**About the Special Issue Editor**

**Kang N. Lee** is a Senior Research Scientist at the NASA Glenn Research Center, Cleveland, OH, USA. He received his Ph.D. from the University of Minnesota, Minneapolis, MN, USA, in 1987. After postdoctoral work at the University of Pennsylvania, Philadelphia, PA, USA, he spent the first 15 years of his career at the NASA Glenn Research Center before moving to Rolls-Royce in 2005. After 11 years at Rolls-Royce, he returned to NASA in 2016. During his first tenure at NASA, he pioneered EBC research while leading the NASA EBC programs that developed the firstand second-generation EBCs, which became the foundations for the current state-of-the-art EBCs. As an Associate Fellow at Rolls-Royce, he provided global leadership for strategy and technology development of high-temperature coatings across the Rolls-Royce product line. Since returning to NASA in 2016, his research has focused on cutting-edge coatings R&D, including next-generation EBCs. He has authored 35 US patents, 20 US patents pending, 80 papers in archival journals, and 4 invited book chapters. He is the recipient of the 2017 Toledo Glass and Ceramic Award, 2010 Rolls-Royce Sir Henry Royce Award, 2003 NASA Public Service Medal, and 2001 R&D 100 Award.

## *Editorial* **Special Issue: Environmental Barrier Coatings**

#### **Kang N. Lee**

NASA Glenn Research Center, Cleveland, OH 44135, USA; ken.k.lee@nasa.gov Received: 12 May 2020; Accepted: 22 May 2020; Published: 27 May 2020

**Abstract:** The global increase in air travel will require commercial vehicles to be more efficient than ever before. Advanced turbine hot section materials are a key technology required to keep fuel consumption and emission to a minimum. Ceramic matrix composites (CMCs) are the most promising material to revolutionize turbine hot section materials because of their excellent high-temperature properties. Rapid surface recession due to volatilization by water vapor is the Achilles heel of CMCs. Environmental barrier coatings (EBCs), which protect CMCs from water vapor, is an enabling technology for CMCs. The first CMC component entered into service in 2016 in a commercial engine, and more CMC components are scheduled to follow within the next few years. One of the most difficult challenges to CMC components is EBC durability because failure of EBC leads to a rapid reduction in CMC component life. Novel EBC chemistries, creative EBC designs, and robust processes are required to meet EBC durability challenges. Engine-relevant testing, characterization, and lifting methods need to be developed to improve EBC reliability. The aim of this Special Issue is to present recent advances in EBC technology to address current EBC challenges.

**Keywords:** EBC; CMC; oxidation; volatility; CMAS; thermomechanical; modeling

#### **1. Introduction**

Ceramic matrix composites (CMCs) are considered a game changer for gas turbine hot section materials because of their excellent high-temperature mechanical properties, oxidation resistance, and a light weight—one-third of nickel-based superalloys [1]. The introduction of CMCs enables a fuel burn reduction up to two percent—few other technology in today's pipeline have this much capability for fuel burn reduction [2]. Additionally, the light weight enables over 50 percent reduction in the turbine component weight [2].

SiC/SiC CMCs derive their oxidation resistance from dense slow-growing silica (SiO2) scale [3]. In combustion environments, however, the protective silica scale volatilizes by reacting with water vapor (H2O), leading to rapid CMC surface recession [4–9]. Alumina (Al2O3) in oxide/oxide CMCs, although not as severe as silica, also suffers surface recession due to water vapor-induced volatilization [10]. External barrier coatings, known as environmental barrier coatings (EBCs), have been developed to protect CMCs from surface recession [11,12]. Over two decades of intense research and development led to the first EBC-coated CMC component-a high-pressure turbine CMC shroud-to enter service in the LEAP engine by CFM International (Cincinnati, Ohio) [2]. More CMC components, including combustor liners and high-pressure turbine vanes, are scheduled to follow in the near future [2].

One of the most difficult challenges to CMC components is EBC durability because failure of EBC leads to a rapid reduction in CMC component life. Key contributors to EBC failure include recession, oxidation, degradation by calcium-aluminum-magnesium silicates (CMAS) deposits, thermal and thermo-mechanical strains, particle erosion, and foreign object damage (FOD) [12,13]. Novel EBC chemistries, creative EBC designs, and robust processes as well as engine-relevant testing, characterization, and lifting methods are required to meet EBC durability and reliability challenges.

#### **2. This Special Issue**

The nine research articles in this special issue present recent advances in EBC technology in the following areas: recession [14]; EBC chemistry and processes [15,16]; oxidation and recession [17]; oxygen permeability and EBC design [18]; CMAS thermal and mechanical properties [19]; CMAS infiltration modeling [20]; mechanical behavior in thermal shock [21]; and thermomechanical reliability: modeling and validation [22].

Smialek et al. [14] investigated the volatility of transient TiO2 and steady-state Al2O3 scales formed on Ti2AlC MAX phase ceramic in 1300 ◦C high velocity (Mach 0.3, 100 m/s) and high pressure (6 atm, 25 m/s) burner rig tests (BRT). Unlike metals, the ceramic was stable at 1300 ◦C. Unlike SiC and Si3N4, neither burner test produced a weight loss, unless pre-oxidized. Volatility of TiO2 was indicated by removal of Ti-rich oxide grains after BRT, while Al2O3 loss, not as pronounced as TiO2, was indicated by grain boundary etching and porosity. Al2O3 net weight loss was observed only on pre-oxidized samples, whose linear loss rate was about one-fifth of SiO2. 7YSZ TBC on Ti2AlC survived for 500 h in the Mach 0.3 burner test with no indication of volatility or spalling.

Vaßen et al. [15] investigated various oxides as an EBC for both oxide/oxide and SiC/SiC CMCs. The goal was to obtain dense and crystalline coatings. APS (atmospheric plasma spraying) Y2O3, Gd2Zr2O7, Y3Al5O12, and Yb2Si2O7 were investigated for oxide/oxide CMCs, while Y2SiO5 and Yb2Si2O7 were investigated for SiC/SiC CMCs. Y2O3 and Gd2Zr2O7 showed promising results for oxide/oxide CMCs: Y2O3 coating showed excellent adhesion due to the formation of chemical bond, while Gd2Zr2O7 coating required laser structuring of the CMC surface for good adhesion. APS Y2SiO5 required a heat treatment to obtain a desirable microstructure for SiC/SiC CMCs. Yb2Si2O7 on SiC/SiC CMCs was investigated using various thermal spray processes (APS, high velocity oxygen fuel (HVOF), very low pressure plasma spraying (VLPPS), and suspension plasma spraying (SPS)). APS coating was dense, but with a high degree of amorphous phase content. HVOF gave higher crystalline phase content than APS, but with some porosity. VLPPS coating gave the highest crystallinity and density. SPS gave a high degree of crystallinity, however segmental cracks could not be avoided. The study demonstrated the feasibility of various thermal spray methods besides APS for EBC processing.

Gatzen et al. [16] investigated YAlO3 as an EBC for Al2O3/Al2O3 CMCs using APS and VLPPS processes. APS resulted in a poor quality coating not suitable as an EBC. VLPPS, on the other hand, resulted in a coating with high crystallinity, high purity, and strong adhesion. The strong adhesion was attributed to the formation of chemical bonding due to the formation of Y3Al5O12 interfacial phase. The coating exhibited excellent thermal cyclic durability and CMAS resistance in laboratory testing, demonstrating its promise as an EBC for oxide/oxide CMCs.

Klemm et al. [17] investigated the oxidation and corrosion (water vapor-induced SiO2 volatilization) of two plasma-sprayed EBC systems (Al2O3/YAG and Si/Yb2Si2O7+SiC/Yb2SiO5) on SiC/SiC CMCs. Two EBC failure mechanisms were observed: spallation due to TGO-induced stresses and corrosion-induced gap at the TGO/EBC interface. Al2O3/YAG EBC, after burner rig test, suffered corrosion at the TGO/EBC interface and inside the Al2O3 bond coat. In Si/Yb2Si2O7+SiC/Yb2SiO5, TGO-induced failure was delayed because SiC particles in the intermediate layer oxidized first, forming a shell surrounding SiC, and thereby served as a getter to reduce the permeation of oxidants (O2, H2O) through EBC. In the burner rig test, small pores developed in the intermediate layer due to SiO2 corrosion. The beneficial gettering function of SiC particles is temporary, disappearing when the SiC particles are fully consumed.

Kitaoka et al. [18] proposed an EBC design based on oxygen permeability calculations to improve oxygen shielding capability and phase stability of mullite coating in Si/mullite EBC for SiC/SiC CMCs. Grain boundary (GB) diffusion of O ions in mullite, from high PO2 mullite surface to low PO2 mullite/Si interface, and simultaneous GB diffusion of Al ions in the opposite direction were calculated. Calculations projected: (i) mullite would decompose at the mullite/Si interface due to the outward Al diffusion and (ii) replacing the Si bond coat with the β'-SiAlON bond coat would increase oxygen shielding capability of mullite—this is because the higher equilibrium PO2 for

the oxidation of β'-SiAlON to form SiO2 compared to the oxidation of Si to form SiO2 suppresses the inward O ion diffusion. Experiments confirmed the mullite decomposition at the mullite/Si interface and the improvement of mullite stability when the Si bond coat was replaced by β'-SiAlON. The latter is because β'-SiAlON serves as a source for Al.

Webster et al. [19] investigated the intrinsic thermal and mechanical properties of a volcanic ash glass obtained from the Eyjafjallajökull eruption of 2010. The properties studied include crystallization, melting temperature, CTE (coefficient of thermal expansion), modulus, hardness, and viscosity. The glass had a low propensity to crystallize in bulk form compared to synthetic sand glasses. Compared to sand glasses, the melting temperature and viscosity were higher, while the CTE and modulus were lower. The higher viscosity was attributed partly to a higher SiO2 and lower CaO content in the volcanic ash. Hardness was similar, but the indentation fracture toughness was about twice that of sand glasses, which was attributed to the presence of Fe3O4 crystallites in the volcanic ash. Understanding the variation of CMAS properties with composition can provide insight into CMAS/coating interactions, which can be highly valuable for the development of CMAS-resistant coatings.

Kabir et al. [20] developed a numerical approach for studying the infiltration kinetics of molten CMAS in EB-PVD TBCs. Detailed analyses of the infiltration kinetics were performed, and the results were verified by experimental infiltration depths of the feathery and coarse microstructures. The study identified key morphological features that are important for infiltration. The rate of longitudinal and lateral infiltration could be minimized by reducing the gap between columns and/or increasing the length of the feather arms. Long feather arms having a lower lateral inclination decreased the infiltration rate and, therefore, reduced the infiltration depth. A key takeaway is that the infiltration kinetics can be altered by careful tailoring of coating microstructure such as feather arm lengths and inter-columnar gaps.

Seo et al. [21] investigated crack formation, crack stability, and crack healing of Si/mullite, Si/(mullite + Yb2SiO5), and Si/Y2SiO5 EBCs on SiC/SiC CMCs under thermal shock cycling between room temperature and 1350 ◦C in air. Mechanical behavior was determined before and after the test via indentation tests. EBCs developed mud cracks on the surface and unidirectional vertical cracks in the top coat after several thermal shock cycles. Mullite-based EBCs exhibited crack healing phenomena after 4000-5000 cycles, where cracks were covered and partially filled with a new phase. Post-test XRD showed only crystalline Al2O3 in addition to the EBC component phases, indicating that the new phase is likely amorphous. The nature of the new phase responsible for cracking healing was not fully understood yet. Si/Y2SiO5 EBC delaminated after 4000 cycles with no crack healing phenomena. After 5000 cycles, mullite-based EBCs showed no significant changes in modulus and hardness due to cracking healing, whereas Y2SiO5 EBC showed a change from elastic to plastic behavior due to delamination.

Kawai et al. [22] investigated the thermomechanical stability of an as-processed EBC by comparing the ERR (energy release rate) for interfacial cracks calculated by FEM analysis and the interfacial fracture toughness determined experimentally. The EBC studied was (SiC/SiAlON/mullite/Yb disilicate-Yb monosilicate graded layer/Yb monosilicates with porous segmented structure). FEM showed thinner SiALON/thicker mullite combinations lower the initial ERR. Fracture in the interface fracture test occurred at the SiC/SiAlON and SiAlON/mullite interfaces, indicating these interfaces are the weak link. ERR by FEM was sufficiently higher than the experimentally determined interfacial fracture toughness, indicating that as-processed EBC possesses sufficient thermomechanical reliability. As-processed EBC with 5–25 μm SiALON and mullite showed no interfacial delamination, confirming the thermomechanical reliability. This study considered only as-processed stresses. Further work is needed to assess the thermomechanical reliability in service environments by incorporating oxidation, CMAS, thermal cycling in temperature gradient, and mechanical stresses.

**Funding:** This research received no external funding.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2020 by the author. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Relative Ti2AlC Scale Volatility under 1300** ◦**C Combustion Conditions**

#### **James L. Smialek** †

NASA Glenn Research Center, Cleveland, OH 44135, USA; Dr.JSmialek@outlook.com † Retired, Oct 29, 2018.

Received: 4 January 2020; Accepted: 3 February 2020; Published: 5 February 2020

**Abstract:** Turbine environments may degrade high temperature ceramics because of volatile hydroxide reaction products formed in water vapor. Accordingly, the volatility of transient TiO2 and steady-state Al2O3 scales formed on the oxidation-resistant Ti2AlC MAX phase ceramic was examined in 1300 ◦C high velocity (Mach 0.3, 100 m/s) and high pressure (6 atm, 25 m/s) burner rig tests (BRT). Unlike metals, the ceramic was stable at 1300 ◦C. Unlike SiC and Si3N4, neither burner test produced a weight loss, unless heavily pre-oxidized. Lower mass gains were produced in the BRT compared to furnace tests. The commonly observed initial, fast TiO2 transient scale was preferentially removed in hot burner gas (~10% water vapor). A lesser degree of gradual Al2O3 volatilization occurred, indicated by grain boundary porosity and crystallographic etching. Modified cubic-linear (growth-volatility) kinetics are suggested. Gas velocity and water vapor pressure play specific roles for each scale. Furthermore, a 7YSZ TBC on Ti2AlC survived for 500 h in the Mach 0.3 burner test at 1300 ◦C with no indication of volatility or spalling.

**Keywords:** MAX phases; scale volatility; burner rigs; EBC

#### **1. Brief Introduction**

Ceramics are widely indicated for higher temperature turbine components. Water vapor attack has become a new element of study because of volatile hydroxides that form by reaction with Al2O3 substrates or with SiO2 scales that form on SiC and Si3N4. The severity of attack is controlled by the thermodynamics of the reaction, velocity, and water vapor pressure [1–4]. The water vapor component of combusted jet fuel is generally indicated at ~10%, with chemical activity increasing with overall system pressure. Typical volatile hydroxides of common oxides observed or projected are CrO2(OH)2, Si(OH)4, TiO(OH)2, and Al(OH)3, in order of decreasing severity [1–4]. While Cr volatiles can occur at low temperatures, 1200 ◦C or more is required for the others to become a noticeable problem. Thus, widely used Al2O3-forming superalloys, typically limited to 1150 ◦C, have not previously exhibited a scale volatility problem. Also, thermal barrier coatings (TBC) allow for higher gas temperatures, but protect the metal and thermally grown oxide (TGO) from both high temperature and high velocity. Al2O3 and SiC ceramic matrix composites (CMC) generally need an environmental barrier coating (EBC) for continuous use above about 1200 ◦C [5]. Furthermore, SiO2 scales are known to have significantly increased growth rates in water vapor, whereas Al2O3 scales show more complex effects on metals than on MAX phases [6–9].

Another class of materials that presents opportunities at intermediate temperatures is that of Al2O3-forming MAX phases, such as Cr2AlC, Ti2AlC, and Ti3AlC2. These compounds are very oxidation resistant, some at or above 1300 ◦C [10], although strength is lacking for unsupported, load-bearing use at this temperature. A broad program evolved at NASA Glenn to examine Type II low temperature hot corrosion resistance of Cr2AlC coatings on superalloys, basic Al2O3 scale kinetics, and extreme durability of yttria-stabilized zirconia (YSZ) TBC on Ti2AlC [11]. Burner rig testing was

also performed that yielded experimental results for TiO2 and Al2O3 scale durability in water vapor. The purpose of this present paper is to catalogue and analyze high pressure (6 atm) and high velocity (100 m/s) burner results with regards to scale volatility issues. Part of the motivation relates to the behavior of Al2O3 scales in high velocity water vapor at 1300 ◦C, an environment not typically used for standard metal alloy tests. The materials and processes were described in detail from a number of studies comprising the source data for this compilation [12–15]. All the results subsequently presented derive from these studies. For further reference, those studies included broad literature surveys and more in-depth discussions.

#### **2. Materials and Methods**

Ti2AlC MAXthal 211® was obtained from Sandvik/Kanthal (Sandviken, Sweden) and EDM machined into ~0.2 cm thick ×1.2 cm wide ×2 cm long furnace samples, ×4 cm long HP-BRT samples, and ×7 cm long Mach 0.3 BRT samples. These were polished through 2400 grit and ultrasonically cleaned in detergent and alcohol. The box furnace tests (Rapid Temp) were conducted in lab air and samples intermittently weighed/inspected over graduated intervals for up to 500 h. Thermogravimetric (TGA) tests were conducted in dry, bottled air for 100 h, using a vertical tube furnace with weights continuously recorded by a thermo-balance (Setaram, Caluire, France). High pressure burner rig tests (6 atm, 25 m/s) were performed in a jet-fuel burner apparatus, completely sealed in a water-cooled stainless-steel chamber, with pressure controlled by an exhaust valve. Test bars were held at the ends in a water-cooled fixture, at 45◦ to the flame. Typical run series achieved 6 h exposure between shutdown and weightings, usually accruing 50 h total test time. Heating and cooldown were generally achieved within 10 min. High velocity Mach 0.3 BRT tests (1 atm., 100 m/s) employed an open jet-fueled burner, with a cantilever gripped, face-on test bar. The face had been coated with 160 μm of Metco 6700 7YSZ TBC by plasma spray-physical vapor deposition (PS-PVD, Sulzer-Oerlikon, 94 kW, 40/80 Ar/He, 1.5 mbar), the backside left uncoated. Cycling (5 h) was obtained by pivoting the burner away from the sample. Heating and cooling were generally complete in 1 min. BRT sample temperatures were monitored by two-color or 8 μm pyrometers, with gas temperature measured by thermocouple. Samples were examined at various intervals and weighed on a Sartorius analytical balance (±0.00005 g). Scales were examined by optical and FEG-SEM microscopy (Hitachi S-4700, Tokyo, Japan), both on the surface and as Ni-plated, metallographically polished cross-sections at the end of the 500 h BRT. Phase contents were estimated from (Brüker 8D) X-ray diffractometer scans and Rietveld refinement, using Jade software (version 6). Complete experimental details are available in the source studies [12–15].

#### **3. Results**

The weight change results from a suite of 1300 ◦C furnace and jet fuel burner tests are presented in Figure 1. Test times were dictated by the specific study they addressed: standard 100 h TGA, extended 300 h furnace pre-oxidation to produce a very slow growing Al2O3 scale, short 50 h HP-BRT to demonstrate kinetics in a labor-intensive, expensive, pressurized apparatus, and 500 h to demonstrate long term TBC durability in the convenient, available, Mach 0.3 high velocity cyclic rig.

The ambient air box furnace (pre-ox) and dry air TGA show similar gains. These are compared to lower curves for the two burner tests. All results indicate a rapid initial uptake in the order of 1 mg/cm<sup>2</sup> within the first hour of exposure. This has often been associated with a rapid transient growth of discontinuous TiO2 scales that are then undercut by a healing layer of slow-growing, steady-state Al2O3. Nearly linear weight loss was observed in the HP-BRT for a pre-oxidized sample, that would otherwise have been masked by the high initial growth rate. Long-term testing was conveniently enabled by the box furnace, used for a pre-oxidation treatment here, and by the accessibility of the open Mach 0.3 burner test. Weight gain of ~2.4 mg/cm<sup>2</sup> was achieved in the Mach 0.3 BRT as compared to about 1.6 mg/cm<sup>2</sup> for the much shorter HP-BRT.

The cubic scaling kinetics were easily treated by correcting (subtracting) the amount of transient TiO2, as graphically interpolated on log-log plots [12,14]. In the case of TGA tests, it was shown that most of the transient growth, w0, took place in the first 10 min. Good linearized fits of (*w* − *w*0) to (*t* − *t*0) <sup>1</sup>/<sup>3</sup> cubic kinetics could be obtained in these well-controlled TGA furnace tests. The cubic scaling constants, extracted from the (*w* − *w*0) offset-corrected mass gain curves show a very well-behaved, single-mechanism Arrhenius dependency, Figure 2. Over the temperature range of 1000 ◦C–1400 ◦C, an activation energy of 334 kJ/mol·K is seen to apply. Accordingly, the same graphical approach was applied to the 1300 ◦C test data for HP-BRT (Figure 3a) and Mach 0.3 BRT (Figure 3b). These yielded *kc* of 0.024 and 0.011 mg3/cm6·h, respectively, compared to 0.212 mg3/cm6·h determined by TGA [12,13]. These BRT reductions reflect losses due to scale volatility effects. The Mach 0.3 test also incorporates protective effects of the YSZ face-coat and lower temperatures (100 ◦C) away from the hot zone.

**Figure 1.** Comparison of 1300 ◦C Ti2AlC furnace and burner oxidation data. Mach 0.3 burner at 1 atm. and 100 m/s; HPBR at 6 atm. and 25 m/s, TGA dry air, and ambient air furnace tests.

**Figure 2.** Arrhenius plot of log cubic oxidation rate constant vs 1/T; TGA furnace tests of Ti2AlC in dry air for 100 h at 1000–1400 ◦C [14].

**Figure 3.** Cubic oxidation kinetics suggested by linearized transient-corrected weight vs t1/<sup>3</sup> behavior for Ti2AlC. (**a**) 6 atm, 25 m/s, 50 h, HP-BRT; and (**b**) 1 atm., 100 m/s, 500 h, Mach 0.3 BRT, with YSZ face-coat [13,15].

Visual confirmation of TiO2 volatility can be surmised from the low-magnification optical micrographs in Figure 4. Here, the scattered, sometimes oriented, initial clusters of light scale phases are seen to coarsen with furnace exposure time and decrease or disappear with HP-BRT exposure time. This effect was semi-quantitatively verified by Rietveld analyses of X-ray diffractometer scans. While TiO2 (rutile) was the primary transient identified at 1200 ◦C and below, the reaction phase of TiAl2O5 was also identified after 1300 ◦C exposures. Here, initial scale quantities of 20% TiO2 and 10% TiAl2O5 were determined after just 0.2 h of furnace exposure. These decreased dramatically to only 0.1% and 1%, respectively, after 80-h HP-BRT exposures (300 h furnace pre-oxidation), the remainder being α-Al2O3 [13].

**Figure 4.** Optical micrographs depicting discontinuous TiO2 scales formed in furnace tests and successive removal in HP-BRT exposures at 1300 ◦C.

The effect of this HP-BRT exposure on the surface structure can be seen in Figure 5. After 300 h pre-oxidation at 1300 ◦C, transient TiO2 and TiAl2O5 bright clusters (T) were retained in (a), but then largely removed by HP-BRT testing for 80 h at 1300 ◦C in (b). Distinct underlying grains of Al2O3 (A) could then be discerned with a much lower Ti EDS signal overall. A linear weight loss rate of 0.012 mg/cm2·h was also measured, as shown by the lower curve in Figure 1. Since this included some modest scale growth, the total removal rate was surmised to be about 0.017 mg/cm2·h. (Pre-oxidation was required to produce a thick scale with a low instantaneous growth rate less than the volatility rate).

**Figure 5.** SEM/BSE surface microstructures of scales formed on Ti2AlC at 1300 ◦C in (**a**) furnace pre-oxidation for 300 h and (**b**) followed by high pressure burner rig (80 h at 6 atm, 25 m/s). T (Ti-rich transient scale); A (Al2O3) [13].

A direct comparison of scales formed in 1300◦C TGA (100 h) and HP-BRT (50 h) in cross-section is presented in Figure 6. The TGA structure shows the Ti-rich remnants of scattered transient scale colonies, with a dense underlayer of Al2O3. The HP-BRT sample exhibits a rather discontinuous surface scale with less distinct Ti-rich regions, if at all. HP-BRT scale volatility is again suggested. The inner Al2O3–Ti2AlC interface is completely intact with no porosity or cracking.

**Figure 6.** SEM/BSE cross sections of scale microstructures formed on Ti2AlC at 1300 ◦C in (**a**) furnace TGA (100 h) [11] and (**b**) high pressure burner rig (50 h at 6 atm, 25 m/s). Ni (plating); T (Ti-rich transient scale); A (Al2O3); M (MAX phase substrate).

Mach 0.3 BRT (1 atm., 100 m/s, 500 h) exposures produced similar effects on surface scale microstructure (uncoated sample backside) (Figure 7). However, since the burner nozzle was about 2.5 cm in diameter centered on the 5 cm long exposed sample length, the sample temperature at the top (a) and bottom (grip end, c) was about 100◦C cooler. This resulted in a less severe attack, with some remnants of bright Ti-rich particles atop Al2O3 grains, the latter exhibiting grain boundary porosity.

In contrast, the hot section (b) showed little vestige of Ti-rich scales, but a highly irregular, open Al2O3 scale. Some grains appeared to be etched crystallographically, forming lamellae, possibly along the hexagonal (0001) basal planes. The platelets retained a slight Ti level; they may have been derived from TiAl2O5 grains where Ti was removed by selective water vapor corrosion. Xrd analyses of the oxidized surface showed ~10% TiO2 after the initial 20 min at 1000 ◦C, then removed by volatile reactions to just 0.1% after the 1300 ◦C exposure, the remainder of the scale being α-Al2O3 [15]. No phase change in the Ti2AlC substrate was apparent other than reduced x-ray diffraction intensity due to absorption from the thickening scale.

In cross-section, Figure 8, little indication of Ti-rich scales remains, and the scale is thicker in the hot zone region (a). The surfaces are very irregular and open, consistent with scale removal by volatile products. Less of this structure remains in the hot zone region compared to the grip end (b). Again, the scale-substrate interface is seen to be completely intact.

**Figure 7.** SEM/BSE surface microstructures of scales formed on Ti2AlC at 1300 ◦C in Mach 0.3 high velocity burner rig (500 h at 1300 ◦C) Uncoated backside at sample (**a**) top, (**b**) hot center, (**c**) grip end. T (Ti-rich transient scale); A (Al2O3); P (etched platelets) [15].

**Figure 8.** SEM/BSE cross-section images of scale microstructures formed on Ti2AlC in Mach 0.3 BRT (1300 ◦C, 500 h, 100 m/s, 1 atm). (**a**) Hot section, uncoated backside; (**b**) cooler grip end, uncoated backside. Ni (plating); A (Al2O3); M (MAX phase substrate) [15].

Figure 9 presents coating structures typifying the as-sprayed (a) and the hot zone region (b) for the coated face of the Mach 0.3 BRT sample. The coating exhibits deposition columns, first textured by the deposition process (a), then by grain growth and surface smoothing during thermal annealing (b). Here, the flame directly impinged on the YSZ coating face, which shows no features of oxide removal by volatility: the zirconia grains and PS-PVD coating columns are basically intact. In the cross-section (Figure 10), the coating, TGO, and MAX phase substrate are also intact with no interfacial porosity, cracks, or delamination. Porosity and metallographic pullout, however, is observed within the scale. Volatility issues have therefore been prevented on the coated face. Coating survival after 500 h testing at 1300 ◦C and 100 m/s is thus indicated on all accounts. Further testing would not be "cost effective" or especially productive as there was little indication of imminent failure, i.e., the same results are expected for 1000 h testing and beyond. Furnace testing has shown similar durability for the YSZ/Ti2AlC system, surviving 2500 h total, including 500 h at 1300 ◦C and scales up to 40 μm thick [14]. While little evidence is seen for detrimental interface reactions, it can be surmised that sustained Al2O3 growth will be limited by the Al reservoir in the MAX phase substrate.

**Figure 9.** SEM microstructure of YSZ coating surface on Ti2AlC after the Mach 0.3 high velocity burner rig test. (**a**) Grip end, 500 h/1200 ◦C test, (**b**) hot zone, 500 h/1300 ◦C test.

**Figure 10.** SEM/BSE hot zone cross-section of YSZ coated Ti2AlC after BRT (Mach 0.3, 500 h/1300 ◦C). (**a**) Full structure, (**b**) detail showing intact YSZ–Al2O3–Ti2AlC interfaces [15].

#### **4. Discussion**

The previous assemblage of results compared the high temperature scaling characteristics of the oxidation resistant Ti2AlC phase under moist, high velocity burner conditions to those from static and dry atmospheres. The distinct appearance and removal of the initial TiO2 transient scale by moisture in a high velocity gas stream was highlighted. The relative susceptibility of TiO2 to TiO(OH)2 formation in water vapor appears preferential compared to that of Al2O3 via Al(OH)3 volatiles. Such a condition has been examined in concert with Jacobson's thermodynamic treatment of various oxides in flowing moist gases [13]. It was predicted that TiO(OH)2 losses would be on the same order as Al(OH)3 losses, but it is now acknowledged that some uncertainties still remain regarding TiO2 volatiles [16,17]. While the Ti-oxides appeared to be removed preferentially, some losses of Al2O3 are also indicated. Critical studies have indeed demonstrated volatile losses and crystallographic etching for bulk Al2O3 [18–20].

The general removal rate of various scales in various water vapor environments can be modeled according to v1/2pH2O*<sup>n</sup>*/ptotal<sup>1</sup>/2, using the original thermochemical-diffusional approach developed by Opila et al. for various oxides [1]. Here, *n* = 1 for TiO2, *n* = 3/2 for Al2O3, and *n* = 2 for SiO2, as dictated by the chemical reaction with water vapor. Accordingly, the relative severity of the Mach 0.3 test (100 m/s, 1 atm) to the HP-BRT (25 m/s, 6 atm) shown here produces relative rig factors of 0.82, 0.33, and 0.14 for the three scales, respectively [15]. Thus, in the Mach 0.3 test, TiO2 is expected to show similar attack severity as in the HP-BRT, while SiO2 is expected to show less attack, with Al2O3 intermediate. This is consistent with efficient removal of TiO2 observed in both tests and more severe removal of SiO2 in the HP-BRT.

An attempt was made to extract volatility kinetics from the weight change curves using a cubic-linear fit (Chen-Tedmon) [21]. This was partially enabled using COSP for Windows originally designed for cyclic oxidation spalling models. For the present case, cubic growth and uniform scale removal is the model selected [13,22]. To account for a decreasing amount of TiO2 with time, the "spall" (removed) thickness exponent (α) is addressed as a negative number (−3). It is recognized that a constant quotient of the "spall" fraction, *Q*0, and cycle duration, τ, yield identical loss rates per hour. (The response basically converges to continuous curves for τ ≤ 1 h and reproduce their analytical expression).

Some solutions are presented in Figure 11 for (a) the HP-BRT and (b) Mach 0.3 BRT results. Both model curves (dashed) are reasonable fits for the experimental data (symbols). The fitting parameters were *kc* = 0.212 mg3/cm6/h, *Q*<sup>0</sup> = 0.220 for the HP-BRT [13] and *kc* = 0.050 mg3/cm6/h, *Q0* = 0.038 for the Mach 0.3 BRT. Both sets fixed a cubic growth exponent, *m* = 3.0, and used a decreasing spall exponent, α = −3. The HP-BRT fit initiated with the same *kc* determined by TGA in dry air. The resulting *Q*<sup>0</sup> was shown to be consistent with a linear volatility weight loss of 0.01–0.02 mg/cm2·h produced for the pre-oxidized sample [13]. The Mach 0.3 BRT, however, maintained face-to-back and center-to-top/bottom temperature gradients. Accordingly, the net Mach 0.3 growth constant was considerably less, being only about one-quarter that of the HP-BRT. Also, the COSP fit projected a low Al volatility weight loss rate of ~0.002 mg/cm2·h after 100 h, converging to just ~0.001 mg/cm2·h after 500 h, or about one-tenth that of the 50 h HP-BRT rates. Remarkably, the Al2O3 scale thickness under the YSZ in the hot zone (at ~1244 ◦C) after 500 h (~20 μm) was basically the same as that produced in the HP-BRT (at ~1300 ◦C) after just 50 h.

**Figure 11.** Cubic-linear fits to burner oxidation results for Ti2AlC: (**a**) high pressure (6 atm, 25 m/s, 1300 ◦C). (COSP for Windows, *kc* = 0.212 mg3/cm6/h, *Q* = 0.220 mg/cm2/h, *m* = 3, <sup>α</sup> = <sup>−</sup>3); (**b**) high velocity Mach 0.3 for YSZ/Ti2AlC (1 atm, 100 m/s, 1300 ◦C). (COSP for Windows: *kc* = 0.050 mg3/cm6/h, *<sup>Q</sup>* <sup>=</sup> 0.038 mg/cm2/h, *<sup>m</sup>* <sup>=</sup> 3, <sup>α</sup> <sup>=</sup> <sup>−</sup>3).

#### **5. Summary**

Results from various thermal exposures of Ti2AlC incorporating water vapor attack in high velocity gas have been examined. Volatility of TiO2 and Al2O3 scales at 1300 ◦C was indicated by the burner rig results, especially when compared to furnace tests. Discontinuous, superficial colonies of Ti-rich oxide grains were essentially cleaned off in both high-pressure (6 atm.) and long-term, highvelocity (100 m/s) burner tests. While moderate weight gains resulted from continuous Al2O3 growth, 300 h pre-oxidation allowed a net weight loss to be observed for a thick slow-growing scale. The linear loss rate was about one-fifth that determined for SiO2 scales formed on SiC in the same exposure. While Al2O3 losses were not as pronounced as TiO2, grain boundary etching and porosity indicated some volatility effects. The PS-PVD YSZ coating on the hot face of the Mach 0.3 test sample showed no volatility effects and was totally protective of the underlying adherent Al2O3 scale and Ti2AlC substrate. No degradation was apparent for 500 h at a temperature (1300 ◦C) well in excess of current alloy system capabilities. Continued testing was unwarranted and impractical since failure would not be expected even for much longer times.

**Funding:** This research received no external funding.

**Acknowledgments:** Substantial contributions were made to the original works by M. Cuy, B. Harder, A. Garg, R. Pastel, J. Buehler, R. Rogers, and N. Jacobson. Those studies were performed under the NASA Fundamental Aeronautics Program.

**Conflicts of Interest:** The author declares no conflict of interest

#### **References**


© 2020 by the author. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).
