Characterization of the high temperature oxidation behaviour of Inconel 625 ® fabricated by additive manufacturing and conventional methods

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Pineda-Arriaga, Javier H. Ramírez-Ramírez, Francisco A. Pérez-González, and 3 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-1728294/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted 7 You are reading this latest preprint version Abstract The differences between the oxidation behaviour of Inconel 625® superalloy fabricated by additive manufacturing and wrought conditions are presented. At 900 and 1000°C, the oxidation kinetics of the specimens exposed to synthetic dry air followed the parabolic rate law, the activation energy being 174 and 139 kJ/mol for the alloys fabricated by additive manufacturing and wrought conditions, respectively. For both alloys, chromia Cr2O3 was the primary oxide found on the surface of the alloy, together with nodules of NiCr2O4 spinel. Metallographic cross-sections of the specimens revealed the development of alumina Al2O3 and rutile TiO2 as internal oxides. The surface and subsurface oxidation of the alloys depended on their microstructures and the presence of the δ-Ni3Nb phase. Superalloys additive manufacturing wrought alloy oxidation Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 Figure 12 Figure 13 Figure 14 Introduction Inconel 625® superalloy exhibits excellent mechanical properties at elevated temperatures. Typically, the microstructural components of this alloy include a g-nickel matrix and secondary phases that, together with a variety of carbides and nitrides, give excellent resistance to high-temperature creep, corrosion, and oxidation [1,2]. When fabricated by conventional casting and forging methods, the metastable-body-centre tetragonal γ''-phase found in the alloy's microstructure transforms into a δ-Ni 3 Nb phase depending on processing times and temperatures. Both microstructural components, γ'' and δ phase, account for the development of additional precipitation-related strengthening mechanisms at elevated temperatures, making it an excellent candidate for the fabrication of components used in engines for aerospace and energy generation applications. Nowadays, advanced technological innovations allow for the fabrication of superalloy components by manufacturing processes different from conventional forging and casting. Investigations show that additive manufacturing practices can produce machinery parts made of Inconel 625®. Despite that by this route, elements with complex geometries can be created, the heterogeneous element distribution promotes high segregation of Nb and Mo in the inter-dendritic limits during solidification, favouring the formation of significant amounts δ-Ni 3 Nb phase. Not surprisingly, the mentioned manufacturing route differs from the conventional ones in the microstructures obtained in each case because the alloys produced by casting and forging develop equiaxed grains with a more homogeneous element distribution [3-5]. Regardless of the manufacturing route selected to fabricate engine components that work at elevated temperatures, the reality is that all of them will experience oxidation [6]. Previous works conducted to understand the high-temperature degradation behaviour of wrought Inconel 625® indicate that the parabolic rate law characterizes the oxidation phenomenon, chromia Cr 2 O 3 , the most stable oxide formed on the surface, and a small amount of Ni-Cr spinel [6-9]. Unfortunately, few works have been dedicated to studying the oxidation behaviour of this and other nickel-based superalloys fabricated by additive manufacturing practices, but the available works should be mentioned. For instance, Ramenate et al. [9] compared the high-temperature oxidation behaviour of Inconel 625® fabricated by Laser Beam Melting (LBM) and conventional techniques at temperatures between 900 and 1050 °C. They found that for both temperatures and manufacturing processes, the main oxide scale was chromia Cr 2 O 3 . However, they pointed out that the formation of a Nb 1.5 Cr 0.5 O 4 oxide was responsible for the greater oxidation experienced by the LBM samples at the highest temperature. Alternatively, Parizia et al. [10] studied the oxidation of alloys fabricated by Laser Powder Bed Fusion (LPBF) in as fabricated and heat-treated conditions. At 900°C and for both specimen conditions, the oxide scale formed was entirely chromia Cr 2 O 3 . Furthermore, they observed that heat-treating of the alloy promoted chromium depletion and precipitation of the δ- Ni 3 Nb phase at the metal-oxide interface and at the grain boundaries, phenomena that enhanced alloy oxidation. Moreover, Sun et al. [11] studied the oxidation behaviour at 900°C of Inconel 625® fabricated by Selective Laser Melting (SLM) using different laser energy densities. Like the investigations already revised, they found that the oxide scale formed included chromia Cr 2 O 3 and a mixture of nonprotective Ni-rich oxides formed at non-specific zones on the surface of the alloy. They indicated that when the energy for melting the alloy is reduced, the formation and spallation of nonprotective oxides were favoured. Finally, Condruz et al. [12] studied the oxidation behaviour of Inconel 625® fabricated by Selective Laser Melting SLM at temperatures between 900 and 1050°C. They found that at 900 °C, the oxide scale included an inner chromia Cr2O3 scale with regions where the spinel (Ni,Fe)Cr 2 O 4 could be found. At 1050°C, they found that the oxide scale formed on the surface of the alloy was entirely (Ni,Fe)Cr 2 O 4 . Private comments from industrial collaborators indicate that Selective Laser Melting (SLM) additive manufacturing techniques have been proposed as convenient and reliable alternatives to undertaking maintenance and repair practices of components used in aerospace and energy generation engines. Unfortunately, this action is conducted between materials with different microstructural characteristics. Therefore, when the alloys are exposed to elevated temperatures, a different oxidation behaviour could be expected between wrought and additive manufacturing alloys, and their behaviour must be analysed and compared. Based on these premises, this work aims to study the differences between the oxidation behaviour of Inconel 625® fabricated by conventional and additive manufacturing routes, emphasizing the microstructure's effect on oxidation. Experimental Procedure For this investigation, samples of Inconel 625® were taken from a billet of the alloy in wrought condition (W) supplied by a local company dedicated to fabricating components of engines used in energy and aerospace applications. Alternatively, the additive manufacturing samples (AM) for the study were fabricated to the final shape using Selective Laser Melting practices (SLM). Samples of 10 × 10 × 20 mm were printed to the final shape in an EOSINT M280 system (EOS GmbH-Electro Optical Systems, Germany) using commercial Inconel 625® powder produced by gas atomization (EOS GmbH, Germany). The selected printing strategy was a stripe laser pattern with a laser rotation of 67° between each solidified layer to minimize printed defects. Printing parameters were 285 W, v = 960 mm/s, h = 0.110 mm, and l = 40 µm, where P [W] is the laser power, v [mm/s] is the scan speed, h [mm] is the hatch distance, and l [mm] is the layer thickness. Argon gas was selected as an inert atmosphere to minimize oxidation during printing. These processing parameters were chosen to guarantee samples' porosity levels below 1% based on the experience of this working group [13]. The nominal chemical compositions of the alloy powder and the wrought alloy are presented in Table 1. It is worth mentioning that the samples of the wrought alloy were machined directly from the billet to the same size as the AM samples using spark erosion techniques. Table 1. Nominal chemical compositions of the Inconel 625 ® alloy powder and the wrought alloy. Alloying elements %wt Ni Cr Mo Nb Fe C Si Al Ti Co Powder Balance 23.78 10.80 4.00 1.2 0.74 0.023 0.43 1.40 0.77 Wrought alloy Balance 23.50 9.75 4.15 2.5 0.53 0.5 0.40 2.85 1 Once obtained, the samples were prepared following standard metallographic procedures for microstructural analysis, oxidation tests, and the cross-section analysis of the oxide scales formed after high-temperature exposure. These practices included grinding the samples with SiC abrasive papers to a 1200 grit surface finish, followed by polishing stages using diamond solutions of 6 mm, 1 mm, and ¼ mm on soft cloths, respectively. Although it is well known that in industrial processing conditions, the surfaces of components are certainly not prepared to a ¼ mm surface finish, this action assisted the analysis of the oxidation phenomenon and its relationship with the microstructure of the alloys. The oxidation behaviour of the alloys was studied by exposure of the samples at 900 and 1000°C using synthetic dry air as an oxidant atmosphere under isothermal conditions in an oxidation reactor. This installation was designed and built for similar studies, and the operation details can be found in the works conducted by Pérez-González et al. [14,15]. The reactor consists of a radiation tube furnace that allows the introduction of inert or oxidant gases to obtain the desired atmosphere. For the current investigation, the furnace was set to the required temperature. Then, the reaction chamber was filled with argon gas which was introduced at 30 cm 3 /min, and either the additive manufacturing or wrought samples were placed close to the hot zone of the furnace using an alumina boat. The endcaps of the tube were closed and sealed, and the samples were then moved to the hot zone of the tube to experience heating to the test temperature under inert conditions, which took ~5 minutes. After this time, the argon gas was replaced by a flow of 60 cm 3 /min of dry air to start with the oxidation tests conducted at different temperatures and for a maximum exposure time of 25 hours. After completion of each oxidation time, the samples were removed from the hot zone of the furnace a let to cool under a natural convection regime. The oxidation kinetics of the samples was determined by gravimetric means measuring the mass of the samples before and after high-temperature exposure using an electronic balance of 1x10 -5 accuracy. The microstructural components of the alloys and the oxide phases resulting from the oxidation tests were determined through X-ray diffraction. A diffractometer coupled with a cobalt K a tube (l=1.79 Å) was used to produce radiation at 50kV and 40mA. High-definition scans were performed over a 2q length of 20 to 95° at a rate of 0.002°/min. The microstructural components and composition of the oxide scales formed relative to the reflections obtained after the scans were determined using the database of standards stemming from most ICDD cards related to the compounds found in the samples. Alternatively, the microstructure of the alloys and the surface characteristics after oxidation were studied using optical (OM) and scanning electron microscopy (SEM). The optical microscopy analysis (OM) did not require much sample preparation. However, for the analysis in the scanning electron microscope (SEM), the samples were coated with a thin layer of gold applied by sputtering means to facilitate electron conduction during the analyses, which were performed using both, secondary and backscattered electron detectors at different acceleration voltages and working distances to give the best image quality. Energy-dispersive X-ray analyses (EDX) were conducted using the detector coupled with the microscope at interest locations. Finally, the microstructure of the oxide scales formed was also studied in the SEM after preparing metallographic cross-sections of the oxides formed. To keep the oxide scale features during the cutting of the samples, these were coated with a layer of a silver dag and let dry for 48 hours. Then, they were sectioned using a cubic boron nitride disc (CBN) in a precision cutting machine. Copious amounts of a water-based cutting fluid were applied during the procedure to avoid sample overheating. Once cut, the samples were mounted in conductive bakelite and prepared following the metallographic procedures described earlier before inspection in the instrument. Results Figure 1 shows the X-ray spectra taken from the samples fabricated by additive manufacturing (AM) and wrought conditions (W). Regardless of the fabrication route, the spectra indicate that the microstructural components of the alloys in both conditions are the g-nickel matrix and the d-Ni 3 Nb phase. Alternatively, Figure 2 a to d shows optical microscopy micrographs of the samples fabricated by additive manufacturing showing distinct orientations concerning the growth direction imposed by the laser during fabrication (z-direction in the Figure). The analysis of these images suggests that homogenous melting occurred with texture imposed by the angle given to the laser during the process (Figure 2 a to c). In addition, Figure 2d is a magnification of a zone in Figure 2c. It is possible to appreciate that inside the different melting pools formed, different growth directions are also established and where elongated-round-shaped dendrites followed orientations related to the solidification direction of the alloy. For comparison purposes, optical micrographs of the microstructures of the wrought alloy (W) and the material fabricated by additive manufacturing (AM) are shown in Figures 3 a and b, respectively, which exhibit the evident difference between the microstructures. A better microstructural analysis of the alloys is shown in Figure 4, which gives backscattered electron micrographs of the samples in wrought (W) (Figure 4a) and additive manufacturing (AM) conditions (Figure 4b). For the alloy in wrought condition (W), equiaxed grains of the g-nickel matrix and TiN and TiC precipitates located at grain boundaries were found together with particles that suggested the presence of the d-Ni 3 Nb phase. Contrarily, for the additive manufacturing (AM) samples, the secondary phases, namely carbides, nitrides, and d-phase, were not easily resolved in the microscope, as is shown in Figure 4b. The energy-dispersive X-ray spectroscopy spectra (EDX) taken from regions of interest for each alloy (points P1 to P3) are also shown to identify with fine detail the microstructural components. Regarding the oxidation behaviour of the alloy, Figure 5 shows the oxidation kinetics plots of the samples for both fabrication conditions. It can be appreciated that, at both temperatures, the alloy fabricated by additive manufacturing (AM) experienced greater oxidation than the alloy in wrought condition (W), suggesting that the former metallic system is more reactive. However, it should be noted that despite the oxidation rates, the scenario with the most significant mass gained included only 1.35 mg/cm 2 after 25 hours of exposure to the oxidation environment selected. Alternatively, the results obtained from the X-ray diffraction analyses taken from the samples after 5-, 15- and 25-hours oxidation for both temperatures and manufacturing conditions are shown in Figures 6 a to d. At 900°C and for all oxidation times, reflections relative to chromia Cr 2 O 3 appear in the X-ray spectra shown in Figures 6a and 6b for the alloys in AM and W conditions, respectively. In addition, the X-ray spectra relative to the additive manufacturing (AM) samples oxidized for 15 and 25 hours also include reflections of the NiCr 2 O 4 spinel (Figure 6a). For the samples in wrought condition (W) at 900°C, reflections of this oxide phase appeared only after 25 hours of oxidation (Figure 6b). It is important to mention that reflections of the alloy matrix and the d-Ni 3 Nb phase were present in the X-ray spectra taken for all de oxidation times studied in the additive manufacturing samples. In contrast, reflections associated with the d-Ni 3 Nb were only revealed after 25 hours of oxidation in the wrought condition samples. A different behaviour was experienced when oxidation occurred at 1000°C. After 5 hours of oxidation and for the additive manufacturing (AM) and wrought conditions (W) samples (Figures 6 c and d, respectively), chromia Cr 2 O 3 and NiCr 2 O 4 spinel were always detected during the analysis. This behaviour differs from the one experienced by the samples in the wrought condition where both oxides were present after 15- and 25-hours oxidation. For the alloy in both manufacturing conditions, reflections of the g-nickel matrix and the d-Ni 3 Nb disappeared after 15 hours of oxidation. Moreover, Figures 7 a to d show backscattered electron micrographs of the surface of the additive manufacturing and wrought alloys oxidized at 900°C for 15 and 25 hours. Particularly for the additive manufacturing samples AM (Figures 7 a and b), regardless of the oxidation time, the surface of the alloy was entirely covered by a homogenous oxide layer that accentuated the grain boundaries, including the presence of oxide nodules. This was not the case for the samples in wrought condition W (Figures 7 c and d) because the grains of the microstructure could still be resolved for both oxidation times, suggesting that a thin oxide scale covered the alloy surface. Nodule formation was also appreciated in this condition, as in the case of the additive manufacture samples, the structures had similar morphologies but distinct size. Furthermore, Figure 8 a to d gives backscattered electron micrographs of the additive manufacturing AM and wrought W samples oxidized for 15 and 25 hours at 1000°C. For the additive manufacturing AM samples (Figure 8 a and b), the growth of the oxides formed at grain boundaries together with the growth of nodules, promoted the formation of a dense oxide layer (Figure 8a) that completely covered the surface of the alloy after 25 hours (Figure 8b). This phenomenon was not appreciated in the samples in wrought condition W (Figures 8 c and d). For both oxidation times, the oxide scale formed on the surface of the alloy was not as dense as in the additive manufacturing samples and only delineated the grain boundaries. Oxide nodule formation was also appreciated. Alternatively, Figures 9 a to d are backscattered electron images taken from the surface of the specimens in both manufacturing conditions after 15 minutes of oxidation at 900 and 1000°C. Figures 9 a and b compare the oxide structures formed in the additively manufactured AM alloy (Figure 9a) and the alloy in wrought condition W (Figure 9b). In both cases, the oxides formed resembled parallelepiped crystals, being the arrangement of these structures denser for the alloy fabricated by additive manufacturing AM. In addition, the energy-dispersive X-ray spectroscopy analysis EDX conducted in area mode suggests that the chemical composition of the oxides is slightly different. This is because, apart from the elements Cr, Ni and O present in spectra taken for both alloys, the spectrum taken from the oxides formed on the surface of the additive manufacturing AM samples also included reflections of the elements titanium Ti and niobium Nb as it is indicated in the EDX analyses shown in Figure 9a and b. Contrarily, differences in the surface morphology of the alloys were present after oxidation at 1000°C, as can be appreciated in Figures 9 c and d for the additive manufacturing AM and wrought W alloys, respectively. The oxide structures found in the additive manufacturing AM alloy included fine and round-shaped structures with a considerable number of voids. The surface morphology was still characterized by parallelepiped crystals for the wrought W alloy, but these were larger than those formed at 900°C. It is worth mentioning that both oxide structures had the same chemical composition as shown in the EDX spectrum relative to each Figure. No titanium Ti nor niobium Nb were detected in the additive manufacturing AM alloy oxides at this temperature. Moreover, Figures 10 a to d are backscattered electron micrographs taken from metallographic cross-sections of the alloys in both conditions and that were oxidized for 15 and 25 hours at 900°C. Particularly Figures 10 a and b correspond to the additive manufacturing AM alloy. It can be appreciated that after 15 hours, a thin and brittle (~5mm) oxide scale grows on the surface, and it is also possible to appreciate zones where nodules were formed at the oxide-gas interface. Additionally, it is also possible to note that a thin-white layer is continuously formed below the metal-oxide interface and in zones where internal oxides were formed. It is essential to mention that the growth of these oxides followed both grain boundaries and the white phase formed. The same behaviour was present for the additive manufacturing AM alloy after 25 hours of oxidation. However, differences were found in the oxide scale thickness (~9mm) and the depth of the oxides formed internally, as can be appreciated in Figure 10 b. The cross-section analysis of the alloy in wrought condition W for these oxidation times and temperature suggests different oxidation behaviours (Figures 10 c and d). For example, for both oxidation times, the oxide scales formed were thinner than those formed in the additive manufacturing AM alloy (~2 and 5 mm after 15 and 25 hours, respectively), and the internal oxidation zones are not as deep and in the case of the additive manufacturing AM samples. However, the formation of a thin-white layer at the metal-oxide interface was also present in the alloy in this condition. As was expected, changes in the oxide scale were experienced when the temperature increased. Figures 11 a to d are backscattered electron micrographs taken from metallographic cross-sections of the alloys in both conditions and that were oxidized for 15 and 25 hours but at 1000°C. Figures 11 a and b correspond to the additive manufacturing AM alloy after 15- and 25-hours oxidation, and, from these figures, it can be seen that the thickness of the oxide scale increases considerably (~9 and 11 mm at 15 and 25 hours respectively) if compared with the samples in this condition but that were oxidized at 900°C. The oxide scale is also thicker than the oxide scale formed on the samples in wrought condition W oxidized at this temperature (Figures 11 c to d). The white layer found at the metal-oxide interface and relative to the d-Ni 3 Nb phase observed in the tests conducted at 900°C seems dissolved at this temperature and for both oxidation times and material manufacturing conditions. Internal oxidation zones were established irrespective of temperature and oxidation times for both alloys. Still, these zones were more prominent for the additive manufacturing AM samples as can be appreciated in the micrographs. Finally, the distribution of elements resulting from the oxidation phenomenon is presented in Figures 12 a and b for the specimens oxidized after 25 hours at 900°C and Figures 13 a and b for the specimens oxidized after 25 hours but at 1000°C. For samples oxidized at 900°C in both conditions, similarities in the element distribution were found. For instance, the oxide scale includes the elements chromium and oxygen, suggesting the formation of chromia Cr 2 O 3 . In addition, at the location where the white layer was found during the earlier analyses, i.e., the metal-oxide interface, the elements nickel and niobium can be found, suggesting that this layer is the delta phase d-Ni 3 Nb which is commonly found in these alloys regardless of the fabrication route. It is worth mentioning that the d-Ni 3 Nb phase was also present next to the internal oxidation zones where the elements aluminium Al, titanium Ti, and oxygen O were also found, suggesting the formation of small amounts of alumina Al 2 O 3 and rutile TiO 2 oxides as in the case of the additive manufacturing AM samples (Figure 13 a) and only alumina Al 2 O 3 in the wrought condition W samples (Figure 13 b). The same behaviour on the composition of the oxide scale and internal oxide formation was present for the additive manufacturing AM and wrought condition W samples oxidized at 1000°C, i.e., the external oxide scale is chromia Cr 2 O 3 . There was also internal oxidation characterized by the presence of alumina Al 2 O 3 and rutile TiO 2 for the additive manufacturing AM alloys and only Al 2 O 3 for the alloy in wrought condition W. The dissolution of the delta phase d-Ni 3 Nb is suggested during this analysis, given that no niobium Nb was found at this temperature. Discussion The analysis of the results suggests that the oxidation kinetics of the alloy is represented by parabolic relationships at both temperatures and regardless of the fabrication route of the alloy. As proposed in several investigations, the parabolic rate law indicates that both the formation of the external layer and the oxide nodules is controlled by the diffusion of ionic species across the oxide scale [16-18]. Therefore, it is possible to state that the growth of the external chromia Cr 2 O 3 layer that is the oxide present at all temperatures and alloy manufacturing conditions is promoted by the diffusion of Cr + and O - ions providing that the reaction between chromium and oxygen is also favoured by thermodynamic equilibrium [19,20]. However, the mass gained by the samples depended on the alloy manufacturing condition, and the samples fabricated by the additive manufacturing AM route resulted in more significant mass gain than the samples in wrought condition W. This can be better appreciated in the plot of Figure 14 which gives the activation energy for the oxidation process for both alloys. It is seen that the additive manufacturing AM samples have a higher activation energy of -174 kJ/mol compared with the activation energy of the wrought condition W samples, which is -139 kJ/mol proving that oxidation of the additive manufacturing AM alloy is more significant than oxidation of the alloy in wrought condition W. The activation energy value obtained for the additive manufacturing AM samples in this study is like the value of -167 kJ/mol obtained by Sharifitabar et al. [21] for the oxidation of the same metallic system suggesting that the experimental route followed in this investigation was adequate despite that testing was only conducted at two temperatures. Apart from the fact that oxidation is a thermally activated process, and the higher the temperature, the greater the oxidation, it is necessary to provide evidence to explain the greater mass gained experienced by the additive manufacturing AM samples at both temperatures. This phenomenon can be related to the microstructure of the alloys in both conditions. As was previously presented, there was a significant difference between the alloys' microstructure and columnar dendrites found in the additive manufacturing AM samples (Figure 2d). Equiaxed grains were found in the alloy in wrought condition W (Figure 3a). The surface analysis of the additive manufacturing AM samples after 15- and 25-hours oxidation at both temperatures (Figures 7 a and c and Figure 8 a) indicated that oxidation occurred at the interface formed between dendrites and given that a considerable amount of these exists in the additive manufacturing AM alloy compared to equiaxed grains in the wrought condition W samples, oxidation is consequently higher because these zones facilitate the diffusion of ionic species and act as nucleation sites for oxidation to occur [10,17]. The surface analysis of the samples in wrought condition W suggests that the oxidation of the alloy is not important at grain boundaries, and the oxidation phenomenon depends only on the oxidation of the surface of the grains. This phenomenon agrees with the work developed by De Sousa et al. [22], which indicated that the growth of the oxide layer in wrought Inconel 625 ® proceeds on the surface of the grains. This process is slower if compared to grain boundary oxidation. Regarding the composition of the oxide scale, the analysis of the X-diffraction spectra of Figure 6 and the energy dispersive X-ray EDX maps presented in Figures 12 and 13, respectively, show that chromia Cr 2 O 3 is the main oxide present favoured the reaction of chromium and oxygen favoured by thermodynamic equilibrium (DG~520 kJ/mol). This information agrees with the findings of works of Sun et al. [11], Sharifitabar et al. [21], De Sousa et al. [22] that investigated the oxidation behaviour of this alloy fabricated by additive manufacturing AM and different routes and being this oxide the most representative for this alloying system. Despite that nickel is the most abundant element in both alloys and its reaction with oxygen is favoured by thermodynamic equilibrium (DG~-280 kJ/mol), its presence in the oxide scale was limited to the formation of NiCr 2 O 4 spinel by the typical reaction between nickel oxide NiO and chromia Cr 2 O 3 as: The energy of formation nickel oxide NiO is less than that of chromia Cr 2 O 3 for the temperatures studied. However, as nickel ions can be mobile, nickel ions are expected to diffuse across the chromia Cr 2 O 3 layer and react at the oxide-gas interface to form nodules whose presence is favoured for long oxidation times, as was shown in the X-ray spectra and the surface analysis of the oxidized samples. In addition, internal oxidation was appreciated at all temperatures and oxidation conditions, but the extent of this phenomenon was more critical in the additive manufacturing AM samples. This could also be related to the alloy's dendritic microstructure in the additive manufacturing AM, which helps oxygen to reach these zones and react to form alumina Al 2 O 3 and lesser amounts of rutile TiO 2 , as suggested during the EDX analysis conducted by elemental mapping on a certain level of chromium depletion was experienced. The reduced amount and Al and Ti in the alloy (as indicated in Table 1), together with the low partial pressure of oxygen found at these locations, favoured the growth of these phases due to their electronic character as n-type semiconductors. These compounds experience increments in the ionic mobility for diffusion when the partial pressure of oxygen is reduced [16]. Because it is more difficult for oxygen to diffuse through the grain boundaries of the alloys and porosity is negligible for the samples in wrought condition W, the development of internal oxidation zones is only limited to a few regions. Finally, during the X-ray diffraction analysis and SEM analysis of the metallographic cross-sections, it was found that the alloy, particularly in additive manufacturing AM conditions, contains a considerable amount of delta phase d-Ni 3 Nb. From the metallurgical point of view, the presence of this phase is essential because, together with other secondary microstructural components like g' and g", it provides additional strengthening mechanisms at elevated temperatures. Then, a question can be posed: What is the delta phase's role during the oxidation of these alloys for both fabrication conditions? The answer to this question is that the presence of this phase is more critical at 900°C than at 1000°C simply because at the highest temperature, this phase is unstable, and its dissolution is expected, as it was appreciated during the cross-section analysis of the samples [23]. Then, the detailed analysis of the micrographs shown in Figures 10 a to d indicates that at 900°C, a continuous layer of the d-Ni 3 Nb phase forms at the metal-oxide interface of the alloys in both manufacturing conditions. This layer does not experience oxidation (as shown by the EDX elemental mapping shown in Figures 12 a and b). Still, for the additive manufacturing AM samples, this layer seems to act as a nucleation site or easy path for oxygen to promote the growth of the internal oxides found in this condition. If the same criterion is applied to the additive manufacturing AM samples oxidized at 1000°C (Figures 11 a and b), it is possible to identify that at the locations where an oxide starts to grow from the metal-oxide interface to the alloy centre, a portion of delta phase d-Ni 3 Nb is always found in these regions acting probably as a nucleation site for oxide growth. Even though this phenomenon has been studied in earlier works for Inconel 625® in wrought conditions W [7], this fact has not been mentioned in investigations related to this alloy fabricated by additive manufacturing AM. Undoubtedly, this phenomenon deserves further analysis, which is being conducted. Conclusions 1.- The oxidation behaviour of the alloys followed the parabolic rate law. Oxidation is more significant for the samples fabricated by additive manufacturing because the boundaries between dendrites function as nucleation sites that promote and enhance oxidation. 2.- For all oxidation temperatures and manufacturing conditions, the oxide layer contains chromia Cr 2 O 3 as its primary oxide. The formation of NiCr 2 O 4 nodules was experienced for long oxidation times. 3. Internal oxidation was also expected for the samples under both manufacturing conditions, including Al 2 O 3 and TiO 2 as principal oxides for this condition. 4.- At 900°C, delta phase d-Ni 3 Nb has an apparent activation effect on the growth of internal oxides, acting as an easy passage for oxygen or a nucleation site for internal oxidation to occur. Even though dissolution of this phase is favoured at 1000°C, the location where it remains also acts as an initiation site for growth of internal oxides. Declarations Acknowledgments The authors would like to thank the lecturer development program (PRODEP), and Universidad Autónoma de Nuevo León for the facilities provided to develop this investigation. Funding declaration This investigation was conducted under funding provided by the National Council for Science and Technology of Mexico (CONACyT) Competing interests statement Not applicable Author Contribution Nelson Garza-Montes-de-Oca wrote the manuscript, designed the experimental route of the investigation, and revised the results obtained. Rafael Colás edited the manuscript and conducted calculations. Juan M. Alvarado-Orozco fabricated the additive manufacturing alloys, edited the manuscript, and conducted calculations. Javier H. Ramírez-Ramírez and Francisco A. Pérez-González conducted the characterization work with experimental techniques and analysed the results. Karen Y. Pineda-Arriaga. Conducted the experiments, analysed, and prepared the results as part of her dissertation work and this manuscript. References M. M. de Oliveira, A. A. Couto, G. F. C. Almeida, D. A. P. Reis, N. B. de Lima, R. 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Stekovic, Chromia layer growth on a Ni-based superalloy: Sub-parabolic kinetics and the role of titanium , Corros. Sci. 75 (2013) 58-66. https://doi.org/10.1016/j.corsci.2013.05.016 N. Ramenatte, A. Vernouillet, S. Mathieu, A, Van de Put, M Vilasi, D. Monceau, A comparison of the high-temperature oxidation behaviour of conventional wrought and laser beam melted Inconel 625, Corros.Sci. 164 (2020) 143-166. https://doi.org/10.1016/j.corsci.2019.108347. S. Parizia, G. Marchese, M. Rashidi, M. Lorusso, E. Hryha, D. Manfredi, S. Biamino, Effect of heat treatment on microstructure and oxidation properties of Inconel 625 processed by LPBF, J.Alloys Compd. 846 (2020) 156418. https://doi.org/10.1016/j.jallcom.2020.156418. Y. Sun, L. Chen, L. Li, X. Ren, High-temperature oxidation behavior and mechanism of Inconel 625 superalloy fabricated by selective laser melting, Opt. Laser Technol. 132 (2020) 106509. https://doi.org/10.1016/j.optlastec.2020.106509 M. Raluca-Condruz, G. Matache, A. Paraschiv, T. Badea, V. Badilita, High Temperature Oxidation Behavior of Selective Laser Melting Manufactured IN 625, Metals 10 (2020) 668. https://doi.org/10.3390/met10050668 M. Li, A. Fang, E. Martinez-Franco,J.M Alvarado-Orozco, Z. Pei, C. Ma, Selective laser melting of metal matrix composites: Feedstock powder preparation by electroless plating, Mater. Lett. 247 (2019) 115-118. https://doi.org/10.1016/j.matlet.2019.03.092 F.A. Pérez-González, N.F. Garza-Montes-de Oca, R. Colás., High temperature oxidation of the Haynes 282© nickel-based superalloy. Oxid. Met. (2014) 82, 145-161. https://doi.org/10.1007/s11085-014-9483-6 F.A. Pérez-González, J.H. Ramírez-Ramírez, M. Terock, N.F: Garza-Montes-de-Oca, U. Glatzel, R Colás (2016). High-temperature oxidation of a nickel base superalloy at different oxygen partial pressures. Corros. Eng., Sci. Technol., 51 (2016) 51 (7), 513-521. https://doi.org/10.1080/1478422X.2016.1158226 O. Kubaschewski, B. E. Hopkins, Oxidation of metals and alloys. Butterworths London (1967). P. Kofstad, Defects and transport properties of metal oxides. Oxid. Met., 44(1995), 3-27. https://doi.org/10.1007/BF01046721 D. J. Young, High temperature oxidation and corrosion of metals (Vol. 1). Elsevier. (2008). O. Kubaschewski, (1977). Metallurgical thermochemistry. International Series on Material Science and Technology, 24 , 478. N. Birks, G.H. Meier, F.S. Pettit, Introduction to the high temperature oxidation of metals. Cambridge University Press (2006) M. Sharifitabar, S. Khorshahian, M.S. Afarani, P. Kumar, N.K Jain, High-temperature oxidation performance of Inconel 625 superalloy fabricated by wire arc additive manufacturing. Corros.Sci. 197 (2022), 110087. https://doi.org/10.1016/j.corsci.2022.110087 A. M De Sousa Malafaia, R. B., de Oliveira, L. Latu-Romain, Y, Wouters, R. Baldan, Isothermal oxidation of Inconel 625 superalloy at 800 and 1000° C: Microstructure and oxide layer characterization. Mater. Charact., 161 (2020), 110160. https://doi.org/10.1016/j.matchar.2020.110160 Chen, Ming-Song, Guan-Qiang Wang, Hong-Bin Li, Y. C. Lin, Zong-Huai Zou, Yan-Yong Ma, Dao-Guang He, and Wei-Dong Zeng. Precipitation and dissolution behaviors of δ phase inside a deformed nickel-based superalloy during annealing treatment, Appl. Phys. A: Solids Surf. 125 (2019): 1-14. https://doi.org/10.1007/s00339-019-2741-3 Additional Declarations No competing interests reported. Cite Share Download PDF Status: Under Review Version 1 posted Editorial decision: Major revision 05 Sep, 2022 Reviews received at journal 04 Jul, 2022 Reviewers agreed at journal 17 Jun, 2022 Reviewers invited by journal 13 Jun, 2022 Editor assigned by journal 13 Jun, 2022 Submission checks completed at journal 07 Jun, 2022 First submitted to journal 05 Jun, 2022 You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-1728294","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":118489286,"identity":"221fde15-8c3e-4717-afa8-e1e41d60cb18","order_by":0,"name":"Karen Y. Pineda-Arriaga","email":"","orcid":"","institution":"Universidad Autónoma de Nuevo León","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Karen","middleName":"Y.","lastName":"Pineda-Arriaga","suffix":""},{"id":118489287,"identity":"127b4228-a03e-4baa-998e-f8b85b82efbd","order_by":1,"name":"Javier H. Ramírez-Ramírez","email":"","orcid":"","institution":"Universidad Autónoma de Nuevo León","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Javier","middleName":"H.","lastName":"Ramírez-Ramírez","suffix":""},{"id":118489288,"identity":"dbd1521c-9a09-4a88-9295-a4ea62e4be33","order_by":2,"name":"Francisco A. Pérez-González","email":"","orcid":"","institution":"Universidad Autónoma de Nuevo León","correspondingAuthor":false,"submittingAuthor":false,"prefix":"","firstName":"Francisco","middleName":"A.","lastName":"Pérez-González","suffix":""},{"id":118489289,"identity":"a1c8a915-13ad-49bd-944f-1762d55f235c","order_by":3,"name":"Juan M. 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Garza-Montes-de-Oca","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA8UlEQVRIiWNgGAWjYFAC5gYE+wMQs7ET1MKI0MI4A6SFmRQtzDxgkoAGc/bG5g8fdzDY9c8+/vCzza9t8nzMDIwfPubg1mLZc7BNcuYZhuQZ5xKSpXP7bhu2MTMwS87chluLwY3ENmbeNoZkhjMMB6Rze24zArWwMfPi03L/YfNnkBb5M4zNvy17btsT1nKDsUEaqMXO4AwzmzTDj9uJhLWcSQT6pU0iwfAMG5tlb8Pt5DZmxmb8fjl++PCHj2029nJn2B/f+PHntu389uaDHz7i0QIFEokNIIqxDUw2EFQPAvYQ6g9RikfBKBgFo2CEAQDjg09l7ydRXQAAAABJRU5ErkJggg==","orcid":"","institution":"Universidad Autónoma de Nuevo León","correspondingAuthor":true,"submittingAuthor":false,"prefix":"","firstName":"Nelson","middleName":"F.","lastName":"Garza-Montes-de-Oca","suffix":""}],"badges":[],"createdAt":"2022-06-05 22:59:11","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-1728294/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-1728294/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":25715747,"identity":"037ecd3d-151b-43a9-80ca-dafdca7add5f","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":81622,"visible":true,"origin":"","legend":"\u003cp\u003eX-ray diffraction spectra of the alloys in \u003cem\u003eas-fabricated\u003c/em\u003e conditions.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/8d30260a5cc2cbe89d548c28.png"},{"id":25715754,"identity":"b9b10030-9ac6-4c14-bfc2-0d4f224de175","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1351694,"visible":true,"origin":"","legend":"\u003cp\u003eOptical microscopy micrographs of the microstructural features of the AM alloy. a) X-Y plane of the microstructure. b) Z-X plane of the microstructure. c) Z-Y plane of the microstructure. d) Magnification of a zone in c).\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/acadafd9f50b782fa7962652.png"},{"id":25716716,"identity":"6318bd08-6286-4966-9fe6-2362e76582f8","added_by":"auto","created_at":"2022-08-26 16:21:56","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":1371974,"visible":true,"origin":"","legend":"\u003cp\u003eOptical microscopy micrographs showing a comparison between the microstructures of the alloys in both conditions. A) W alloy. b) AM alloy.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/510dceb2aa2919af4be3a327.png"},{"id":25716295,"identity":"9b78d282-f7b5-4de7-995a-b95cca12de50","added_by":"auto","created_at":"2022-08-26 16:16:56","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":502243,"visible":true,"origin":"","legend":"\u003cp\u003eBackscattered electron micrographs of the microstructure of material in \u003cem\u003eas-received \u003c/em\u003econdition. a) W alloy showing equiaxed grains and microstructural components. b) AM alloy showing dendrite distribution.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/7a0c5e1b912f61fed4a76926.png"},{"id":25715748,"identity":"f683b906-cd01-4322-9db0-690ef90b0e0d","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":88345,"visible":true,"origin":"","legend":"\u003cp\u003eOxidation kinetics of the alloys for both manufacturing conditions at 900 and 1000°C.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/f9e4c1fdf53734a1fde8cfd9.png"},{"id":25716296,"identity":"012214f3-cd6c-4da9-bd61-ad199e4bd488","added_by":"auto","created_at":"2022-08-26 16:16:56","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":229618,"visible":true,"origin":"","legend":"\u003cp\u003eX-ray diffraction spectra taken from the samples after oxidation. a) AM alloy at 900°C. b) W alloy at 900°C. c) AM alloy at 1000°C. d) W alloy at 1000°C.\u003c/p\u003e","description":"","filename":"6.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/82fc522d7802ba9c34223bb3.png"},{"id":25715760,"identity":"a7abcd13-6085-4d22-aa80-f8744aa84565","added_by":"auto","created_at":"2022-08-26 16:11:57","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":358915,"visible":true,"origin":"","legend":"\u003cp\u003eSurface morphology of the samples after oxidation. a) AM alloy after 15 hours oxidation at 900°C. b) AM alloy after 25 hours oxidation at 900°C. c) W alloy after 15 hours oxidation at 900°C. b) W alloy after 25 hours oxidation at 900°C.\u003c/p\u003e","description":"","filename":"7.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/272d6431df3b75b2aff43dad.png"},{"id":25716299,"identity":"26b67da3-7e55-4ccd-a601-920d2330a337","added_by":"auto","created_at":"2022-08-26 16:16:56","extension":"png","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":356644,"visible":true,"origin":"","legend":"\u003cp\u003eSurface morphology of the samples after oxidation. a) AM alloy after 15 hours oxidation at 1000°C. b) AM alloy after 25 hours oxidation at 1000°C. c) W alloy after 15 hours oxidation at 1000°C. b) W alloy after 25 hours oxidation at 1000°C.\u003c/p\u003e","description":"","filename":"8.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/91e7d767981997a532c145ff.png"},{"id":25715750,"identity":"952272c5-ddc6-4ef1-977d-1a5424407ff8","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":491076,"visible":true,"origin":"","legend":"\u003cp\u003eSurface morphology of the samples after 15 hours oxidation. a) AM alloy 900 °C, b) W alloy at 900°C. c) AM alloy at 1000 °C. d) W alloy at 1000 °C.\u003c/p\u003e","description":"","filename":"9.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/35cdde74b0760728c0cbdee8.png"},{"id":25716864,"identity":"8da200c2-042c-4f21-a964-ee8eac95a833","added_by":"auto","created_at":"2022-08-26 16:26:56","extension":"png","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":360357,"visible":true,"origin":"","legend":"\u003cp\u003eBackscattered electron micrographs of the metallographic cross-section of the samples after oxidation. a) AM alloy after 15 hours oxidation at 900°C. b) AM alloy after 25 hours oxidation at 900°C. c) W alloy after 15 hours oxidation at 900°C. b) W alloy after 25 hours oxidation at 900°C.\u003c/p\u003e","description":"","filename":"10.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/50eecfa7fffb614e0f80e3b2.png"},{"id":25715759,"identity":"503cf4d8-3475-4e14-979c-8edeb5a8187e","added_by":"auto","created_at":"2022-08-26 16:11:57","extension":"png","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":379189,"visible":true,"origin":"","legend":"\u003cp\u003eBackscattered electron micrographs of the metallographic cross-section of the samples after oxidation. a) AM alloy after 15 hours oxidation at 1000°C. b) AM alloy after 25 hours oxidation at 1000°C. c) W alloy after 15 hours oxidation at 1000°C. b) W alloy after 25 hours oxidation at 1000°C.\u003c/p\u003e","description":"","filename":"11.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/6379701f6f4c0dc76515ad8b.png"},{"id":25716301,"identity":"71483605-dd6a-4bc4-8a9f-93e695bc0edd","added_by":"auto","created_at":"2022-08-26 16:16:56","extension":"png","order_by":12,"title":"Figure 12","display":"","copyAsset":false,"role":"figure","size":711939,"visible":true,"origin":"","legend":"\u003cp\u003eEnergy dispersive X-ray spectroscopy element mapping of the samples after 25 hours oxidation at 900°C. a) AM alloy. b) W alloy.\u003c/p\u003e","description":"","filename":"12.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/ad4b3e34fe3ef276057835e7.png"},{"id":25715757,"identity":"1dde6d28-bdc8-450c-8bb0-d71a0657ea74","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":13,"title":"Figure 13","display":"","copyAsset":false,"role":"figure","size":711711,"visible":true,"origin":"","legend":"\u003cp\u003eEnergy dispersive X-ray spectroscopy element mapping of the samples after 25 hours oxidation at 1000°C. a) AM alloy. b) W alloy.\u003c/p\u003e","description":"","filename":"13.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/cb9e544e7771b8765f430c7c.png"},{"id":25715756,"identity":"7af996e1-464c-42cb-a8aa-a2bf43f06df4","added_by":"auto","created_at":"2022-08-26 16:11:56","extension":"png","order_by":14,"title":"Figure 14","display":"","copyAsset":false,"role":"figure","size":77157,"visible":true,"origin":"","legend":"\u003cp\u003eActivation energy of the alloys.\u003c/p\u003e","description":"","filename":"14.png","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/71ba3f9d94f69925e88198d9.png"},{"id":25716865,"identity":"462e3dd3-b416-42d0-aafc-3fa265319fba","added_by":"auto","created_at":"2022-08-26 16:27:00","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":776015,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-1728294/v1/cff99ff2-0ab7-470e-b6ce-07b15211dc45.pdf"}],"financialInterests":"No competing interests reported.","formattedTitle":"Characterization of the high temperature oxidation behaviour of Inconel 625 ® fabricated by additive manufacturing and conventional methods","fulltext":[{"header":"Introduction ","content":"\u003cp\u003eInconel 625\u0026reg; superalloy exhibits excellent mechanical properties at elevated temperatures. Typically, the microstructural components of this alloy include a g-nickel matrix and secondary phases that, together with a variety of carbides and nitrides, give excellent resistance to high-temperature creep, corrosion, and oxidation [1,2]. When fabricated by conventional casting and forging methods, the metastable-body-centre tetragonal \u0026gamma;\u0026apos;\u0026apos;-phase found in the alloy\u0026apos;s microstructure transforms into a \u0026delta;-Ni\u003csub\u003e3\u003c/sub\u003eNb phase depending on processing times and temperatures. Both microstructural components, \u0026gamma;\u0026apos;\u0026apos; and \u0026delta; phase, account for the development of additional precipitation-related strengthening mechanisms at elevated temperatures, making it an excellent candidate for the fabrication of components used in engines for aerospace and energy generation applications.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eNowadays, advanced technological innovations allow for the fabrication of superalloy components by manufacturing processes different from conventional forging and casting. Investigations show that additive manufacturing practices can produce machinery parts made of Inconel 625\u0026reg;. Despite that by this route, elements with complex geometries can be created, the heterogeneous element distribution promotes high segregation of Nb and Mo in the inter-dendritic limits during solidification, favouring the formation of significant amounts \u0026delta;-Ni\u003csub\u003e3\u003c/sub\u003eNb phase. Not surprisingly, the mentioned manufacturing route differs from the conventional ones in the microstructures obtained in each case because the alloys produced by casting and forging develop equiaxed grains with a more homogeneous element distribution [3-5].\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eRegardless of the manufacturing route selected to fabricate engine components that work at elevated temperatures, the reality is that all of them will experience oxidation [6]. Previous works conducted to understand the high-temperature degradation behaviour of wrought Inconel 625\u0026reg; indicate that the parabolic rate law characterizes the oxidation phenomenon, chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e, the most stable oxide formed on the surface, and a small amount of Ni-Cr spinel [6-9].\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eUnfortunately, few works have been dedicated to studying the oxidation behaviour of this and other nickel-based superalloys fabricated by additive manufacturing practices, but the available works should be mentioned. For instance, Ramenate et al. [9] compared the high-temperature oxidation behaviour of Inconel 625\u0026reg; fabricated by Laser Beam Melting (LBM) and conventional techniques at temperatures between 900 and 1050 \u0026deg;C. They found that for both temperatures and manufacturing processes, the main oxide scale was chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e. However, they pointed out that the formation of a Nb\u003csub\u003e1.5\u003c/sub\u003eCr\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e oxide was responsible for the greater oxidation experienced by the LBM samples at the highest temperature.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eAlternatively, Parizia et al. [10] studied the oxidation of alloys fabricated by Laser Powder Bed Fusion (LPBF) in as fabricated and heat-treated conditions. At 900\u0026deg;C and for both specimen conditions, the oxide scale formed was entirely chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e. Furthermore, they observed that heat-treating of the alloy promoted chromium depletion and precipitation of the \u0026delta;- Ni\u003csub\u003e3\u003c/sub\u003eNb phase at the metal-oxide interface and at the grain boundaries, phenomena that enhanced alloy oxidation.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eMoreover, Sun et al. [11] studied the oxidation behaviour at 900\u0026deg;C of Inconel 625\u0026reg; fabricated by Selective Laser Melting (SLM) using different laser energy densities. Like the investigations already revised, they found that the oxide scale formed included chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and a mixture of nonprotective Ni-rich oxides formed at non-specific zones on the surface of the alloy. They indicated that when the energy for melting the alloy is reduced, the formation and spallation of nonprotective oxides were favoured.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eFinally, Condruz et al. [12] studied the oxidation behaviour of Inconel 625\u0026reg; fabricated by Selective Laser Melting SLM at temperatures between 900 and 1050\u0026deg;C. They found that at 900 \u0026deg;C, the oxide scale included an inner chromia Cr2O3 scale with regions where the spinel (Ni,Fe)Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e could be found. At 1050\u0026deg;C, they found that the oxide scale formed on the surface of the alloy was entirely (Ni,Fe)Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e.\u003c/p\u003e\n\u003cp\u003ePrivate comments from industrial collaborators indicate that Selective Laser Melting (SLM) additive manufacturing techniques have been proposed as convenient and reliable alternatives to undertaking maintenance and repair practices of components used in aerospace and energy generation engines. Unfortunately, this action is conducted between materials with different microstructural characteristics. Therefore, when the alloys are exposed to elevated temperatures, a different oxidation behaviour could be expected between wrought and additive manufacturing alloys, and their behaviour must be analysed and compared. Based on these premises, this work aims to study the differences between the oxidation behaviour of Inconel 625\u0026reg; fabricated by conventional and additive manufacturing routes, emphasizing the microstructure\u0026apos;s effect on oxidation.\u003c/p\u003e"},{"header":"Experimental Procedure","content":"\u003cp\u003eFor this investigation, samples of Inconel 625\u0026reg; were taken from a billet of the alloy in wrought condition (W) supplied by a local company dedicated to fabricating components of engines used in energy and aerospace applications. Alternatively, the additive manufacturing samples (AM) for the study were fabricated to the final shape using Selective Laser Melting practices (SLM). Samples of 10 \u0026times; 10 \u0026times; 20 mm were printed to the final shape in an EOSINT M280 system (EOS GmbH-Electro Optical Systems, Germany) using commercial Inconel 625\u0026reg; powder produced by gas atomization (EOS GmbH, Germany). The selected printing strategy was a stripe laser pattern with a laser rotation of 67\u0026deg; between each solidified layer to minimize printed defects. Printing parameters were 285 W, v = 960 mm/s, h = 0.110 mm, and l = 40 \u0026micro;m, where P [W] is the laser power, v [mm/s] is the scan speed, h [mm] is the hatch distance, and l [mm] is the layer thickness. Argon gas was selected as an inert atmosphere to minimize oxidation during printing. These processing parameters were chosen to guarantee samples\u0026apos; porosity levels below 1% based on the experience of this working group [13]. The nominal chemical compositions of the alloy powder and the wrought alloy are presented in Table 1. It is worth mentioning that the samples of the wrought alloy were machined directly from the billet to the same size as the AM samples using spark erosion techniques.\u003c/p\u003e\n\u003cp\u003eTable 1. Nominal chemical compositions of the Inconel 625 \u0026reg; alloy powder and the wrought alloy.\u003c/p\u003e\n\u003ctable border=\"1\" cellpadding=\"0\" cellspacing=\"0\"\u003e\n \u003ctbody\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" width=\"11.904761904761905%\"\u003e\n \u003cp\u003e\u0026nbsp;\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd colspan=\"10\" valign=\"top\" width=\"88.0952380952381%\"\u003e\n \u003cp\u003e\u003cstrong\u003eAlloying elements %wt\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" width=\"11.904761904761905%\"\u003e\n \u003cp\u003e\u0026nbsp;\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"10.54421768707483%\"\u003e\n \u003cp\u003e\u003cstrong\u003eNi\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"9.183673469387756%\"\u003e\n \u003cp\u003e\u003cstrong\u003eCr\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e\u003cstrong\u003eMo\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"12.07482993197279%\"\u003e\n \u003cp\u003e\u003cstrong\u003eNb\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.673469387755102%\"\u003e\n \u003cp\u003e\u003cstrong\u003eFe\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.482993197278912%\"\u003e\n \u003cp\u003e\u003cstrong\u003eC\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e\u003cstrong\u003eSi\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.333333333333334%\"\u003e\n \u003cp\u003e\u003cstrong\u003eAl\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.653061224489796%\"\u003e\n \u003cp\u003e\u003cstrong\u003eTi\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.142857142857143%\"\u003e\n \u003cp\u003e\u003cstrong\u003eCo\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" width=\"11.904761904761905%\"\u003e\n \u003cp\u003ePowder\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"10.54421768707483%\"\u003e\n \u003cp\u003eBalance\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"9.183673469387756%\"\u003e\n \u003cp\u003e23.78\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e10.80\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"12.07482993197279%\"\u003e\n \u003cp\u003e4.00\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.673469387755102%\"\u003e\n \u003cp\u003e1.2\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.482993197278912%\"\u003e\n \u003cp\u003e0.74\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e0.023\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.333333333333334%\"\u003e\n \u003cp\u003e0.43\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.653061224489796%\"\u003e\n \u003cp\u003e1.40\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.142857142857143%\"\u003e\n \u003cp\u003e0.77\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" width=\"11.904761904761905%\"\u003e\n \u003cp\u003eWrought alloy\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"10.54421768707483%\"\u003e\n \u003cp\u003eBalance\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"9.183673469387756%\"\u003e\n \u003cp\u003e23.50\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e9.75\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"12.07482993197279%\"\u003e\n \u003cp\u003e4.15\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.673469387755102%\"\u003e\n \u003cp\u003e2.5\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.482993197278912%\"\u003e\n \u003cp\u003e0.53\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.503401360544217%\"\u003e\n \u003cp\u003e0.5\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"8.333333333333334%\"\u003e\n \u003cp\u003e0.40\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.653061224489796%\"\u003e\n \u003cp\u003e2.85\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" width=\"7.142857142857143%\"\u003e\n \u003cp\u003e1\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003c/tbody\u003e\n\u003c/table\u003e\n\u003cp\u003e\u0026nbsp;Once obtained, the samples were prepared following standard metallographic procedures for microstructural analysis, oxidation tests, and the cross-section analysis of the oxide scales formed after high-temperature exposure. These practices included grinding the samples with SiC abrasive papers to a 1200 grit surface finish, followed by polishing stages using diamond solutions of 6 mm, 1 mm, and \u0026frac14; mm on soft cloths, respectively. Although it is well known that in industrial processing conditions, the surfaces of components are certainly not prepared to a \u0026frac14; mm surface finish, this action assisted the analysis of the oxidation phenomenon and its relationship with the microstructure of the alloys.\u003c/p\u003e\n\u003cp\u003eThe oxidation behaviour of the alloys was studied by exposure of the samples at 900 and 1000\u0026deg;C using synthetic dry air as an oxidant atmosphere under isothermal conditions in an oxidation reactor. This installation was designed and built for similar studies, and the operation details can be found in the works conducted by P\u0026eacute;rez-Gonz\u0026aacute;lez et al. [14,15]. The reactor consists of a radiation tube furnace that allows the introduction of inert or oxidant gases to obtain the desired atmosphere. For the current investigation, the furnace was set to the required temperature. Then, the reaction chamber was filled with argon gas which was introduced at 30 cm\u003csup\u003e3\u003c/sup\u003e/min, and either the additive manufacturing or wrought samples were placed close to the hot zone of the furnace using an alumina boat. The endcaps of the tube were closed and sealed, and the samples were then moved to the hot zone of the tube to experience heating to the test temperature under inert conditions, which took ~5 minutes. After this time, the argon gas was replaced by a flow of 60 cm\u003csup\u003e3\u003c/sup\u003e/min of dry air to start with the oxidation tests conducted at different temperatures and for a maximum exposure time of 25 hours. After completion of each oxidation time, the samples were removed from the hot zone of the furnace a let to cool under a natural convection regime. The oxidation kinetics of the samples was determined by gravimetric means measuring the mass of the samples before and after high-temperature exposure using an electronic balance of 1x10\u003csup\u003e-5\u003c/sup\u003e accuracy.\u003c/p\u003e\n\u003cp\u003eThe microstructural components of the alloys and the oxide phases resulting from the oxidation tests were determined through X-ray diffraction. A diffractometer coupled with a cobalt \u003cem\u003eK\u003c/em\u003ea\u0026nbsp;tube (l=1.79 \u0026Aring;) was used to produce radiation at 50kV and 40mA. High-definition scans were performed over a 2q\u0026nbsp;length of 20 to 95\u0026deg; at a rate of 0.002\u0026deg;/min. The microstructural components and composition of the oxide scales formed relative to the reflections obtained after the scans were determined using the database of standards stemming from most ICDD cards related to the compounds found in the samples.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp;Alternatively, the microstructure of the alloys and the surface characteristics after oxidation were studied using optical (OM) and scanning electron microscopy (SEM). The optical microscopy analysis (OM) did not require much sample preparation. However, for the analysis in the scanning electron microscope (SEM), the samples were coated with a thin layer of gold applied by sputtering means to facilitate electron conduction during the analyses, which were performed using both, secondary and backscattered electron detectors at different acceleration voltages and working distances to give the best image quality. Energy-dispersive X-ray analyses (EDX) were conducted using the detector coupled with the microscope at interest locations.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp;Finally, the microstructure of the oxide scales formed was also studied in the SEM after preparing metallographic cross-sections of the oxides formed. To keep the oxide scale features during the cutting of the samples, these were coated with a layer of a silver dag and let dry for 48 hours. Then, they were sectioned using a cubic boron nitride disc (CBN) in a precision cutting machine. Copious amounts of a water-based cutting fluid were applied during the procedure to avoid sample overheating. Once cut, the samples were mounted in conductive bakelite and prepared following the metallographic procedures described earlier before inspection in the instrument.\u003c/p\u003e"},{"header":"Results","content":"\u003cp\u003eFigure 1 shows the X-ray spectra taken from the samples fabricated by additive manufacturing (AM) and wrought conditions (W). Regardless of the fabrication route, the spectra indicate that the microstructural components of the alloys in both conditions are the\u0026nbsp;g-nickel matrix and the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase.\u003c/p\u003e\n\u003cp\u003eAlternatively, Figure 2 a to d shows optical microscopy micrographs of the samples fabricated by additive manufacturing showing distinct orientations concerning the growth direction imposed by the laser during fabrication (z-direction in the Figure). The analysis of these images suggests that homogenous melting occurred with texture imposed by the angle given to the laser during the process (Figure 2 a to c). In addition, Figure 2d is a magnification of a zone in Figure 2c. It is possible to appreciate that inside the different melting pools formed, different growth directions are also established and where elongated-round-shaped dendrites followed orientations related to the solidification direction of the alloy. For comparison purposes, optical micrographs of the microstructures of the wrought alloy (W) and the material fabricated by additive manufacturing (AM) are shown in Figures 3 a and b, respectively, which exhibit the evident difference between the microstructures.\u003c/p\u003e\n\u003cp\u003eA better microstructural analysis of the alloys is shown in Figure 4, which gives backscattered electron micrographs of the samples in wrought (W) (Figure 4a) and additive manufacturing (AM) conditions (Figure 4b). For the alloy in wrought condition (W), equiaxed grains of the\u0026nbsp;g-nickel matrix and TiN and TiC precipitates located at grain boundaries were found together with particles that suggested the presence of the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase. Contrarily, for the additive manufacturing (AM) samples, the secondary phases, namely carbides, nitrides, and d-phase, were not easily resolved in the microscope, as is shown in Figure 4b. The energy-dispersive X-ray spectroscopy spectra (EDX) taken from regions of interest for each alloy (points P1 to P3) are also shown to identify with fine detail the microstructural components.\u003c/p\u003e\n\u003cp\u003eRegarding the oxidation behaviour of the alloy, Figure 5 shows the oxidation kinetics plots of the samples for both fabrication conditions. It can be appreciated that, at both temperatures, the alloy fabricated by additive manufacturing (AM) experienced greater oxidation than the alloy in wrought condition (W), suggesting that the former metallic system is more reactive. However, it should be noted that despite the oxidation rates, the scenario with the most significant mass gained included only 1.35 mg/cm\u003csup\u003e2\u003c/sup\u003e after 25 hours of exposure to the oxidation environment selected.\u003c/p\u003e\n\u003cp\u003eAlternatively, the results obtained from the X-ray diffraction analyses taken from the samples after 5-, 15- and 25-hours oxidation for both temperatures and manufacturing conditions are shown in Figures 6 a to d. At 900\u0026deg;C and for all oxidation times, reflections relative to chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e appear in the X-ray spectra shown in Figures 6a and 6b for the alloys in AM and W conditions, respectively. In addition, the X-ray spectra relative to the additive manufacturing (AM) samples oxidized for 15 and 25 hours also include reflections of the NiCr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e spinel (Figure 6a). For the samples in wrought condition (W) at 900\u0026deg;C, reflections of this oxide phase appeared only after 25 hours of oxidation (Figure 6b). It is important to mention that reflections of the alloy matrix and the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase were present in the X-ray spectra taken for all de oxidation times studied in the additive manufacturing samples. In contrast, reflections associated with the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb were only revealed after 25 hours of oxidation in the wrought condition samples.\u003c/p\u003e\n\u003cp\u003eA different behaviour was experienced when oxidation occurred at 1000\u0026deg;C. After 5 hours of oxidation and for the additive manufacturing (AM) and wrought conditions (W) samples (Figures 6 c and d, respectively), chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and NiCr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e spinel were always detected during the analysis. This behaviour differs from the one experienced by the samples in the wrought condition where both oxides were present after 15- and 25-hours oxidation. For the alloy in both manufacturing conditions, reflections of the g-nickel matrix and the d-Ni\u003csub\u003e3\u003c/sub\u003eNb disappeared after 15 hours of oxidation.\u003c/p\u003e\n\u003cp\u003eMoreover, Figures 7 a to d show backscattered electron micrographs of the surface of the additive manufacturing and wrought alloys oxidized at 900\u0026deg;C for 15 and 25 hours. Particularly for the additive manufacturing samples AM (Figures 7 a and b), regardless of the oxidation time, the surface of the alloy was entirely covered by a homogenous oxide layer that accentuated the grain boundaries, including the presence of oxide nodules. This was not the case for the samples in wrought condition W (Figures 7 c and d) because the grains of the microstructure could still be resolved for both oxidation times, suggesting that a thin oxide scale covered the alloy surface. Nodule formation was also appreciated in this condition, as in the case of the additive manufacture samples, the structures had similar morphologies but distinct size.\u003c/p\u003e\n\u003cp\u003eFurthermore, Figure 8 a to d gives backscattered electron micrographs of the additive manufacturing AM and wrought W samples oxidized for 15 and 25 hours at 1000\u0026deg;C. For the additive manufacturing AM samples (Figure 8 a and b), the growth of the oxides formed at grain boundaries together with the growth of nodules, promoted the formation of a dense oxide layer (Figure 8a) that completely covered the surface of the alloy after 25 hours (Figure 8b). This phenomenon was not appreciated in the samples in wrought condition W (Figures 8 c and d). For both oxidation times, the oxide scale formed on the surface of the alloy was not as dense as in the additive manufacturing samples and only delineated the grain boundaries. Oxide nodule formation was also appreciated.\u003c/p\u003e\n\u003cp\u003eAlternatively, Figures 9 a to d are backscattered electron images taken from the surface of the specimens in both manufacturing conditions after 15 minutes of oxidation at 900 and 1000\u0026deg;C. Figures 9 a and b compare the oxide structures formed in the additively manufactured AM alloy (Figure 9a) and the alloy in wrought condition W (Figure 9b). In both cases, the oxides formed resembled parallelepiped crystals, being the arrangement of these structures denser for the alloy fabricated by additive manufacturing AM. In addition, the energy-dispersive X-ray spectroscopy analysis EDX conducted in area mode suggests that the chemical composition of the oxides is slightly different. This is because, apart from the elements Cr, Ni and O present in spectra taken for both alloys, the spectrum taken from the oxides formed on the surface of the additive manufacturing AM samples also included reflections of the elements titanium Ti and niobium Nb as it is indicated in the EDX analyses shown in Figure 9a and b.\u003c/p\u003e\n\u003cp\u003eContrarily, differences in the surface morphology of the alloys were present after oxidation at 1000\u0026deg;C, as can be appreciated in Figures 9 c and d for the additive manufacturing AM and wrought W alloys, respectively. The oxide structures found in the additive manufacturing AM alloy included fine and round-shaped structures with a considerable number of voids. The surface morphology was still characterized by parallelepiped crystals for the wrought W alloy, but these were larger than those formed at 900\u0026deg;C. It is worth mentioning that both oxide structures had the same chemical composition as shown in the EDX spectrum relative to each Figure. No titanium Ti nor niobium Nb were detected in the additive manufacturing AM alloy oxides at this temperature.\u003c/p\u003e\n\u003cp\u003eMoreover, Figures 10 a to d are backscattered electron micrographs taken from metallographic cross-sections of the alloys in both conditions and that were oxidized for 15 and 25 hours at 900\u0026deg;C. Particularly Figures 10 a and b correspond to the additive manufacturing AM alloy. It can be appreciated that after 15 hours, a thin and brittle (~5mm) oxide scale grows on the surface, and it is also possible to appreciate zones where nodules were formed at the oxide-gas interface. Additionally, it is also possible to note that a thin-white layer is continuously formed below the metal-oxide interface and in zones where internal oxides were formed. It is essential to mention that the growth of these oxides followed both grain boundaries and the white phase formed. The same behaviour was present for the additive manufacturing AM alloy after 25 hours of oxidation. However, differences were found in the oxide scale thickness (~9mm) and the depth of the oxides formed internally, as can be appreciated in Figure 10 b.\u003c/p\u003e\n\u003cp\u003eThe cross-section analysis of the alloy in wrought condition W for these oxidation times and temperature suggests different oxidation behaviours (Figures 10 c and d). For example, for both oxidation times, the oxide scales formed were thinner than those formed in the additive manufacturing AM alloy (~2 and 5 mm after 15 and 25 hours, respectively), and the internal oxidation zones are not as deep and in the case of the additive manufacturing AM samples. However, the formation of a thin-white layer at the metal-oxide interface was also present in the alloy in this condition.\u003c/p\u003e\n\u003cp\u003eAs was expected, changes in the oxide scale were experienced when the temperature increased. Figures 11 a to d are backscattered electron micrographs taken from metallographic cross-sections of the alloys in both conditions and that were oxidized for 15 and 25 hours but at 1000\u0026deg;C. Figures 11 a and b correspond to the additive manufacturing AM alloy after 15- and 25-hours oxidation, and, from these figures, it can be seen that the thickness of the oxide scale increases considerably (~9 and 11\u0026nbsp;mm at 15 and 25 hours respectively) if compared with the samples in this condition but that were oxidized at 900\u0026deg;C. The oxide scale is also thicker than the oxide scale formed on the samples in wrought condition W oxidized at this temperature (Figures 11 c to d). The white layer found at the metal-oxide interface and relative to the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase observed in the tests conducted at 900\u0026deg;C seems dissolved at this temperature and for both oxidation times and material manufacturing conditions. Internal oxidation zones were established irrespective of temperature and oxidation times for both alloys. Still, these zones were more prominent for the additive manufacturing AM samples as can be appreciated in the micrographs.\u003c/p\u003e\n\u003cp\u003eFinally, the distribution of elements resulting from the oxidation phenomenon is presented in Figures 12 a and b for the specimens oxidized after 25 hours at 900\u0026deg;C and Figures 13 a and b for the specimens oxidized after 25 hours but at 1000\u0026deg;C. For samples oxidized at 900\u0026deg;C in both conditions, similarities in the element distribution were found. For instance, the oxide scale includes the elements chromium and oxygen, suggesting the formation of chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e. In addition, at the location where the white layer was found during the earlier analyses, i.e., the metal-oxide interface, the elements nickel and niobium can be found, suggesting that this layer is the delta phase\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb which is commonly found in these alloys regardless of the fabrication route. It is worth mentioning that the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase was also present next to the internal oxidation zones where the elements aluminium Al, titanium Ti, and oxygen O were also found, suggesting the formation of small amounts of alumina Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and rutile TiO\u003csub\u003e2\u003c/sub\u003e oxides as in the case of the additive manufacturing AM samples (Figure 13 a) and only alumina Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e in the wrought condition W samples (Figure 13 b). The same behaviour on the composition of the oxide scale and internal oxide formation was present for the additive manufacturing AM and wrought condition W samples oxidized at 1000\u0026deg;C, i.e., the external oxide scale is chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e. There was also internal oxidation characterized by the presence of alumina Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and rutile TiO\u003csub\u003e2\u003c/sub\u003e for the additive manufacturing AM alloys and only Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e for the alloy in wrought condition W. The dissolution of the delta phase\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb is suggested during this analysis, given that no niobium Nb was found at this temperature.\u003c/p\u003e"},{"header":"Discussion","content":"\u003cp\u003eThe analysis of the results suggests that the oxidation kinetics of the alloy is represented by parabolic relationships at both temperatures and regardless of the fabrication route of the alloy. As proposed in several investigations, the parabolic rate law indicates that both the formation of the external layer and the oxide nodules is controlled by the diffusion of ionic species across the oxide scale [16-18]. Therefore, it is possible to state that the growth of the external chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e layer that is the oxide present at all temperatures and alloy manufacturing conditions is promoted by the diffusion of Cr\u003csup\u003e+\u003c/sup\u003e and O\u003csup\u003e-\u003c/sup\u003e ions providing that the reaction between chromium and oxygen is also favoured by thermodynamic equilibrium [19,20]. However, the mass gained by the samples depended on the alloy manufacturing condition, and the samples fabricated by the additive manufacturing AM route resulted in more significant mass gain than the samples in wrought condition W. This can be better appreciated in the plot of Figure 14 which gives the activation energy for the oxidation process for both alloys. It is seen that the additive manufacturing AM samples have a higher activation energy of -174 kJ/mol compared with the activation energy of the wrought condition W samples, which is -139 kJ/mol proving that oxidation of the additive manufacturing AM alloy is more significant than oxidation of the alloy in wrought condition W. The activation energy value obtained for the additive manufacturing AM samples in this study is like the value of -167 kJ/mol obtained by Sharifitabar et al. [21] for the oxidation of the same metallic system suggesting that the experimental route followed in this investigation was adequate despite that testing was only conducted at two temperatures.\u003c/p\u003e\n\u003cp\u003eApart from the fact that oxidation is a thermally activated process, and the higher the temperature, the greater the oxidation, it is necessary to provide evidence to explain the greater mass gained experienced by the additive manufacturing AM samples at both temperatures. This phenomenon can be related to the microstructure of the alloys in both conditions. As was previously presented, there was a significant difference between the alloys\u0026apos; microstructure and columnar dendrites found in the additive manufacturing AM samples (Figure 2d). Equiaxed grains were found in the alloy in wrought condition W (Figure 3a). The surface analysis of the additive manufacturing AM samples after 15- and 25-hours oxidation at both temperatures (Figures 7 a and c and Figure 8 a) indicated that oxidation occurred at the interface formed between dendrites and given that a considerable amount of these exists in the additive manufacturing AM alloy compared to equiaxed grains in the wrought condition W samples, oxidation is consequently higher because these zones facilitate the diffusion of ionic species and act as nucleation sites for oxidation to occur [10,17]. The surface analysis of the samples in wrought condition W suggests that the oxidation of the alloy is not important at grain boundaries, and the oxidation phenomenon depends only on the oxidation of the surface of the grains. This phenomenon agrees with the work developed by De Sousa et al. [22], which indicated that the growth of the oxide layer in wrought Inconel 625 \u0026reg; proceeds on the surface of the grains. This process is slower if compared to grain boundary oxidation.\u003c/p\u003e\n\u003cp\u003eRegarding the composition of the oxide scale, the analysis of the X-diffraction spectra of Figure 6 and the energy dispersive X-ray EDX maps presented in Figures 12 and 13, respectively, show that chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e is the main oxide present favoured the reaction of chromium and oxygen favoured by thermodynamic equilibrium (DG~520 kJ/mol). This information agrees with the findings of works of Sun et al. [11], Sharifitabar et al. [21], De Sousa et al. [22] that investigated the oxidation behaviour of this alloy fabricated by additive manufacturing AM and different routes and being this oxide the most representative for this alloying system. Despite that nickel is the most abundant element in both alloys and its reaction with oxygen is favoured by thermodynamic equilibrium (DG~-280 kJ/mol), its presence in the oxide scale was limited to the formation of NiCr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e spinel by the typical reaction between nickel oxide NiO and chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e as:\u003c/p\u003e\n\u003cp\u003e\u003cimg src=\"data:image/png;base64,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\"\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003e\u0026nbsp;\u003c/strong\u003eThe energy of formation nickel oxide NiO is less than that of chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e for the temperatures studied. However, as nickel ions can be mobile, nickel ions are expected to diffuse across the chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e layer and react at the oxide-gas interface to form nodules whose presence is favoured for long oxidation times, as was shown in the X-ray spectra and the surface analysis of the oxidized samples.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eIn addition, internal oxidation was appreciated at all temperatures and oxidation conditions, but the extent of this phenomenon was more critical in the additive manufacturing AM samples. This could also be related to the alloy\u0026apos;s dendritic microstructure in the additive manufacturing AM, which helps oxygen to reach these zones and react to form alumina Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and lesser amounts of rutile TiO\u003csub\u003e2\u003c/sub\u003e, as suggested during the EDX analysis conducted by elemental mapping on a certain level of chromium depletion was experienced. The reduced amount and Al and Ti in the alloy (as indicated in Table 1), together with the low partial pressure of oxygen found at these locations, favoured the growth of these phases due to their electronic character as n-type semiconductors. These compounds experience increments in the ionic mobility for diffusion when the partial pressure of oxygen is reduced [16]. Because it is more difficult for oxygen to diffuse through the grain boundaries of the alloys and porosity is negligible for the samples in wrought condition W, the development of internal oxidation zones is only limited to a few regions.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eFinally, during the X-ray diffraction analysis and SEM analysis of the metallographic cross-sections, it was found that the alloy, particularly in additive manufacturing AM conditions, contains a considerable amount of delta phase\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb. From the metallurgical point of view, the presence of this phase is essential because, together with other secondary microstructural components like\u0026nbsp;g\u0026apos; and\u0026nbsp;g\u0026quot;, it provides additional strengthening mechanisms at elevated temperatures. Then, a question can be posed: What is the delta phase\u0026apos;s role during the oxidation of these alloys for both fabrication conditions? The answer to this question is that the presence of this phase is more critical at 900\u0026deg;C than at 1000\u0026deg;C simply because at the highest temperature, this phase is unstable, and its dissolution is expected, as it was appreciated during the cross-section analysis of the samples [23]. Then, the detailed analysis of the micrographs shown in Figures 10 a to d indicates that at 900\u0026deg;C, a continuous layer of the\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb phase forms at the metal-oxide interface of the alloys in both manufacturing conditions. This layer does not experience oxidation (as shown by the EDX elemental mapping shown in Figures 12 a and b). Still, for the additive manufacturing AM samples, this layer seems to act as a nucleation site or easy path for oxygen to promote the growth of the internal oxides found in this condition. If the same criterion is applied to the additive manufacturing AM samples oxidized at 1000\u0026deg;C (Figures 11 a and b), it is possible to identify that at the locations where an oxide starts to grow from the metal-oxide interface to the alloy centre, a portion of delta phase\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb \u0026nbsp;is always found in these regions acting probably as a nucleation site for oxide growth. Even though this phenomenon has been studied in earlier works for Inconel 625\u0026reg; in wrought conditions W [7], this fact has not been mentioned in investigations related to this alloy fabricated by additive manufacturing AM. Undoubtedly, this phenomenon deserves further analysis, which is being conducted.\u003c/p\u003e"},{"header":"Conclusions","content":"\u003cp\u003e1.- The oxidation behaviour of the alloys followed the parabolic rate law. Oxidation is more significant for the samples fabricated by additive manufacturing because the boundaries between dendrites function as nucleation sites that promote and enhance oxidation.\u003c/p\u003e\n\u003cp\u003e2.- For all oxidation temperatures and manufacturing conditions, the oxide layer contains chromia Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e as its primary oxide. The formation of NiCr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e4\u003c/sub\u003e nodules was experienced for long oxidation times.\u003c/p\u003e\n\u003cp\u003e3. Internal oxidation was also expected for the samples under both manufacturing conditions, including Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e and TiO\u003csub\u003e2\u003c/sub\u003e as principal oxides for this condition.\u003c/p\u003e\n\u003cp\u003e4.- At 900\u0026deg;C, delta phase\u0026nbsp;d-Ni\u003csub\u003e3\u003c/sub\u003eNb has an apparent activation effect on the growth of internal oxides, acting as an easy passage for oxygen or a nucleation site for internal oxidation to occur. Even though dissolution of this phase is favoured at 1000\u0026deg;C, the location where it remains also acts as an initiation site for growth of internal oxides.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAcknowledgments\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors would like to thank the lecturer development program (PRODEP), and Universidad Aut\u0026oacute;noma de Nuevo Le\u0026oacute;n for the facilities provided to develop this investigation.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFunding declaration\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis investigation was conducted under funding provided by the National Council for Science and Technology of Mexico (CONACyT)\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests statement\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eNot applicable\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor Contribution\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eNelson Garza-Montes-de-Oca wrote the manuscript, designed the experimental route of the investigation, and revised the results obtained.\u003c/p\u003e\n\u003cp\u003eRafael Col\u0026aacute;s edited the manuscript and conducted calculations.\u003c/p\u003e\n\u003cp\u003eJuan M. Alvarado-Orozco fabricated the additive manufacturing alloys, edited the manuscript, and conducted calculations.\u003c/p\u003e\n\u003cp\u003eJavier H. Ram\u0026iacute;rez-Ram\u0026iacute;rez and Francisco A. P\u0026eacute;rez-Gonz\u0026aacute;lez conducted the characterization work with experimental techniques and analysed the results.\u003c/p\u003e\n\u003cp\u003eKaren Y. Pineda-Arriaga. Conducted the experiments, analysed, and prepared the results as part of her dissertation work and this manuscript.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eM. M. de Oliveira, A. A. Couto, G. F. C. Almeida, D. A. P. Reis, N. B. de Lima, R. Baldan, Mechanical Behavior of Inconel 625 at Elevated Temperatures, Metals 9 (2019) 1-13. https://doi.org/10.3390/met9030301\u003c/li\u003e\n\u003cli\u003eF. Xu, Y. Lv, Y, Liu, B. Xu, P. He, Effect of heat treatment on microstructure and mechanical properties of Inconel 625 alloy fabricated by pulsed plasma arc deposition. 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Baldan, Isothermal oxidation of Inconel 625 superalloy at 800 and 1000\u0026deg; C: Microstructure and oxide layer characterization. Mater. Charact., 161 (2020), 110160. https://doi.org/10.1016/j.matchar.2020.110160\u003c/li\u003e\n\u003cli\u003eChen, Ming-Song, Guan-Qiang Wang, Hong-Bin Li, Y. C. Lin, Zong-Huai Zou, Yan-Yong Ma, Dao-Guang He, and Wei-Dong Zeng. Precipitation and dissolution behaviors of \u0026delta; phase inside a deformed nickel-based superalloy during annealing treatment, Appl. Phys. A: Solids Surf. 125 (2019): 1-14. https://doi.org/10.1007/s00339-019-2741-3\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"oxidation-of-metals","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"oxid","sideBox":"Learn more about [Oxidation of Metals](http://link.springer.com/journal/11085)","snPcode":"11085","submissionUrl":"https://submission.nature.com/new-submission/11085/3","title":"Oxidation of Metals","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"Superalloys, additive manufacturing, wrought alloy, oxidation","lastPublishedDoi":"10.21203/rs.3.rs-1728294/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-1728294/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"The differences between the oxidation behaviour of Inconel 625® superalloy fabricated by additive manufacturing and wrought conditions are presented. At 900 and 1000°C, the oxidation kinetics of the specimens exposed to synthetic dry air followed the parabolic rate law, the activation energy being 174 and 139 kJ/mol for the alloys fabricated by additive manufacturing and wrought conditions, respectively. For both alloys, chromia Cr2O3 was the primary oxide found on the surface of the alloy, together with nodules of NiCr2O4 spinel. Metallographic cross-sections of the specimens revealed the development of alumina Al2O3 and rutile TiO2 as internal oxides. The surface and subsurface oxidation of the alloys depended on their microstructures and the presence of the δ-Ni3Nb phase.","manuscriptTitle":"Characterization of the high temperature oxidation behaviour of Inconel 625 ® fabricated by additive manufacturing and conventional methods","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2022-08-26 16:11:54","doi":"10.21203/rs.3.rs-1728294/v1","editorialEvents":[{"type":"communityComments","content":0},{"type":"decision","content":"Major revision","date":"2022-09-06T02:32:12+00:00","index":"","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2022-07-04T16:46:19+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"db7df6ac-9146-4c79-9915-d5f98a6a9ae4","date":"2022-06-17T07:18:36+00:00","index":"hide","fulltext":""},{"type":"reviewersInvited","content":"","date":"2022-06-14T02:52:29+00:00","index":"","fulltext":""},{"type":"editorAssigned","content":"","date":"2022-06-14T02:40:42+00:00","index":"","fulltext":""},{"type":"checksComplete","content":"","date":"2022-06-07T11:39:44+00:00","index":"","fulltext":""},{"type":"submitted","content":"Oxidation of Metals","date":"2022-06-05T22:49:42+00:00","index":"","fulltext":""}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"oxidation-of-metals","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"oxid","sideBox":"Learn more about [Oxidation of Metals](http://link.springer.com/journal/11085)","snPcode":"11085","submissionUrl":"https://submission.nature.com/new-submission/11085/3","title":"Oxidation of Metals","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"d6d4a7d2-d51a-4ea7-b97e-d943a667807b","owner":[],"postedDate":"August 26th, 2022","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[],"tags":[],"updatedAt":"2022-10-01T21:14:13+00:00","versionOfRecord":[],"versionCreatedAt":"2022-08-26 16:11:54","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-1728294","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-1728294","identity":"rs-1728294","version":["v1"]},"buildId":"7rjqhiLT3MXkJMwkYKINL","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}