Effect of the powder particle structure and substrate hardness during vacuum cold spraying of Al2O3

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Effect of the powder particle structure and substrate hardness during vacuum cold spraying of Al2O3 ⁎

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Li-Shuang Wang1, Heng-Fu Zhou1, Ke-Jie Zhang, Yu-Yue Wang , Cheng-Xin Li , Xiao-Tao Luo, Guan-Jun Yang, Chang-Jiu Li State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi'an Jiaotong University, Xi’an, Shaanxi Province 710049, PR China

A R T I C L E I N F O

A BS T RAC T

Keywords: Vacuum cold spray Alumina coating Powder particle structure Substrate hardness

In this paper, the effect of the powder particle structure and substrate hardness during vacuum cold spraying (VCS) of Al2O3 is investigated. Our results help understand the underlying deposition mechanism during VCS in more detail and enable the tailoring and improving of the resulting coatings. Two structurally different alumina feedstocks were used for this study. We find that the loosely agglomerated powder bonds to the substrate primarily through coordinated deformation of the nano-sized powder particles. The sintered powder, on the other hand, bonds to the substrate through severe fracture and deformation of the particles. High-resolution transmission electron microscopy (HR-TEM) was employed to observe details in the interfacial microstructure of the coatings on the two substrates with differing hardness. The hard steel substrate facilitates particle fracture, which leads to cohesive particle/particle-bonding in the coating region close to the substrate. The softer aluminum substrate leads to strong interfacial coating/substrate-bonding because the particles are embedded into the substrate. In summary, the fracture and deformation of the feedstock as well as the substrate hardness affect both adhesion (coating/substrate bonding) and cohesion (particle/particle bonding) considerably.

1. Introduction Ceramic coatings are widely used as thermal-resistant coatings [1], corrosion-resistant coatings [2], wear-resistant coatings [3], and others. Typically, these coating are produced using chemical vapor deposition [4], magnetron sputtering [5], electrophoretic deposition [6], or thermal spray coating [7]. Among these methods, thermal spray coating is one of the most promising methods because it is also a fast, flexible, and cost-effective technique. However, due to the larger feedstocks used in thermal spraying (15–100 µm), the thickness of the obtained coatings typically exceeds 50 µm. Vacuum cold spray (VCS) is a variation of thermal spray technology. It is also known as aerosol deposition method (ADM) [8], and has been widely investigated as an emerging coating-deposition method that is especially suited for depositing thin and dense ceramic coatings with submicronand nano-sized powder, at room temperature [9,10]. During VCS,

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powder particles are accelerated and injected into a vacuum chamber with the help of a carrier gas. The resulting high speed of the particles ensures strong impact when they hit the substrate. After the particles were fractured and deformed, many desirable ceramic coatings are obtained together with the compaction of any residual powder. Upon the high-speed impact, the particles are fractured and deformed. With the successive compaction of particles, the vacuum cold sprayed ceramic coatings are built up layer by layer. All this can be accomplished without additional heating [8]. Unlike thermally sprayed ceramic coatings, the ceramic feedstock is deposited at room temperature during VCS. This makes it easier to produce nano-sized ceramic coatings and to expand the range of available suitable substrates (ceramics, metals, polymers). Moreover, VCS is an excellent flexible method to produce ceramic films with thicknesses ranging from several to tens of micrometers. This is possible thanks to the option to use nano- or micro-sized feedstocks [11,12]. Additionally, the deposition

Corresponding authors. E-mail addresses: [email protected] (Y.-Y. Wang), [email protected] (C.-X. Li). These authors contributed equally to this work, and should be considered as co-first authors.

http://dx.doi.org/10.1016/j.ceramint.2016.12.085 Received 31 August 2016; Received in revised form 22 November 2016; Accepted 15 December 2016 0272-8842/ © 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Wang, L.-S., Ceramics International (2016), http://dx.doi.org/10.1016/j.ceramint.2016.12.085

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2. Materials and method

rate of VCS is about hundred nanometers to several micrometers per minute, which is considerably higher than for CVD [4], sputtering [5], and other common methods used to produce ceramic films [6]. Furthermore, these advantages extend the application of VCS to produce a wide range of other products. For example, VCS made it possible to produce porous TiO2 photo-anode coatings for dye-sensitive solar cells [13,14], and porous (La, Sr) (Co, Fe) O3-δ cathode layers for solid oxide fuel cells [15]. VCS, through feedstock fragmentation under intensive compaction, makes it possible to prepare dense ceramic coatings,. The technique can be used to prepare lead zirconate titanate (PZT) [16] coatings with excellent ferroelectric properties for microelectromechanical systems, alumina coatings for electronic components [17], electrolyte layers for solid oxide fuel cells [18], solid-state Li-ion batteries [19], and corrosion-resistance coatings for metal substrates [20]. All of these VCS-prepared dense coatings show excellent properties in their respective applications. Despite the various intended uses of these VCS-prepared coatings, cohesion and interfacial adhesion are the two main research topics in previous studies [9,10,21]. A typical material used in VCS, α-Al2O3, is commonly used to explore the compacting and bonding mechanism during the VCS process. Akedo et al. [22] studied the deposition mechanism through experiments and numerical simulation. Their results demonstrate that the crystallite size that appears in the deposited layers is much smaller than in the starting powder. Additionally, the estimated maximum shock energy of feedstocks was similar to the fracture toughness of the powder material. This suggests that fracture of the particle or plastic deformation occurs during deposition. Moreover, Chun et al. [23] suggested that the built-up of VCS coating was mainly the result from particle fragmentation and consolidation. Although the deposition mechanism of VCS has not been understood well until now, it is clear that both the powder structure and the substrate properties are crucial for the preparation of the desired coatings [16,17,21,24–26]. Nam et al. [16] prepared a dense alumina coating using VCS on glass and feedstock with a diameter of 400 nm. They also found that chalk-like coatings can grow on glass with hard alumina agglomerating powder. This confirms that the structure of the feedstock affects the deposition results for alumina coating with VCS. Kim et al. [27] studied the effect of metal substrates on Y2O3 coatings produced by VCS, as well as interface anchoring and the densification of coatings. They discovered that Sn with its lower melting temperature produces better interfacial bonding than stainless steel. On the other hand, a much denser coating was obtained on stainless steel because of its higher hardness compared to Sn. Lee et al. [26] investigated the effect of substrate hardness on the growth of VCS alumina coatings by characterizing the microstructure of Al2O3 particles at the initial deposition stage. The results show that the deposition rate and the hardness of the obtained coatings are proportional to the hardness of the substrates. The aforementioned investigations indicate that both the powder structure and the substrate properties have strong effects on adhesion and densification. However, it is still unclear how these factors affect the deposition and the resulting microstructure i.e. the particle/particle interface and coating/substrate interface. The intrinsic differences of deposition behavior resulting from feeding powder with different structures and substrate properties have not been fully understood. They may be the key, however, to reveal details of the deposition mechanism during VCS. In this study, two types of alumina powder, a loosely agglomerated nano-crystal powder and submicron sized sintered powder were used to study the effect of the feedstock structure. In addition, some mechanical properties of the obtained coatings were evaluated to help understand the internal microstructural differences between these two coatings. Furthermore, two types of metal substrates with different hardness values were also used to explore the effect of substrate hardness on the obtained microstructure, phase transition, and interfacial bonding.

2.1. Materials Two types of commercially available α-Al2O3 powder (Showa Denko, Japan) were used as the starting powder for the coatings. One is loosely agglomerated nano-crystal powder with a diameter of 80 nm (powder A). The other is sintered submicron powder with a diameter of 400 nm (powder B). We used glass as a substrate to investigate the deposition behavior for these two different powder types. In order to investigate the effect of substrate hardness on the deposition behavior of alumina coating, two types of metal substrates, Al and stainless steel (SUS304), were used. According to the previously published study [28], smooth substrate surfaces aid the deposition. Therefore, these two substrates with different hardness values were polished to obtain a similar surface roughness. Before deposition, the substrates were cleaned in acetone with ultrasound, for ten minutes. 2.2. Sample preparation The alumina coatings were deposited using a VCS-2000 vacuum cold spray system developed at Xi’an Jiaotong University, as shown in Fig. 1. As reported previously [14], the system consists of a vacuum chamber, an aerosol chamber, a particle-accelerated nozzle, a carrier gas unit, a two-dimensional worktable, and a control unit. During the deposition process, α-Al2O3 powder was carried using helium gas to produce an aerosol flow. The particles were accelerated by the nozzle again, injected into the vacuum chamber, before they hit the substrate at high speed. The deposition parameters are shown in Table 1. 2.3. Coating characterization The morphology and phase structure of both the feedstocks and VCS-prepared coatings were characterized using scanning electron microscopy (SEM, TESCANMIRA 3 LMH, Czech) and X-ray diffraction (XRD, X'pert PRO, PANalytical, The Netherlands). The microstructure of the coating/substrate interface was examined using high-resolution transmission electron microscopy (HR-TEM, JEM-2100F, JEOL, Tokyo, Japan). The preparation details for a cross-section of a coating layer for TEM observation are the following: After face-to-face gluing of the coatings, the sample was polished from both sides with a grinding machine to obtain about 20 µm. Then, the sample was ion-milled with an ion mill (Gantan 691, Ganta, Oxford, UK) to obtain a suitable thickness for TEM examination. More details of the preparation of the coating sample for TEM observation can be found elsewhere [29]. The coating hardness was measured by nano-indentation using the NanoTest (Nano Indenter G200, MTS). The hardness was derived

Fig. 1. Schematic diagram of the vacuum cold spray setup.

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porous, and thus the formed coatings still contain nano-sized pores and gaps. Consequently, the deposition rate for powder A was higher than for powder B. The nanoindentation test suggests that the coatings with powder A and powder B have a mean hardness of 3.5 ± 0.8 GPa and 7.8 ± 1 GPa, respectively, as shown in Fig. 6. Li et al. [33] reported that the mechanical properties of coatings were significantly influenced by the internal microstructure of the thermal spray coatings. The difference in coating hardness is attributed to the fact that the particles in powder B had a higher kinetic energy, which leads to an easier fracture, firmer compacting, or stronger deformation of the particles during successive deposition. According to a previous study [17], the severely agglomerated powder can absorb part of the kinetic energy during VCS process. This is due to the cushion effect of agglomeration, which leads to the formation of a chalk-like film. In our study, the agglomerated powder A also reduced the velocity of particles. In other words, a stronger “tamping effect” occurred in the coating with powder B, which has strong particle/particle bonding and higher hardness.

Table 1 Deposition parameters for vacuum cold spraying. Parameter

Value

Unit

Chamber pressure He gas flow rate Distance from nozzle exit to substrate Nozzle traversal speed Nozzle orifice size Spray passes

80–180 3–6 10 10 2.5×0.2 1,10

Pa L min−1 mm mm s−1 mm×mm Times

from the load-depth penetration curve of a Berkovich indenter [30]. For each sample, at least 10 indentations were made on its polished cross section. 2.4. Numerical analysis of Al2O3 particle deposition during VCS The commercially available software Ansys-fluent 16.0 (Ansys Inc.) was used to examine the deposition of Al2O3 particles during VCS. In order to improve the accuracy of our results, a three-dimensional axissymmetric model was employed to simulate the gas flow and the Al2O3 particle velocity in the self-developed converging nozzle with an orifice size of 2.5×0.2 mm2. A schematic diagram of the nozzle in this study is shown in Fig. 2. The simulation domain was meshed to 930,516 grids to obtain an almost mesh-size-independent solution. For simplification, the Al2O3 particles were assumed spherical. In addition, only the dynamic drag force of the gas was considered in this study. The energy equation together with the K-ε turbulent model was used in the simulation, and discrete particle modeling (DPM) was employed to compute the acceleration of the Al2O3 particles.

3.2. Investigation of the deposition units of the two powder types \To reveal differences during deposition of the two different powder types, the deposition units for each powder were collected by scanning the substrate in one pass using the same parameters as coating deposition (Table 1). Fig. 7 shows the morphology of α-Al2O3 deposition units deposited by the powder A and powder B, respectively. It can be found that the deposition units formed from these two types of powder were significantly different. For powder A, particle clusters (marked with a white rectangle shown in Fig. 7a) of the original nano-crystal grains were found on the substrate. This indicates that the nano-sized grains in powder A are indeed agglomerates. As a result, a weak bond between the nano-crystal particles in porous agglomeration would remain in the VCS coatings. Furthermore, when the agglomerated particles hit the glass, coordinated deformation among the nano-grains consumes part of their kinetic energy. Hence, the shock energy cannot fully convert into bonding energy, and cohesion between the particles as well as the adhesion between coating and substrate would be weak. For powder B, which has 400 nm diameter particles, individual particles (not agglomerates) are the deposition units, as shown in Fig. 7b. The individual particles consist of many smaller fragments, which suggests that the particles break apart upon hitting the substrate surface. Some separated smaller particles (marked with open circles in Fig. 7b), which are smaller than 100 nm, can be seen on the substrate. This suggests that some small particles of powder B were also deposited. Moreover, as the white arrows in Fig. 7b indicate, craters were formed on the substrate surface due to the high impact force of the submicron Al2O3 particles during the deposition. Lee et al. [26] reported that interfacial bonding between VCS alumina coatings and their substrates relies on plastic deformation and fragmentation of the submicron particles. In our study, as can be seen from the deposition unit of powder B, plastic deformation and fragmentation occurred too, which facilitates strong bonding at the interface. The differences between these two powder types for the deposition are shown schematically in Fig. 8. Fig. 8a shows that loosely agglomerated particles, which consist of many nano-grains, hit the substrate. As shown in Fig. 8b, fragmentation and deformation occur after sintered submicron Al2O3 particles hit the substrate. Fig. 9 shows the result of our numerical simulation for helium gas and alumina particle velocity. Fig. 9a suggests that the gas velocity begins to increase at the throat of the nozzle, and then rapidly increases to about 1750 m/s due to the low chamber pressure. Subsequently, the gas velocity decreases to zero because of the blocking substrate. Fig. 9b shows the variation of particle velocity in the direction of the nozzle axis. We find that the particle velocity increases first before it decreases to several hundred meters per second, depending on the particle size. For particles of 400 nm diameter, the average impact velocity was about 430 m/s, which is higher than the critical velocity for alumina particles (~150 m/ s) [22].

3. Results and discussion 3.1. Characterization of the alumina coatings obtained from two types of powder The morphology of the two types of α-Al2O3 powder is shown in Fig. 3. Fig. 3a reveals that powder A has a loosely agglomerated morphology with a feature size of about 80 nm. This structure was mainly formed by nano-size grains. Fig. 3b shows that the morphology of powder B is facet and irregular with a diameter of about 400 nm. Fig. 4 shows the morphology of the surface and the cross-section of αAl2O3 coatings deposited on glass with these two types of powder. Both coating types show little roughness and strong interface bonding. Furthermore, there are no visible large pores and cracks due to the “tamping effect” caused by the continuous impacts of particles [31]. Apart from these similarities, more flattened particles appeared in the coating surface with powder B, as shown in Fig. 4b. This indicates that the impact force for powder B is larger than for powder A. The boundaries between the particles became ambiguous in both surface and cross section (Fig. 4b and d), suggesting that strong particle/ particle bonding occurred in the coating with powder B. Although the two types of powder were deposited under the same conditions, the deposition rate is different, as shown in Fig. 5. According to a previous study [32], ultrafine grains tend to form loose agglomerations before or during the VCS process, which is consistent with what we see here, as shown in Fig. 3a. The agglomerations are

Fig. 2. Schematic diagram of the nozzle.

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Fig. 3. Morphology of α-Al2O3 powder at a high magnification: (a) powder A, (b) powder B.

Fig. 4. Morphology of VCS-prepared α-Al2O3 coatings with powder A (a, c) and powder B (b, d). (a, b) surface morphology, (b, d) cross-section.

For the loosely agglomerated powder A, the deposition process can be described as follows: (i) the nano-grains maintain their original size; (ii) coatings are produced through the impact of successive agglomeration on the substrate i.e. successive accumulation of many coating layers. Some nano-sized pores appear due to the intrinsic porosity and insufficient kinetic energy of successive agglomeration. In a previous study [13], agglomerated TiO2 powder with a primary particle size of 25 nm was used to deposit coatings via VCS. The result was a porosity of about 43–49%, and pore diameters ranging from about several nanometers to one hundred nanometers. In other words, it is relatively difficult to obtain a dense coating using agglomerated powder, and it does not matter whether agglomeration is loose or hard. For the sintered submicron powder B, because of the high velocity, severe fragmentation, and deformation occurs when the sintered particles hit the substrate. In addition, a weak bond between fragments and nano-sized pores forms. Particles can also sometimes collide with previous deposits at high speed, which improves the cohesion between fragments. This also strengthens the adhesion between coating and substrate. Overall, a coating with a well-bonded coating/substrate

Fig. 5. Effects of the gas flow rate on the deposition rate of coatings deposited with powder A and powder B, respectively.

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different substrates. Powder B was characterized using TEM, as shown in Fig. 11. The bright-field TEM image (BF-TEM) shows that the particle diameters in powder B are about 400 nm, as shown in Fig. 11a. The phase structure of powder B was identified using selected area electron diffraction (SAD) analysis. The SAD pattern for the powder marked “A” in Fig. 11a is shown in Fig. 11b. The SAD pattern shows a single crystal α-Al2O3 structure with a clear hexagonal lattice (hcp), which is consistent with the X-ray diffraction results of alumina powder, as shown in Fig. 11c. Details of the microstructure for the interface between the Al2O3 VCS coating and the Al substrate using HR-TEM are shown in Fig. 12. The alumina coating bonded well with the underlying Al substrate according to the BF-TEM image in Fig. 12a and c. The SAD pattern of region “A” in Fig. 12a is shown in Fig. 12b. The phase structure of the coating is polycrystalline hexagonal α-Al2O3. It is the same as for the original powder, but the grain size was reduced to tens of nanometers, as shown in Fig. 12a. Alumina particles were embedded into the aluminum substrate – see the black arrows in Fig. 12a. This confirms that the penetration effect strengthens the bonding between coating and substrate. We also found that a weakly bonded interface between particles appears in the alumina coating close to the substrate – see the white arrows in Fig. 12a and c. The weakly bonded interface in the alumina coating is attributed to insufficient fracture and compaction of

interface and cohesive interface between particles is produced, layer by layer, through successive particle compaction. 3.3. Microstructural characterization of the VCS alumina coating of two metal substrates Based on the discussion above, powder B was chosen as the starting powder to deposit coatings on different metal substrates using VCS at room temperature. Aluminum (Al) and stainless steel (SUS 304) were chosen as substrates because of their different hardness values. Alumina coatings were deposited using the parameters shown in Table 1. The hardness of stainless steel (~200 GPa) is significantly higher than for aluminum (~72 GPa). The microstructure of the cross-section of the coatings deposited on the two metal substrates is shown in Fig. 10. Anchoring layers of different thicknesses form between the coating and different substrates. The thickness of the anchoring layers was about 700 nm and 200 nm for Al and SUS 304, respectively, see the white arrows in Fig. 10. The thickness of the anchoring layer decreased with increasing substrate hardness. Fig. 10a shows that the coating/substrate interface was relatively rough when the particles hit the substrate at a high speed. This is because alumina particles were embedded in the Al substrate with a lower hardness. In contrast, the SUS 304 with its higher hardness can better resist the impact of the high velocity particles, which results in a smoother interface, as shown in Fig. 10b. As reported previously [9], the formation of an anchoring layer is beneficial for coating/substrate bonding. In addition, a better coating/substrate bond was achieved with the lower-hardness substrate. 3.4. Effect of substrate hardness on the microstructure and interface bonding The coating/substrate interface, including its detailed morphology and interface bonding, was further investigated with HR-TEM for the

Fig. 8. Schematic diagram of the deposition behavior of the two powder types: (a) powder A, (b) powder B.

Fig. 6. Hardness of coatings deposited with powder A and powder B, respectively.

Fig. 7. Deposition units for the two types of powder: (a) powder A, (b) powder B. In (b), the open white rectangle and the open white circle correspond to the fractured Al2O3 particles and small Al2O3 particles, respectively.

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Fig. 9. (a) Gas velocity in axial direction of the computational domain; (b) The velocity of different sizes of Al2O3 particles in axis direction of the computational domain.

alumina particles fractured more severely when they hit SUS 304. The SAD pattern of region “A” in Fig. 13b confirms that the phase structure of this region is still polycrystalline hexagonal α-Al2O3. The HR-TEM image corresponding to region “(d)” in Fig. 13c is clearly visible in Fig. 13d. The FFT pattern “B” was inserted in Fig. 13d, which corresponds to region “B”. The diffused halo rings in the FFT pattern suggest that this region is essentially amorphous. In other words, the amorphous phase was formed locally in the region close to the coating/ SUS 304 interface. For metal and ceramic coatings, it was reported that the transition from crystalline to amorphous phase (C-A transition) occurs typically during rapid quenching, high pressure, or a high strain rate [34–36]. For example, an amorphous phase was observed in conventional plasma sprayed fully melted Al2O3 splats, because of fast cooling during splat quenching [34]. Paulo et al. [35] studied the phase transition of a Ni nanowire under high strain-rates by studying the molecular dynamics at room temperature. Their results show that the C-A transition also occurs when the strain-rate was above the critical value at room temperature. Chen et al. [36] observed the amorphization of ceramic materials for larger pressure and higher temperature. In a previous paper [37], localized amorphization of a cold-sprayed BN/NiCrAl nano-composite coating was induced by a high strain rate. This was confirmed by HR-TEM images and theoretical calculations. Park et al. [38] observed amorphous alumina in vacuum cold sprayed alumina coating if alumina particles are deposited at higher flow rates. Moreover, boundary amorphization of vacuum cold sprayed TiN coating was observed [39]. Therefore, it is necessary to identify the conditions that induce the C-A transition during VCS. Dun et al. [23] simulated the deposition behavior of alumina particles with a diameter of 300 nm that hit alumina substrates at 350 m/s during vacuum cold spray. The maximum impact pressure and temperature at the interface between particle and substrate were 5.7 GPa and 600 K, respectively. They

Fig. 10. Cross-section of the interface between alumina coatings and different substrates: (a) Al, (b) SUS304.

particles. The HR-TEM image shown in Fig. 12d, which corresponded to the region marked with “(d)” in Fig. 12c, show clear crystal fringes in the alumina coating. The fast Fourier transform (FFT) pattern (inset B in Fig. 12d) suggests the presence of a crystalline structure of hexagonal Al2O3 (hcp) at the coating/substrate interface. Fig. 13 shows typical cross-section TEM morphology of the Al2O3 coating on the SUS 304 substrate. It can be seen in Fig. 13a and c that the alumina coating bonds well with the SUS 304 substrate, and the interface is not very distinct. The grain size in the alumina coating is much smaller than that in the alumina coating on the Al substrate (Fig. 12a). In addition, the obtained coating was so dense that the particle interface cannot be seen clearly, which confirms that the

Fig. 11. TEM images of the starting powder Al2O3: (a) BF-TEM image; (b) SAD pattern corresponding to region “A” in (a). (c) X-ray diffraction pattern of the starting powder Al2O3.

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Fig. 12. TEM images of the Al2O3 VCS coating on the Al substrate: (a) low magnification of the interface between Al2O3 coating and substrate, (b) SAD pattern of region “A” in (a); (c) high magnification of the interface between Al2O3 coating and substrate; (d) HR-TEM image corresponding to region “(d)” in (c). (Inset B is a FFT image corresponding to region “B”). The white dashed lines indicate the boundaries between Al and the Al2O3 coating.

substrate. Well-bonded particle interfaces appear close to the interface between coating and substrate in the harder SUS 304 substrate. In other words, both the microstructure and the phase structure at the coating/substrate interface are significantly different for different substrate hardness values. Our results confirm that the mechanical properties of the substrate play an important role for both cohesive and adhesive bonding during vacuum cold spraying at room temperature.

occurred about 1.6 ns after particle impact. The maximum temperature was significantly lower than the melting point of Al2O3 (~2303 K). Therefore, the possibility of rapid quenching induced amorphization of the alumina coating could be eliminated in this study. According to the simulation results shown in Fig. 9, the impact velocity of particles for powder B is 430 m/s. As a result, the impact pressure could be greater than 5.7 GPa according to the paper by Dun et al. [23]. Based on the analysis above, we conclude that the observed localized amorphization of alumina at the interface between coating and substrate is primarily due to the high impact pressure combined with the high strain rate due to short impact time during vacuum cold spraying. According to the previous study [40], the kinetic energy of the alumina particles (before impact) converts to both particle fracture energy and bonding energy (particle/particle or particle/substrate). In this study, the difference in kinetic energies between the alumina particles for Al and SUS 304 substrates should be negligible because we used identical deposition parameters. The deformation of the Al substrate consumes kinetic energy and thus decreases the impact pressure, which results in a weakly-bonded interface for the coating (Fig. 12a). In addition, an anchoring layer with a thickness of ~700 nm was formed at the interface between coating and substrate, as shown in Fig. 10, which in turn strengthens interfacial bonding. The impact pressure was higher for SUS 304, which has a higher hardness, compared to the Al substrate. This led to the formation of an anchoring layer of only about 200 nm (Fig. 10b) and the localized amorphization of the alumina coating close to the interface between coating and

4. Conclusions Differences during the deposition of VCS alumina coatings were investigated using structurally different powders and substrates with different hardness values. The loosely agglomerated nano-sized alumina powder was deposited at high speed with the deposition unit of the agglomerates containing nano-size pores. They then bonded to the substrate through coordinated deformation of the original nano-grains. The sintered submicron powder, on the other hand, was deposited with deposition units being individual particles. These single particles underwent serious fracture upon impact on the substrate. The microhardness of the coating with the sintered powder was two times that of the coating with the loosely agglomerated powder. Better particle/ particle bonding was observed for the sintered submicron powder. For the softer Al substrate, a thicker anchoring layer, facilitating coating/ substrate interface bonding, was created at the Al2O3/Al interface. More severe fracturing of alumina particles occurs when the particles hit SUS 304 because of the higher impact pressure. In conclusion, better coating/substrate adhesive bonding can be achieved in softer 7

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Fig. 13. TEM images of the Al2O3 VCS coating on the SUS304 substrate: (a) BF-TEM images of the interface between Al2O3 coating and substrate; (b) SAD pattern corresponding to region “A” in (a); (c) high magnification of the interface between Al2O3 coating and substrate; (d) HR-TEM images corresponding to region “(d)” in (c). Inset B is a FFT image showing the presence of the amorphous phase. The white dashed lines indicate the boundaries between SUS304 and the Al2O3 coating. [7] A. Feuerstein, J. Knapp, T. Taylor, A. Ashary, A. Bolcavage, N. Hitchman, Technical and economical aspects of current thermal barrier coating systems for gas turbine engines by thermal spray and EB-PVD: a review, J. Therm. Spray Technol. 17 (2) (2008) 199–213. [8] J. Akedo, M. Lebedev, A. Iwata, H. Ogiso, S. Nakano, Aerosol deposition method (ADM) for nano-crystal ceramics coating without firing, Mat. Res. Soc. Symp. Proc. 778 (2003) 289–294. [9] J. Akedo, M. Lebedev, Microstructure and electrical properties of lead zirconate titanate (Pb(Zr0.52/Ti0.48)O3) thick films deposited by aerosol deposition method (9B)Jpn. J. Appl. Phys. 1 (38) (1999) 5397–5401. [10] J. Akedo, Aerosol deposition of ceramic thick films at room temperature: densification mechanism of ceramic layers, J. Am. Ceram. Soc. 89 (6) (2006) 1834–1839. [11] Y.Y. Wang, Y. Liu, C.J. Li, G.J. Yang, J.J. Feng, K. Kusumoto, Investigation on the electrical properties of vacuum cold sprayed SiC-MoSi2 coatings at elevated temperatures, J. Therm. Spray Technol. 20 (4) (2011) 892–897. [12] X.L. He, G.J. Yang, C.J. Li, C.X. Li, S.Q. Fan, Room temperature cold sprayed TiO2 scattering layer for high performance and bending resistant plastic-based dyesensitized solar cells, J. Power Sources 251 (2014) 122–129. [13] S.Q. Fan, C.J. Li, G.J. Yang, L.Z. Zhang, J.C. Gao, Y.X. Xi, Fabrication of nano-TiO2 coating for dye-sensitized solar cell by vacuum cold spraying at room temperature, J. Therm. Spray Technol. 16 (5–6) (2007) 893–897. [14] S.Q. Fan, C.J. Li, C.X. Li, G.J. Liu, G.J. Yang, L.Z. Zhang, Preliminary study of performance of dye-sensitized solar cell of nano-TiO2 coating deposited by vacuum cold spraying, Mater. Trans. 47 (7) (2006) 1703–1709. [15] J.J. Choi, J.H. Choi, J. Ryu, B.D. Hahn, J.W. Kim, C.W. Ahn, W.H. Yoon, D.S. Park, Low-temperature fabrication of nano-structured porous (La,Sr)(Co,Fe)O3-δ cathodes by aerosol deposition, J. Alloy Compd. 545 (2012) 186–189. [16] J. Akedo, M. Lebedev, Powder preparation in aerosol deposition method for lead zirconate titanate thick films, Jpn. J. Appl. Phys. 41 (1) (2002) 6980–6998 (11B). [17] S.M. Nam, N. Mori, H. Kakemoto, S. Wada, J. Akedo, T. Tsurumi, Alumina thick films as integral substrates using aerosol deposition method, Jpn. J. Appl. Phys. 43 (1) (2004) 5414–5418 (8A). [18] S.F. Wang, Y.F. Hsu, C.H. Wang, C.T. Yeh, Solid oxide fuel cells with Sm0.2Ce0.8O3-δ electrolyte film deposited by novel aerosol deposition method, J. Power Sources

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