E L S E V I E R Thin Solid Films 280 (1996) 188-198

Microstructure of plasma-sprayed Ni-A1 alloy coating on mild steel

H.C. Chen a,b, E. Pfender b,. "The Second Department of Mechanical Engineering, South China University of Technology, Guangzhou 510641, People's Republic of China

b ERCfor Plasma-Aided Manufacturing, Department of Mechanical Engineering~ University of Minnesota, Minneapohs, MN 55455, USA

Received 13 July 1995; accepted 18 October 1995

Abstract

Microstructure and fracture surface morphologies were characterized for a cross-sectional plasma-sprayed Ni-A1 alloy coating-mild steel substrate system by scanning electron microscopy, electronic probe microanalysis (EMPA) and transmission electron microscopy. The plasma-sprayed Ni-AI coating-steel substrate interface was found to be of pure metallurgical nature. The bonding layer consisting of Fe3AI closely relates to exothermic reactions from Ni and AI during spraying. Ni3A1, the product from the Ni-A1 reaction, and Fe3A1 kept a fixed crystallographic orientation relationship. From the coating-substrate interface toward the coating surface, the fracture mode changes from preferential interlamellar fracture, to intersplat cleavage fracture, and finally te quasi-cleavage fracture. The changes are thought to be related to the contact condition between splats. Microstructural changes observed through the coating thickness seem to be caused by the non-uniform cooling rate distribution.

Keywords: Coatings; Interfaces; Plasma processing and deposition; Transmission electron microscopy

1. Introduction

Plasma-spraying of Ni-A! alloy coatings is widely used as bond coats under ceramic or metal coatings. This improves both adhesion of the ceramic overlay with the substrate, and protects the substrate from oxidation by air or other oxidizing atmospheres, diffusing through the pores of the ceramic over- lay. Plasma-sprayed Ni-AI alloy coatings are also used as working coatings because of their high bond strength with substrates, good hot corrosion resistance and low oxide inclu- sions. This kind of coatings are very suitable for restoring dimensions of worn parts due to their excellent machinability and they can be used at elevated temperatures ranging up to 800 °C. Ni-AI alloy spray powder has been proven to be economic and of interest for studying in-flight trajectories of particles and plasma-particle interactions [ 1-3].

Over the past years, many studies have focused on plasma spraying parameter optimization and plasma-particle inter- actions. Recently, the solidification process of splats on impact with the substrate and the coating microstructure, especially the coating-substrate interface, characterized by transmission electron microscopy (TEM) and scanning elec- tron microscopy (SEM), have been receiving more attention, because they can provide important information useful for

* Corresponding author.

0040-6090/96/$15.00 © 1996 Elsevier Science S.A. All rights reserved SSDi 0040-6090 ( 95 ) 0 8195- X

modifying the plasma spray parameters and for clarifying the coating-substrate interfacial bonding. Unfortunately, earlier SEM and TEM observations were mostly based on planar sections of coatings, parallel to the substrate surface. They could neither identify grain orientation nor characterize changes in the coating microstructure through the coating cross-section. This may have been due to difficulties in pre- paring cross-sectional samples for TEM and SEM, especially samples containing the coating-substrate interface.

In this paper, the authors present a systematic characteri- zation of the coating microstructure, identification of the coat- ing-substrate interface, and an explanation of the mechanisms of reactions occurring at this interface.

2. Experimental methods

The commercial 95/5 (Ni/Al:95/5 wt.%) Ni-AI alloy powder used in the experiments is spherical in shape with a size distribution from 50 to 100 lxm, a mean diameter of 85 Ixm and a standard deviation of 12 p,m [ 2]. A cross-sectional view of the particles is shown in Fig. 1, indicating that the particles consist of a Ni core and a A1 cladding on the exterior of the powder particles as they are received from the manu- facturer. The cladding prepared in this way produces an exo- thermic reaction of the powder, thereby enhancing the

H.C Chen, E. Pfender /Thin Solid Films 280 (1996) 188-198 189

: ~ ~ . . ' • . ~ o . - .

Fig. 1. Cross-sectional view of the spray particles.

bonding between the coating and the substrate. This makes it possible to successfully prepare coating-substrate interfacial toils for transmission electron microscopy (TEM) studies. Before plasma spraying, 55 × 30 × 2 mm mild steel substrates were ultrasonically cleaned and then grit blasted with 60 mesh ( 100 i~m) A1203 under pressures around 0.4 MPa for approx- imately 30 s to produce a surface roughness of about 1.4 p.m. The substrate was attached to a rotating 305 mm diameter disk and repeatedly traversed through the plasma jet with a velocity of 0.56 m s- 1. Coatings were formed using a Miller SG- 100 plasma spray torch for a 2 min spraying time without substrate cooling or preheating. The Miller SG-100 plasma spray torch used in this experiment was operated in the sub- sonic mode (anode 2083-175, cathode 1083A-129, gas injec- tor 2083-113). The arc gas injector is located at the base of the cathode with four holes (i.d., 1,98 mm) drilled tangen- tially into the gas injector walls. The powder injector, located about 10 mm before the torch exit, is straight with respect to the torch axis. The operating conditions for this experiment are summarized in Table 1.

Specimens for metallographic examination were cut per- pendicular to tke coating surface. Composition profiles and X-ray maps were analyzed for a cross-sectional interfacial region using a computerized JEOL JXA-8900 Superprobe.

The structure of the cross-sectional coating was analyzed by SEM after controlled fracture of the coating-substrate

system. A slot 1 mm wide cut into the back face of the substrate opposite to the coating allows fracture of the sample under a controlled bend lo~d. Such a fracture process exposes a large area of the coating-substrate interfacial region which has been characterized by SEM (JEOL 840-//). This method can provide more information about the coating-substrate interface compared with the controlled fracture of a coating removed from the substrate.

Techniques for preparation of cross-sectional sample~ of the coating-substrate interface are similar to those described in Refs. [ 4,5 ]. Two thin slides 1 mm thick were first sectioned along the direction normal to the coating surface and put together face to face with epoxy, and then thinned mechani- cally and polished to 30--50 p~m. Finally the sample thickness has been further reduced with a ion-thinning machine. The samples containing the coating-substrate interface were examined with the JEOL 1210 transmission e;ectron microscope.

3. Results and analys~s

3.1. Composition distributions of eiements over a cross- section of the coating

Composition distributions were analyzed on a polished cross-section of the coating by using a computerized micro- probe analyzer. Fig. 2 shows X-ray maps and ccmposition profiles of AI, Ni and Fe of the same area of the coating- subslxate interfacial region. Distribution of AI in the coating is non-uniform, discontinuously piling up at the interfaces between iamellae, especially at the coating-substrate inter- face. Such a pile-up of AI is due to preferential melting of the low-melting-point A! layer on the particle surface during spraying and penetration into the pores of the regions men- tioned above. Inter-diffusion between Ni, AI and Fe took place in the coating-substrate interface. The diffusion depth of Ni and Ai from the coating to the substrate is 1.5-2 I~m, while Fe diffusion from the substrate to the coating, is up to 6 I~m.

3.2. Fractography

Table 1 Operating conditions

Parameters Settings

Arc current 800 A Primary gas flo~, rate (Ar) 45 slm Secondary gas flow rate (He) 16 slm Carrier gas flow rate (At) 4.5 sim Powder feed rate 39 g min- i Standoff distance 90 mm Substrate velocity 0.56 m s- 1 Spray time 2 rain Torch nozzle i.d. 8 mm Powder injector tube i.d. 2 mm

Detailed information about the microstructure can be obtained from observation of a controlled fracture surface of the coating-substrate system. Fig. 3 shows a typical fracture surface near the coating-substrate interfacial region. A layer of island-like structures, 5-15 Ixm thick, contacts with the coating-substrate interface. Away from this island-like layer a columnar structure layer is observed. Higher magnification (Fig. 3(b)) shows that each island-like region actually con- sists of many fine grains of different sizes. Cracks occurred at the boundaries between island-like regions. It is interesting to note that a bonding layer (annotated as the interfacial region in Fig. 3 (b)) exists between the coating and the sub- strate. In some local interfacial regions, there is no island-

190 H.C Chen, E. Pfender l Thin Solid Films 280 (1996) 188-198

Fig. 2. X-ray maps and composition profiles of elements in the coating-substrate interracial region, showing inter-diffusion between Ni, AI and Fe, and AI pile- up in the interfaces between both lamellae and the coating and substrate.

like structure layer and the columnar structure layer directly contacts with the coating-substrate interface, as shown in Fig. 3(c). Moreover, cracks, causing coating spallation, did not propagate along the coating-substrate interface line, but propagated along the boundary between the island-like layer and the columnar layer or along the interfaces between lamel- lae (see Fig. 3(a) and Fig. 4). Fracture in this area is pref- erentially interlameUar fracture (see Fig. 4).

Fig. 5 shows a typical fracture surface in the center of the coating. The fracture surface morphology is much different from that near the coating-substrate interface. The character- istics of the preferential interlamellar fracture is lost. It is replaced by so-called intersplat cleavage fracture, with the cracks propagating along boundaries between splats (Fig. 5(a) and 5(b)). Although the grain structure within a splat still remains columnar, the columnar direction is no longer simply normal to the lamellar interface. The structure is com- posed of columnar grains with different growth directions, as shown in Fig. 5(c). In the boundary layer between splats, equiaxed fine grains with different sizes can be observed (see Fig. 5(d)) . Only in a small area where the fracture path follows preferentially the lines of the lamellar boundaries ( shown by the arrow in Fig. 5 (a)), the columnar grain grow- ing direction remains normal to the lamellar interface (Fig. 5(e) ). The change in fracture mode shows improvement of interlamellar adhesion, as demonstrated by interlamellar

bridges existing in the fracture surface, as shown in Fig. 6. In the following section, the interlamellar bridges are further characterized by TEM.

In the area close to the coating surface, the fracture mode changed again, with a quasi-cleavage fracture being observed, as shown in Fig. 7. The fracture surface shows that the coating underwent some degree of deformation before fracture. This must be due to further improvement of inter- lamellar and intersplat adhesion.

3.3. TEM observations on the coating-substrate interface

The samples prepared in this study allow one to directly observe the coating-substrate interfacial region by TEM. A typical TEM of the morphology of the coating-substrate interfacial region is seen in Fig. 8(a). Two layers exist between the coating and the substrate. One is close to the substrate and is identified by TEM to be face-centered cubic FeaAI, as shown in Fig. 8(a). Its corresponding dark-field image is shown in Fig. 8(b). This layer is generally 0.1-0.3 I~m thick, but locally up to 0.8 I~m representing the bonding layer connecting the coating and the substrate. Another layer close to the coating is about 0.2 ~m thick, identified by TEM as tetragonal NiaAI, which is the product from a Ni-AI reac- tion during spraying. Fe3AI and Ni3AI have a fixed crystal- lographic orientation relationship:

H. C. Chen, E. Pfender / Thin 6olid Films 280 (1996) 188-198 191

Fig. 3. A typical fracture surface near the coating-substrate interface (SEM): (a) island-like structure layer, directly in contact with the coating-substrate interface, and co'lumnar structure layer; (b) higher magnification image of the island-like structure, consisting of fine grains changing in size, and close-up of the coating-subs.trate interfacial region; (c) columnar structure layer in contact directly with the coating-substrate interface (no island-like structure layer).

(110) Ni3Ai / i (100 ) Fe3AI

[ 111 ] Ni3AI/J / [011 ] Fe3AI

i.e. the bonding layer (Fe3AI) forms in close relation with NiaAI, whica is further proven by Fig. 9. The bonding layer in this a~rea is very thin (25 nm) and discontinuous, because there is no NiaAI. Moreover, Fig. 9 also shows the transition area of Fe:~AI growing into the coating.

3.4. TEM microstructural characterization throughout the coating thickness

TEM shows that plasma-sprayed Ni-AI alloy coatings are composed of Ni, AI, intermetallic compounds such as Ni3AI and NiAI, a~nd oxides. The coating is built up layer by layer in the form of lamellae. The thickness of the lamellae changes from place to place and from lamella to lamella. In some areas, the thiickness of the lameUae is from 2.5 to 5.5 p.m, but in other areas, only 0.8-1.0 p~m (Fig. 10(a)). Interlamellar and intersplat porosity (shown by arrows) can be observed, which is l~lieved to arise from an intermittent contact

between lamellae and between splats. Fig. 10(b) shows a region where two individual lamellae have different thick- ness, 1.5 and 0.6 lxm, respectively. The larneUae are generally thicker at their center and thinner at their periphery, as shc, wn in Fig. 10(c).

The microstructural morphologies of the coating changefi not only from location to location and from lamella to lamella, but also from one region to another within a lamella. The microstructure in the interlamellar or intersplat boundary regions is similar to that near the coating-substrate interface. It consists of fine- or micro-crystals and amorphous phases, as shown in Fig. 9 and Fig. 11 (a). The amorphous film at the interlamellar or at the intersplat boundary regions is very thin and preferentially etched away during ion-beam thinning or during TEM observation, leaving gaps (Fig. 11 (a)) . At the periphery of the splat, the microstructure often consists of fine grains and thin plate-like crystals (shown by arrow in Fig. 11 (b)) . Within a lameUa, the microstructure consists of columnar grains, as shown in Fig. 12(a), the columnar grains are 0.3--0.5 ~m thick, and approximately 1.5 ~m long. How- ever, the columnar grains are not always parallel to each other. Some of them nucleate on both sides of the splat and grow

192 H, C Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198

trast in the lath boundaries can be observed, which was caused by deformation when the particle impacted with the substrate, as shown in Fig. 14(c).

Fig. 4. SEM fracture surface near the coating-substrate interface, showing the preferential interlamellar fracture and cracks propagating along the boundary between island-like structure layer and columnar structure layer and along interlamellar boundaries.

into the center where they intersect. Especially in the area from the center to the surface of the coating, the columnar structure is not entirely composed of one-directional growing grains, but it is a combination of multi-directional growing grains (see Fig. 12(b) ).

It is, in-general, thought that the columnar grains do not grow across the lameUar interfaces. But in this study, TEM shows some interlamellar bridges in some regions, especially far away from the coating-substrate interface. Fig. 13 shows such a film, 0.2 Ixm thick, in which the fine- or micro-crystals on the amorphous matrix nucleate at both sides of the lamellae and grow into the center of the film finally intersecting to form an inter-growing zone and thus a bridge which connects two lamellae. This is consistent with SEM observations of fracture surfaces.

For partially melted particles, the melted portion on the particle surface will undergo a spread and solidify upon impact with the substrate, forming a columnar structure (Fig. 14(a)), while the unmelted core (G region in Fig. 14(a)) is embedded into the coating, displaying a lath- shaped structure (Fig. 14(b)). Twin boundary fringe con-

4. D i s c u s s i o n

4.1. Formation mechanism of the bonding layer between the coating and the substrate.

TEM shows that the bonding layer, i.e. Fe3AI, is often accompanied by a layer ofNi3Al. This bonding layer becomes very thin and discontinuous in the absence of Ni3AI. There- fore, the formation of the bonding layer depends to a large degree on Ni3AI, i.e. on the reaction between Ni and Ai.

During plasma spraying of Ni-AI alloy powder, particles are injected into the plasma and heated. The AI cladding on the Ni core will first melt as the surface temperature of the particle reaches the aluminum melting point of 993 K. As more of the AI cladding melts the A! layer thickens when the distance to the substrate decreases. Ito et al. demonstrated by differential thermal analysis (DTA) that the reaction between AI and Ni begins at temperatures near the aluminum melting point [ 6]. In the case of plasma-spraying of Ni-AI alloy, the average measured particle temperatures are in excess of 1900 K [ 3 ], which implies that the Ni-AI reactions on the droplet surface take place during flight of the particles from the torch exit to the steel substrate. The reactions continue up to the moment of droplet impact with the substrate. According to thermodynamic calculations, the Ni-A! reactions forming intermetallic compounds are all exothermic, and there are two kinds of reactions based on TEM phase identification:

Ni + AI ~ NiAI - AHt ( 1 )

3Ni + AI --* Ni3Al - AH2 (2)

where AH~ and AH2 represent the heat released from the reactions, which depend on the reaction temperature. At 1600 K, for example, AHI and AH2 are 139.63 and 176.05 kJ mol- ~ respectively. Therefore, when the droplets impact on the steel substrate, the latter will receive heat by conduction from the molten droplet, and more importantly heat from the exothermic Ni-AI reactions.This heat rapidly raises the local substrate surface temperature to levels which speed up the diffusion of elements such as Ni, AI and Fe in the droplet- substrate interface, and thus form an inter-diffusion zone between Ni, AI and Fe (Fig. 1). Since the affinity between Fe and AI is greater than that between Ni and AI, Fe diffusing from the substrate will react with AI in Ni3AI and form Fe3AI, resulting in a metallurgical bonding between the coating and the substrate. In the absence of exothermic reactions, the substrate would receive less heat from the droplet, resulting in a very thin and discontinuous bonding layer, because dif- fusion of Fe is less promoted in this case.

The fact that AI is mainly present in the coating-substrate interfacial region and interiamellar boundary regions is due

H. C. Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198 193

Fig. 5. SEM fracture surface in the center of the coating: (a) showing intersplat cleavage fracture, with a small area where the fracture path follows the intedamellar boundaries, manifesting the columnar structure (shown by arrow); (b) higher magnification image of cracks propagating along the intersplat boundaries; (c) columnar grains with different growth directions within a lamella; (d) equiaxed fine grains in the intersplat boundary layers; (e) columnar grains with the growth direction normal to the interlamellar interface.

to AI melting on the droplet surface and its penetration into pores of the these regions during droplet impact with the substrate. It should be pointed out that the penetration of the melted AI will cause, to some extent, a liquid-sintering effect between lamellae, which improves interlamellar adhesion.

4.2. Changes in fracture mode of the coating

The SEM photographs show that the fracture mode of the coating changes from the preferential interlamellar fracture to intersplat cleavage and finally to quasi-cleavage fracture

194 H.C. Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198

Fig. 6. gEM fracture surface showing bridges connecting two lamellae.

Fig. 7. $EM fracture surface close to the coating surface, showing a quasi- cleavage fracture.

when going from the coating-surface interface toward the coating surface. This is attributed to changes in the degree of contact of droplets with the substrate or previously deposited layers, resulting in changes in interlamellar and intersplat adhesion.

As a liquid droplet impinges on the substrate or on the previously deposited layers, the liquid droplet spreads out on the surface, cools rapidly and solidifies to form a lamella. It is also possible that a liquid droplet fills a void on the substrate or on the previously deposited layers as it impinges. Thus, mechanical locking may take place not only at the interface between the bottom of the coating and the grit blasted sub- strate but also at every interface between the lamellae. The degree of droplet spreading and the void-filling during impact depends on the contact condition of droplets with the sub- strate or previously deposited layers, and also on droplet temperature, particle size, velocity, viscosity, wetting behav- ior, and rate and direction of heat removal. As soon as the droplet temperature drops below the melting point of the deposited material, the solidification begins, preventing fur- ther spread of the droplet and a lameila will be formed.

When the molten droplet impinges on the cold substrate, it may explode. In spite of the relatively poor thermal contact between the splat and the cold substrate, the large temperature difference between them leads to a rapid cooling and solidi- fication of the splat. Therefore, at the area in the coating- substrate interface, the spreading process of each droplet and void-filling process of the droplets is diminished, resulting in decrease of mechanical locking effect between lamellae. Thus the poor adhesion between lamellae is predictable, which is responsible for the preferential interlamellar fracture. With increasing coating thickness, the substrate temperature will rise, and the contact conditions of the droplet with the pre- viously deposited splats will be improved. Therefore, spread- ing and void-filling processes of the droplets become more likely, which means mechanical locking between lamellae is enhanced, and thus interlamellar adhesion will be improved, as observed by changes of the fracture mode of the coating. Because of the better contact of droplets with splats, inter- growing zones or bridges are observed between lamellae, as revealed by SEM and TEM micrographs. Improvement of interlamellar adhesion is responsible for the change of frac- ture mode of the coating from preferential interlamellar frac- ture to the intersplat cleavage fracture in the coating center. Toward the coating surface, the fracture mode changes from the intersplat cleavage fracture to the quasi-cleavage fracture because the better contact of droplets with splats, resulting in further enhancement of the spreading and the void-filling processes of the droplet and further improvement of interla- mellar and intersplat adhesion.

4.3. Microstructural changes through the coating thickness

The microstructure of the lamella depends mainly on the cooling rate of the splat, and thus on the contact conditionc, between splats. It should be pointed out that the degree of superheat, the solidification temperature and temperature range as well as the initial temperature of the splat will influ- ence the resulting cooling rate [7 ].

In spite of the relatively poor thermal contact between the splat and the cold substrate, the large temperature difference between them leads to a rapid cooling and solidification of the splat. As a consequence, nuclei form homogeneously at extremely high rates and grow forward towards the liquid, forming a layer of micro-crystals. Because the undercooling condition in front of the substrate is non-uniform, the under- cooled liquid conditions are limited to small local regions, resulting in island-like structures (Fig. 2(a) ). As the crystals grow, the temperature of the liquid ahead of the crystals may even increase due to release of the heat of fusion, thus sup- pressing nucleation and growth of the crystals. As a result, columnar grain growth will continue along the direction of highest heat removal which is usually normal to the substrate surface. As the coating thickness increases, the substrate tem- perature will rise, and the better contact of splats with the previously deposited layers wiU lead to rapid heat removal from splats in a direction perpendicular to the contact inter-

H. C. Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198 195

t ; Ni3AI [ ] FeaAI (0

Fig. 8. A typical TEM morphology of the coating-substrate inteffacial region and corresponding electron diffraction patterns. (a) Inteffacial region consisting of two layers, the bonding layer, FeaAI, and the NiaA! layer; (b) dark-field image of the bonding layer (layer Fe3AI), showing uneven thicl~.ess; (c) selected area electron diffraction pattern (SAEDP) of the face-centered cubic (f.c.c.) 7-Fe suhstrate along [ 112] direction; (d) SAEDP of f.c.c. Ni along [011] direction in the coating; (e) SAEDP of layer Fe3AI and layer Ni3AI; (f) schematic representative indexing of SAEDP in (e).

Fig. 9. TEM micrographs showing a very thin and discontinuou~ bonding layer between the coating and the substrate. (a) Transition area of Fe3AI growing into the coating; (b) higher magnification image showing very thin and discontinuous layer of Fe3Al.

196 H. C Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198

Fig. 10. TEM cross-section showing (a) a series of lamellae, (b) two different thicknesses of lamellae in local regions, and (c) uneven thickness of lamella.

~,i ~i~iii

Fig. 11. 'gEM micrographs showing microstmctures of (a) the interlamellar and intersplat boundary layer, (b) the periphery of the splat.

iiiiiii~

~i!~:~ ,̧ ~!!!~i ~ . . . . . . . . ~̧ ~̧

Fig. 12. TEM micrographs showing (a) the columnar structure within a lameUa, noting that the columnar grains nucleated at both sides of the splat, (b) the columnar grains with different growth directions in the center of the coating.

H.C. Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198 197

face between splats, resulting in columnar crystal growth. In the case of rough substrates, columnar grains will grow in different directions, as shown in Fig. 5(c) and Fig. 12(b).

At the periphery of a splat, undercooling conditions of the liquid are also expected because of the poor thermal contact. Therefore equiaxed fine grains arising from homogeneous nucleation are observed, while in the contact area between splats, the solidification process is similar to that of a thin layer near the coating-substrate interface.

5. Conclusions

1. Plasma-sprayed Ni-A1 alloy coatings cannot be separated mechanically from the steel substrate

2. The bonding between the coating and the substrate is purely metallurgic. The bonding layer, generally 0.1-0.3 pom thick, consists of Fe3Al, formed by interfacial reac- tions between the coating and the substrate. Its thickness and structure depends on exothermic reactions between Ni and Al during spraying. Without these exothermic reac- tions, the bonding layer becomes very thin and discontinuous.

3. The bonding layer and the layer produced by Ni-Al exo- thermic reactions keep a fixed crystallographic orientation relationship:

Fig. 13. TEM micrograph showing a thin film consisting of amorphous phase and micro- or fine-crystals, noting bridge (shown by arrow) formed as a result of inter-growth of the crystals from two sides of the lamellae.

( 110)NiaAI//(100) F¢3AI

[ 111 ] Ni3kl / / [ 011 ] F~3AI

4. From the coating-substrate interfitce toward the coating surface, the fracture mode changes from preferential inter- lamellar fracture to intersplat cleavage fracture and finally to quasi-cleavage fracture. This is attributed to the change in the contact conditions between splats.

5. The microstructure of the coating changes from location to location, and from lamella to lamella and even from region to region within a lamella, which is attributed to the non-uniform cooling rates throughout the coating.

°°

Fig. 14. TEM micrographs showing a semi-melted particle. (a) The liquid on the surface of the particle solidified as the columnar structure; (b) unmelted core developed into lath-shaped morphology; (c) twin boundary fringe contrast appeared in the lath boundaries, resulting from, deformation of the particle when impacting with the substrate.

198 H.C Chen, E. Pfender / Thin Solid Films 280 (1996) 188-198

Acknowledgements

The authors thank Mr. K. Leung for help in preparing plasma-sprayed coatings. This work has been supported by NSF ECD-87-21545.

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