Multifunction high temperature coating system based on aluminium particle technology

Transcription

1 Multifunction high temperature coating system based on aluminium particle technology V. Kolarik a, M. Juez-Lorenzo b, M. Anchústegui. H. Fietzek Fraunhofer-Institute for Chemical Technology ICT Joseph-von-Fraunhofer Str 7, Pfinztal, Germany a vladislav.kolarik@ict.fraunhofer.de; b maria.juez-lorenzo@ict.fraunhofer.de Keywords: Multifunction, thermal barrier coating, diffusion coating, aluminium particle technology Abstract. Spherical Al particles sized in the range of 2 to 5 µm were deposited with an organic binder by brushing on the austenitic steel X6 CrNi (Alloy 304H). The coated samples were annealed in air at 400 C for 1 h in order to expel the binder. For studying the oxidation behaviour in air, isothermal experiments were performed at 700 C and 900 C with oxidation times of 5 h, 100 h and 1000 h. The oxide formation was studied in situ by high temperature X-ray diffraction (HT- XRD) up to 100 h. Field emission scanning electron microscopy (FE-SEM) was applied to investigate the surface and the cross-section of the particle coating. During oxidation, the stable α-al 2 O 3 was identified in situ by HT-XRD on all studied samples at both temperatures. No meta-stable alumina phases were found. In the initial state, 2 h at 900 C, the Al particles are completely oxidised to hollow alumina spheres, controlled predominantly by the reaction due to the small particle size and relatively high surface portion. Simultaneously, the Al-rich diffusion layer is formed in the substrate. On further exposure, a thin protective alumina scale continues growing on the top of the diffusion layer. After exposure to both 700 C and 900 C, a coating structure was encountered, which consists of a quasi-foam top coat from conjoint hollow spherical alumina particles and an Al-rich diffusion layer below. The quasi-foam top coat has the potential to effectuate as thermal barrier by gas phase insulation, while the diffusion layer below serves as protective coating against oxidation. The approach by particle size processing opens a potential for obtaining a complete thermal barrier coating system in one manufacturing step. The coating properties can be adjusted by parameters like selection of source metal/alloy, particle size, substrate, binder and heat treatment. Introduction Raising the operating temperature is mostly the essential key parameter for achieving a higher efficiency of energy conversion as well as a reduction of pollutants and of the CO 2 emission in practically all processes running at high temperatures. Several research programmes, like the COST Actions 522 and 536 for example, have been dedicated to this matter. Efficient protection of components at high temperatures against aggressive environments can only be achieved by coating systems, as the structural material itself must be designed to meet the mechanical loads. Higher operating temperatures mostly lead the existing coatings to their limits and imply the demand for more advanced coating systems. Furthermore, increasing material costs and a severe competition on the global market raise the demand for economic approaches. In land based gas turbines, metallic overlay coatings are used, which protect the blades and vanes from oxidation and corrosion, as well as thermal barrier coatings (TBC) protecting the metal against high temperatures and/or reducing the effort for inner cooling. The overlay coatings consist usually of an MCrAlY alloy, where M is Ni and/or Co, and are deposited mostly by thermal spraying methods, such as LPPS or HVOF [1,2]. Thermal Barrier Coatings (TBC) from yttria partially stabilized ZrO 2 (YPSZ) are deposited by plasma spraying methods and Electron Beam-Physical

2 Vapour Deposition (EB-PVD) [1,2]. Bond coats with similar compositions like the overlay coatings are applied between substrate and TBC. While the oxidation mechanisms of the overlay MCrAlY coatings are presently well understood, research and development are now focussed on the TBC systems. Their spectrum of failure mechanisms is more complex and comprises effects, such as spallation, phase transitions, sintering, inter-diffusion and bond coat oxidation [3,4]. In boilers, heat exchangers, super-heaters and reactors, easily weldable ferritic and austenitic steels are widely used in the high temperature components. The operating temperatures are lower than in gas turbines, but the environments are often much more aggressive. In order to develop advanced protective coatings for boiler and heat exchanger materials, different coating compositions deposited by pack-cementation, aluminisation and fluidized bed have been investigated [5-8]. Detailed models supported by long-term experimentation were developed in order to understand the diffusion processes and their impact on the lifetime [6]. Numerous research was also dedicated to the mismatch of thermal expansion coefficients of Fe-Al phases and the substrate material [7,8]. Special attention has been paid on slurry diffusion coatings for application in boilers as well as in steam and gas turbines, aiming at aluminising the surface in order to achieve the formation of a protective alumina scale at high temperatures [9-11]. These coatings are produced by depositing particles of aluminium, aluminides or mixtures of Al with other metals like Ti mostly by brushing and pasting with a subsequent curing at temperatures around 350 C and a diffusion heat treatment at 700 C. The demand for developing high performance and economic protective coatings as well as TBCs however, is huge, although commercial thermal spray coatings as well as slurry coatings are currently offered [12-14]. Recent research on the oxidation of nano- and micro-sized metal particles revealed that spherical Al particles in the size range of 2 µm to 20 µm do not form meta-stable alumina phases at high temperatures. They form rather hollow spheres consisting of the stable α-al 2 O 3 filled with air [15]. The discovered particle size window was the initial motivation for a research work based on the leading idea to deposit spherical Al particles with an appropriate size by brushing, spraying, dipping or rolling, and obtaining by an adequate heat treatment a complete coating system comprising an adherent top coat from conjoint hollow oxide spheres and a diffusion layer below it. The thermal barrier effect could be achieved by gas phase insulation in the hollow oxide spheres yielding more flexibility in the source metal selection. The possibility to manufacture such a coating in one procedure could provide multi-functionality for a wide spectrum of applications. Such a new concept would in fact offer breaking new ground in coating technology, but extensive research work is still necessary. The purpose of the present work was to investigate the basic viability and potential for the development of economic coating systems based on particle size processing. Experimental Spherical Al particles sized in the range of 2 to 5 µm were deposited with silica oil as binder by brushing on the austenitic steel X6 CrNi (Alloy 304H) obtaining a coating with an average thickness of 60 µm. The Al-particles were produced by SIBTHERMOCHIM in Russia. The composition of X6 CrNi (material number ) is Fe (bal), C, 1.0 Si, 2.0 Mn, P, S, Cr and Ni. The samples were cured at 400 C for 1 h in order to expel the binder. For studying the oxidation behaviour in air, isothermal experiments were performed at 700 C and 900 C with oxidation times of 5 h, 100 h and 1000 h. The oxide formation at 900 C was studied in situ by high temperature X-ray diffraction (HT-XRD) for 100 h, recording a series of diffraction patterns with time intervals of 2 h, and in the initial state up to 5 h with time intervals of 30 min. The experimental set-up for the HT-XRD consists of an X-ray diffractometer, a high temperature device with a programmable temperature controller and a linear detector. Isothermal measurements and freely selectable temperature programmes can be performed between room tempera-

3 ture and 1600 C under oxidizing conditions. Phase changes and the formation of new products are detected continuously during the experiment. For the oxide growth, the intensities of the oxide peaks from a series of in situ X-ray diffraction patterns are determined as a function of time by a summing method. This procedure calculates the peak intensities summing the counts of each channel in the range of the peak. The background is subtracted to eliminate the influence of changing background intensity. The resulting intensity curves iz(t) for each oxide and its modification show their formation as a function of time taking into account the absorption of the X-ray beam in the growing oxide layer [16]. Field emission scanning electron microscopy (FE-SEM) with energy dispersive element analysis (EDX) was applied to study the surface of the particle coating as well as the cross-section. Results After deposition and curing for 1 h at 400 C, the raw coating consists of spherical Al particles with an average size of 2 to 5 µm forming a homogeneous surface structure (Fig. 1). In Fig. 2 the series of HT-XRD patterns recorded in situ during the first 100 h of oxidation at 900 C is plotted. The first pattern in Fig. 2 was recorded at room temperature before heating and represents the Al particle coating in the state as shown in Fig. 1 after curing and identifying the reflexes from aluminium as well as those from the austenitic substrate below. The series of HT-XRD patterns recorded in situ during the first 100 h at 900 C, reveals the formation of α-al 2 O 3 from the first pattern at room temperature. No meta-stable alumina phases were found by HT-XRD. After 8 h of oxidation small amounts of α-fel 2 O 3 are detected. In the X- ray diffraction pattern recorded after cooling at room temperature, reflexes from the alloy are observed besides those of the oxides (Fig. 2). At 700 C, α-al 2 O 3 is the only oxide that is identified in situ during the whole oxidation time of 100 h. The Pt reflexes stem from the sample holder. Fig. 1 Surface of the raw coating after curing for 1 h at 400 C

4 t Fig. 2 Series of HT-XRD patterns recorded in situ during the oxidation at 900 C for 100 h. The first pattern is recorded before and the last one after the oxidation. For evaluating the oxidation kinetics, the intensities of the XRD reflexes of α-al 2 O 3 were analysed as a function of time, yielding the intensity curve iz(t). At 900 C the intensity curve iz(t) of α-al 2 O 3 climbs in the first 2 h strongly, and on further exposure until 100 h it grows slower following a more parabolic course (Fig. 3). The HT-XRD measurement in the initial state up to 5 h reveals the change of the oxidation kinetics with a higher time resolution. A quasi-linear course of the intensity curve iz(t) is observed in the first 2 h, and then only a slight increase follows (Fig. 4). The morphology of the surface after oxidation at 900 C is for the investigated oxidation times of 5 h, 100 h and 1000 h comparable and shows no evident difference for longer exposure times. Alumina spheres with a clean, smooth surface are found, obviously consisting of the stable phase α-al 2 O 3. Approximately 50 % of the spheres however, show square-edged patterns on their surface indicating the plate-like morphology typical for θ-al 2 O 3, while maintaining their overall spherical shape in the average range of 3 to 7 µm. Even thin whiskers in nano-scale are found on some particles (Fig. 5). After the oxidation at 700 C a similar morphology composed of spheres with smooth and square-edged surface is observed for the studied oxidation times of 5 h, 100 h and 1300 h (Fig. 6). No θ-al 2 O 3 was identified by XRD after the oxidation at both 700 C and 900 C. In cross-section after exposure to both 700 C and 900 C, a coating structure was encountered that consists of a quasi-foam top coat from hollow spherical alumina particles and a diffusion layer below that forms a protective alumina scale. Fig. 7 shows such a coating formed in 5 h at 900 C. The quasi-foam alumina top coat has an average thickness of 60 µm and is adherent to the metal. The diffusion layer consists of two iron aluminide phases, a matrix with 38 wt% Al (light grey in Fig. 7) and a phase with 53 wt% Al distributed in it with increasing content towards the interface with the top coat, where it forms a thin coherent layer (dark grey in Fig. 7). A nickelenrichment is found in the substrate below the aluminium diffusion layer. For investigating to which extent the spherical shape of the Al particles is of importance, mechanical pressure was applied to the coating surface before the curing process in order to deform

5 the particles. The oxidation experiments were then performed under the same conditions as described above. With the slightly deformed Al particles a different surface morphology is obtained for both studied temperatures and all oxidation times. Already after 5 h at 700 C the alumina particles show a broke-open structure, which on closer inspection reveals that one hemisphere collapsed into the interior of the other, forming some kind of double-walled half-shell (Fig. 8). This collapse effect occurs as well after 1000 h at 900 C, although the spheres show holes rather than a collapsed hemisphere and intact particles are still observed (Fig. 9). The collapse effect demonstrates that the top coat is really consisting of hollow oxide spheres as clearly evident from Figs. 8 and 9. Sample V 900 C 100h Sample U. 900 C 5hr Intensity iz(t) Intensity iz(t) Al2O3 Alfa-Al2O Time [h] 100 0,0 1,0 2,0 3,0 4,0 5,0 6,0 Time [h] Fig. 3 Intensity curve iz(t) of α-al 2 O 3 in the first 100 h at 900 Fig. 4 Intensity curve iz(t) of α-al 2 O 3 in the first 5 h at 900 C a) b) c) Fig. 5 Surface of the coating after oxidation at 900 C in air (10000 : 1) a) 5 h b) 100 h c) 1000 h 2 µm

6 Quasi-foam structured alumina Diffusion zone Substrate Fig. 6 Surface of the coating after 1300 h oxidation at 700 C (20000 : 1) Fig. 7 Quasi-foam alumina coating with Alrich diffusion zone on an austenitic substrate after 5 h at 900 C Fig. 8 Surface of the coating after 5 h at 700 C with Al particles deformed prior to exposure (10000 : 1) Fig. 9 Surface of the coating after 1000 h at 900 C with Al particles deformed prior to exposure (5000 : 1) Discussion The results confirm that on spherical Al particles with an average size of 2 to 5 µm the formation of meta-stable alumina phases is suppressed. The relatively high surface portion of the Al particles in this size range probably facilitates the reaction controlled oxide formation, which is reflected in the enhanced quasi-linear kinetics during the oxidation of the Al particles (Figs. 3 and 4). With a particle radius of 1 to 2.5 µm, the diffusion paths are short and allow furthermore a quick transport of Al to the surface. Diffusion along grain boundaries can be neglected with the given particle size leading to a homogeneous transport rate of Al to the surface. It is possible that under these conditions the formation of the stable α-al 2 O 3 is favoured. As alumina scales grow predominantly by outward diffusion of Al, hollow alumina spheres are formed from the spherical Al particles. The local morphologies typical for θ-al 2 O 3 (Fig. 5) could be due to lattice defects in the Al particles, which influence the diffusion paths. The HT-XRD patterns however, do not detect any

7 meta-stable alumina phases in situ. Consequently, if θ-al 2 O 3 forms initially on the sphere surface, the transformation to the stable α-al 2 O 3 is too fast for being detected by the HT-XRD. The platelike morphology however remains after the phase transition, an effect that was demonstrated elsewhere [17]. An evaluation of the corresponding particle sizes didn t reveal any size dependencies. The deposition of spherical Al particles with a diameter of 2 to 5 µm on an austenitic substrate and the subsequent subjection to high temperatures according to the described parameters, leads to the formation of a structure, which has the potential for a complete coating system that consists of a top coat from hollow alumina spheres and an Al-rich diffusion layer below it (Fig. 7). From the intensity curves iz(t) in Figs. 3 and 4 it can be deduced that the formation of the quasifoam alumina top coat occurs at 900 C within the first 2 h, when the Al particles are oxidised to alumina spheres. Simultaneously, an Al-rich diffusion layer begins to form in the substrate. With further exposure, the intensity curve iz(t) grows more moderately with a quasi-parabolic course, monitoring the formation of a protective alumina scale on the top of the diffusion layer. The top coat with its structure of conjoint hollow alumina spheres has the potential to act as a thermal barrier by gas phase insulation. Further research however, is needed and currently being performed for investigating the optimum parameters in regard to the deposition process, the heat treatment and the thickness for achieving a mechanically stable and adherent top coat. Also the mechanical performance and potential application fields are subject of further research. The diffusion layer simultaneously formed below the top coat yields protection against oxidation and corrosion. Only a smaller part of the Al diffuses into the substrate, as the majority of the Al particles are oxidised forming the top coat. The Al content in the diffusion layer is therefore limited, and the formation of the Al-rich brittle iron aluminides, such as Fe 2 Al 5, is local. The diffusion zone consists mainly of iron aluminides with lower Al content, in view of the phase diagram lower than 40 wt%. The formation of small amounts of hematite detected in situ at 900 C indicates that iron aluminides with lower Al-contents must be predominant during exposure to temperature. In experiments deforming the Al source particles prior to oxidation, it was shown that already a slight deformation from the spherical shape prevents the particles from forming stable hollow alumina spheres by oxidation. During the exposure to high temperature, one hemisphere collapses into the other resulting in a double-walled half-shell. The following mechanism is proposed for the collapse effect: Subjected to the oxidation temperature, the Al particle is quickly covered by an alumina scale before it can melt and change its shape. Molten Al portions are kept in the alumina shell. The alumina scale is growing predominantly by Al diffusion outward, rather than by oxygen diffusion inward, leading to the formation of a void inside the particle. With a growing alumina scale, the void formation inside the particle is proceeding continuously. As the inward diffusion of air is slow, the gas pressure inside the sphere is very low. Due to the spherical coving of the particle, the alumina scale can withstand the external atmospheric pressure until the void is gradually filled with air. With a deformed shape, however, the alumina shell cannot take advantage of a homogeneous coving to stand the outer atmospheric pressure and collapses at the deformed position. The mechanism of the collapse effect is plotted graphically in Fig. 10. Fig. 10 Model of the collapse effect mechanism in the case of deformed Al spheres

8 The collapse effect is more pronounced at 700 C than at 900 C (Figs. 8 and 9). At 900 C the alumina scale on the particle grows faster and reaches the maximum possible scale thickness earlier and therewith the maximum mechanical strength to counteract the outer atmospheric pressure. The collapse effect causes more often holes at the deformation positions (Fig. 9). The results with deformed Al particles and the understanding of the collapse effect imply that the spherical shape of the Al particles is essential for the formation of the top coat from hollow alumina spheres. Conclusions The approach by particle size processing using spherical source metal particles with a defined diameter opens a potential for obtaining a complete thermal barrier coating system with an Al diffusion layer below it in one manufacturing step. The coating properties can be adjusted by parameters like selection of source metal/alloy, particle size, substrate, binder and heat treatment. The thermal effect is achieved by gas phase insulation in the hollow oxide spheres, leading to more flexibility in material selection. References [1] C. Coddet: Mater. Sci. Forum Vols (2004) p. 193 [2] M.G. Hocking, V. Vasantasree, P.S. Sidky, Metallic and ceramic coatings: production, high temperature properties and applications., Longman Scientific & Technical, London (1989) [3] O.A. Adesanya, K. Bouhanek, F.H. Stott, P. Skeldon, D.G. Lees, G.C. Wood: Mater. Sci. Forum Vols (2201) p. 639 [4] S. Saunders, J. Banks, S. Osgerby, D. Rickerby, C. Chunnilall: Mater. Sci. Forum Vols (2001) p. 775 [5] V. Rohr, M. Schütze, E. Fortuna, D.N. Tsipas, A. Milewsha, F.J. Perez: Materials and Corrosion Vol. 56 (2005) p. 874 [6] Y. Zhang, A.P. Liu, B.A. Pint: Materials and Corrosion Vol. 58 (2007) p. 751 [7] Y. Zhang, B.A. Pint, G.W. Garner, K.M. Cooley, J.A. Haynes: Surf. Coat. Technol. Vols (2004) p. 35 [8] B.A. Pint, Y. Zhang, P.F. Tortorelli, J.A. Haynes, I.G. Wright: Materials at High Temperatures Vol. 18 (2001) p185 [9] A. Aguëro, R. Muelas, M. Gutiérrez, R. Van Vulpen, S. Osgerby, J.P. Banks: Surf. Coat Technol. Vol. 201 (2007) p [10] A. Agüero, R. Muelas, A. Pastor, S. Osgerby: Surf. Coat. Technol. Vol. 200 (2005) p [11] K. Murakami, N. Nishida, K. Osamura, Y. Tomota: Acta Materialia Vol. 52 (2004) p [12] Information on [13] Information on [14] Information on [15] N. Eisenreich, H. Fietzek, M. Juez-Lorenzo, V. Kolarik, V. Weiser, A. Koleczko in: Proc. Int. Conf. Microscopy of Oxidation, edited by G.J. Tatlock and H.E. Evans, Science Reviews (2005) pp [16] V. Kolarik, W. Engel and N. Eisenreich: Mater. Sci. Forum Vols (1993) p. 563 [17] V. Kolarik, A. Kolb-Telieps, H. Hattendorf, M. Juez-Lorenzo, H. Fietzek, R. Hojda in: Materials Aspects in Automotive Catalytic Converters, edited by Hans Bode, Willey-VCH, ISBN , (2001) pp