γ-Ni matrix with grain boundaries as the preferential nucleation site. A few isolated planar arrangements of dislocations within the matrix phase were also noticed. Examination of irradiated samples brought out the existence of radiation induced defect clusters, dislocation loops, network dislocations, ion-vacancy complexes and to a limited extent He stabilized voids. In addition, nearly complete dissolution of the native alpha-Cr precipitates accompanied with concomitant increase in Cr and Mo content of the matrix upon irradiation were confirmed through TEM-EDS analysis. The size of voids was observed to be in the range of ~ 2 to 6 nm. An increase in the size of the defect clusters/small dislocation loops with irradiation dose was also noticed. The density of irradiation induced defects showed a tendency of saturation with increase in irradiation dose. SRIM calculation indicated the damage depth to extend up to 350 nm from the irradiated surface which was complemented by nanohardness measurement study which showed a clear increase in hardness up to a depth of ~ 300 nm from the surface. GIXRD and PAS studies also established lattice swelling as an irradiation effect in this alloy probably due to He-stabilized voids. Other microstructural parameters such as domain size, microstrain and dislocation density were estimated by performing line profile analysis of GIXRD data and these microstructural observations were used to rationalized the irradiation induced hardening phenomenon.

Acknowledgements

The authors extend their sincere gratitude to the Director, Variable Energy Cyclotron Centre (VECC), India, and to Prof. P.Y. Nabhiraj, Head, ECR Ion Source Facilities Section of this centre, for generously providing access to the irradiation facility. The authors also express their sincere appreciation to the Head, Department of Metallurgical Engineering and Materials Science, IIT Bombay for their assistance in facilitating XRD measurements for this study. Furthermore, the authors wish to convey their gratitude to Dr. R. Tewari, Associate Director, Materials Group, Bhabha Atomic Research Centre, India for his unwavering interest and invaluable insights that significantly enriched the progression of this study. This study has been funded by the Department of Atomic Energy, Government of India.

Author contributions: Not applicable

Conflicts of interest or competing interests: There is no conflicts of interest of any author of this manuscript.

Data and code availability: Experimental data will be made available on demand by the corresponding author.

Supplementary information: Not applicable

Ethical approval: Not applicable

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Understanding the Impact of He Ion Irradiation on the Microstructure and Mechanical Properties of a Ni-Cr-Mo Alloy: A Comprehensive Study

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Abstract

Figures

1. INTRODUCTION

2. EXPERIMENTAL

3. RESULTS

3.1. Microstructure prior to irradiation

3.2. X-Ray Diffraction Line Profile analysis

3.2.1 Modified Williamson Hall Method

3.3. Transmission electron microscopy of irradiated samples

3.4. Positron annihilation spectroscopy study

3.5. Nano indentation

4. DISCUSSION

CONCLUSIONS

Declarations

References

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