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# Advanced time-resolved characterization of Stress Assisted Grain Boundary Oxidation of 718 Ni superalloy

To cite this article: P D Resende et al 2022 IOP Conf. Ser.: Mater. Sci. Eng. 1249 012046

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# Advanced time-resolved characterization of Stress Assisted Grain Boundary Oxidation of 718 Ni superalloy

P D Resende<fnref n="1" />,<fnref n="2" />, N Viganò<fnref n="1" />, D Bardel<fnref n="3" />, J Réthoré<fnref n="4" />, and W Ludwig<fnref n="1" />,<fnref n="2" />

E-mail: pedro.damas-resende@esrf.fr

**Abstract.** Nickel superalloys are used for harsh condition application cases as they have high chemo-thermomechanical stability. However, they can suffer from embrittlement due to stress corrosion cracking. This effect is difficult to observe as it can take place at long time scales. Here we propose a feasible experiment to study stress assisted grain boundary oxidation, a phenomenon that has similar mechanisms involved that can take place in laboratory compatible time scales. We observed the event using phase and diffraction contrast tomography while applying monotonic loading of the sample at 650 °C. This initial analysis shows that the experimental setup is a good candidate for the study of such degradation mechanism.

## 1. Introduction

Nickel superalloys (SA) are widely known for withstanding harsh environments while maintaining good chemo-mechanical stability, key properties for applications in nuclear reactors, turbines and others [1–3]. The IN718 alloy is mainly composed by Ni-Cr-Fe with additions of elements like Nb and Mo. This alloy is constituted by a γ matrix that is hardenable by γ'/γ'' precipitates. The crystal structure of the matrix is an FCC solid solution, while the γ' has a ordered cubic structure of type L1₂ with approximate composition of Ni₃(Al,Ti) and γ'' has an ordered BCT structure of type D0₂₂ [2,4] and approximate composition of Ni₃Nb. Both phases are coherent with the matrix, while γ'' is the main hardening precipitate [4,5]. The occurrence of δ phase is also possible with an orthorhombic structure [6]. MC carbides can also be found in the microstructures [4].

Regarding the application in nuclear power plants, premature failure of components made with 718 alloy was found in literature due to the Stress Corrosion Cracking (SCC) phenomenon [1]. The main mechanism of failure in reactor conditions (~350 °C, mechanical loading, and irradiation) is Stress Corrosion Cracking (SCC). Leonard *et al.* [1] studied leaf springs that were subjected to a commercial reactor exposure for 2 years. In their study it was found that the samples showed extensive grain boundary (GB) cracking in spite of its low content of δ phase. They also found that cracks tend to propagate preferentially along random high angle GBs.

SCC is caused by a coupling between oxygen diffusion and mechanical loading [7,8] that leads to embrittlement. This susceptibility of the material to embrittlement in such application is related to crystallographic orientation relationships within the grain structure of the material [9,10]. In the literature the most common type of study relies on observation of *post-mortem* samples, and/or surface-limited observations. There is a clear need to complement this type of surface analysis with in-

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<footnote n="1">¹ European Synchrotron Radiation Facility (ESRF), Grenoble, France</footnote>

<footnote n="2">² MATEIS, INSA Lyon, Université de Lyon, UMR5510 CNRS, France</footnote>

<footnote n="3">³ Framatome, 2 rue Professeur Jean Bernard, 69007 Lyon, France</footnote>

<footnote n="4">⁴ GeM - Génie Civil et Mécanique - UMR 6183 CNRS - Centrale de Nantes - Université de Nantes</footnote>
situ bulk observations in order to gain a better understanding of the kinetics and to account for the three-dimensional character of the damage process.

Observing SCC *in-situ* is a difficult task as the phenomena can take a long time and never reach failure [11]. Stress Assisted Grain Boundary Oxidation (SAGBO) observation arises as a sound method to mimic SCC within laboratory consistent time-scales. SAGBO occurs at conditions with combined high temperature, oxygen partial pressure and mechanical loading over a limited range of strain rates [7,8,12]. Andrieu, Max and Viguier showed in their work [8] that at high strain rates (~10⁻² s⁻¹) intragranular cracking and embrittlement is not observed even at higher temperatures in air atmosphere. For lower strain rates (< 10⁻⁴ s⁻¹), it is possible to see a clear transition from the transgranular fracture regime, accompanied by Portevin-Le Chatelier effect, to an intergranular failure, depending on the temperature. Therefore, a condition to carry out SAGBO observation can be defined around 650 °C and strain rates in the order of ~10⁻⁵ s⁻¹.

In the present work we present an experimental setup for *in-situ* observation of SAGBO that allows to correlate 3D grain orientation maps with the evolution of a crack network by means of combining phase contrast and diffraction contrast tomography. Whereas the former allows to visualize the 3D crack network during *in-situ* testing at high temperature, the former can be used to map the 3D microstructure of the material prior to the *in-situ* experiment. For heat treatments promoting precipitation of the delta phase, it may also be possible to use the phase contrast data to calculate local strain fields by means of digital volume correlation (DVC) [13,14]. However, the weak difference in electron density implies poor image contrast between matrix and the precipitate phase and careful optimization of the scan acquisition and processing steps is required to provide suitable contrast for this type of analysis. Hereafter we discuss incipient processing details to achieve such image quality.

## 2. Experimental procedures

The *in-situ* experiment was conducted at the ID11 Materials Science beamline at ESRF. Miniaturized samples of IN718 alloy were cut in dog-bone shape specifically designed for the Nanox [15] tensile rig. The rig was modified to accommodate a longer shaft allowing SAGBO *in-situ* observations inside a tubular furnace by means of phase contrast tomography (PCT). The samples, already mounted in the stress-rig, were initially grain-mapped by means of diffraction contrast tomography (DCT) before straining and/or heating. All measurements were carried out on the 3DXRD endstation of ID11.

### 2.1. Diffraction contrast tomography (DCT)

The data acquisition for the grain map of the undeformed sample was performed using a 2048 x 2048 pixels sCMOS detector system with effective pixel size of 2.44 µm and sample-detector distance, *d* = 7 mm. 50.25 keV x-rays (λ = 0.24673 Å) were used with an incident beam of 0.7 x 0.3 mm in cross-section. The diffraction spots images were measured with angular steps (Δω = 0.1°) and exposure time of 0.05 s over a full turn of the sample. The DCT data processing pipeline is implemented in Matlab is based on finding and indexing Friedel pairs among the segmented spots and then selecting the spots for shape reconstruction [16,17].

### 2.2. Phase Contrast Tomography (PCT)

The PCT imaging was performed with effective pixel size of 0.61 µm and similar beam conditions. The reconstructions were performed using an implementation from ESRF's data analysis group. Denoising and ring removal routines were also applied to assure the best conditions for digital volume correlation (DVC).

*In-situ data acquisition.* The cylindrical furnace was placed over the Nanox sample using a motorized translation stage. The temperature was controlled by measurements from a thermocouple near the sample region. The dwell temperature was set to 923 K and remained stable for at least 1 h before the test. The force applied to the sample at a strain rate of order of ~10⁻⁶s⁻¹ was recorded through the experiment allowing the determination of the tensile behavior as function of time.
*PCT data processing pipeline.* The PCT data was mainly processed using ESRF's Nabu software for tomographic reconstruction [18] using a filtered back-projection (FBP) algorithm. The experimental setup in tomographic acquisitions can introduce image artifacts, such as rings, streaks and error-motion originated defects. Different methods have been proposed to suppress rings [19–22]. In this work we chose to apply the ring corrections suggested by Vo, Atwood and Drakopoulos in their paper [21,23]. The error motion was calculated using a method based on the tomographic consistency premise [24] and the resulting translation vectors were applied before final reconstruction.

### 3. Results and Discussion

#### 3.1. *Initial reconstructions*

Figure 1 depicts a vertical slice reconstruction of different time steps of the SAGBO test without corrections. One can notice the evolution of shape/features in the middle slice as the sample is deformed, this is due to the relative movement imposed by the tensile rig together with cross-section diminishing. It is also possible to notice that the dark spots on the slices, which correspond to MC carbides, present self-similar star-like artifacts. These artifacts are due to small movements of the sample relative to the rotation axis, caused by periodic temperature fluctuations during a full rotation of the sample in our experimental setup. Ring artifacts are prominent in all the images and high frequency noise can be observed.

<figure><img src="image_5.png" /><figcaption>**Figure 1.** Vertical slice reconstruction of 7 different time steps along the in-situ SAGBO test.</figcaption></figure>

#### 3.2. *Deringed and error-motion corrected reconstructions*

Figure 2 shows a separation between the effects of error-motion and error-motion + ring correction steps in the 2 minutes acquisition. Comparing figure 2(a) and (b) it's possible to see that the amplitude of the star-like artifacts is reduced. In figure 2(c) one can see that finer rings are removed, although some weaker, new ones have been introduced at locations which seem correlated with the presence of larger carbides.

In figure 3 we present the ring and error-motion corrected slices. It's possible to note that smaller carbide precipitates have enhanced contrast and boundary definition in the reconstructed slice, while bigger particles maintain residual streak artifacts, most likely due to strong phase contrast at the operating distance. The ring removal procedure added/enhanced streak artifacts on the reconstruction as previously reported in [20] potentialized by interaction with cracks in latter time steps.
<figure><img src="image_6.png" /><figcaption>**Figure 2.** Comparison between slices with (a) no corrections, (b) error-motion correction, and (c) ring and error-motion correction.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Figure 3.** Middle slice reconstruction of 7 different time steps along the in-situ SAGBO test with ring and error-motion correction.</figcaption></figure>

### 3.3. Effects of filtering on the quality of FBP reconstructed slice

In order to use tomography data to calculate 3D strain fields it is necessary that the contrast found on the images are related to local characteristics of the sample. DVC is based on the assumption of optical flow, in which each point of the sample produces a level of grey within its neighbourhood that will be maintained as the sample is strained. High frequency noise, as well as ring artifacts, is detrimental to DVC, as it introduces random contrast that does not represent, nor follows, the deformation of the sample. In figure 4 we present different frequency cut-off values used with a Hann filter [25] for the reconstruction of the slices. A divergent colormap was chosen to accentuate the differences in pixel value. It is possible to note that even though the filtering produced an overall smoothing of the image the internal gray level fluctuation remain unchanged.

<figure><img src="image_8.png" /><figcaption>**Figure 4.** Slices obtained by using different frequency cut-off values of (a)0.5 to (f)1.0 with increments of 0.1 between pictures. Blue color represents lower intensity values in the reconstructed image.</figcaption></figure>
### 3.4. Diffraction contrast tomography

In figure 5 we present the DCT reconstruction of three mid-height centered superimposed on the respective radiograph of the same view of the undeformed sample. It is possible to notice that the sample has a comparatively large grain size with respect to its lateral dimensions, with one grain extending over the full cross-section of the sample on the upper half of the image.

On the referred grain it is possible to see a twin forming an angle of 49.6 degrees with the loading axis of the sample. On the right-hand side of the image a radiograph of the deformed sample, prior to big crack opening, is presented. One can infer from this image that the deformation of the sample is dictated by the rotation of the big grain (purple) in the clockwise direction. On the radiograph it is also possible to see faint extinction spots, with orientation angle of 54.2 degrees, that corroborate this observation.

<figure><img src="image_9.png" /><figcaption>**Figure 5.** DCT map of the middle slice through the thickness superposed to the radiograph of the sample.</figcaption></figure>

### 4. Conclusions

In this work we presented an overview of a first attempt of in-situ, time-lapse observation of SAGBO in 718 Ni superalloy. We showed that a combination of phase and diffraction contrast tomography can be used to follow the evolution of damage during a continuous tensile test and we succeeded in the direct observation of SAGBO while monotonically loading the sample at appropriate strain rates. However, the challenging experimental conditions give rise to various artifacts in the reconstructed images. Further improvement of the experimental conditions is required to reach the level of data quality required for quantitative analysis of the evolving strain fields by means of DVC. Nevertheless, our preliminary results confirm the feasibility of joint PCT & DCT characterization, enabling crystallographic analysis of the damage initiation sites.

### Acknowledgments

This project has received funding from the European Union's Horizon 2020 research and innovation programme under the Marie Skłodowska-Curie grant agreement No 847439.

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