Effect of the Noble Metals Addition on the Oxidation Behavior of Ni<sub>3</sub>Al

Ramos-Hernandez, J. J.; Porcayo-Calderon, J.; Gonzalez-Rodriguez, J. G.; Mayén, Jan; Lara-Rodriguez, G.; Salinas-Bravo, V. M.; Martinez-Gomez, L.

doi:https://doi.org/10.1155/2018/1616340

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Research Article | Open Access

Volume 2018 | Article ID 1616340 | https://doi.org/10.1155/2018/1616340

Effect of the Noble Metals Addition on the Oxidation Behavior of Ni₃Al

J. J. Ramos-Hernandez,¹J. Porcayo-Calderon,²J. G. Gonzalez-Rodriguez,²Jan Mayén,³G. Lara-Rodriguez,⁴V. M. Salinas-Bravo,⁵and L. Martinez-Gomez^6,7

Academic Editor: Amit Bandyopadhyay

Received06 Jul 2017

Revised12 Oct 2017

Accepted01 Nov 2017

Published17 Jan 2018

Abstract

This paper discusses the effect of the addition of noble metals on the microstructure, mechanism, and oxidation kinetics of the Ni₃Al intermetallic alloy. Ni₃Al was doped with 1% (atomic percent) of Au, Ag, Pd, and Pt. Oxidation behavior of the alloys was evaluated at 900, 1000, and 1100°C in O₂ for 24 hours. XRD analysis showed that the addition of noble metals favored the oxide growth on preferential crystallographic planes. In addition, the preferential substitution of the noble metals in the Ni₃Al structure modifies the surface composition by increasing the Al/Ni ratio. It was observed that most of the alloys showed a subparabolic behavior, and only the intermetallic base and the alloy doped with Ag show a parabolic behavior at 900°C. The developed oxides were analyzed both superficially and in cross section by scanning electron microscopy (SEM) and energy dispersive X-ray analysis (EDXA). It was evident that only the intermetallic base showed the formation of a duplex oxide scale (Al₂O₃/NiO). The alloys doped with noble metals showed the oxide growth practically of pure Al₂O₃. This was due to a decrease in the diffusion of the Ni cations because of the presence of the noble metals in the crystalline structure.

1. Introduction

Ni₃Al intermetallic compounds are excellent candidates for high-temperature structural applications due to their thermochemical and mechanical properties. They can perform beyond conventional temperatures because of their excellent resistance to corrosion and oxidation [1]. During solidification of an alloy with stoichiometry close to that of Ni₃Al, there is in the liquid phase a competition for phase stability between β_NiAl, γ′_Ni3Al, and γ_Ni with stable eutectic (β-γ′), metastable (β-γ), and peritectic (γ′-γ) phases which have common borders in the Ni-Al binary diagram. Because the region of phase stability γ′ is very narrow, any deviation in the Ni/Al ratio favors the stability of any of the aforementioned phases; that is, an increase in the Ni content tends to favor the stability of γ′-γ phase. According to Verin [2], the microstructure of the Ni-23Al-0.001Ce-0.004Co (at %) is composed of a combination of dendritic phases γ′ and γ, where the matrix is the phase γ′ and the interdendritic phase is the phase γ; similar results were obtained by Chiba and Hanada [3] for the alloy Ni-23Al (at %) and Jozwik et al. [4] for the alloy Ni-22Al-0.002Zr-0.001B (at %). Moreover, Hunziker and Kurz [5] mentioned that the directional solidification of intermetallic Ni-25Al (at %) shows a microstructure where the phases β, γ, and γ′ are found. In this case, the γ′ phase is formed during cooling of metastable eutectic β-γ. It has been shown that there are subsequent transformations after the solidification process, which complicates the interpretation of the phases produced.

On the other hand, the solidification rate is a factor that influences the microstructure of the alloy. When the solidification is slow, the convection in the liquid phase favors formation of a single γ′-phase. However, in a rapid solidification process, there is little convection in the liquid phase and it promotes the formation of a structure with phases γ′’ and β [6]. The addition of a third element to the Ni-Al system may stabilize or destabilize a given phase. According to Inoue et al. [7], the addition of Cr, Mn, Fe, Co, or Si destabilizes the phase γ′, and according to Goetzinger et al. [8], the addition of Ti stabilizes the phase. Sun et al. [9] indicate that Pd, Ag, and Cu are elements forming the γ phase, while Pint et al. [10] show that Pt promotes the reduction of the temperature of the phase transformation from β to γ′ during the process of cooling and phase transformation from γ′ to β during heating.

During the oxidation process, β and γ′ phases play an important role due to their content of Al [11]. According to Haerig and Hofmann [12], in the early stages of the oxidation process, NiO and Al₂O₃ are formed simultaneously due to the diffusion of cations towards the outside of the alloy. Bardi et al. [13] show that the growth of the oxide is in the form of islands. Al consumption owing to the formation of Al₂O₃ diminishes the aluminum content below 17% (at %), which is the content needed to form the outer protective layer [14]. This favors the growth of NiO onto the Al₂O₃ layer. This pattern occurs because of the conducting properties of the formed oxides: Al₂O₃ is an amphoteric oxide (movement of cations and anions is possible) allowing outward diffusion of Ni⁺ cations. Al₂O₃ forms internally due to inward diffusion of O²⁻ anions, while NiO is a p-type conductor, which preferably promotes conduction of cations [15–17]. Al₂O₃ growth onto alloys such as Ni₃Al can adopt different crystallographic arrangements: between 427 and 927°C, the γ-Al₂O₃ phase is stable; until 1077°C, the θ-Al₂O₃ phase is stable; and above 1077°C, the α-Al₂O₃ phase is stable [18]. The θ → α transformation is accompanied by a change in the transport properties of the scale. The θ-Al₂O₃ scale mainly grows by outward diffusion of cations, forming cavities in the scale/alloy interface, resulting in the deterioration of the adhesion layer on the substrate. In contrast, the growth of α scale is dependent on the diffusion of the O²⁻ anion inwards, so that the oxide fills the preformed cavities. This improves the adhesion of the layer on the substrate. Therefore, it is desirable for the growth of an α-Al₂O₃ layer [1, 14]. It has been observed that the only growth of an outer scale of Al₂O₃ is possible under the following conditions, by increasing both the temperature and concentration of Al [12], where the oxidation process is carried out at oxygen partial pressures below the equilibrium of NiO formation [13, 19–21] and also in alloys which were implanted with B⁺ and Cr⁺ ions [14]. Increasing the oxygen partial pressure and temperature, it is possible to form the spinel NiAl₂O₄ at the NiO/Al₂O₃ interface [22]. Garza et al. [23] attributed this behavior to the difference between the free energies of formation of Al₂O₃ and NiO which is four times larger than that between Al₂O₃ and NiAl₂O₄. Velon and Olefjord [15] mentioned that the spinel could also be formed in the initial stages of oxidation at low temperatures (300°C).

A key aspect to increase the oxidation resistance of an alloy is to improve the adherence of the protective oxide mainly because most applications implicate high-temperature thermal cycles [13, 24]. The most effective method is the addition of elements or compounds such as Hf, Y, Zr, Cr, Pd, Pt, and SiC [25–31], which can change the localized concentration of aluminum and reduce its growth rate. Gleeson et al. [32] indicate that the addition of Pt has beneficial effects in promoting the formation of a stable Al₂O₃ scale. Pt is an unreactive element whose addition decreases the chemical activity of Al in the γ′ phase, and it increases the Al/Ni ratio, which promotes the formation of Al₂O₃ in certain crystallographic planes and increases the Al diffusion into the formed scale. It has been calculated that the preferential site of substitution of noble metals like Ag, Au, Pd, and Pt in the Ni₃Al structure occurs in locations where there are normally nickel atoms [33–37]; therefore, the addition of these elements can modify the oxidation behavior of the alloy. Therefore, this paper evaluates the effect of the addition of these elements in low concentration (1 at.%) on the oxidation behavior of the Ni₃Al intermetallic alloy.

2. Experimental Procedure

The intermetallic alloys with nominal composition Ni₇₅Al₂₅ and Ni₇₅Al₂₄M₁ (where M = Ag, Au, Pd, and Pt) were manufactured in quartz crucibles in an induction furnace under inert atmosphere using elements with purity higher than 99.99%. The ingots were prepared by the fusion of stoichiometric amounts of the high-purity elements (solid rods), and the melt was homogenized by the convective currents generated by the induction furnace. After melting the metal elements, the ingots were allowed to cool to room temperature inside the induction furnace. From the ingots, samples were cut with dimensions of 20 × 5 × 2 mm to carry out the tests.

XRD analysis of alloys was carried with a Bruker D8 Advance diffractometer equipped with a Cu Kα (k = 1.5418 Å; 40 kV) monochromatic source and a LYNXEYE^TM 1D positional detector. Scans were performed from 2θ = 5° to 100° with a 0.02° step and an irradiation time of 0.2 seconds.

In order to reveal the microstructure of the intermetallic alloys, they were etched by immersing in HF-HNO₃ solution (1 : 1 molar ratio) for 5 to 10 seconds at room temperature.

For the oxidation test, samples were grinded with a silicon carbide paper to 600-grit size, washed with water, and ultrasonically cleaned with acetone. In order to have a larger reaction area, two samples were employed in each test: one was used for surface analysis and the other for cross-sectional analysis.

Oxidation tests were carried out in pure oxygen at 900, 1000, and 1100°C. This range was chosen since at these temperatures, there is a change in the stability of the different phases of Al₂O₃; at 900°C, transition between γ-phase and θ-phase begins; at 1000°C, θ-phase is stable; and at 1100°C, α-phase is stable. Before each test, the electric furnace was preheated to the experimental temperature. The record of the weight gain curves as a function of exposure time was done with a thermobalance CI MK2-M5 Head LM 2-01 and the LabWeigh software. Analyses were carried out by means of a scanning electron microscope (Zeiss model DSM 960) and the energy dispersive X-ray analysis (EDXA).

3. Results and Discussion

3.1. XRD Analysis of Alloys

The XRD profiles of the alloys with and without noble metal additions are shown in Figure 1. The XRD profile of the Ni₃Al alloy shows the characteristic peaks (111), (200), (220), (311), and (222) corresponding to the L₁₂ superstructure (γ′-phase), and their relative intensities are consistent with previous reports [38–40], which indicate a complete transformation from metals to the Ni₃Al intermetallic alloy.

The peaks observed in the XRD profiles of the doped alloys are wide, suggesting a small crystallite size. In addition, the intensities of some peaks increase and decrease depending on the added noble metal. The Ni₇₅Al₂₄Ag₁ profile shows highest intensities in (100) and (200), and the intensity of the peak decreased in (111); this suggests the crystal growth in two preferential planes. This preferential growth was also observed in the Ni₇₅Al₂₄Pd₁ alloy, where the profile shows a slight increase in the (222) peak, indicating a preferential unidirectional growth [41]. The XRD profile for the Ni₇₅Al₂₄Au₁ alloy shows a notable increase in the intensity of the (311) peak and a slight increase in the intensity of the (200) peak, and the characteristic peak (111) shows no change. On the other hand, the XRD profile of the Ni₇₅Al₂₄Pt₁ alloy shows an increase in the intensity of all its peaks, and an unidentified peak at 2θ = 29.56 is observed. Considering that the positions of the atoms in the unit cell affect the intensities but not the directions of the diffracted beams [42], the change in peak intensities can be completely attributable to the addition of noble metals.

Properties such as hardness show a significant dependence on crystallographic planes: for Ni₃Al, the highest value occurs in (100), followed by (111) and (110) [43]. However, several studies about the Ni₃Al oxidation indicate that there are no significant differences in the oxidation behavior between planes (100), (110), and (111) [44–48]. Taking into account that the planes (100) and (110) consist of a series of two layers: the most stable having a composition of 50% Ni and 50% Al and the other with 100% Ni and the plane (111) having a composition of 75% Ni-25% Al [49, 50]; then due to it, it is possible the presence of zones with different chemical compositions onto alloy surface (i.e., change in the Al/Ni ratio). In addition, substitution at preferential sites by the addition of noble metals can further change the surface chemical composition, thereby greatly influencing its oxidation behavior.

3.2. Microstructural Analysis

Figure 2 shows the microstructure of the intermetallic Ni₇₅Al₂₅ alloy, where the presence of two phases is observed. The EDXA showed that the dendrite is a phase with high concentration of Al, so that the (γ-γ′) phase is in the matrix and the (β-γ′) phase is in the dendrite core as indicated by Hunziker and Kurz [51] and Rosas et al. [52], respectively. During the solidification process, there is no cooling rate sufficiently slow to maintain equilibrium in the melt; that is, there is insufficient time for the composition gradients to be diffusion removed and the dendrites retained [53]. Its formation also may be attributable to the difference between the melting points of the constituent phases. According to Lee et al. [11, 17], when the solidification process begins, it is favored by the formation of the dendritic structures (β-γ′), and during the cooling process, the γ-phase decreases while the γ′-phase increases.

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Figure 3 shows the microstructure of intermetallic alloys evaluated, and it is observed that there are notable changes in the microstructure of the doped alloys, which can be explained according to the preferential growth plane of each alloy (Figure 1) and an increase in the solidification rate. The dendrite structure of Ni₇₅Al₂₅ (Figure 3(a)) has a spacing between secondary arms of 47 μm, and this value is typical in a conventional solidification process [54]. The space between the secondary arms of the dendrites is inversely related to the cooling rate during the solidification process, which has been correlated according to the following equation [55]:where d₂ is the secondary dendrite arm spacing, v is the solidification cooling rate, and k is a constant of the alloy.

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Since the solidification cooling rate is inversely proportional to the space between the secondary arms of the dendrites, from Table 1, it can be noticed that the solidification cooling rate in the Ni₇₅Al₂₄Pd₁, Ni₇₅Al₂₄Pt₁, and Ni₇₅Al₂₄Ag₁ alloys was faster than in the Ni₇₅Al₂₅ alloy, so that there is a lower space between the secondary dendrite arms; moreover, several directions of dendrites in the Ni₇₅Al₂₄Pt₁ alloy are congruent with its preferential growth plane. This same affirmation is consistent with the Ni₇₅Al₂₄Pd₁ alloy with almost only one preferential growth plane. The solidification cooling rate in the Ni₇₅Al₂₄Au₁ alloy was slower than the undoped alloy such that the dendrites almost disappeared.

It is observed that the addition of Ag favors an increase in the ratio γ-γ′/β due to the reduction in the number and volume of the interdendritic phase (Figure 3(b)). Because of the low solubility of Ag in Ni [56], it is possible to observe the presence of Ag-rich phases. According to Zhou and Guo [57, 58], the presence of Ag at grain boundaries begins during the solidification process since this element segregates at the solid-liquid interface, and at the same time, the Ag solubility decreases causing their precipitation into the grains. The microstructure of the Au-doped intermetallic alloy showed a honeycomb appearance (Figure 3(c)); that is, there is not a tendency to form precipitates as in the case of Ag-doped intermetallic alloy. Sheng et al. [59] found that the effect of addition of 1% Au on the microstructure of the NiAl intermetallic alloy causes a grain size reduction and the formation of α-AlAu₂ precipitates. It can be assumed that the absence of the α-AlAu₂ phase is due to the lower content of Al in Ni₃Al compared to that in NiAl. Figures 3(d) and 3(e) show the microstructure of the intermetallic alloy doped with Pd and Pt, respectively. The Pd-doped alloy has dendritic structures of bigger area; however, the relationship γ-γ′/β is much greater than in the alloy doped with Ag and Au, indicating that, at room temperature, the γ-γ′ phase is stable. The Pt-doped alloy has very fine dendritic formations in which there is no enrichment of Al. In both cases, a uniform distribution of Pd and Pt is observed as in the case of the Au-doped intermetallic alloy. Meng et al. [60] indicate that the addition of Pd into Ni₃Al intermetallic alloy generates an increase of delocalized bonding electrons, and this causes a strengthening of the grain boundaries and improves the ductility of the alloy. Furthermore, Li et al. [61, 62], Oquab and Monceau [63], and Zhang et al. [64] mention that Pd plays an important role in stabilizing the β-phase in NiAl at high temperatures; the same effect is attributed to the addition of Pt [10]. No precipitates were observed in the interdendritic region of the alloys, indicating that the phases found there are results of a eutectic reaction, contrasting with the work of Lee et al. [11], which presents cell interdendritic microstructures that Hunziker and Kurz [5] defined as peritectic structures.

3.3. Oxidation Kinetics

Figure 4 shows the curves (solid lines) of weight gain per unit area as a function of exposure time during the process of oxidation of the intermetallic alloys between 900 and 1100°C. The curves were fitted (dashed lines) to an equation of the form [24, 65]:where Δm is the weight gain per unit area, k is the oxidation rate constant, n is the exponent of velocity, and t is the time of exposure. The values of k and n to fit the experimental results are shown in Table 2. Adjustments were done using a nonlinear least-squares regression method (Gauss–Newton method) [66]. According to the n values obtained, it can be observed that only at 900°C, the oxidation kinetics of the Ni₃Al alloy and the Ag-doped alloy follows a parabolic behavior, all other n values correspond to a subparabolic behavior. The oxidation mechanism of a protective oxide that grows with parabolic kinetics mainly consists of the diffusion of cations from the metal-oxide interface to the oxide-gas interface; however, when the oxidation mechanism approaches to cubic kinetics, some authors attribute this behavior to the oxide grain growth [67, 68], although other authors claim that the formation of a protective layer like α-Al₂O₃ is more a rule than an exception [69]. An additional characteristic of the subparabolic kinetics is that at the beginning of the oxidation process, a great weight gain is observed with respect to that observed in a parabolic kinetic, being this a characteristic of the transient stage of the oxidation process.

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At 900°C, Ag- and Au-doped alloys had the highest weight gain throughout the exposure period with respect to the undoped alloy, whereas Pd- and Pt-doped alloys had the lower weight gain. During first two hours, the Au-doped alloy has a significantly greater weight gain, while Ag-, Pd-, and Pt-doped alloys have a similar and slightly higher weight gain than the undoped alloy. At 1000°C, all doped alloys (except Pd-doped alloy) have the greater weight gain with respect to the undoped alloy, in a decreasing order of Au-, Ag-, and Pt-doped alloys. This behavior may be because the doped alloys stabilize the θ-Al₂O₃ phase and therefore reduce the α-Al₂O₃ phase transformation as Pt does according to Cadoret et al. [70] and Hayashi and Gleeson [71]. The θ-Al₂O₃ phase generates more pathways for diffusion because of its smaller grain size compared to α-Al₂O₃ phase. The k values at 1000°C show magnitudes that explain the high increase in weight gain for doped alloys. At 1100°C, there is a clear grouping in terms of weight gain between the Pd- and Pt-doped alloys (lowest) and the Ag- and Au-doped alloys (higher) compared to undoped alloy, similar as at 900°C. The Pt-doped alloy has the lower weight gain than the undoped alloy throughout the exposure period, while the Pd-doped alloy has a higher weight gain than the undoped alloy only in the first six hours. All alloys except Pd-doped alloy have a remarkably high weight gain in the first hour.

The weight gain curves of Figure 4 give an impression about the effect of the addition of noble metals in the oxidation process, but it is also necessary to know the rate of oxidation. Considering strictly a parabolic behavior n = 0.5, for comparative purposes (2) can be rewritten aswhere k_p is the parabolic rate constant. The linearity of (3) allows graphically displaying and estimating the k_p value considering the slope of the Δm² − t relationship. In Figure 5, changes in the parabolic rate constant showing two k_p values can be identified, the first k_p value being greater than the second one; this is due to the fact that it is mathematically forcing a curve with a subparabolic behavior to be adjusted to a parabolic behavior [69], so the k_p value that will be taken into account to characterize the oxidation process will be that of the second slope which is presented in Table 3.

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At 900°C, undoped and Ag-doped alloys show a constant k_p value throughout the exposure period. The k_p value of Pd- and Pt-doped alloys is similar, and it is lower than that for the undoped alloy. At 1000°C, the k_p value of the Au-doped alloy decreases at three hours after starting the test, the k_p value of Ag, Pd, and Pt doped alloys decreases at four hours after starting the test, while for the undoped alloy, its k_p value remains constant. At 1100°C, the k_p value of the undoped and doped alloys decrease at the first half an hour after starting the test, except the Pt-doped alloy, for which the k_p value remains constant. In most alloys, at all temperatures, it is observed at most only one change in k_p, except the Pd-doped alloy at 1100°C with two changes: half an hour and two hours after starting the test. At 1000°C, the k_p value of the Au-doped alloy decreases at three hours after starting the test, and for Ag-, Pd-, and Pt-doped alloys, the k_p value decreases at four hours after starting the test, while for the undoped alloy, its k_p value remains constant. At 1100°C, the k_p value of all alloys (except Pt) decreases after 30 min, and for Pt-doped alloy, the k_p value remains constant.

When it is not possible to obtain accurate measurement from the slopes, it is necessary to consider the average of the instantaneous parabolic constants, which results in differentiating (3) versus time [65],

The instantaneous parabolic constant (k_p) of the oxidation tests is shown in Figure 6. This type of representation of the experimental data provides more information regarding the evolution of oxidation rates. At 900°C (Figure 6(a)), in the first two hours, the k_p values of the Au-doped alloy decrease seven times its equilibrium value, while in the Pd- and Pt-doped alloys, the k_p values decrease two times their equilibrium value; on the other hand, the Ag-doped alloy and undoped alloy maintain a constant rate. At 1000°C (Figure 6(b)), it is observed that the values of the instantaneous parabolic constant k_p of doped alloys tend to decrease in the first 6 hours, and then they remain constant, and for the undoped alloy, the k_p values remain constant. For the Ag-doped alloy, its k_p value decreases six times its equilibrium value, and for other doped-alloys, their k_p values decrease two times their equilibrium value. Finally, at 1100°C (Figure 6(c)), the Pt-doped alloy shows a stable behavior, and the other alloys show a fast decrease in their k_p values in the first hour.

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In order to visualize the effect of the temperature on the rate constants, it is necessary to determine the activation energy of the oxidation process from the Arrhenius equation.where Q is the activation energy, R is the universal gas constant, T is the temperature, k_p is the rate constant, and k₀ is the preexponential factor. Plotting ln k_p versus −(RT)⁻¹ from the slope of the lines, the Q value can be determined. Perez et al. [65] suggest assuming a parabolic kinetics in the whole temperature range; therefore, the average value of instantaneous k_p from Figure 6 was calculated. Table 4 shows the average values of instantaneous k_p.

In order to compare the results of this work, Figure 7 shows the instantaneous k_p values obtained along with those reported in other studies. At 900 and 1000°C, the k_p values of both doped and undoped alloys tend to k_p values where the formation of the θ-Al₂O₃ phase may be possible in NiAl as reported by Brumm and Grabke [72]. At 1100°C, the k_p values are concentrated at a point between the formation of θ-Al₂O₃ and α-Al₂O₃; an explanation for this behavior could be that, at this temperature, the stability of the α-Al₂O₃ begins [18] and this tends to decrease the weight gain [70, 71]. At 900°C and 1000°C, the k_p value of the undoped alloy is similar to results reported by Lee and Kim [26] and Choi et al. [73], and at 1100°C, it is similar to results of Kuenzly et al. [74]. The Pt-doped alloy has k_p values similar to the undoped alloy, Ag- and Au-doped alloys have higher k_p values, and the Pd-doped alloy has lower k_p values.

Activation energy is an important parameter that characterizes the predominant oxidation mechanism [75]; for example, the results presented by Ochiai et al. [75] and Kuenzly et al. [74] suggest that, in both cases, the oxidation mechanism is similar, although the experimental conditions were different, that is, cyclic oxidation tests and continuous oxidation tests. On the other hand, analyzing the results for the formation of α-Al₂O₃ scale [72] and NiO scale [76], it is observed that the preexponential factor is inversely proportional to the protective capacity of the protective oxide. Table 5 shows the activation energies and preexponential factors obtained from the Arrhenius graph (Figure 7). The activation energy value of Ni₃Al is similar to those values reported by Lee and Kim [26] and Cao et al. [77], but almost three times lower than that reported by other authors [73, 74]. This difference is because the intermetallic alloy used in this study was not thermally treated. Activation energy values of the doped alloys show that the additions of the noble elements affect the Ni₃Al oxidation mechanism in the following order: Pt > Au > Ag > Pd. On the other hand, the Pd-doped alloy showed the lowest value of the preexponential factor.

3.4. SEM Analysis

SEM analysis showed that the oxide scale formed on all alloys had nonuniform thicknesses. In the case of Ni₃Al, the average thicknesses observed were between 2.44 and 2.86 μm in the temperature range studied. At 900°C (Figure 8), a double-layer oxide was developed, an outer Ni-rich layer and an internal Al-rich layer, with similar thicknesses. Other authors have described similar results [78]. However, at higher temperatures, only the presence of an Al-rich oxide was detected.

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In general, the thickness of the oxide layer developed onto doped alloys was lower than that observed onto Ni₃Al. At 900°C, the average thicknesses were between 0.98 and 1.5 μm, at 1000°C, between 2.13 and 2.20 μm (except the Pt-doped alloy, its average thickness was 4.61 μm), and at 1100°C, between 1.43 and 2.55 μm (except the Ag-doped alloy, its average thickness was 4.52 μm). Figure 9 shows the thickness of the oxide formed on the intermetallic alloys oxidized to 900°C; similar aspects were observed at higher temperatures. At all test temperatures, the Pd-doped alloy showed the oxide layer with the lowest thickness (Figure 10(a)). In general, all doped alloys developed an Al-rich monolayer oxide (Figures 10(a)–10(c)), as has been reported in other studies [79]. In the case of alloys doped with Pt (at 1000°C) and with Ag (at 1100°C), zones with multilayer oxides and thicknesses up to 20 μm were observed (Figures 10(d) and 10(g)). In all cases, an Al-depleted zone was observed at the interface alloy oxide (Figures 8 and 10). This is due to the consumption of this element for the formation of protective oxide, Al₂O₃ [80]. Similar depletion processes have been reported in studies of corrosion by molten salts [81–85].

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Figure 11 shows the surface morphology of the oxides formed at 900°C. The formation of oxides in the form of needles can be observed only onto nondoped alloy and onto those alloys doped with Ag and Au. This type of morphology is associated with the presence of the metastable phases of alumina [86–88], whose transformation from γ-Al₂O₃ to θ-Al₂O₃ occurs between 800 and 950°C [72]. It is worth mentioning that the alloys that presented this type of morphology had a greater weight gain (Figure 4(a)).

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Figure 12 shows the surface morphology of the oxide scale at 1000°C. It is observed that, at this temperature, there is a greater presence of oxides in the form of needles (θ-Al₂O₃) both in the nondoped alloy and in those alloys doped with Ag, Au, and Pt. The alloy doped with Pd showed a lower presence of this phase, possibly due to a decrease in the growth rate of the θ-Al₂O₃ phase. Again, the surface morphology is correlated with the observed weight gain (Figure 4(b)); that is, the greater the presence of the θ-Al₂O₃ phase, the greater the mass gain.

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On the other hand, the morphology of the oxides formed at 1100°C was completely different (Figure 13). In this case, the oxides formed show wrinkle shaped, which is characteristic of the α-Al₂O₃ phase [72]. This morphology is more evident in the nondoped alloy and in those doped with Ag and Au, and in the alloys doped with Pd and Pt, a more compact and dense morphology was observed. These morphological differences indicate a difference in the growth rates of the oxide layers. In all cases, the weight gain at 1100°C was lower than that observed at 1000°C (Figure 4). This suggests that oxide layers rich in α-Al₂O₃ are more protective than those rich in θ-Al₂O₃.

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Figure 14 shows the typical morphology of the Ni₃Al oxidized in oxygen for 24 hours. In general, a granular scale involving both heterogeneous Ni and Al oxides, mainly Ni oxides, as reported by Cao et al. [77] were observed. At 900°C, there are well-defined regions rich in Al₂O₃ which largely coincide with the interdendritic regions (Figure 14(a)). This suggests that the lateral growth of NiO could not fully cover the Al₂O₃, and the initial stage of oxidation occurred preferentially in the interdendritic phases [11]. At 1000 and 1100°C, the Ni₃Al surface was oxidized uniformly in comparison with the oxidized alloy at 900°C, but the surface mainly contains Al₂O₃ (Figures 14(d) and 14(g)). In some regions, the oxide layer has been cracked, probably during cooling, and it was possible to observe the presence of cavities into the scale-alloy interface (Figure 15). These sorts of cavities are formed by the consumption of Al and Ni in the course of formation of the oxide scale, and they have an impact on the adhesion of the scale [70].

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The difference in the oxidation process between the dendritic and the interdendritic phases lies in the thermochemistry of alumina formation between the phases; that is, in the NiAl-rich dendritic phase, the following reaction may occur:

On the other hand, in the interdendritic phase where the Ni3Al phase predominates, the following reaction occurs:

Thermochemically, the dendritic phase is more favored than the interdendritic phase. Considering also that in the doped alloys there is a considerable decrease in the dendritic phase (Section 3.2), it can be established that there exists a relation between the microstructural change and the thermochemical equilibrium conditions for the formation of an alumina layer. On the other hand, the absence of NiO in the surface of the doped alloys can also be attributed to the increase in the grain size of the Al₂O₃; that is, to say, a decrease in the diffusion of Ni through the grain boundaries of the Al₂O₃ layer, which is related to the kinetics of subparabolic growth.

4. Conclusions

The XRD analysis of the alloys showed that the addition of 1% (atomic percent) of noble metals to the Ni₃Al intermetallic alloy promotes the preferential growth planes and increases the solidification rate by changing the microstructure of the base alloy. The oxidation behavior of doped alloys may be attributable to the change in the surface composition, taking into account that every growth plane has different compositions and that the preferential substitution site of the noble metals in the structure of Ni₃Al increases the Al/Ni ratio. In general, the addition of noble metals promotes the formation of oxide layers rich in Al₂O₃ and is evident that the diffusion of the Ni cations through Al₂O₃ scale is decreased, since the formation of NiO during the test time was not evident. Ag and Au additions increase the value of the parabolic rate constant k_p, while Pd and Pt decrease, the k_p values tend to values where the formation of θ-Al₂O₃ and α-Al₂O₃ is possible. The noble metal addition increased the activation energy value of the oxidation process compared to the nondoped Ni₃Al, except Pd. The low solubility of Ag and its low melting point caused microstructural localized failures into Ag-doped alloy.

Conflicts of Interest

The authors declare that they have no conflicts of interest regarding the publication of this paper.

Acknowledgments

Financial support from Consejo Nacional de Ciencia y Tecnología (CONACYT, México) (Project no. 159898 and Ph.D. scholarship to J. J. Ramos-Hernandez, Registration no. 202898) is gratefully acknowledged.

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Effect of the Noble Metals Addition on the Oxidation Behavior of Ni3Al

Abstract

1. Introduction

2. Experimental Procedure

3. Results and Discussion

3.1. XRD Analysis of Alloys

3.2. Microstructural Analysis

3.3. Oxidation Kinetics

3.4. SEM Analysis

4. Conclusions

Conflicts of Interest

Acknowledgments

References

Copyright

Effect of the Noble Metals Addition on the Oxidation Behavior of Ni₃Al