Study on Preparation of Nano-CeO₂ Modified Aluminized Coating by Low Temperature Pack Aluminizing on γ-TiAl Intermetallic Compound

Abstract

TiAl alloy offers advantages including low density, high specific strength and stiffness, and excellent high-temperature creep resistance. It is widely used in the aerospace, automotive, and chemical sectors, as well as in other fields. However, at temperatures of 800 °C and above, it forms a porous oxide film predominantly composed of TiO₂, which fails to provide adequate protection. Applying high-temperature protective coatings is therefore essential. Oxides demonstrating protective efficacy at elevated temperatures include Al₂O₃, Cr₂O₃, and SiO₂. The Pilling–Bedworth Ratio (PBR)—defined as the ratio of the volume of the oxide formed to the volume of the metal consumed—serves as a critical criterion for assessing oxide film integrity. A PBR value greater than 1 but less than 2 indicates superior film integrity and enhanced oxidation resistance. Among common oxides, Al₂O₃ exhibits a PBR value within this optimal range (1−2), rendering aluminum-based compound coatings the most extensively utilized. Aluminum coatings can be applied via methods such as pack cementation, thermal spraying, and hot-dip aluminizing. Pack cementation, being the simplest to operate, is widely employed. In this study, a powder mixture with the composition Al:Al₂O₃:NH₄Cl:CeO₂ = 30:66:3:1 was used to aluminize γ-TiAl intermetallic compound specimens via pack cementation at 600 °C for 5 h. Subsequent isothermal oxidation at 900 °C for 20 h yielded an oxidation kinetic curve adhering to the parabolic rate law. This treatment significantly enhanced the high-temperature oxidation resistance of the γ-TiAl intermetallic compound, thereby broadening its potential application scenarios.

1. Introduction

Three intermetallic compounds in the Ti–Al binary system have attracted significant research interest: Ti₃Al, TiAl, and TiAl₃ [1]. Among these, TiAl alloys have emerged as highly promising lightweight and high-temperature structural materials due to their advantageous properties, including high melting point, low density, high specific strength, excellent high-temperature creep resistance, and robust oxidation resistance at elevated temperatures [2]. These alloys are commonly employed in critical components such as gas turbine blades, engine exhaust valves, and compressor impellers. Rapid advancements in aviation, energy, and defense industries necessitate continuous improvements in energy utilization efficiency, driving progressive increases in turbine inlet gas temperatures. Combustion chamber temperatures in advanced aero-engines and large-scale ground-based gas turbines now reach 1700 °C [3]. However, Wu et al. [4] report that when TiAl alloys operate above 800 °C, they form predominantly non-protective rutile TiO₂ scales or mixed TiO₂/Al₂O₃ scales, which severely limit their extended applicability. Consequently, the application of high-temperature protective coatings is essential in practical implementations [5].

Typical high-temperature protective coatings include aluminum-based coatings, modified aluminum coatings, cladding coatings (MCrAlY coatings, where M = Fe, Ni, Co, or combinations thereof), and thermal barrier coatings [6]. Tang et al. [7] selected Ti-50Al-10Cr alloy as a coating material, depositing it onto γ-TiAl alloy via magnetron sputtering. After oxidizing the TiAlCr coating at 900 °C for 1000 h, cross-sectional analysis revealed a dense Al₂O₃ scale on the coating surface with minimal interdiffusion between the coating and substrate. Mengis et al. [8] aluminized Ti-48Al-2Cr-2Nb alloy through pack cementation at 600 °C for 4 h, obtaining a ∼10 μm thick coating. The coated sample exhibited a final mass gain at 700 °C approximately one-third that of the uncoated specimen. Following 1000 h of cyclic oxidation, a thin, adherent Al₂O₃ scale and an Al-rich interdiffusion zone formed on the aluminized surface, with internally precipitated TiO₂ displaying α-Al₂O₃ and rutile structures. Swadzba et al. [9] synthesized silicon-modified aluminide coatings on Ti-45Al-5Nb-0.2B-0.2C using pack cementation (850 °C, 6 h) with a penetrant mixture of Si powder, Al powder, Al₂O₃ filler, and NH₄Cl activator. The resulting ∼43 μm coating exhibited a bilayer structure: an external TiAl₃ matrix and an intervening ∼400 nm equiaxed TiAl₂ layer at the substrate interface, accompanied by dispersed Ti₅Si₃ phases. Isothermal oxidation at 950 °C for 3000 h produced a continuous, uniform ∼2.5 μm α-Al₂O₃ scale. Hao et al. [10] fabricated Al₂O₃–Y₂O₃ composite coatings on γ-TiAl alloy via electrophoretic deposition (EPD). These coatings significantly enhanced the alloy's oxidation and scale spallation resistance by suppressing outward Ti diffusion from the substrate while promoting selective Al oxidation within the γ-TiAl alloy.

Among high-temperature protective coatings, aluminide coatings are widely employed industrially due to mature processing technology and low production costs. However, traditional aluminizing temperatures typically range from 850 to 1050 °C. These elevated temperatures have a limitation: when applied at high temperatures, they can cause damage and degradation to the substrate material [11]. Combined with thermal stresses, this promotes internal cracking and deformation, further exacerbating mechanical property degradation [12]. Lower aluminizing temperatures also offer significant energy savings. Current research on low-temperature aluminizing primarily focuses on formula development, achieving reduced processing temperatures through addition of infiltration promoters, active elements, and compositional adjustments. Nevertheless, temperatures rarely fall below 650 °C. In this study, TiAl alloy served as the substrate. Through optimized formula design and incorporation of 0-1% rare earth oxide CeO₂ [13,14], aluminizing was successfully performed at 600 °C. This process yielded an aluminum-based coating exhibiting excellent high-temperature oxidation resistance.

2. Materials and Methods

2.1. Coating Process

The substrate material was a γ-TiAl-based alloy supplied by the Titanium Alloy Department, Institute of Metal Research, Chinese Academy of Sciences.

Table 1. The nominal chemical composition of γ-TiAl intermetallic compounds.

Elements	Al	V	Cr	Nb	N	O	Mn	Ti
Nominal	32.3	≤1.5	≤1.0	≤0.20	≤0.05	≤0.015	6.32	Bal.
EPMA detected	33.0	0.53	0.97	0.17	0.01	0.01	6.01	Bal.

presents the nominal composition and actual composition (determined by electron probe microanalysis, EPMA-1610, Shimadzu, Kyoto, Japan) of the γ-TiAl intermetallic compound (wt.%). Specimens were sectioned into 15 × 10 × 2 mm sheets via wire-electrode cutting, followed by sequential preparation: mechanical grinding (120−800# SiC abrasive paper), edge chamfering, ultrasonic cleaning in organic solvents (acetone and ethanol), drying, and storage. All substrates underwent surface sandblasting prior to processing.

In the experiments of this article, the method of embedding with aluminizing was chosen, and the schematic diagram is shown in Figure 1. Prior to experimentation, the pack cementation powder (excluding the NH₄Cl activator) was uniformly blended in a ball mill according to the designated weight percentages (wt.%). The mixed powder was then placed in a stainless steel crucible and dried at 400 °C for 2 h under an inert atmosphere. Subsequently, the NH₄Cl activator was added in the prescribed proportion and thoroughly mixed. The dried powder mixture was transferred to a corundum crucible. Substrates were positioned within the powder bed at one-third to one-half of the crucible height, ensuring complete powder envelopment. The powder was compacted firmly to maximize substrate-powder contact. The crucible was heat-treated in a muffle furnace at 600 °C for specified durations, followed by air cooling. Samples were then extracted, rigorously cleaned, and prepared for testing. For cross-sectional analysis, a protective nickel layer was electrodeposited onto coated specimens using a plating solution containing 200 g/L NiSO₄·6H₂O, 40 g/L H₃BO₃, 40 g/L NiCl₂, and 120 g/L Na₃C₆H₅O₇·2H₂O (trisodium citrate dihydrate) to prevent alumina layer damage during metallographic preparation. The pack cementation composition is detailed in

Table 2. The formula components and the proportion of each component of the diffusion agent in the powder-embedding aluminizing process.

Raw Materials of Penetrant	Al	Al2O3	NH4Cl	CeO2	Al
Formula 1 (wt.%)	30	66	3	1	30
Formula 2 (wt.%)	30	67	3	0	30

.

2.2. Isothermal Oxidation

To evaluate the oxidation resistance of coated specimens, isothermal oxidation tests were conducted in static air at 800 °C and 900 °C for 20 h, respectively. Sample mass changes were continuously monitored using a thermobalance. Post-oxidation mass measurements were recorded with a precision of 0.00001 g.

2.3. Analyzing Methods

Phase analysis of the coatings was performed using X-ray diffraction (XRD) (Panalytical, B.V., Almelo, The Netherlands). Microstructural morphology (surface and cross-section) of specimens before and after ablation testing was characterized by field-emission scanning electron microscopy (FE-SEM) (FEI Co., Hillsboro, OR, USA) (with coupled energy-dispersive spectroscopy (EDS) (Oxford instruments Co., Oxford, UK).

3. Results

3.1. The Influence of the Composition of the Diffusion Accelerator on the γ-TiAl-Intermetallic-Compound-Powder-Embedded Aluminized Coating

Figure 2 presents the microstructural morphology of aluminide coatings deposited on γ-TiAl intermetallic compounds via pack cementation with 0% and 1% nano-CeO₂ additions. The results demonstrate nano-CeO₂’s significant promotive effect. As shown in Figure 2A,D, CeO₂-modified coatings exhibit greater thickness (~40 μm) than unmodified counterparts, with enhanced density, continuity, and absence of porosity or penetration features. This improvement stems from rare earth doping strengthening coating–substrate adhesion [15]. Conversely, conventional aluminide coatings display interfacial inhomogeneity and cracking (Figure 2A,B). The CeO₂-modified coating surface appears uniformly smooth. Figure 3 describes the diffusion process of aluminizing on Ti₆Al₄V alloy. Similar to this paper, throughout the diffusion process, active Al atoms deposited on the surface diffuse into the matrix. Meanwhile, due to the high Ti content in the substrate, Ti atoms inevitably migrate from regions of high chemical potential (substrate) to low chemical potential (coating) to approach chemical equilibrium, consistent with Fick's diffusion law [16]. Given that Al diffusion rates in nickel-based systems exceed those of Ti [17], the coating thickness depends primarily on Al diffusion. It can be seen that the coating doped with CeO₂ (~40 μm) is thicker than its undoped counterpart (~19 μm), indicating indirectly that Ce presence enhances Al interdiffusion into the substrate.

This is because rare earth elements have very strong chemical reactivity. NH₄Cl decomposes into HCl and Cl₂ at 450 ℃. HCl reacts with CeO₂ as follows [18]:

2CeO₂ + 2HCl = Ce₂O₃ + Cl₂ + H₂O

(1)

Ce₂O₃ + 6HCl = 2CeCl₃ + 3H₂O

(2)

Ce₂O₃ + 3Cl₂ = 2CeCl₃ + 3/2O₂

(3)

CeCl₃ diffuses onto the substrate surface and is reduced to active Ce atoms. The deposited rare earth Ce activates the coating surface. As Ce atoms possess a larger atomic radius than Ti atoms, their incorporation induces significant lattice distortion while reducing surface free energy [19]. This phenomenon enhances Al adsorption and diffusion, rapidly forming a high-Al-concentration surface layer. The resultant concentration gradient facilitates rapid inward Al diffusion.

3.2. The Influence of Different Aluminizing Times

Using Formula 1, coatings were heat-treated at 600 °C for 1, 3, and 5 h. Figure 4 presents the principal elemental composition and dominant phases on aluminized coating surfaces at varying durations. Following substrate sandblasting pretreatment, pack cementation at 600 °C yielded TiAl₃ intermetallic compound coatings. EDS analysis (Figure 4E) confirmed a substantial increase in surface aluminum content from 32.2% to 78.78% post-treatment.

Figure 5 presents cross-sectional morphologies of aluminide coatings obtained after pack cementation durations of 1, 3, and 5 h at 600 °C. Coating thickness increased progressively from ∼12 μm (1 h) to ∼28 μm (3 h), reaching ∼31 μm after 5 h. While thickness continuously grew, the growth rate decelerated with extended processing time. This behavior arises because temperature dominantly governs diffusion coefficients. When aluminizing temperature remains constant and duration exceeds a critical threshold, atomic diffusion approaches completion. At this stage, increased coating thickness reduces atomic diffusivity, thereby diminishing growth kinetics. Consequently, in this experiment, we chose 5 h as the embedding aluminizing duration when the aluminizing temperature is 600 ℃

3.3. Surface Morphology Analysis

Figure 6 presents the surface morphology of the aluminide coating deposited on γ-TiAl intermetallic compound via 600 °C pack cementation for 5 h. As shown in Figure 6A, the coating surface exhibits uniform topography devoid of visible cracking. Higher-magnification imaging (Figure 6B) reveals a microstructure comprising irregularly sized blocky particles with adhered fine particulates. Interparticle voids are evident, while the blocky particles demonstrate non-uniform spatial distribution lacking discernible orientation.

Figure 6D presents EDS elemental analysis of representative regions. Selected areas at 100× and 5000× magnification revealed a composition dominated by Al, Ti, and Mn. Comparative analysis of regions 2 and 3 (5000×) showed minimal variation in Al/Ti/Mn weight percentages, confirming compositional homogeneity throughout the coating.

Figure 6C presents the cross-sectional SEM morphology and elemental depth profile of the aluminide coating deposited on γ-TiAl intermetallic compound via 600 °C pack cementation for 5 h. Line scanning analysis (left-to-right: coating surface to substrate) reveals a uniform ∼30 μm thick coating with a sharp coating–substrate interface. The coating exhibits full density without through-thickness cracking. EDS line scan analysis indicates a significant decrease in Al concentration at 29 μm from the surface, coinciding with increased Ti intensity. Combined with XRD analysis, the coating was identified as primarily TiAl₃ phase.

3.4. Research on High-Temperature Oxidation Resistance

Coatings prepared per Formula 1 (600 °C, 5 h) underwent isothermal oxidation at 800 °C and 900 °C for 20 h, respectively. Oxidation kinetics curves were derived from unit area mass gain measurements, with subsequent microstructural analysis performed via SEM/EDS.

3.4.1. Oxidation Kinetics Curve

Figure 7 presents oxidation kinetic curves for γ-TiAl intermetallic compound substrates and corresponding pack-cementation aluminide coatings after 20 h isothermal oxidation at 800 °C and 900 °C. As shown in Figure 7A, both materials exhibit mass gain during initial 800 °C exposure (<10 h). However, the uncoated substrate experienced oxide scale spallation after 10 h, accompanied by mass loss, indicating complete loss of oxidation resistance. In contrast, aluminized specimens demonstrated rapid initial mass gain followed by parabolic kinetics stabilization. This behavior signifies prompt formation of a continuous, dense Al₂O₃ scale that impedes oxygen permeation and suppresses further oxidation. Consequently, aluminized specimens exhibited minimal net mass gain with oxidation rates substantially lower than the substrate. Figure 7B reveals distinct oxidation mechanisms at 900 °C. While the substrate followed linear oxidation kinetics (indicating no protective scale formation), aluminized specimens maintained parabolic kinetics. This confirms sustained formation of protective Al₂O₃ scales that effectively isolate the substrate from oxygen ingress. Thus, aluminized coatings again demonstrated significantly reduced oxidation rates compared to the unprotected substrate. From Figure 7C, it can be seen that during the mid-stage oxidation, the oxidation kinetics curve shows a slow increase/exhibits a gradual upward trend.

3.4.2. Analysis of the Oxidation Resistance of Aluminized Coating After Constant Temperature Oxidation at 800 ℃ for 20 H

Following 20 h isothermal oxidation at 800 °C in static air, pack-cementation aluminide samples exhibited a color transition from light gray to dark gray. Figure 8 presents post-oxidation microstructural characteristics. As shown in Figure 8A, the oxide scale remains intact without cracking or spallation. Surface morphology at 5000× magnification (Figure 8B) reveals coarsened blocky particles with increased surface particulates compared to pre-oxidation states (Figure 7). EDS analysis (Figure 8D) confirms regions 1−4 consist primarily of Al₂O₃ with minor TiO₂. Cross-sectional examination (Figure 8C) shows no distinct oxide layer, attributed to the thin (<1 μm) scale intertwining with the electrodeposited nickel layer during metallographic preparation, rendering it SEM-indistinguishable. As indicated by the EDS line scan in Figure 8C, aluminum content peaks at ~70% on the coating surface and progressively decreases toward the substrate. XRD analysis (Figure 8E) identifies TiAl₃ and δ-Al₂O₃ as dominant surface phases.

3.4.3. Analysis of the Oxidation Resistance of Aluminized Coating After Constant Temperature Oxidation at 900 ℃ for 20 H

Following 20 h of isothermal oxidation at 900 °C in static air, pack-aluminized samples transitioned from light gray to dark gray, with Figure 9 characterizing the post-oxidation microstructure as exhibiting: (i) an intact oxide scale without cracking/spallation (Figure 9A), (ii) agglomerated block-like oxides under 10,000× magnification, (iii) predominant Al₂O₃-TiO₂ composition via EDS (Figure 9D,E), (iv) a sharp coating–substrate interface with minimal oxide thickness (Figure 9C,D), (v) Kirkendall voids at the oxide–coating boundary (magnified inset in Figure 9D) formed by vacancy accumulation due to divergent Ti/Al diffusion kinetics during oxide consumption—where outward Al diffusion depletes coating aluminum faster than inward Ti migration, and (vi) surface-phase dominance of TiAl and α-Al₂O₃ confirmed by parallel-light spectroscopy (Figure 9F), all corroborated by EDS line-scan data showing an elevated Al concentration at the coating surface, progressive Al depletion toward the substrate, and corresponding Ti enrichment (Figure 9D, the inset profiles).

4. Discussion

Thermodynamically, the Gibbs free energy for Al₂O₃ formation is marginally lower than for TiO₂ formation (ΔG difference: 218 kJ/mol at standard pressure [20]). This proximity enables competitive oxidation. Consequently, TiAl intermetallic compounds exposed to air form mixed Al₂O₃/TiO₂ scales rather than exclusive Al₂O₃ layers. Kinetically, Ti diffusion through TiO₂ substantially exceeds Al diffusion through Al₂O₃, accelerating TiO₂ growth. These factors collectively drive dual-oxide formation on TiAl alloys during air oxidation.

Numerous studies [21,22,23] have examined the high-temperature oxidation mechanism of TiAl intermetallic compounds. Research indicates that during initial oxidation, Al₂O₃ and TiO₂ form discontinuously as island structures. TiO₂ exhibits bidirectional growth (lateral and vertical), consuming surrounding Ti while locally enriching Al, thereby promoting Al₂O₃ nucleation. Consequently, the initial oxide scale comprises mixed Ti_xO_y (various valence states) and Al₂O₃ phases. A thin Ti-rich oxide overlayer forms due to Al's significantly lower oxidation kinetic tendency compared to Ti. As oxidation progresses, outermost Ti_xO_y phases transform into stable TiO₂ upon prolonged air exposure. Kirkendall porosity develops at the TiO₂/(Al₂O₃ + TiO₂) interface due to differential cation diffusion. Furthermore, rapid growth of the non-protective porous TiO₂ scale creates weak adhesion between the mixed oxide layer and TiAl substrate, ultimately causing scale spallation and loss of protective function. However, as aluminizing produces an aluminum-rich TiAl₃ phase on γ-TiAl intermetallic compounds, the coating surface achieves 85.63 wt.% aluminum content. During oxidation, this enables formation of a continuous, dense, and adherent protective Al₂O₃ scale. This barrier effectively inhibits inward oxygen diffusion and outward metal ion transport, significantly reducing oxidation kinetics. A TiAl₂ interlayer persists between the aluminide coating's TiAl₃ phase and the γ-TiAl substrate during high-temperature oxidation.

The literature [23,24] demonstrates that above 800 °C, TiAl₂ exhibits greater stability than TiAl₃ and Ti. The diffusion rate of Al follows TiAl > TiAl₃ > TiAl₂. Consequently, TiAl₃ persists throughout intermediate oxidation stages. After sustained exposure > 800 °C, the scale primarily consists of Al₂O₃ and residual TiAl₃, consistent with XRD analysis. After oxidation at 900 °C for 20 h, a small amount of TiO₂ phase formed in the oxide layer. This occurred because the rapid oxidation rate during the initial stage at 900 ℃ led to partial TiO₂ formation. As oxidation proceeded, a continuous and dense α-Al₂O₃ film developed, effectively inhibiting oxygen ingress and trapping residual TiO₂ on the alumina surface. Concurrently, elevated temperature accelerated the decomposition of TiAl₃ and enhanced Al diffusion, promoting TiAl phase formation.

During prolonged oxidation, inherent brittleness of the aluminide coating and continuous Al depletion for Al₂O₃ formation generate internal stresses. These induce microcracking, wrinkling, and eventual spallation in the oxide scale, leading to progressive thinning. Consequently, mid-stage oxidation exhibits diminished mass gain, reflected in the kinetics curve's attenuated slope. As oxidation persists, spalled regions expose the substrate. However, the coating's self-healing capacity facilitates continuous Al₂O₃ reformation, maintaining substrate protection.

5. Conclusions

This formulation, namely the use of the 30%Al-66%Al₂O₃-3%NH₄Cl-1%CeO₂ composition, successfully deposited aluminide coatings on γ-TiAl intermetallic compounds at 600 °C for 5 h. The coating thickness increased progressively with aluminizing duration, though growth decelerated over time. Beyond a critical period, thickness asymptotically approaches a maximum value. The optimal coating thickness was achieved after 5 h at 600 °C.

The addition of the rare earth oxide CeO₂ promoted the diffusion of aluminum atoms, achieving an oxide film approximately 31.30 µm thick at a temperature of 600 °C for 5 h.

The oxidation kinetics of both γ-TiAl intermetallic substrates and low-temperature aluminized specimens conformed to parabolic laws after 20 h at 800 °C. At 900 °C, however, the substrate exhibited linear oxidation kinetics, demonstrating no high-temperature oxidation resistance. In contrast, aluminized specimens maintained parabolic kinetics. These results confirm the coating's significant enhancement of substrate oxidation resistance.