Influences of Al concentration and Nb addition on oxidation behavior of Ti₂AlC ceramics at high temperatures

ABSTRACT

A titanium aluminum carbide, Ti₂AlC, which is classified among MAX phase ceramics, was studied as a potential candidate for various mechanical components used in high-temperature applications. The impact of its chemical composition on its high-temperature oxidation process was determined. Ti₂AlC powders with various Al contents and with or without Nb addition were synthesized via a conventional reaction technique followed by 16 h of annealing in vacuum at 1300°C. The synthesized Ti₂AlC powders were consolidated by means of 15 min of pulsed electric current sintering at a die temperature of 1300°C in vacuum under a uniaxial pressure of 30 MPa. In the presence of an aluminum reservoir in the form of TiAl₃, Ti₂AlC has excellent resistance against high-temperature oxidation. The promising results concerning the addition of Nb to TiAl provided the rationale for a similar modification of Ti₂AlC. The results of oxidation tests on Nb-doped Ti₂AlC likewise showed excellent oxidation resistance. Alloying with Nb can improve the oxidation resistance of Ti₂AlC with low Al content, allowing the formation of a protective Al₂O₃ scale and inhibiting the growth of TiO₂.

KEYWORDS:

• MAX phase
• Ti₂AlC
• oxidation resistance
• Nb

1. Introduction

MAX phase ceramics are layered ternary transition metal carbides and nitrides represented by the general formula M_n+1AX_n, where M stands for a transition metal (Sc, Ti, V, Cr, Mn, Fe, Y, Zr, Nb, Mo, Lu, Hf, Ta or W), A is a group A element (Al, Si, P, S, Cu, Zn, Ga, Ge, As, Pd, Cd, In, Sn, Ir, Au, Tl, Pb or Bi) and X represents C or N [1]. MAX phase ceramics have both metallic and ceramic features because of the presence of both metallic and covalent bonds in the crystal [2]. Since Ti₂AlC belongs to this group, it also exhibits both metallic and ceramic properties, including high electrical and thermal conductivity, high thermal shock resistance, lower density than Ni alloys and steels, good high-temperature oxidation resistance and high hardness [3,4]. Ti₂AlC has been studied as a potential candidate for application in various mechanical components used at high temperatures as well as room temperature. The mechanical strength of Ti₂AlC decreases sharply at around 900°C, at this point, Ti₂AlC undergoes the brittle-to-ductile transition (BDT) [5]. The application of Ti₂AlC ceramics would therefore be limited at temperatures higher than the BDT temperature.

When it comes to structural elements exposed to high temperatures, one of their most important properties is high-temperature oxidation resistance as well as high-temperature mechanical strength and creep resistance. However, there is no well-established understanding of the oxidation properties of Ti₂AlC at high temperatures. Some sources consider Ti₂AlC to have improved resistance to oxidizing conditions [6–8], while others do not [9–11]. Certain results indicate linear oxidation behavior occurred with the formation of a non-protective TiO₂ scale on the surface of Ti₂AlC after long-term oxidation [8]. To promote the applications of Ti₂AlC in high-temperature components, the oxidation process of Ti₂AlC must be stabilized.

The chemical composition of Ti₂AlC may affect its oxidation behavior. Because Ti₂AlC is a line phase [12], every commercial Ti₂AlC powder contains other phases, such as TiC, Ti₃AlC₂ or TiAl₃ [13]. This makes it impossible to study the oxidation behavior of Ti₂AlC using samples with a controlled phase ratio. Because the protective scale that forms on Ti₂AlC is Al₂O₃, its continued growth requires a sufficient Al supply during high-temperature oxidation. Al-rich phases like TiAl₃ may therefore improve the oxidation resistance of Ti₂AlC.

High-temperature oxidation of TiAl with poor oxidation resistance leads to the formation of oxide scale including the rutile polymorph of TiO₂ and Al₂O₃ [14]. This oxidation behavior is very similar with Ti₂AlC with poor oxidation resistance. Improvements of oxidation resistance on TiAl have been studied with several techniques. As one of the common methods for improving the oxidation resistance of TiAl, addition of Nb is applied [15], which would inhibit the growth of the TiO₂ scale on the surface of TiAl and subsequently facilitates the formation of a continuous Al₂O₃ scale. The oxidation resistance of Ti₂AlC doped with Nb was studied by Donchev et al. [16], and they reported a decreased oxidation rate, indicating the protective action of a Nb addition. However, the mechanism underlying the beneficial effect of Nb on the oxidation properties of Ti₂AlC is unclear in some cases. Adding Nb to Ti₂AlC with low Al content may not necessarily improve its oxidation resistance to a very significant degree, since this compound is in itself highly susceptible to oxidation [17].

In this study, the effects of varying the Al content and Nb addition on the oxidation behavior of Ti₂AlC in air at 800°C was investigated to establish a material design strategy that would optimize the high-temperature oxidation resistance of Ti₂AlC ceramics.

2. Experimental procedure

Commercial Ti, Al, and C powders were mixed in Ti:Al:C molar ratios of 2:1:0.9, 2:1.2:0.9, and 2:1.5:0.9, with the corresponding samples receiving the designations Al1, Al1.2 and Al1.5, respectively. Some Al1 and Al1.2 samples had 5% of their Ti content replaced with Nb; these samples are referred to as Al1Nb and Al1.2Nb. The particle size and purity of the raw powders were as follows: 38 μm and 99.9% for Ti, 3 μm and 99.9% for Al, 1–2 μm for C and 45 μm for Nb (Kojundo Chemical Laboratory). The powder mixture was placed in an alumina container and annealed for 16 h in vacuum at 1300°C. The reacted powder underwent phase identification by means of X-ray diffraction (XRD, Rigaku corporation MiniFlex 600, using monochromatic CuKα radiation). Al1 is categorized as Al-poor Ti₂AlC, and Al1.2 and Al1.5 are categorized as Al-rich Ti₂AlC. Table 1 shows the symbol and category of each sample investigated in the present study. The synthesized powder was consolidated over 15 min using a pulsed electric current sintering technique with a graphite die at 130°C, in vacuum, and under a uniaxial pressure of 30 MPa. The phases present in the sintered samples were also identified with XRD. The relative density of the samples reached at least 98% of the theoretical one. Samples for oxidation experiments were cut into a small quarter-circle shape with a radius of 7.5 mm and a thickness of 3 mm using a diamond saw. All samples for the oxidation tests were ground with SiC abrasive paper with varying grit (up to 3000), and then subsequently polished with 3 and 1 μm diamond slurry to ensure that the surface of each sample had undergone the same preparation. The samples were heated for times ranging from 2 to 14 d in an electric furnace with laboratory air at 800°C. For thermogravimetric analyses (TGA), the samples were cut into small rectangles with dimensions of 10 × 5 × 1 mm³ using a diamond saw. The oxidation kinetics were evaluated at 800°C in ambient air, over a period of 150 h using a thermogravimetric apparatus which enabled the continuous monitoring of the weight with an accuracy of to 10^–6 g (MK2 Vacuum Head Microbalance, CI Electronics Ltd, UK). The phase compositions of the samples were analyzed by means of XRD. The samples were mounted in resin and cut with a diamond cutting wheel for cross-sectional observations by means of scanning electron microscopy (SEM), performed using a FEI Nova NanoSEM 200 microscope coupled with an EDAX Genesis XM2 system required for the analysis of chemical composition via energy-dispersive X-ray spectroscopy (EDS).

Table 1. Sample symbols and categories.

Ti:Al:C:Nb	Symbols	Category name
2:1:0.9:0	Al1	Al-poor
2:1.2:0.9:0	Al1.2	Al-rich
2:1.5:0.9:0	Al1.5
1.9:1:0.9:0.1	Al1Nb	Nb-doped
1.9:1.2:0.9:0.1	Al1.2Nb

3. Results

Figure 1 shows the XRD patterns of the sintered samples in all the groups. The XRD graphs with the bottom portion enlarged were smoothed to remove noise. The main peaks in these XRD patterns were identified as originating from Ti₂AlC. In the case of Al-poor Ti₂AlC, peaks attributed to Ti₃AlC₂, TiAl₃ and TiC were also observed. The patterns recorded for the remaining samples had minor peaks originating from Ti₃AlC₂ and TiAl₃. The surface images of Ti₂AlC and Nb-doped Ti₂AlC samples with different Al molar ratios after oxidation at 800°C for up to 14 d in air are shown in Figure 2. Yellowish scales had formed on the surface of Al-poor Ti₂AlC. After the oxidation test, the surface of the Al1Nb sample differed in color compared to its counterpart without Nb addition. Oxidation time did not affect the appearance of any of the Al-rich and/or Nb-doped samples to a significant degree.

Figure 1. X-ray diffraction patterns recorded for bulk samples: Al1, Al1.2, Al1.5, Al1Nb and Al1.2Nb.

Figure 2. Magnified images of Ti₂AlC and Nb-doped Ti₂AlC samples with various Al molar ratios after 2 − 14 d of isothermal oxidation in laboratory air at 800°C.

Figure 3 shows the XRD patterns obtained for sintered samples from all groups after 14 d of oxidation. The main peaks in XRD patterns of Al-rich and Nb-doped oxide samples were identified as originating from Ti₂AlC. The minor peaks were attributed to rutile-type TiO₂ and different polymorphs of Al₂O₃. Because peaks observed for α-Al₂O₃, θ-Al₂O₃, and γ-Al₂O₃ are similar and they all have low intensity, it was impossible to precisely determine the type of Al₂O₃. The EDX element mapping of the surface of Al-rich and Nb-doped samples after 14 d of oxidation clearly shows aluminum and oxygen (Figure 4). The conducted XRD, SEM, and EDX observations indicate that an Al₂O₃ scale with a thickness of less than 1 μm was formed on these samples. The predominant phase in the oxide scales formed in the case of all Al-rich and Nb-doped samples was Al₂O₃. Figure 5 shows SEM and EDS images of the surface of Al1.2Nb after 150 h of isothermal oxidation in laboratory air at 800°C. Needle-like growth of Al₂O₃ was clearly observable only on the surface of Al1.2Nb. Figure 6 shows the cross-sectional SEM and EDX images of the Al-poor sample after oxidation for 14 d. The results of XRD, SEM, and EDX investigations indicate that rutile-type TiO₂ formed on the scale’s surface, while Al₂O₃ and TiO₂ had formed underneath it. In the case of this sample, thickness increased linearly with oxidation time, becoming approximately 120 μm after 14 d (Figure 7). The oxide scale formed on all Al-poor samples was determined to consist of rutile-type TiO₂ with poor oxidation resistance.

Figure 3. X-ray diffraction patterns recorded for samples Al1, Al1.2, Al1.5, Al1Nb and Al1.2Nb after 14 d of isothermal oxidation in laboratory air at 800°C.

Figure 4. SEM and EDS images of the cross-sections of samples Al1.2, Al1Nb and Al1.2Nb after 14 d of isothermal oxidation in laboratory air at 800°C.

Figure 5. SEM and EDS images of the surface of Al1.2Nb after 14 d of isothermal oxidation in laboratory air at 800°C.

Figure 6. SEM and EDS images of the cross-sections of Al1 after 14 d of isothermal oxidation in laboratory air at 800°C.

Figure 7. Oxide scale thickness (x) versus time (t) for Al1 oxidized at 800°C for 14 d.

Figure 8 shows the kinetics of isothermal oxidation for all of the sintered samples in a linear plot representing mass gain per unit area in a function of time. The data show that the oxidation kinetics of the samples approximately obeys the parabolic rate law. The slowest constituent process that determines the oxidation rate is therefore the diffusion of reagents in the oxidation product. The largest increase in the mass of the corrosion product was observed for the Al1 sample, while the Al1.2Nb sample was characterized by the smallest increase. The parabolic oxidation rate constants (k_p) of the samples determined using the Pilling-Bedworth equation are listed in Table 2 [18]. As can be seen, the parabolic rate constant of oxidation decreases with increasing aluminum content. The oxidation rate constants of the Al1 and Al1.5 samples are different by as much as one order of magnitude. On the other hand, the addition of niobium reduced the k_p values by another order of magnitude, further improving the oxidation resistance of the Al1Nb and Al1.2Nb samples.

Figure 8. Oxidation kinetics for all investigated samples over 150 h of isothermal oxidation in laboratory air at 800°C, presented in a linear plot.

Table 2. Parabolic rate constants of oxidation (k_p) of all samples determined after 150 h of isothermal oxidation in laboratory air at 800°C.

Sample	kp/g^2cm^−4s^−1	Category name
Al1	1.6 × 10^−12	Al-poor
Al1.2	3.1 × 10^−13	Al-rich
Al1.5	1.0 × 10^−13
Al1Nb	6.4 × 10^−14	Nb-doped
Al1.2Nb	2.7 × 10^−14

4. Discussion

The Al₂O₃ oxide scale formed on the surface of Al-rich Ti₂AlC such as Al1.2 and Al1.5 and Nb-doped one can be considered a protective layer, even though its thickness was below 1 μm even after 14 d of oxidation. This is demonstrated by the lower oxidation resistance of Al-poor Ti₂AlC, which had a main TiO₂ scale on its surface and a non-continuous Al₂O₃ scale beneath it. It is the evidence that the dispersion of an Al-rich phase and doping with Nb are effective at improving the oxidation resistance of Ti₂AlC ceramics.

Brumm and Grabke stated that the oxidation rate depends on the oxide phase formed on the alloy; γ-Al₂O₃, θ-Al₂O₃, and α-Al₂O₃ have been observed, and the growth of each of these phases was associated with different k_p values [19]. Figure 9 shows the k_p values of Ti₂AlC determined in the present study compared with those of NiAl reported by Brumm and Grabke [19]. The k_p value for Al₂O₃ growth on Ti₂AlC in the cited study is similar to those for γ- or θ-Al₂O₃. The k_p of Al1.2 was slightly higher than that of other Al-rich and Nb-doped samples because of partial non-protective oxidation of the Al1.2 sample’s corners. Non-protective oxidation tends to occur at the corners of Ti₂AlC samples [10]. The supply of Al is limited at this location because of the small area of the lower part. When a continuous Al₂O₃ scale is formed on a protruding part, such as a corner, the Al₂O₃ scale may crack due to growth stress [13]. Needle-like Al₂O₃ was observed on the surface of Al1.2Nb oxidized at 800°C for 150 hours. The needlelike Al₂O₃ were sparsely present in some places and did not form dense colonies. Mechanism of the formation of needlelike shape is explained by the stress generated by the transformation from γ- to θ-Al₂O₃. When γ-Al₂O₃ transforms into θ-Al₂O₃, the sample volume decreases by 1.7% [20]. The reduced volume causes tensile stress in the θ-Al₂O₃ layer and compressive stress in the γ-Al₂O₃ layer formed on top. Ando et al. reported that a possible mechanism is that the <111> plane shifts in the <112> direction to form twins that relieve compressive stress generated in the γ-Al₂O₃ grain near the transition interface from γ- to θ-Al₂O₃ [20]. The twin boundary serves as a preferential diffusion path for Al, and the twinned part preferentially grows outward, resulting in a needle shape.

Figure 9. Arrhenius plot showing the parabolic rate constants determined for the oxidation of Ni₅₀AI₅₀ [19] and the k_p values determined in the present experiment (Table 2).

Needle-like crystals of γ- or θ-Al₂O₃ may form on the surface when Al₂O₃-forming materials are oxidized at low temperatures of 500–800°C as well [21–23]. The shape of the Al₂O₃ grains and the comparison of k_p indicates that the Al₂O₃ scale formed on Ti₂AlC at 800°C should be regarded as being γ- or θ-Al₂O₃.

When γ-Al₂O₃ transforms into θ-Al₂O₃, the sample volume decreases by 1.7%; the transformation from θ-Al₂O₃ to α-Al₂O₃ is associated with a decrease of 12.3% [20]. The volume reduction caused by the change of the Al₂O₃ phase causes tensile stress and leads to the cracking of the scale. This phenomenon may cause decrease in the oxidation resistance of Ti₂AlC. The formation mechanism of Al₂O₃ should therefore be considered when designing Ti₂AlC with excellent oxidation resistance in mind.

The following tendencies can be observed during the oxidation of Al-rich Ti₂AlC. A continuous Al₂O₃ scale forms under a non-continuous TiO₂ scale as the material starts to oxidize. Since the diffusion coefficient of Al in Ti₂AlC is much faster than that of Ti, the Al₂O₃ scale on top of Ti₂AlC becomes thin and non-continuous [24]. When continuous Al₂O₃ does not form in the initial oxidation stage, its subsequent formation is accompanied by the formation of TiO₂ [25]. The rate at which this non-continuous TiO₂ scale grows is faster than the growth rate of Al₂O₃ [26], and its underlying mechanism is the outward diffusion of Ti through the grain boundaries of the discontinuous Al₂O₃ scale, made possible because the diffusion coefficient of O in the TiO₂ particles is smaller than that of Ti [27]. TiO₂ forms on the surface, and Al₂O₃ forms underneath it. In order for the Al₂O₃ scale to be continuous, the supply of Al from Ti₂AlC needs to be sufficient. The formation of a continuous Al₂O₃ scale suppresses the formation of the TiO₂ scale.

The oxidation of Al-poor Ti₂AlC takes a different course, similar to that observed for Al-poor TiAl. The oxide scale that forms in this case consists mostly of rutile-type TiO₂ and Al₂O₃ located underneath it, and it does not provide significant protection against oxidation [14]. The supply of Al to the surface is insufficient to allow the formation of the Al₂O₃ scale. Ti is oxidized into TiO₂ and its volume increases, which destroys the Al₂O₃ scale. In order for the oxidation resistance of Ti₂AlC to be improved, it is therefore necessary to ensure that the supply of Al is sufficient and that the growth of the TiO₂ layer is inhibited.

Al-rich Ti₂AlC had high oxidation resistance owing to the development of an Al₂O₃ scale. The Al-rich phase (TiAl₃) supplied the needed Al, allowing the continuous Al₂O₃ scale to form and improve the material’s resistance to oxidation. In addition to increasing the amount of Al, the supply of Al for Ti₂AlC can also be improved by shortening the Al supply path. Al atoms in Ti₂AlC can migrate more easily in the outward direction along grain boundaries and the base planes than other atoms during high-temperature oxidation [24]. Yu et al. reported that obtaining finer Ti₂AlC grains facilitates Al diffusion and thus the development of the Al₂O₃ scale, which is yet another way in which the oxidation resistance of Ti₂AlC can be improved [28]. There are various ways to reduce grain size in Ti₂AlC, including the use of a finer starting powder or pinning with fine impurity particles. In addition to using a finer starting powder, the preparation process can also be optimized by applying a lower sintering temperature, which yields Ti₂AlC with a high concentration of Al.

Some results concerning the oxidation of Ti₂AlC show that linear kinetics lead to poor oxidation resistance after long-term oxidation [8]. In such conditions, Al-depletion zones in Ti₂AlC develop as the Al is drained by the formation of the continuous Al₂O₃ scale, and after a sufficiently long time of oxidation or if cycles are involved Al-rich Ti₂AlC may be affected in much the same way as the material with low Al content. An additional method to prevent non-protective oxidation of Ti₂AlC is required for long-term use.

This was precisely the rationale for the addition of Nb. When Nb⁵⁺, which has a higher valence than Ti⁴⁺, is added to TiO₂, the number of oxygen vacancies in TiO₂ decreases [29], and the growth of TiO₂ is suppressed. After the solid dissolution of Nb₂O₅ in TiO₂, the Nb concentration at which the oxygen ion vacancies in TiO₂ disappear is estimated to be Ti:Nb = 1:0.04 [15]. For the Nb-doped sample in this experiment, Nb was added in an amount equal to 5% of Ti, an amount considered sufficient to fill the oxygen ion vacancies in TiO₂. The cause of non-protective oxidation was the destruction of the Al₂O₃ scale due to the growth of TiO₂. Inhibiting TiO₂ growth by adding Nb is one of the factors that improve oxidation resistance.

Adding more Nb (10%) to TiAl than the solid solution limit in TiO₂ may further improve oxidation resistance [15,30]. However, when Nb is added to TiAl in large amounts (15% or more), Nb₂O₅ and TiNb₂O₇ scales are formed, and oxidation resistance is reduced [15,30]. The oxidation resistance of Ti₂AlC may also be further improved by increasing the amount of Nb added; the optimal amount of Nb is yet to be determined.

When combined, the change in the amount of Al and the addition of Nb have a pronounced effect on the oxidation resistance of the Ti₂AlC ceramics, which is a property of fundamental importance from the perspective of its eventual application. Such modifications are therefore necessary for the stable operation of components manufactured from this material in conditions that involve high temperatures.

5. Conclusions

The oxidation behavior of Ti₂AlC with various Al concentrations and with or without Nb addition was investigated in air at 800°C in periods ranging from 2 to 14 d. Ti₂AlC powder samples were synthesized for different molar ratios of Al and amounts of added Nb. Al-poor Ti₂AlC was characterized by low oxidation resistance due to the formation of a TiO₂ scale that did not reduce the oxidation rate. Nb-doped Ti₂AlC and samples with higher Al contents exhibited far superior resistance to oxidation owing to the formation of a protective Al₂O₃ scale. The addition of Nb improved oxidation resistance even in the case of the Ti₂AlC sample with the lowest Al content.

Acknowledgments

This work was supported partly by the Osawa Scientific Studies Grants Foundation.

Disclosure statement

No potential conflict of interest was reported by the author(s).