Oxidation resistance of Si-modified aluminide coatings on Mar-M246 superalloy using the halide activated pack cementation technique

Abstract

In the present work, the Halide Activated Pack Cementation (HAPC) was adopted to form protective aluminide coatings due to its versatility and low cost. Codeposition of Al and Si on Mar-M246 was performed by HAPC using two different cements constituted of 75 wt.% Al₂O₃ and 25 wt.% of a mixture formed in different weight ratios of (Al: Si): PM I (50:50) and PM II (72.5:27.5). The coatings are composed in outer part of NiAl and Ni₂Al₃ phases with significant presence of Si on PM I. After 200 h of oxidation test at 1000 °C, the samples coated by PM I and PM II exhibited the best performance than uncoated one, with mass gains near 0.3 mg/cm², which is half times the mass gained by uncoated alloy. All coated samples were characterized before and after the oxidation test by SEM, EDS and XRD techniques. The exposure of the coated to high temperature led to the formation of Al₂O₃ and a modification of the coating morphology by the dissolution of the Ni₂Al₃ phase.

Keywords:
Mar-M246 Superalloy; Thermodynamic Analysis; Halide Activated Pack Cementation; Aluminizing and Siliconizing; High Temperature Oxidation

1. Introduction

Mar-M246 is a well-known Ni-based superalloy that is largely used for high temperatures application in aerospace engines as and in automotive turbocharger. This alloy operates at temperatures in the range 800-1000 °C, combined with high stress loads. For the MAR-M246, the values of the oxidation rate constant (Kp) are in the 10^-14-10^-12 g²cm^-4s^-1 range, from 800 up to 1000 °C respectively, which is typical for alumina forming alloys (Denzine, et al. 1976; Johnston e Parr, 2014). Thus, long term creep and oxidation by combustion gases are the main source of damage that the alloy must endeavor (Denzine, et al. 1976; Johnston e Parr, 2014). Recent investigations have shown that intergranular oxidation and degradation of carbides in the subsurface region at temperatures near 950 °C has deleterious effects on the creep behavior of this class of materials (Denzine, et al. 1976). Thus, many investigations have shown that coatings can be applied on the surface of Ni-based superalloys to increase their resistance at high temperatures in aggressive environments and enhance the creep behavior during utilization.

The application of surface coatings enables a potential increase in operating temperature and/or service life, consequently enhancing the performance of parts and components manufactured from the respective alloy. This proves to be critical in power generation plants, allowing for reduced fuel consumption and emissions of pollutants such as CO₂ and NO₂, while also facilitating improved efficiency in thermal processes (Yan et al., 2014). The most efficient way to achieve a good protection is the use of coatings, which can be obtained by different processes: PVD (Physical Vapour Deposition), CVD (Chemical Vapour Deposition), HAPC (Halide Activated Pack Cementation), Thermal Spray Coating and Laser Surface Alloying (Tsipas et al., 2007; Liu et al., 2008; Prasad e Paul, 2011).

The Halide Activated Pack Cementation (HAPC) is versatile and easy-to-implement process, used for covering parts of complex geometry and produces adherent and uniform layers with smooth surfaces and controlled thickness. Several investigations show that this process is efficient for protecting superalloys and steels for different type of industrial applications (Xiang et al., 2003; Pistofidis et al., 2006; Cheng e Park, 2013; Yuan et al., 2013; Lange et al., 2015). The deposition occurs at high temperature (650-1200 °C) using a powder mixture called cement, composed of an inert filler material (Al₂O₃, for example), powder of the metallic elements to be deposited (e.g., Al, Si, Cr or a intermetallic one) called masteralloy, material that will be deposited on the surface of the substrate, and a halide salt (NaCl, NaF, NH₄Cl) as activator (Bianco e Rapp, 1993). Protection is achieved when the coating oxidizes at high temperatures, leading to the formation of protective oxides (e.g., Al₂O₃, SiO₂, Cr₂O₃) that inhibit the accelerated progression of oxidative processes on the surface.

For the Mar-M246 superalloy, the aluminization process using pure aluminum as the masteralloy was introduced in the work by Glória et al. (2021). In this study, the proposal is to coat the Mar-M246 superalloy using the aluminization process with a masteralloy composed of aluminum and silicon (Al-Si). In this context, the effect of silicon in the HAPC process for coating formation and post-oxidation tests will be analyzed.

2. Materials and experimental methods

2.1 Thermodynamic analysis

A Thermodynamic analysis can be employed as a tool to predict the possibility of depositing the coating elements contained in the masteralloy (Levine e Caves, 1974). In this case, the chemical equilibrium occurring between the surface of the alloy to be coated and the deposition gases generated in the chemical reactor by the gaseous metal halides, the latter being formed in the reaction between the masteralloy and the activating salt, is assessed. Once this equilibrium is established, the partial pressure of the gaseous halides can be determined as a function of the process temperature. It is known that higher partial pressures of gaseous halides suggest a greater potential for depositing the coating elements, and the minimum deposition pressure is in the order of 10^-6 to 10^-5 atm. Several software programs can be used to assist in thermodynamic analyses (Oates, 1987; Smith, 1996; Bale et al., 2016), and in this study, HSC Chemistry 6.0 was used to generate results to identify the primary gaseous halides formed and evaluate their deposition capacity as a function of their respective partial pressure and the process temperature used in the HAPC experiments.

2.2 Procedures for the HAPC process and oxidation testing

Discs of the Mar-M246 alloy (Ni_bal-10W-10Co-9Cr-5.5Al- 2.5Mo-1.5Ti-1.5Ta-0.15C-0.015B in wt%), of 7 mm diameter and 1.5 mm thick, were cut from vacuum arc melted bars. Prior to the pack cementation process, the substrates were grinded with SiC abrasive papers up to P2400 grit and polished with 0.05 μm alumina suspension. The samples were cleaned ultrasonically in isopropanol, cleaned with water, and then dried with hot air. For the aluminization process, two powder mixtures (PM I, PM II) composed of 75 wt.% of Al₂O₃ and 25 wt.% of a mixture of Al (99.75% pure) and Si (99.999% pure) elementary powders were used. The proportions of Si and Al in the 25 wt.% of PM were:

PM I: 50 wt.% of Si and 50 wt.% of Al;

PM II: 27.5 wt.% of Si and 72.5 wt.% of Al.

Because of its low thermodynamic stability, NH₄Cl was used as the activating salt. The amount of NH₄Cl used to activate the cement was approximately 15 mg to generate a reactive gas with approximately 1 atm. Samples of the MAR-M246 alloy to be coated were introduced together with cements in silica tubes sealed under primary vacuum. The HAPC processes were performed for 1, 9, 16 and 25 h at 1000 °C using a muffle furnace. After cooling, the reactors were removed and broken, the coated substrates were cleaned in ultrasonic bath with isopropanol and dried under hot air.

For the cyclic oxidation tests, the samples were placed on an alumina crucible and introduced in a tubular furnace under static laboratory air. The samples were oxidized under laboratory air at 1000 °C for a duration of 240 h. For mass change monitoring, the specimens were frequently removed from the furnace within a 24 h interval, cooled down up to room temperature, and hand weighted using an analytical balance with 0.1 mg of accuracy. Three specimens of each composition were utilized.

2.3 Metallographic characterization

The phases present in the coated substrates and the oxidation products formed after the oxidation tests were identified by XRD, on the surface of the specimens. The measurement were realized using Cu-Kα radiation, diffraction angle (2θ ) ranged from 10 to 90º, step size of 0.02º and a counting time of 1s per step. Identification of the different phases was performed using the PowderCell software (Kraus e Nolze, 1996) and a database (Villars e Calvert, 2007).

For the metallographic characterization, the coated samples and the oxidized ones were cold mounted in epoxy resin and prepared using conventional metallographic methods. The characterizations of the coating composition, thickness and microstructure were conducted by Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS).

3. Results and discussion

3.1 Thermodynamic analysis

Figure 1 shows the Gibbs Free Energies of formation of some fluorides and chlorides between 0 and 1200 °C (Glória et al., 2021). Chloride-based activators are the most indicated due their low thermodynamic stability at high temperatures. Consequently, chlorides can provide greater amounts of the gaseous species of the donor elements forming a reactive rich in Si and Al. Among the chlorides most used, NH₄Cl has the lowest stability, therefore it was chosen to activate the PMs of the cement used in the present work.

Figure 1
Free Gibbs energy formation for some activators considering 1 mol of Cl₂(g)/F₂(g).

Figure 2 shows the composition of the gas phases generated by (a) PM I and (b) PM II as a function of log of partial pressures and temperatures which varies between 700-1200 °C. In both cases, it is evident that the aluminum deposition process has a high probability of occurring on the alloy's surface. This is attributed to the formation of gaseous aluminum-based metal halides with partial pressures above the minimum deposition pressure (10^-6-10^-5 atm), considering the temperature of 1000 °C used in the HAPC experiments. However, the formation of gaseous silicon-based halides was considerably lower in terms of quantity and pressure when compared to the aluminum-based metal halides. Therefore, it is likely that the silicon deposition process does not occur with the same intensity as the aluminization process. Additionally, it only has a significant chance of occurring under condition PM I, as in this condition, only the presence of the SiCl₂(g) metal halide was observed near the minimum deposition pressure.

Figure 2
Formation of activated halides in the reaction of pure Al and Si with the activator NH₄Cl in (a) PM I and (b) PM II.

Below, the primary gaseous metal halides identified as potential candidates for the deposition of aluminum and silicon on the alloy's surface under PM I and PM II conditions have been listed. In both cases, the observed pressures are either above or close to the minimum required for the deposition process in HAPC experiments. It is essential to note, however, that not all gaseous metal halides effectively deposit on the surface, even when they exhibit sufficient partial pressures. In such cases, more complex thermodynamic analyses involving the interaction of the substrate with the deposition gas need to be conducted (Levine e Caves, 1974).

PM I: AlCl₃(g), AlCl₂(g), AlCl(g), Al₂Cl6(g) and SiCl₂(g);

PM II: AlCl₃(g), AlCl₂(g), AlCl(g), Al₂Cl₆(g).

Based on these data, it can be assumed that codeposition de Al e Si may take place by using the PM I, since aluminum and silicon halides have partial pressures of the same order of magnitude, which is a sine qua non condition for codeposition as demonstrated previously by Bianco and Rapp (Bianco e Rapp, 1993).

3.2 Morphology of the coated substrates

Figure 3 shows the surface x-ray patterns of the 2 coated samples. The highest intensity peaks near 43°, 18°, 25° and 85° indicate the presence of NiAl and Ni₂Al₃ phases. In addition, peaks near 30°, 35° and 40° indicate the presence of Al₂O₃ that can be residual particles sintered on the surface of the coating, or particles incrusted in the structure of the external layer caused a possible external growth. In PMs I and II, it is possible to observe a series of smaller peaks in the region of 35 - 50° that indicate the formation of an Al-Cr intermetallic that occurs due to the high activity of aluminum (Tong et al., 2010).

Figure 3
X-ray diffractogram of the surface formed by the coatings.

Coatings obtained by PM I and PM II are quite similar in terms of morphology and chemical composition as showed in the Figure 4 and Figure 5. The coatings are composed of a single layer having approximately 50 at.% of Al. The coating obtained by the PM I cement is approximately 140 µm thick while that obtained by the PM II cement is almost 170 µm tick. It is worth to note the presence of small precipitates in both coatings, evenly distributed through the coating thickness. Because of their small size, it was not possible to identify their nature. However, since they are characterized by higher chemical contrast, one can suppose that they are rich in refractory elements, originally present in the substrate, that do not dissolve into the NiAl and Ni₂Al₃ aluminides.

Figure 4
a) Micrograph of the coating formed by PM I b) Atomic distribution profile in the coating formed by PM I c) Si atomic distribution profile in PM I.

Figure 5
a) Micrograph of the coating formed by PM II b) Atomic distribution profile in the coating formed by PM II.

As observed, silicon deposition was only feasible under the PM I condition, in contrast to the non-deposition of this element under the PM II condition. The obtained results are in good agreement with the predictions from the thermodynamic analyses conducted, which indicated that only under the PM I condition is there a genuine possibility of Si deposition by the SiCl₂(g) metal halide. However, the silicon concentration throughout the coating remained below 7 at.%, as shown in the Figure 4.c.

Furthermore, it is worth nothing that the concentration profile and morphology of these coatings are very similar to the results obtained by Glória et al. (2021) for the aluminization of the Mar-M246 Alloy using pure Al powder as masteralloy in a cement activated also by NH4Cl. Although the similarities observed for these coatings, it seems reasonable to state that the presence of Si in the PM I and the PM II relatively reduces the chemical activity of Al, since there was a great difference in the final thickness of layers obtained PM I and PM II in this work, and the use of pure Al in cement leading to a coating of 290 µm for the same conditions of time and temperature (Glória et al., 2021). The presence of carbides in the microstructure of the coatings from the PM I and PM II indicates an inward growth, i.e. a higher mobility of Al compared to Ni during the growth of the layer.

3.3 Oxidation test

Two specimens coated under the PM I and PM II conditions, along with the uncoated Mar-M246 alloy, were exposed to high temperatures in laboratory air to assess their oxidation resistance. Figure 6 presents the specific mass gain versus the exposure time. It can be observed that samples coated under the PM I and PM II conditions exhibited excellent oxidation resistance. Similarly, the uncoated sample notably demonstrated good oxidation resistance, albeit inferior to the coated ones.

Figure 6
Mass gain by surface area in an oxidation test carried out for 240 hours at 1000°C.

In other studies, oxidation tests on the uncoated conventional Mar-M246 superalloy were also performed at 1000 °C. These tests were distinguished by the execution mode, involving quasi-isothermal tests (Alkmin et al., 2021) and thermal cycling tests (Cunha et al., 2021). In both cases, the oxidation tests predicted a good performance for the uncoated Mar-M246 superalloy, with the stabilization of oxidative processes, as observed in this study. However, according to other work (Gloria et al., 2021) under the same oxidation test conditions and referred to as a reference (Figure 6), the Mar-M246 superalloy exhibited a significant mass gain up to 72 hours, followed by a decrease in mass mainly due to the spalling of the oxides scale. Thus, it is suggested that temperatures around 1000 °C can be highly critical in certain situations for the uncoated Mar-M246 alloy, leading to pronounced oxidation of the Mar-M246 alloy.

The results of the XRD patterns of PM I and PM II performed on the surface of the oxidized samples are shown in Figure 7. Specific peaks associated to NiAl and Al₂O₃ can be identified in the diffractogram. In this case, it is suggested that during the oxidation of the coatings, an induced phase transformation of the previously existing Ni₂Al₃ phase into NiAl and Al₂O₃ occurred according to the following reaction:

Figure 7
XRD patterns of PM I and PM II oxidize surfaces.

$$2{\mathrm{Ni}}_2{\mathrm{Al}}_3+\frac{3}{2}{\mathrm{O}}_2\to 4\mathrm{NiAl}+{\mathrm{Al}}_2{\mathrm{O}}_3\mathrm{\Delta }{\mathrm{G}}^0=-1189.8\frac{\mathrm{kj}}{\mathrm{mol}}$$

In both coated conditions (PM I and PM II), surface protection was promoted by the formation of the Al₂O₃ oxide layer, originating from the oxidation of the aluminized coating. In the EDS analysis of the PM I coating (Figure 4b), small amounts of silicon were observed in the coating. However, selective oxidation of silicon was not achievable following the oxidation test, resulting in the formation of SiO₂ oxides or more complex combinations involving this element. Additionally, the addition of silicon to the Al-Si masteralloy could have the following key effects on coating formation and oxidation tests.

I. Reduction of the activity of pure aluminum in the Al-Si masteralloy;

II. Enrichment of the coating by Si during deposition;

III. Formation of silicon-based protective oxides or the selective induction of other protective elements through the presence of silicon.

Considering the proportions used for the masteralloys (PM I and PM II), the most significant effect of silicon in the Al-Si masteralloy was its ability to reduce the activity of pure aluminum compared to other studies (Glória et al., 2021). This can be of interest when aiming to control growth, defects, and phases present in the coating, as observed in similar studies conducted with the Al-Cr masteralloy for steels (Wang et al., 2010). However, the deposition of silicon along with aluminum becomes significantly hindered under the conditions employed for PM I and PM II. In a similar vein, small concentrations of silicon in the coating appear to indicate that they do not induce its selective oxidation for the formation of protective oxides, as indicated by the obtained results in this work.

Figure 8 and Figure 10 show the cross sections as well as the measured concentration profiles for the oxidized PM I and PM II coatings respectively; Figure 9 and Figure 11 show the EDS elemental mapping for the oxidized PM I and PM II, respectively, cross sections. The SEM analysis reveals major microstructural modifications of the coatings morphology after the oxidation test. Two regions were identified on the cross sections of the oxidized samples: i) an external part with a thickness of approximately 145 μm and 200 µm for the PM I and PM II coatings respectively. For the PM I coating, the concentration of Al is close to 40 at. % near the subsurface region reaching values near 35 at. % close to the interdiffusion region. For the PM II coating, the concentration of Al is close to 45 at. % near the subsurface region reaching values near 35 at. % close to the interdiffusion region. ii) For both coatings, an internal region called IZ was formed with a thickness of 65 μm approximately, characterized by the presence small lamellar precipitates. These phases are rich in refractory elements (Ta, W and Ti) and are known as Topologically Close-Packed (TCP) phases. They appear because of Ni diffusion to the outer layer during the exposure of high activity aluminide coatings to elevated temperatures (Kearsey et al., 2004).

Figure 8
PM I a) SEM micrograph and b) atomic distribution profile of the coating obtained after 240 hours of oxidation test.

Figure 9
PM I EDS elemental mapping after 240 hours of oxide.

Figure 10
PM II a) SEM micrograph and b) atomic distribution profile of the coating obtained after 240 hours of oxidation test.

Figure 11
PM II EDS elemental mapping after 240 hours of oxidation test.

Remarkably, following the oxidation tests, the coating thickness increased under PM I conditions (from 140 µm to 200 µm) and PM II conditions (from 170 µm to 270 µm), while the aluminum concentration profile tended to decrease locally. Due to the high temperature of the oxidation test and the extended time, an interdiffusion phenomenon may have occurred (Bates et al., 2009), where aluminum in the coating tends to diffuse by solid-state diffusion on two fronts, being consumed at the surface through selective oxidation, and also diffusing into the substrate region due to the concentration difference. The interdiffusion phenomenon can be highly critical for coatings with low aluminum concentration and ultra-thin characteristics when exposed to high temperatures for extended periods (>20,000 h). In contrast, the coatings obtained under PM I and PM II conditions demonstrated favorable conditions to withstand the interdiffusion.

The coating layer of PM I was the only one chosen to determine the growth kinetics since it has presented modification of its composition with Si and demonstrated oxidation resistance similar to PM II. The growth kinetic was investigated at 1000 °C for durations of 1, 4, 9 and 16 h. The Table 1 presents the total thickness values for the different times of aluminization.

Table 1
Coating thicknesses obtained by PM I for 1, 4, 9 and 16 h of HAPC process at 1000 °C.

Time (h)	Thickness (µm)
1	45.7
4	127.9
9	131.2
16	231.2

The kinetic law involved in the growth of the coating can be described by a general Equation 1 as proposed by Kofstadt (1988):

(1)$$\mathrm{\Delta }{x}^n=K\times t$$

With: ∆x is the coating thickness; K is the growth constant; t is the time; n is the kinetic order.

The kinetic order can be obtained from the double logarithmic plot log(x) versus log(t) according to the Equation 2:

(2)$$\log \mathrm{\Delta }x=n^{-1}\log (t)+C$$

The value of “n” calculated from the log (x) versus log(t) as shown in the Figure 12 is 1.85, very close to the parabolic law (n=2). Thus, it can be assumed that the growth of the layer occurs by solid-state diffusion according to the Equation 3:

Figure 12
Logarithm profile of the thickness variation of the coating as a function of time for PM II.

(3)$$\mathrm{\Delta }{x}^2=K_p\times t$$

The parabolic constant can be determined from the plot of the thickness variation versus the square root of time as presented in the Figure 13. The value for K_p was 8.33 x 10^-9 cm²s^-1, smaller than that obtained at 1000 °C (3.14 x 10^-8 cm²s^-1) (Gloria, 2021) by pure Al PM, indicating the drop in Al's reactivity when mixed with Si in the PM, already discussed above.

Figure 13
Linearized profile of parabolic growth of the coating as a function of time root for the HAPC process using PM II.

4. Conclusions

A coating layer was successfully created by utilizing HAPC with Al and Si powders on the surface of MAR-M246 nickel superalloy. The resulting layers were cohesive, adherent, devoid of cracks, measuring between 140 to 170 microns thick. They principally composed of NiAl and Ni₂Al₃ phases, as revealed using XRD. Two mixtures, designated PM I e II, were tested where only PM I demonstrated satisfactory co-deposition characteristics for both elements, which corresponded well with thermodynamic simulations presented within this study. Upon evaluating coatings performance during oxidation testing against uncoated samples showcased mass gains that surpassed slightly more than half that observed from non-coated specimens (0.5 mgcm^-2 compared to Uncoated Samples' value, 0.9 mgcm^-2. Those coated under PM II gained even less weight, 0.3 mgcm^-2). The predominance of NiAl and Al₂O₃ phases on the surface after the oxidation tests was also observed. Lastly, the layer growth for PM II during the cementation process was calculated, confirming that the presence of Si in the mixture decreased the activity of Al.

Acknowledgements

The authors gratefully the financial support of CAPES. Also the Alcoa Alumínio S/A Aluminum Powder Plant for supplying the aluminum powder used in this work.

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Publication Dates

• Publication in this collection
  24 Mar 2025
• Date of issue
  2025

History

• Received
  02 May 2024
• Accepted
  06 Aug 2024