SiO2 thermal barrier coatings

Surface & Coatings Technology 262 (2015) 154–165

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Effect of high temperature on the surface morphology and mechanical properties of nanostructured Al2O3–ZrO2/SiO2 thermal barrier coatings Venkatachalam Rajendran ⁎, Arumugam Karthik, Saturappan Ravisekaran Srither, Sundarmoorthy Arunmetha, Palanisamy Manivasakan Centre for Nanoscience and Technology, K.S. Rangasamy College of Technology, Tiruchengode 637 215, Tamil Nadu, India

a r t i c l e

i n f o

Article history: Received 24 March 2014 Accepted in revised form 16 December 2014 Available online 24 December 2014 Keywords: Composite Sol–gel coating Thermal barrier Oxidation Hot corrosion Nanoindentation

a b s t r a c t Nanostructured Al2O3–ZrO2/SiO2 multilayer coating was applied on 316L SS specimens using dip-coating method. Then, the coated specimens on heat treatment to 800, 900, and 1000 °C for 100 h showed the formation of a stable structural phase of t-ZrO2-toughened Al2SiO5 along with ZrSiO4. The microstructural images of the heattreated specimens observed the level of porosity and that of BHJ-pore size distribution was decreased in the coating along with crack splats. Then, the coating was subjected to cyclic oxidation and hot corrosion test for 20 cycles. The obtained results showed that the rate of cyclic oxidation resistance enhanced when compared to that of hot corrosion resistance. The results of X-ray diffraction analysis showed hot corrosion resistance in molten salt, with non-leaching t-ZrO2 structural phase of ceramic coating. The mechanical properties of the nanostructured coating improved with an increase in the heat treatment. The maximum hardness (H) and reduced elastic modulus (Er) values of the coating were approximately 18.43 ± 0.62 and 68.30 ± 1.35 GPa, respectively. Spallation and oxidation resistance of the coating were explored through mechanical properties, in situ scanning probe microscopy, and atomic force microscopy. The results showed that silicate-based thermal barrier coating provided better protection to SS specimen from high-temperature corrosion. © 2014 Elsevier B.V. All rights reserved.

1. Introduction In recent years, stainless steel alloys have attracted considerable attention because of their potential usage in a wide range of industrial applications to withstand harsh environments. This seems possible due to their important properties such as high strength, toughness, corrosion resistance, and thermal stability [1]. The durability of stainless steel alloys at extreme operating conditions such as high-temperature cyclic oxidation and hot corrosion is one of the crucial properties to be considered. An attempt has been made to develop a new surface on the metal alloy substrates by incorporating a desirable feature through structural ceramic coatings [2]. The structurally insulated materials were used to decrease the interface temperature between the substrate and top ceramic coating and to increase the resistance of the component against high-temperature oxidation [3]. Applications of ceramic-based nanostructured thermal barrier coatings (TBCs) are more dominant than those of any other conventional materials. This is because of their exceptional thermophysical properties at extreme operating conditions such as high temperature and additional mechanical loading [4,5]. Nevertheless, failure of TBCs occurs due to inadequate adhesion between the substrate and ceramic coatings [6].

⁎ Corresponding author. Tel.: +91 4288 274741 4; fax: +91 4288 274880. E-mail address: [email protected] (V. Rajendran).

http://dx.doi.org/10.1016/j.surfcoat.2014.12.039 0257-8972/© 2014 Elsevier B.V. All rights reserved.

The bond coat layer is responsible for the formation of thermally grown oxide (TGO) layers when TBC is subjected to high temperatures. It causes damages such as coating crack, spallation, and oxidation on ceramic top coat layer due to increased thermal stress mismatch between the TGO layers and metallic substrate. When the thickness of microscale ceramic layer exceeds limitation, tunneling of cracks and brittleness occur across the bond coat with base metallic substrate [7,8]. The use of controlled-thickness nanoscale multilayer coating is gaining importance as it increases the lifespan of the product by arresting the crack extension during high-temperature oxidation and corrosion. The existing multilayer coating is capable of arresting thermal loads and rates of chemical reaction across the interface of coatings, resulting in reduced thermal conductivity [9,10]. The advanced refractory materials, such as alumina–zirconia (Al2O3–ZrO2), are used as an insulative ceramic candidate for TBCs because of the metastable tetragonal (t) structural phase of ZrO2 with an active phase of Al2O3 [11–13]. The combination of Al2O3 and ZrO2 binary materials significantly yields the required physicochemical properties [14,15]. In addition, ceramic materials with silica matrix are proposed as a coating material to eliminate the possibility of the advanced cracks, delamination, and spallation of coating at high temperatures. Moreover, the filling of Al2O3 and ZrO2 particles in silica matrix is found to produce a better refractory material [16–18]. Certain amount of porosity and density helps accommodate existing thermal stress within the porous region [19,20]. The mechanical

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behavior of TBCs is an important factor to be studied to explore the mechanical properties, such as hardness and elastic Young's modulus, by nanoindentation [21–23]. Recently, Al2O3–ZrO2-filled silica coating on mild steel (EN3) has been studied as a potential candidate for TBC applications with regard to temperature [13]. In this study, nanostructured Al2O3–ZrO2/SiO2 multilayer coating was applied on 316L SS specimens and its thermal behavior was studied extensively. The structural phase, thermal stability, surface area, pore size distribution and oxidation resistance of the coatings were explored through high-temperature oxidation and hot corrosion studies. The evaluation of microstructural and mechanical properties of the coating was carried out over a wide range of temperatures (from 800 to 1000 °C) for 100 h. 2. Experimental 2.1. Powder production Alumina-toughened zirconia binary (Al2O3–ZrO2) nanoparticles were produced from nitrate precursors [Al(NO3)3∙ 9H2O, ZrO(NO3) 2∙ 2H2O] using well-established hot-air spray pyrolysis technique in our laboratory. The prepared Al2O3–ZrO2 nanocomposite was discussed in our earlier study [13]. Fig. 1. XRD pattern of nanostructured Al2O3–ZrO2-filled silica barrier coating on SS specimen.

2.2. Substrate preparation Identical 316L SS specimens with 2.5 × 1.5 × 0.2 cm dimensions were used in this study. All the specimens were polished using different grades (i.e., 400, 600, 800, and 1000 grit) of silicon carbide sheet and the surface roughness of around 0.3606 μm was kept (Table 1). The polished specimens were washed in an acetone bath under sonication followed by washing in deionized water to remove debris and ions present on their surface. Finally, the specimens were sonicated in 0.1 M nitric acid bath to create micro-scale pore fillers on their surface, enabling them to mechanically anchor with top coating material along with an existing surface roughness. Surface-cleaned specimens of 316L SS (hereafter termed as SS) were used for further studies. 2.3. Nanostructured multilayer coating on SS specimens The homogeneous Al2O3–ZrO2-filled silica sol prepared in our earlier study [13] was used to coat the polished SS specimens by dip-coating method at constant dipping and withdrawing speed of 0.8–1 mm s−1 [24,25]. The coated specimens were air-dried at ambient temperature for 1 h and then preheated at 400 °C for 30 min. This procedure was repeated till six layers were coated on the SS specimens. The nanostructured multilayer-coated specimens have been hereafter termed as SSC specimens. The thermal stability of the SSC specimens was investigated using direct heat treatment in a furnace at 800, 900, and 1000 °C for

100 h (hereafter termed as SSC800, SSC900, and SSC1000, respectively) at the rate of 5 °C min−1. Similarly, the uncoated SS specimen heat treated at 800 °C is termed as SS800. The SS and SSC specimens were subjected to cyclic oxidation and hot corrosion resistance test at the previously mentioned three oxidation temperatures for 20 cycles. 2.4. High-temperature oxidation resistance test The oxidation resistance studies were carried out on both SS and SSC specimens using a programmable high-temperature furnace. The SS specimen was placed inside the furnace and then temperature was raised from room temperature to 800 °C at the rate of 5 °C min−1. The temperature was maintained at 800 °C for 5 h and then the specimen was allowed to cool down to room temperature under atmospheric conditions. The weight of the cooled SS specimen was measured using a digital electronic balance (CP225D; Sartorius, Göttingen, Germany) with an accuracy of ±0.01 mg. This process yielded one cyclic oxidation. The SS specimen was then subjected to 20 such cycles for 100 h. Using this procedure, we carried out high-temperature cyclic oxidation experiments on three identical SSC specimens at 800, 900, and 1000 °C with a total oxidation period of 100 h for each sample.

Table 1 Nanostructured TBCs properties and its corrosion resistance at different temperatures. Specimens

Hardness (H) GPa

Reduced modulus (Er) GPa

Depth displacement (h) nm

Density g cm−3

Specific surface area (SSA) m2 g−1

Average surface roughness (Ra)

Oxidation and hot corrosion 100 h for 20 cycles

SS SS800/SS at 800 °C SSC SSC800/SSC at 800 °C SSC900/SSC at 900 °C SSC1000/SSC at 1000 °C

8.02 ± 0.31 0.901 ± 0.025 3.74 ± 0.22 15.49 ± 0.46 17.69 ± 0.34 18.43 ± 0.62

157.58 ± 1.43 46.77 ± 0.89 107.37 ± 1.10 75.43 ± 1.29 68.33 ± 1.23 68.30 ± 1.35

81.7 223.5 92.6 81.3 82.3 85.9

– – 0.91 1.07 1.25 1.33

– – – 21.2 2.49 1.55

SPM 10 μm

AFM 50 μm

Weight loss g cm−2

0.3603 1.3537 0.6690 0.0189 0.2143 0.2664

– – – 0.171 0.353 0.969

– 0.0238 – 0.0001 0.0013 0.0066

– 0.1553 – 0.0015 0.0034 0.0724

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2.5. High-temperature hot corrosion resistance test The SS and three identical SSC specimens were subjected to hot corrosion test. A mixture of corrosive salts [50 V2O5 + 25 Na2SO4 + 25 NaCl (wt.%)] was deposited onto the SS specimens. Analytical grade V2O5 (98%; Sigma-Aldrich), Na2SO4 (98%; Merck), and NaCl (99.5%; Merck) were used without any further purification. Before the deposition of the salt, all the samples were preheated in furnace at 250 °C for 30 min. After preheating, aqueous solution of crystalline salt was sprayed on the specimens (SS and the three identical SSC specimens) using an ultrasonic atomizer (model VCX 134PB; Sonics & Materials). The saltdeposited samples were subjected to a hot-air oven at 100 °C for 3 h [26]. The specimens were then cooled down to room temperature and the initial weight of the samples was measured. Then, the hot corrosion test was carried out on all the samples for 20 cycles using the oxidation test procedure. The weight of the samples was measured after each cycle. The spallation resistance of the SS and SSC specimens was determined considering changes from the measured weight. 2.6. Characterization The peak positions and structural phase of SS and SSC specimens at different temperatures were analyzed using an X-ray diffractometer (XRD; X'Pert PRO; PANalytical, Almelo, the Netherlands). The diffractometer equipped with Cu Kα radiation source (λ = 1.5406 Å) was operated at 40 kV (30 mA) with a scanning range of (2θ) 10°–80° and rate of 1° s−1 to obtain the XRD pattern of all the samples. The microstructural surface feature and cross-sectional coating were characterized using field-emission scanning electron microscope (FE-SEM; SU6600; Hitachi) and scanning electron microscope (SEM; JSM-6390LV; JEOL,

Tokyo, Japan) studies. A surface area analyzer (BET Autosorb AS-1MP; Quantachrome, Boynton Beach, FL) with multipoint BJH (Barret– Joyner–Halenda) method was used to measure the specific surface area, pore size distribution, and pore volume of SSC coating after the heat treatment. The coated part was peeled out, degassed at 295 °C for 3 h, and then subjected to nitrogen adsorption/desorption isotherms. The chemical compositions of the coating were identified using energy-dispersive spectroscopy (EDS; JSM-6390LV; JEOL). Archimedes's principle was used to measure the density of the SSC specimen before and after heat treatment [27]. The mass gained by the specimen was weighed using a digital electronic balance. The thickness of the nanostructured coating was measured using a profilometer (model SJ-301; Mitutoyo, Japan). To determine the hardness (H) and reduced modulus (Er), we carried out nanoindentation both before and after the heat treatment of the specimens (SS, SS800, SSC800, SSC900, and SSC1000). The nanomechanical test of all the samples was carried out using a Berkovich triangular pyramid diamond nanoindenter (TI 700 Ubi; Hysitron, USA). Initially, a silica quartz fuse was used to calibrate the instrument. The typical three-point indentations were carried out for each specimen during each loading and unloading with an interval of 5 s at maximum force of 1000 μN. The average of H and Er values was given for each specimen (Table 1). The mechanical profile for all the specimens was derived automatically by software (quasi-static analysis tool) from load–displacement curve using the following equation: P max : ð1Þ Aðhc Þ The hardness was described as the ratio of the applied maximum indentation load (Pmax) to the projected area of the hardness impression (A) with function of contact depth (hc). Then, the hardness of the

H¼

Fig. 2. FE-SEM images of microstructure of TBCs on SS specimens: (a) SSC800, (b) SSC900 and (c) SSC1000.

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Fig. 3. BET specific surface area analysis of TBCs. a) Adsorption and desorption isotherms and b) BJH-pore size distribution.

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Fig. 4. SEM cross-sectional image of TBCs on SS specimens: (a) SSC800, (b) SSC900 and (c) SSC1000.

coating was measured using Eq. (1). Er was expressed from the elastic properties of the specimen and using Eq. (2), which defined the peak indentation load and penetration depth for specimen: 1 1−ν 2s 1−ν ¼ þ Er Ei Es

mullite (Al2SiO5) along with SS800. These diffraction patterns are matched with JCPDF nos. 75-1564, 80-2155, and 79-1456 for ZrSiO4, tZrO2, and Al2SiO5, respectively. The new ZrSiO4 and Al2SiO5 phases

ð2Þ

where E and ν represent the Young's modulus and Poisson's ratio of specimen and indenter, respectively. Er of materials is denoted as follows: pffiffiffi π Er ¼ pffiffiffiffiffiffiffiffiffiffiffiffi  S 2 Aðhc Þ

ð3Þ

where S is the elastic contact stiffness from the slope of unloading curves (S = dP/dh) [28,29]. The surface topography of all the samples was determined using in situ scanning probe microscope (SPM) equipped with a nanoindenter. In addition, the surface topography of SSC800, SSC900, and SSC1000 specimens was analyzed using an atomic force microscope (AFM; version 7; Innova, USA). 3. Results and discussion 3.1. Microstructural evaluation of TBCs on the SS specimens Nanostructured multilayer-coated specimens before and after the heat treatment along with base SS are shown in Fig. 1. The diffraction pattern of SS specimen shows their simple face-centered phase [30] and that of the SSC specimen shows amorphous coating with the presence of only SS peaks [13]. However, nanostructured multilayer-coated specimens (SSC800, SSC900, and SSC1000) showed the formation of distinct and complex structural phases of zircon (ZrSiO4), t-ZrO2, and

Fig. 5. EDS spectra of Al2O3–ZrO2 filled silica TBCs on SS specimens.

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exist due to the effect of heat treatment, wherein diffused Si4 + ions from SiO2 react with Al2O3–ZrO2 matrix [31,32]. It is noteworthy that solid-state reaction between SiO2 and Al2O3–ZrO2 of specimen SSC1000 shows the same crystalline phase as that of specimen SSC800. Al2 O3 –ZrO2 þ TEOS→ Al2 O3 –ZrO2 þ SiO2 Amorphous

Diffusion

→ ZrSiO4 þ t‐ZrO2 þ Al2 SiO5 Crystallization ð800–1000



CÞ

Many researchers have reported the formation of ZrSiO4 and Al2SiO5 structural phases from the crystalline phases of t-ZrO2, m-ZrO2, and αAl2O3 with SiO2 network (micro/nanomaterials) at temperatures in the range of 1200–1500 °C [32–36]. Nevertheless, this observation

159

(Fig. 1) shows that ZrSiO4 and Al2SiO5 phases are formed even at a low temperature of 800 °C, which is possibly due to the availability of amorphous Al2O3–ZrO2 content with a narrow particle size distribution (10–55 nm). In our previous study, the crystalline ZrSiO4 phases were observed at 600 °C in Al2O3–ZrO2/SiO2 [13], which supports the present observation. Thus, it is evident that the crystalline ZrSiO4 and Al2SiO5 phases are formed at 800 °C onward while using the amorphous Al2O3–ZrO2 phase. The FE-SEM images of SSC800, SSC900, and SSC1000 specimens are shown in Fig. 2. The observed microstructural features show the dense surface coating is free from cracks, spallation, and delamination. However, after the heat treatment (800, 900, and 1000 °C), the porosity of

Fig. 6. High-temperature cyclic corrosion resistance test of TBCs: (a) oxidation and (b) hot corrosion in static air for 100 h period of 20 cycles.

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the coating is gradually decreased (Fig. 2a–c), which is clear from the porosity distribution of void-free coating. In that case, the isotropic coating shrinkage is expected during the heat treatment while the existence of porosity volume is not burnout throughout coating. Moreover, the observed individual pores are noninteracting with neighboring pores, which is evident from the observed images (Fig. 2a–c). The figure also shows the presence of segmentation crack splats along with semimolten state (Fig. 2c). BET surface area and multipoint BJH desorption results for nanostructured coated specimen after heat treatment are shown in Fig. 3. Specific surface areas of coated specimens are found to decrease in magnitude at elevated temperatures, as shown in the desorption/adsorption hysteresis cure (Fig. 3) and the same is reported in Table 1. The decrease in surface area of coating at elevated temperatures confirms the shrinkage of coating as well as silica diffusion in Al2O3–ZrO2/SiO2 matrix. In addition, it may also be due to the existence of intermediate phase transformation as evident in XRD pattern (Fig. 1) The corresponding isotherm desorption branch of BJH method shows that the pore size distribution in SSC800, SSC900, and SSC1000 specimens decreases over a wide range as obtained in BJH isotherm (Fig. 3) the same is evident in FE-SEM (Fig. 2). The cross-sectional SEM images of SSC800, SSC900, and SSC1000 specimens are shown in Fig. 4, which clearly shows that the porosity distribution in the coating decreases with an increase in temperature (800, 900, and 1000 °C; Fig. 3a–c). This observation is in line with that evident from the FE-SEM images (Fig. 2). The observed

porous structural multilayer coating along with crack segment in turn may enhance thermal insulation, hence protecting SS specimen from high-temperature oxidation without coating compliant [4]. The average thickness of multilayer coating was measured as 2.5 μm using the profilometer. The presence of top ceramic layer with limited thickness on SS specimen is observed as crack-free coating, which is more evident from the FE-SEM and SEM images (Figs. 2 and 4). The density of all the specimens (SSC, SSC800, SSC900, and SSC1000) is shown in Table 1. It is clear from the observed density that nanostructured coating becomes dense as temperature increases. The EDS spectra of the base SS specimen along with SSC and SSC1000 specimens are shown in Fig. 5. The mapping of EDS spectra of SS specimen shows the presence of base elements (Fe, Cr, Mn, and Ni), as shown in Fig. 5. Further, the EDS spectra show the presence of intense peaks for Al, Zr, and Si after coating on 316L alloy with its base elements. Further, SSC1000 specimen confirms the existence of similar elements as that of SSC. 3.2. Cyclic oxidation resistance of nanostructured coating on SS specimen The nanostructured multilayer-coated specimens after subjecting to different cyclic oxidation temperatures (800, 900, and 1000 °C) are shown in Fig. 6a along with the uncoated SS specimen at 800 °C. The base SS specimen undergoes maximum oxidation at the rate of 0.0238 mg cm−2 at the end of 20 cycles after 100 h. However, the SSC

Fig. 7. SEM image of TBCs hot corrosive resistance: (a) SS at 800, (b) SSC at 800, (c) SSC at 900 and (d) SSC at 1000 °C for 100 h after 20 cycles along with EDS spectra.

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Fig. 7 (continued).

at 800, 900, and 1000 °C specimens show spallation of 0.0001, 0.0013, and 0.0066 mg cm−2, respectively, which show that the rate of spallation increases with the increase in temperature (Fig. 6a). Nevertheless, the rate of spallation for SSC specimen at 1000 °C is 261% lower than that for SS specimen at 800 °C. However, the rate of spallation for SSC specimen at 800 °C is negligible when compared with that of specimens oxidized at 900 and 1000 °C. These results indicate that SSC specimen at 800 °C has better oxidation resistance than SSC specimens oxidized at 900 and 1000 °C. This also shows that SSC specimens help control the high-temperature oxidation. Furthermore, it is noted that the surface oxidation resistance of coating seen in the FE-SEM images (Fig. 2) may support the observations made from the cyclic oxidation test.

3.3. Cyclic hot corrosion resistance of nanostructured coating on SS specimen The rate of spallation of SS specimen at 800 °C is approximately 0.1553 mg cm−2 whereas that of SSC specimens at 800, 900 and

1000 °C is 0.0015, 0.0034, and 0.0724 mg cm−2, respectively (Fig. 6b). It indicates that the rate of spallation for SSC specimen at 800 °C is the lowest when compared with that for SSC specimens at 900 and 1000 °C. Meanwhile, the rate of spallation for SSC at 1000 °C is 53% lower than that for SS specimen at 800 °C. However, the observed hot corrosion resistance for SSC specimen at 1000 °C (Fig. 6b) is 90% lower than that of oxidized SSC specimen (Fig. 6a), whereas the same for the SS specimen at 800 °C is 84%. The observed increase in the rate of spallation in hot corrosion resistance test (Fig. 6b) is due to the formation of molten salts sodium vanadate (NaVO3) and sodium sulfate (Na2SO4) during oxidation of the SS and SSC specimens. The formation of Na2SO4 from sulfation (SO2) of sodium chloride (NaCl) is as follows: V2 O5 þ Na2 SO4 þ NaCl→2NaVO3 þ SO3

ð4Þ

2NaCl þ SO2 þ 2O2 →Na2 SO4 þ Cl2 :

ð5Þ

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It is well known that the mixture of corrosive salts reduces the melting point, which instantly increases their flexural corrosion rate [37,38]. The SEM images of SS and SSC specimens subjected to hot corrosion resistance test after 20 cycles along with the EDS spectra are shown in Fig. 7. SS specimen (Fig. 7a) at 800 °C shows a vigorous surface leaching in the form of needlelike crystals. After hot corrosion at 800, 900, and 1000 °C, the microstructural images of SSC specimens (Fig. 7b–d) show them to be covered with fused and molten salts along with top ceramic coating. However, SSC at 800 °C and SSC at 900 °C specimen show better surface stability than SSC specimen at 1000 °C (Fig. 7d). Moreover, spallation and crack propagation in the top surface layer is observed for SSC specimen at 1000 °C (Fig. 7d). The microstructural analysis indicates that infiltration of corrosive salts is more in SSC specimen at 1000 °C than in SSC specimens at 800 and 900 °C. Thus, the top surface layer degrades due to the presence of NaVO3 and Na2SO4 molten salts. It is evident that the corresponding EDS spectra (Fig. 7) which shows that intensity peaks for top ceramic coating consist of continuous oxides scale of Fe, Cr, Mn, Ni, Co, Cl, Na, V, and S along with ceramic Al, Zr, and Si. In this case, the close observation of microstructural SEM images (Fig. 7b–d) shows the porous coating favors the increase in hot corrosion oxidation rate at high temperatures, as reported earlier [39,40]. The XRD pattern of SSC specimens after hot corrosion in 50 V2O5 + 25 Na2SO4 + 25 NaCl (wt.%) environment at different cyclic temperatures is shown in Fig. 8. During the first cycle at 800 °C for 5 h, the coated specimen shows fragile phase of SS alloy (Fig. 8a). When it is extended to the cyclic temperatures of 900 and 1000 °C for 5 h, ZrSiO4 phases are only observed with the distribution of oxidized alloy peaks. On further increasing cyclic temperature to 1000 °C for 10 h (Fig. 8d), t-ZrO2 and Al2SiO5 phases are observed along with ZrSiO4. In addition, three identical SSC specimens with deposition of corrosion salt are further subjected into cyclic hot corrosion at 800, 900, and 1000 °C for 100 h (20 cycles). The observed results are shown in Fig. 8e–g for comparison. The sample after hot corrosion oxidation (20

Fig. 8. XRD pattern of TBC coating after cyclic hot corrosion with molten salt deposition (50 V2O5 + 25 Na2SO4 + 25 NaCl (wt.%)): a) SSC at 800 for 5 h, (b) SSC at 900 for 5 h and (c) SSC at 1000 for 5 and (d) SSC at 1000 °C for 10 h: at 100 h period of 20 cycles, (e) SSC at 800, (f) SSC at 900 and (g) SSC at 1000 °C.

cycles) shows peaks for t-ZrO2, Al2SiO5, and ZrSiO4 phases in addition to the peaks observed for SS oxidized specimen. The peak for t-ZrO2 phase after 20 cycles of heat treatment is absent during the first cycle at 5 h, as evident from Fig. 8a–c, while the major SS alloy species was also disappeared. During the hot corrosion, the fused deposited salts (V2O5, Na2SO4, and NaCl) combine to form molten NaVO3 and Na2SO4 salts [41]. In this case, the molten salt serves as catalyst whereas oxygen carrier readily oxidizes SS specimen through the surface-coated ceramic layer. This results in the formation of corrosion products of Cr2O3, Fe2O3, Fe3O4, and NiCr2O4 [42–44] along with the molten salts as evident from Fig. 8. Moreover, the observed rich ZrSiO4, t-ZrO2, and Al2SiO5 structural phases after 20 cycles (Fig. 8e–g) in the SSC specimens, which are the same the structural phase observed from (Fig. 1). It is emphasized that during the heat treatment and hot corrosion test (Figs. 1 and 8), solidstate silicate-based barrier coating is formed on SS specimen. This indicates the presence of structural phases across the coating may lead to molten salt corrosion resistance.

3.4. Mechanical properties and surface analysis of TBCs Fig. 9 shows the surface mechanical properties of SS, SS800, SSC, SSC800, SSC900, and SSC1000 specimens. Fig. 9a shows the load– displacement curve and topography images of a polished blank SS specimen. The H and Er values of SS specimen are 8.02 ± 0.31 and 157.58 ± 1.43 GPa, respectively, with maximum depth displacement of 81.7 nm. The observed mechanical value is in close agreement with that of SS specimen [45]. After heat treatment of SS800 specimen (Fig. 9b), a substantial change in the values of H and E r (0.901 ± 0.025 and 46.77 ± 0.89 GPa, respectively) is observed with a maximum depth displacement of 223.5 nm. The difference in the mechanical property (Fig. 9a and b) is due to microstructural surface changes caused by high-temperature treatment. The surface changes on SS and SS800 specimens are visualized through in situ SPM (2D and 3D images). In case of nanostructured ceramic-coated SSC specimen (Fig. 9c), the values of H and Er are 3.74 ± 0.22 and 107.37 ± 1.10 GPa, respectively, with a coating displacement of approximately 92.6 nm. This low mechanical value of amorphous SSC specimen indicates soft nature of ceramic coating, which is evident from the observed topographic images (Fig. 9c). The SSC specimens after heat treatment (SSC800, SSC900, and SSC1000) are shown in Fig. 9d–f. The hardness of heat-treated ceramic coating is found to increase with an increase in the operating temperature, whereas the modulus is found to decrease significantly, as reported in Table 1. Nevertheless, depth displacement of coating is increased at elevated temperatures. However, the overall mechanical properties of barrier coatings are not altered. The maximum hardness and elastic Young's modulus values of SSC1000 specimen are found to be 18.43 ± 0.62 and 68.30 ± 1.35 GPa, respectively. The increase in resultant hardness and modulus of coating after heat treatment is due to the densification of the coating with respect to the sintering temperature, as reported earlier [21,46–48]. The obtained maximum hardness value of heat-treated SSC specimens in this study is higher than that of earlier reported Al2O3–ZrO2 [23] and Y2O3–ZrO2 composites [49]. It states that Al2O3–ZrO2-filled silica coating enhances composite's mechanical properties, which are associated with crystalline size effect of Al2O3–ZrO2 nanocomposite and existing stable structural phases of Al2SiO5, ZrSiO4, and t-ZrO2 (Fig. 1). This study confirms that nanostructured multilayer coating resists deformation, scratch, and erosion at high temperatures [50]. The corresponding 2D and 3D projections of SSC specimen after heat treatment (Fig. 9d–f) show that the nanostructured coatings are structurally well bonded with metallic SS specimen at high temperatures [51]. In addition, SSC800 and SSC900 specimens maintain their structural surface stabilities without major surface changes, whereas surface

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changes occur for SSC1000 specimen because of the high-temperature oxidation. AFM images are used to show oxidation and spallation resistance of SSC specimen at different temperatures, as shown in Fig. 10. SSC800 and SSC900 specimens show coating is protected from high-temperature oxidation (Fig. 10a and b). However, small oxidation pits rarely emerge in both specimens. When the specimen is heated at 1000 °C, the level of oxidation pits increases randomly, as shown in Fig. 10c. Further, it is clear that the coating remains free of cracks and deformation at high temperatures (Fig. 10). Moreover, the level of oxidation on coatings is measured from its surface roughness by in situ SPM and AFM image analyses, which is found to increase with respect to increase in temperature, as reported in Table 1.

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4. Conclusions Nanostructured Al2O3–ZrO2-filled silica multilayer coating that served as thermal barrier for SS specimens with an effect of temperature shows the complex crystalline phase of t-ZrO2-toughened silicates (Al2SiO5 and ZrSiO4). The FE-SEM and SEM images of surface feature and cross section of heat-treated structural SSC specimens show a better surface-protective coating against high temperature, due to porous and segment crack splats. The surface area and BJH-pore size distribution of heat treated SSC specimen are decreased in wide range of elevated temperatures. The cyclic oxidation shows an improved oxidation resistance when compared to that of hot corrosion test. The observed results confirm that the existence of porous coating was more favorable to

Fig. 9. Nanoindentation test images of TBCs on SS specimens: (a) SS, (b) SS800, (c) SSC, (d) SSC800, (e) SSC900 and (f) SSC1000.

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Fig. 9 (continued).

accommodate thermal stress and enhance cyclic oxidation resistance. After hot corrosion, the XRD pattern shows stable high-temperature structural phases of silicates and t-ZrO2 without their monoclinic phase under molten salt deposition. Further, the improved mechanical properties of nanostructured coating are because of densification of the coating with respect to increasing temperatures as well as stable structural phases of t-ZrO2, Al2SiO5, and ZrSiO4. Thus, depth displacement, surface roughness, oxidation pits, and spallation resistance of multilayer coating were observed using nanoindentation, in situ SPM, and AFM analyses. This study confirms that nanostructured Al2O3–

ZrO2/SiO2 coating can be considered as a potential candidate for hightemperature TBC applications.

Acknowledgments A. Karthik acknowledges the Senior Research Fellowship (SRF) (08/570(0002)/2013)-EMR-I provided by the Council of Scientific and Industrial Research (CSIR), New Delhi, to carry out this research project.

V. Rajendran et al. / Surface & Coatings Technology 262 (2015) 154–165

165

Fig. 10. AFM surface topography 3D image of TBCs on SS specimen: (a) SSC800, (b) SSC900 and (c) SSC1000.

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