A study of high velocity oxy-fuel thermally sprayed tungsten carbide based coatings. Part 1: Microstructures

Materials Science and Engineering A246 (1998) 11 – 24

A study of high velocity oxy-fuel thermally sprayed tungsten carbide based coatings. Part 1: Microstructures C. Verdon, A. Karimi, J.-L. Martin * De´partement de Physique, Institut de Genie Atomique, Ecole Polytechnique Fe´de´rale de Lausanne, PH-B Ecublens, 1015 Lausanne, Switzerland Received 15 September 1997; received in revised form 6 November 1997

Abstract The microstructures of two tungsten carbide–cobalt (WC – Co) coatings, deposited using high velocity oxy-fuel (HVOF) thermal spraying method in different conditions, are studied. They are compared with that of the WC – Co powder grains injected in the flame, in an attempt to understand the transformations that occur during deposition. For this purpose, various imaging and analytical techniques in electron microscopy are used, in addition to global characterization methods such as X-ray diffraction and fluorescence. These methods reveal that the coatings are made of distinct islands, elongated along the substrate direction, which exhibit a nano-crystalline matrix containing tungsten, cobalt and carbon. The fraction of WC grains in the coating is smaller than that in the powder and fluctuates throughout the coating. A net loss in carbon is evidenced in the coatings as compared to the powder grains. New phases, W2C and W, appear in specific locations in the microstructure in relation with the local composition of the matrix. Very little metallic cobalt is retained. The extent of the transformation is related to the spraying conditions. Some processes that account for the change in microstructure and composition during spraying are proposed. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Tungsten carbide; Coatings; Thermal spraying; Microstructure; TEM

1. Introduction Tungsten carbide – cobalt (WC – Co) alloys are used in a wide range of wear resistant applications in a variety of conditions which include sliding, fretting, abrasion and erosion. The erosion resistance of such composites is well-studied both on sintered bulk specimens [1,2] and on coatings obtained by different methods [3,4]. It has been shown that for thermally sprayed coatings, decarburization occurs during the spraying process that affects the microstructure with a concomitant decrease in the hardness and the wear resistance of the materials. This phenomenon depends on deposition conditions which are characterized by high temperature, a generally oxidizing atmosphere and high cooling rates, so that the sprayed material is subjected to complex physical and chemical transformations. In particular, several studies based on global characterization methods have revealed that during the spraying of * Corresponding author. Tel.: +41 21 6933371; fax: + 41 21 6934470; e-mail: [email protected] 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(97)00759-4

WC–Co, a part of WC is transformed into WC1 − x, W2C and metallic tungsten [5], while another part is dissolved in the cobalt matrix to form M6C or M12C carbides [6]. More recently, a study based on X-ray diffraction (XRD), differential thermal analysis (DTA) and transmission electron microscopy (TEM), claims the formation of a amorphous compound [7]. The extent of these transformations depends on the spraying method, the operating parameters [8] and the powder characteristics [9]. New processes, such as high velocity oxy-fuel (HVOF) thermal spraying, have been designed to retain a larger fraction of WC in the coating [10]. In this method, the hypersonic velocity of the flame shortens the time of interaction between the powder and the flame, and in conjunction with the relatively low temperature in the latter, limits WC decomposition. In addition, the high kinetic energy acquired by the powder particles ensures a good cohesion in the coatings and allows for the production of materials free of any significant porosity. However, the transformations mentioned above still occur [11], but to a limited extent.

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Several studies [6,12,13] have attempted to relate the spraying parameters and the powder characteristics to the microstructural features and the wear resistance of the coatings. Such an approach aims at optimizing the spraying conditions to obtain an increased wear resistance. Nevertheless, the specific role played in the wear mechanisms by each component of the microstructure is not clear. Moreover the coating features were mostly studied by macroscopic means such as XRD and chemical analysis or scanning electron microscopy (SEM). These are straightforward methods, but do not provide enough information about how and where phase transformations take place during the deposition process. The characterization, at a fine scale, of the microstructure of thermal sprayed WC – Co, as well as the localization of the phases formed during spraying, are unavoidable steps in understanding how wear takes place in such complex microstructures. The present investigation is focused on the microscopic description and analysis of the microstructural components in HVOF sprayed WC – Co. A comparison between the microstructures in the powder grains and the coating is presented in an attempt to understand the phenomena that occur during spraying and in particular how decarburization takes place. The second part of this article describes the wear behavior of HVOF WC – M coatings (where M= Co and CoCr) in the case of slurry erosion, in the light of this first part.

2. Experimental procedures

2.1. Material The starting powder was of nominal composition WC 88 wt.%–Co 12 wt.%. Its size distribution, determined with a laser diffusion sizer, was found to be in the range from 20 to 70 mm, with a mean value of 40 mm. The coatings were deposited by HVOF thermal spraying onto an austenitic stainless steel substrate. Two coatings, hereafter named 1 and 2 respectively, were produced with the above powder using two sets of spraying parameters which are summarized in Table 1. The most significant difference between them is the fuel gas used, namely H2 or C3H8. The thickness of the coatings were about 300 mm and their average Vickers hardness was 1141 and 972 N mm − 2 for coatings 1 and 2, respectively. The latter values are the average of 12 measures realized with an applied load of 1.6 N. As expected, SEM observation of the impressions showed that this system is prone to cracking, thus complicating the interpretation of the results. This matter will be discussed in more detail in the second part of this article.

2.2. Microstructural characterization To study the microstructure at its different levels of heterogeneity, several techniques were used: classical imaging with secondary or backscattered electrons (respectively SE and BSE) was carried out on a Cambridge S360 SEM, while a Philips CM20 TEM was used for conventional observations of thin foils, selected area diffraction (SAD) and micro-diffraction experiments. The two latter methods were used for local phase identification. Some high resolution observations were also performed in TEM. Quantitative image analysis was performed on SEM micrographs to measure average grain size or to quantify the relative abundance of the phases. For global structure analysis, the XRD technique was also used. XRD patterns were obtained on a u –2u diffractometer using the Cu Ka radiation with a diffracted beam monochromator. Spectra were recorded for 2u in the range of 20–90° with a beam diameter of the order of 1 mm. In addition, the relative proportion of phases were quantified using the Rietveld method [14] in the same way as already done in previous work by the authors [15]. Since some amount of nano-crystalline phase had been detected, the crystallinity of the samples was evaluated by an indicator I which is defined as the ratio between the areas of the Bragg peaks (crystalline material) and the total area of the spectrum for u comprised between 30 and 55°. This indicator varies from 0 to 100% and it is clear that it increases monotonically witha the actual mass fraction of the crystalline phases, but as the structure factor could not be determined for the nano-crystalline phase, the quantitative relation between this indicator and the actual mass fraction is not known.

2.3. Analytical electron microscopy For local chemical analysis, a field emission gun TEM Hitachi HF-2000 equipped with energy dispersive Table 1 Sets of HVOF spray parameters used in the present investigation

Fuel gas Fuel gas flow rate Oxygen flow rate Powder feed rate Carrier gas Carrier gas flow rate Spray distance Oxygen excess in the flamea a

Coating 1

Coating 2

H2 420 l min−1 450 l min−1 30 g min−1 N2 20 l min−1 225 mm 2.1

C3H8 55 l min−1 420 l min−1 40 g min−1 N2 20 l min−1 300 mm 2.2

Ratio between the experimental oxygen to fuel ratio and the oxygen to fuel ratio necessary for full combustion (assuming formation of CO).

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spectroscopy (EDS) and parallel electron energy loss spectroscopy (PEELS) was used. This type of microscope allows for a chemical analysis in areas of a few nm2 only. In order to determine the concentrations of carbon, cobalt, tungsten of a given area, EDS and PEELS spectra were recorded simultaneously. The relative elemental abundance is then expressed by the two following ratios: W/Co (measured by EDS) and C/Co (measured by PEELS). The detailed procedure as well as the problems encountered during these measurements are already described elsewhere [16]. Complementary global chemical analysis was also performed: X-ray fluorescence (XRF) and a combustion method were used to determine the metallic elements (W, Co) and the carbon contents respectively.

2.4. Specimen preparation The following procedure was used to prepare TEM samples. The coatings were spark machined to produce 100 mm thick discs parallel to the substrate. They were then dimpler ground down to 40 mm and ion milled with an 8 kV Ar beam under an incidence of about 4° until they were perforated. Then final thinning was performed at 4 kV under a more grazing incidence by applying a retarding potential, in order to prevent selective sputtering of the different phases [17]. Large observable areas were obtained. In order to prepare cross-section TEM samples, a small piece of the coating (250× 300×700 mm) was embedded in a sintered WC– Co holder and subsequently thinned in the same way as above [18].

3. Results

3.1. Microstructural identification 3.1.1. X-ray diffraction Fig. 1 presents a comparison between the XRD pattern of the starting powder and those of the coatings. In the powder spectrum, the intensity peaks could be related to WC and cubic cobalt, but other weak peaks corresponding to trigonal W2C and Co6W6C are also observed. The latter carbides were probably produced during powder manufacturing. Only small amounts of those compounds are present as revealed by the Rietveld quantification of the XRD pattern in Table 2. Because XRF of Co is excited by the Cu radiation, its absorption coefficient is much higher than that of WC [19]. This prevents a quantitative analysis of the phases containing Co, in particular the matrix, but still allows quantification of the other phases, one relative to the other, assuming that the absorption by the matrix is the same for all of them. Hence, the value which appears in Table 2 for the Co content of the

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Fig. 1. XRD spectrum of the powder, coatings 1 and 2. Humps on the coatings spectra, which are situated between the Bragg peaks and the base lines (dotted lines), reveal a nano-crystalline phase.

starting powder is not taken from XRD measurements but from XRF analysis. According to the XRD pattern of coating 1, transformations occur during spraying. WC is still present, but the amount of W2C has increased as compared to the powder, metallic tungsten has appeared while neither cobalt nor Co6W6C are detectable anymore. It is to be noted that there is no free carbon (graphite, main peak at 2u=26.38°) in the coatings. In addition, two large humps are present, which are centered respectively at 2u= 40° and 2u=72°. They have been previously identified [7] and can be connected to a nano-crystalline phase [15] containing cobalt and tungsten, which corresponds to the matrix. The fact that the cobalt peaks are detected in the powder particles and not in the coating, suggests that most of the cobalt is retained in the nano-crystalline binder phase. Nevertheless, the TEM observations (see Section 3.1.4) will show that very small amounts of metallic cobalt dendrites are produced in a few areas. As can be seen from the XRD pattern of coating 2, similar phases are present, but the extent of the transformations is lower than in coating 1. Table 2 summarizes the quantitative results of the Rietveld analysis. The amounts of W2C and metallic tungsten as compared to that of WC are smaller in coating 2. Moreover the fraction of the nano-crystalline binder phase, shown by the crystallinity index I, is also smaller in coating 2.

3.1.2. SEM obser6ations Fig. 2a shows the structure of the starting WC–Co powder. It was manufactured by agglomeration and densification that results in a nearly spherical shape of the powder grains. The microstructure of the inner part of a single powder particle is better seen on a polished section (Fig. 2b). The particle is made up of interconnected angular carbide grains with a mean size of about 1.1 mm, which exhibit bright contrast, embedded in the

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Table 2 X-ray Riedvelt analysis of the relative quantities of crystalline phases in the powder and in the coatings Phases

WC wt.%

W2C wt.%

W wt.%

Co wt.%

Co6W6C wt.%

I%

Starting powder Coating 1 Coating 2

85 60 73

2 25 21

— 15 6

12a — —

1 — —

100 42 73

I is the index of crystallinity. The cobalt content is calculated from XRF analysis.

a

cobalt binder phase which has been etched during polishing and therefore exhibits a dark contrast. The volume fraction of the binder as compared to WC is about 20 vol.% which is in agreement with the nominal weight composition of the powder (12 wt.% Co). Except for a few pores, the powder is dense. According to the X-ray diffraction analysis already described (Section 3.1.1), the raw powder contains only small amounts of secondary phases besides the two major ones, WC and Co. These characteristics make the microstructure of the powder very near that of sintered WC–Co and close to an ideal two phase material. For comparison, Fig. 2c and d show a cross section of coating 1 and 2 respectively. They exhibit a splat pattern-like structure which consists of islands which are elongated in directions parallel to the substrate surface. Some of them contain carbide grains which have a bright contrast while other islands appear to be more homogeneous. In the regions which are rich in carbides, the corresponding grains have lost their angularity. A comparison of the powder particle and the coatings microstructures, shows that in the latter ones, the volume fraction of the matrix has increased at the expense of the carbide content, due to the dissolution of WC in cobalt. In contrast to the powder grains, there are no longer carbide – carbide contacts in the coatings. SEM image analysis showed that the matrix phase covers about 65 vol.% of the microstructure in coating 1. In coating 2 the matrix covers only about 40 vol.% of the microstructure, that means that the WC grain dissolution is more restricted as compared to coating 1. In addition coating 2 cross section exhibits more pores than coating 1. According to their angular shape, they probably result from carbide pull out during polishing [20]. A close inspection of Fig. 2c,d shows that the matrix exhibits a bright or dark contrast in areas where the carbide fraction is small or large respectively. In the BSE imaging mode, the Z contrast predicts that the brighter zones have a higher tungsten content [21]. This means that the tungsten content of the matrix is higher in areas where carbide dissolution was more pronounced as confirmed by the microanalytical observations (see Section 3.2.1).

3.1.3. Chemical analysis Table 3 shows the chemical compositions of the starting powder and the coating respectively. No big changes are observed in the relative tungsten and cobalt contents of both coatings as compared to the powder, implying that cobalt was not lost preferentially during deposition, as it may be the case for plasma spraying [13]. In contrast, the coatings exhibit a substantial carbon loss which is more pronounced for coating 1 (− 45%) than for coating 2 (− 19%). 3.1.4. Electron microscope obser6ations To complete this macroscopic analysis and investigate morphology and development of the new phases, a TEM and SEM study was carried out. The different features constituting the microstructure of the coatings will be presented successively. The following description applies to both coatings 1 and 2, as they are qualitatively similar and differ only by quantitative aspects (as described previously and in Section 4.4). The differences will be discussed where necessary. 3.1.4.1. The nano-crystalline phase. The binder phase was found to be nano-crystalline all over the section. It can be easily identified since its contrast does not change through tilting experiments and the diffraction diagrams exhibit a typical diffuse ring pattern (Fig. 3 in insert). The high resolution TEM micrograph (Fig. 3) shows the nano-crystallites that constitute the matrix. They are only observable at the edge of the hole of the TEM sample, where the foil is thin enough to contain only one crystallite layer. Indeed, lattice fringes are clearly visible, which correspond to end-on-crystallographic planes of a crystallite. In thicker parts, the superimposition of several crystallites prevents the observation of phase contrasts. The matrix is not textured, at least at the scale of the smallest SA aperture i.e. 1 mm, as revealed by the isotropic ring pattern. The size of the crystallites on the image is in the range from 2 to 8 nm and a quantitative analysis of the ring broadening [18] gives a mean value of about 4 nm. 3.1.4.2. Di-tungsten carbide. WC, W2C and W grains have been differentiated by systematic microdiffraction experiments along several zone axes in TEM, since it is

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Fig. 2. SEM micrograph showing the morphology and structure of the original powder and the coatings. SE images of (a) the powder grains and (b) a polished section of a powder particle. Cross section BSE images (Z contrast) of (c) coating 1 and (d) coating 2.

not always possible to distinguish between these three phases by indexing one microdiffraction pattern only. An example of a W2C crystallite is presented in Fig. 4. The TEM micrograph shows a characteristic corerim structure of W2C (B) which surrounds a WC grain (A). Both compounds are in epitaxy as illustrated by the diffraction patterns in inserts: the two unit cells have all their a-, b-, c-axes parallel one to another. The c-axes are in the plane of the micrograph and their common direction is parallel to the diffraction vector. Since (001) tungsten planes have different spacing in WC and W2C, i.e. 0.284 and 0.236 nm respectively, there is an important mismatch between the two lattices. This gives rise to moire´ fringes at the Table 3 Chemical analysis by XRF (Co, W) and combustion method (C) of the starting powder and the coatings

Starting powder Coating 1 Coating 2

W wt.%

Co wt.%

C wt.%

82.2 91 83.8 91 81.39 1

12.49 1 13.39 1 14.39 1

5.37 9 0.02 2.94 9 0.01 4.339 0.01

interface (M), perpendicular to diffraction vector. The interface between W2C and the nano-crystalline phase (N) is not diffuse but rather sharp and exhibits a smooth curvature. In addition, contrary to WC which contains structural defects such as dislocations and stacking faults, W2C is almost free of defects as observed through tilting experiments. Its morphology suggests that during carbide dissolution decarburization also takes place at the surface of WC grains. Carbon diffusion towards the outside of the grain leads to sub-carbide formation.

3.1.4.3. Tungsten precipitates. Fig. 5 shows a typical tungsten precipitate. It is : 100 nm in size i.e. much smaller than WC and W2C grains. In addition, the image and the diffraction pattern of the entire grain show that it is made of a cluster numerous small crystals about 10 nm each, while WC and W2C grains were single crystals. These tungsten crystallites have been observed only in the vicinity of island boundaries (arrow in Fig. 5) where carbides are no longer present. 3.1.4.4. Cobalt dendrites. In a few areas where the density of remaining carbides is high, dendritic struc-

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Fig. 3. High resolution TEM micrograph of the nano-crystalline binder phase showing the crystallites (coating 1). In insert, the related SAD pattern shows that the crystallites are randomly oriented.

tures are observed. Fig. 6 shows an example of such dendrites which have grown on WC grain and are surrounded by the nano-crystalline phase. The microdiffraction experiments reveal, as expected, that each dendrite is a single crystal. Combined with EDS measurements, they also reveal that the dendrites are made of a solid solution of tungsten in cobalt (8 – 25 at.%) with a cubic structure. The tungsten concentration in the nano-crystalline binder in the region was found to be low e.g. W/Co B 0.3. These cobalt dendrites are seldom encountered in the coating microstructure. Their small volume fraction is confirmed by the absence of cobalt peaks in the XRD patterns of Fig. 1.

3.1.4.5. General morphology. Fig. 7 illustrates how the phases described above appear when observed in SEM (except for Co). They are identified through their characteristic size and shape. Two different islands separated by a boundary (C) can be seen. Within the one located on the left hand side, the spatial distribution of the phases is well defined. Far from the island boundary, WC dissolution is limited and WC is directly surrounded by the nano-crystalline binder. Closer to the boundary, one can distinguish the typical WC (A)– W2C (B) structure with the core-rim morphology similar to that of Fig. 4. Regions corresponding to W2C are preferentially etched to the WC during polishing, allowing for an easy recognition of that phase on SE image. Different configurations of the core-rim structure can be seen depending on the plane of sectioning. Very close to the boundary, small tungsten grains (W) are present in some areas of the adjacent island.

3.2. Microanalysis results 3.2.1. Nano-crystalline matrix Considering the heterogeneity of the microstructure, a spatial distribution of the matrix composition was expected. Therefore the points of analysis were randomly chosen in the nano-crystalline phase along the hole’s edge. Cross section TEM samples were used so as to explore many different islands through the coating thickness. Fig. 8a,b show the distribution of the composition, expressed by the C/Co atomic ratio measured by PEELS as a function of the W/Co atomic ratio measured by EDS, through the nano-crystalline binder. The data points indicate that the matrix is enriched in tungsten and carbon as compared to the starting Co matrix. Since there is no loss of tungsten and cobalt contents during deposition (see Section 3.1.3), the W/Co ratio measured in the matrix is directly related to the quantity of the original carbides which have dissolved in the matrix during the spraying process. The variation of this ratio through the nano-crystalline phase points out that this dissolution is heterogeneous. The mean values of the W/Co ratios extracted from these diagrams are 2.0 and 0.9 for coatings 1 and 2, respectively. For comparison the calculated nominal value of W/Co is 2.3 for the powder. This indicates that the dissolution of WC grains is more pronounced for coating 1 than for coating 2. Moreover, most points of these diagrams are situated below a dashed line which corresponds to W/C =1, which is the nominal value of the W/C ratio in WC–Co powder grains. This shows that only part of the carbon

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Fig. 4. Observation of W2C formation in coating 1. TEM image of the core-rim structure of WC (A) and W2C (B). The related SAD patterns in insert show the perfect epitaxy of the two phases. Moire´ fringes at the interface (M) are due to the difference in d001 values. N indicates the nano-crystalline binder phase.

atoms resulting from WC dissolution is retained in the matrix. This local loss of carbon and the formation of phases containing less carbon than WC (W2C and W) account for the global carbon loss which was measured by chemical analysis.

3.2.2. Island boundaries in the coating Some of the boundaries between the different islands are visible in the SEM micrograph of the cross section of coating (Fig. 2c,d, Fig. 7). They do not look the same everywhere, some of them being hardly distinguishable while others show a sharp line contrast. Fig. 9 is a dark field image of a boundary between two islands in coating 1. Each island exhibits the characteristic nano-crystalline contrast in dark field which allows crystallites to be seen. The boundary itself consists of an interlayer of about 8 nm wide which is amorphous [16]. Its chemical composition differs strongly from that of the adjacent nano-crystalline phase as revealed by TEM microanalysis (Table 4). The nano-crystalline phase composition is roughly the same on either side of the interface. In this phase and near the interface, the atomic W/Co ratio (about 2.5) is higher than the mean value measured through the sample (2.1) and is still higher in the interlayer (3.3). The C/W ratio is low (about 0.4) in the matrix on both sides as compared to values measured across the coating (Fig. 8). This points out that decarburization is more important near the boundaries. Some oxygen was also detected in these areas, as opposed to in the matrix far from the boundaries.

Depending on their location in the microstructure, these interlayers show different widths and compositions. Their width was observed to vary from a few up to 50 nm and their corresponding C/Co and O/Co ratios can be multiplied by a factor of 3–50 for both elements as compared to their value in the adjacent nano-crystalline phase. Despite of a systematic examination of many interlayers, no correlation between width and composition has been observed.

4. Discussion

4.1. Summary of the obser6ations All the microstructural features that we have presented above are summarized in Fig. 10. This figure schematically represents the microstructure on a cross section of the coating. It is made up of elongated islands parallel to the substrate, which contain various amounts of WC phase according to the extent of WC dissolution in cobalt. This dissolution produces a nanocrystalline binder phase after cooling. Therefore the composition of the binder fluctuates through the microstructure and it was previously shown that the tungsten concentration in the binder increases at the expense of the volume fraction of the remaining WC grains. In addition, new crystalline phases appear during the spraying process as described later. The tungsten precipitates are formed in the vicinity of some island boundaries, in areas which do not

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Fig. 5. Observation of W precipitates in coating 1. Typical TEM micrograph of a W polycrystalline precipitate with the corresponding indexed SAD pattern in the insert. The arrow points to an island boundary. N indicates the nano-crystalline binder phase.

contain any WC grains and exhibit low carbon concentration. The WC grains have completely dissolved, in agreement with the high tungsten concentration in the nano-crystalline phase in the area. The cobalt dendrites, which were seldom encountered in the above samples, were only observed in areas where a relatively small fraction of WC grains has dissolved with a concomitant low W/Co ratio in the matrix nearby. The W2C crystals form the characteristic core-rim structure already reported. Areas where they are observed seem to be less specific than the other phases. Colonies of W2C/WC grains are present for a variety of local WC phase volume fractions, except for extreme values which correspond to complete dissolution on the one hand or to a low amount of dissolution on the other hand. Finally, some islands are bounded by an amorphous interlayer of varying composition and width. Its chemical composition is different from that of the neighboring nano-crystalline matrix.

4.2. Comparison with pre6ious results A straight forward comparison is not possible for the following reason. The thermally sprayed WC – Co coatings which have been studied by different authors, were obtained through various techniques such as air plasma spray, vacuum plasma spray, high energy plasma and HVOF. In addition, for the same technique, the spraying parameters differ. The comparison of coating 1 and 2 illustrates the influence of the former parameters on the microstructural features. The powder characteristics (composition, manufacturing processes, grain size and

shape), were shown to have a great influence on the composition and properties of the resulting coating [8]. Nevertheless, some common features can be demonstrated in the coating microstructures which reveal that a few similar phenomena take place during various spraying processes. The WC and W2C phases which have been identified here using electron micrographs, micro-diffraction and microanalytical results, agree with previously published XRD results [6–8,11,22] as well as the almost complete absence of any metallic cobalt, except for cobalt contents \ 16 wt.% [9]. W2C seems to be the first phase formed when carbon loss occurs. In contrast, the noticeable amount of metallic tungsten detected in our samples, was confirmed in some studies [8,11] but not observed in others [6,9]. As far as the nano-crystalline phase is concerned, its existence is clearly established in Section 3 above on the basis of high resolution images and dark field observations in TEM. The size of the crystallites as well as a large spatial variation of the phase composition over the sample, explain the peak broadening of the various phases contained which accounts for the humps observed on the XRD pattern of Fig. 1. Some works [7,23] report the presence of similar features on XRD patterns and by comparison with bright field TEM observations and DTA data conclude that the matrix has an amorphous structure.

4.3. Microstructure formation during spraying 4.3.1. Structure modifications Most of the above observation are consistent with the following description which is sketched in Fig. 11. During the HVOF spraying process, the WC –Co pow-

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Fig. 6. TEM image of the dendrites (D) around WC grains (A). The corresponding SA diffraction pattern in insert can be indexed as a dilated fcc Co structure. N indicates the nano-crystalline binder phase.

der particles are accelerated and heated as they travel through the flame (Fig. 11a). They melt except for a fraction of the WC grains they contain. The coating is built up by the piling up of the impacting droplets which are flattened by the acceleration forces and rapidly cooled down. The islands within the coating result from the collapse of the droplets on the material which has just been deposited (Fig. 11b). Smaller islands could correspond to smaller droplets or droplet fractions which result from particle division, either in

the flame [24] or when impacting the coating [25]. In addition, chemical reactions and phase transformations take place during the entire process. The temperature and the velocity of the particles when they impact the substrate depend on their size. However, computer simulations of the HVOF process under relatively standard conditions [26,27] show that the temperature reached at the center of WC–Co powder grains remains between the temperature of the eutectic in the W–Co–C system (1320°C, WC 65

Fig. 7. Micrograph showing the appearance of the different phases observed in SEM. WC grains (A) are partially dissolved and may be surrounded by W2C rim (B). They are embedded in a nano-crystalline binder phase (N). Colony of W grains (W) are present in the vicinity of an island boundary (C).

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wt.%–Co 35 wt.% [28]), and the temperature of WC peritectic decomposition (2785°C [29]) for the range of particle size simulated (10 – 44 mm in diameter). In this range of temperatures, the ternary W – Co – C diagram (see [30] and Appendix A) shows that the higher the temperature, the higher the WC solubility in the liquid. This liquid gives rise to the nano-crystalline binder during rapid cooling as the droplets impact the substrate, while the non-dissolved WC grains are retained in the coating. This process accounts for the presence of cobalt, tungsten and carbon in the binder phase of the coating as well as for the above correlation between the local composition of the matrix and the local volume fraction of remaining WC. According to this description, the observed microstructural heterogeneity is mainly due to three factors: 1. the size distribution of the powder particles, 2. their various paths in the flame, which are responsible for a large temperature and velocity spectrum

Fig. 8. Chemical analysis of the matrix by EDS and PEELS techniques in TEM. The distribution of the nano-crystalline phase composition is plotted on a diagram of the C/Co atomic ratio measured by PEELS versus the W/Co atomic ratio measured by EDS. (a) for coating 1 and (b) coating 2.

for the powder grains as they impact the substrate, and 3. thermal equilibrium is not reached, which allows for a temperature gradient inside the grains (Fig. 11a). It is to be noted that if the effect of the two first factors can be limited by careful design of the powder and the spraying equipment, the latter is unavoidable if the primary goal is to retain WC in the coating.

4.3.2. Carbon loss mechanism The experimental observations show a global carbon loss and shed new light on the decarburization processes. It is generally admitted that the loss of carbon is due to its interaction with oxygen during spraying. It has been claimed that carbon loss occurs through direct reaction of WC with the oxygen contained in the flame [22]. For other plasma sprayed coatings, the proposed loss mechanism includes WC decomposition into W2C and carbon, followed by oxidation of the sub-carbide [8,11]. However, all the above observations suggest that during spraying, WC grains liberate carbon (and tungsten) into the surrounding liquid matrix. Carbon is oxidized at high temperature at the surface of the droplet and then eliminated in the form of carbon oxide. Thus the surface of the droplets acts as a sink for carbon, that induces carbon diffusion through the liquid toward the surface thereby allowing for the decarburization of the inner part of the droplets. Carbon could also reach the surface of the droplet by convection in the liquid. Besides, creation of new surfaces through in flight particle fragmentation [24] can enhance the process. This description accounts for the fact that not only the outer shell of the droplets which was in direct contact with oxygen is decarburized, but also the inner part of the islands. Notably, W2C was formed far from the island boundaries (Fig. 7) and the entire matrix has a deficiency of carbon compared to the amount of tungsten in solution (C/W B1, Fig. 8a,b). According to these mechanisms, the rate at which carbon is removed during spraying depends on the three following mechanisms which act in series: 1. the dissolution WC grains, 2. the diffusion of carbon toward the surface of the particles, and 3. carbon oxidation at the surface. The temperature distribution influences the three mechanisms. In addition, the first one depends on the size of carbide grains in the powder. In particular powder containing fine carbides should lead to a faster dissolution. The third step notably depends on the surface to volume ratio of the powder particles, that is their shape, size and the fragmentation they can experience in the flame. The slowest of the three controls the process.

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Fig. 9. Boundary between two different islands in coating 1. Dark field TEM micrograph (using the first diffraction ring) of the nano-crystalline phase belonging to two different adjacent islands and separated by an interlayer. The chemical analysis at points 1 – 3 may be found in Table 4.

This description is supported by the experiments conducted by Basinska-Pampuch and Gibas [31] who noted that pure WC powders experience a lower carbon loss than WC–Co powders sprayed in the same conditions. It also explains why tungsten carbide–cobalt powder containing W2C is less subjected to carbon loss [9]: owing to different eutectic temperatures for WC and Co (1320°C) and W2C and Co (\ 1800°C [18]), W2C can only dissolve in Co at higher temperature. As a result, for a same thermal history, W2C dissolves less than WC, which limits the potential carbon loss. The detail of the process by which carbon and oxygen react at the droplet surface is not known. However, the C and O rich interlayers which were found at some island boundaries, could result from this process. These interlayers can not be attributed to carbonated residues resulting from the gas combustion, since they were found in the microstructure even when H2 was used as fuel gas (coating 1). It is also obvious that such interlayer control the bond strength between the splats which constitute the coating. Table 4 EDS and PEELS analysis of the elemental content (atomic ratio) at a boundary between two islands in coating 1 No.

W/Co

C/Co

O/Co

1 2 3

2.4 3.3 2.6

0.93 10.5 1.1

0.29 5.1 0.45

Numbers correspond to those of Fig. 9.

4.3.3. New phase formation The nano-crystalline structure of the binding phase shows that solidification occurred under high cooling rates, that are typical for thermal spray processing [32]. Because of the lack of data concerning the cooling stage for the coatings studied or even for similar systems, it is difficult to ascertain the conditions under which the different phases are formed during solidification. In particular the nano-crystalline character of the matrix can be explained in two ways. The cooling rate could be high enough to provide for a high level of undercooling which would result in a high rate of homogeneous nucleation [33]. Alternatively the liquid could experience a glass transition followed by crystallization. A study [23] has shown that recrystallization of the binder phase takes place through annealing at a relatively low temperature (671°C). Such conditions can be experienced by the already deposited material during spraying. This latter case would mean that a mixture of W, Co and C is prone to form a metallic glass in a composition range. Similarly, the way W2C appears within the coating can not be deduced solely from the final microstructure. The coexistence of the crystalline phases with the nanocrystalline binder suggests that they were not formed at the same time. Moreover W2C decomposes below 1250°C and must be formed before quenching. W2C could result from direct decarburization of WC, a process which is concurrent to carbide dissolution during heating.

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Fig. 10. Schematics of the coating microstructure perpendicular to the substrate. Each phase is represented by a different pictogram. (A) WC carbides. (N) Nano-crystalline matrix. (B) WC + W2C. W – W crystallites. (D) Co dendrites surrounding WC grains. (C) Splat boundaries. The different phases and crystallites are not drawn at the same scale for the sake of clarity.

The presence of W precipitates can be understood by considering the temperature evolution during the particle flight. Within the spraying distance, the powder particle temperature first increases, reaches a maximum, and then decreases as it moves toward the substrate [34]. This could induce an in-flight solidification of the observed W crystallites before the droplets impact the substrate and prior to the rapid solidification of the liquid part of the droplet into the nano-crystalline binder. The special location of the small W crystallites, only in the vicinity of island boundaries, suggests that the precipitation is not only driven by temperature. According to the C-rich corner of the W – Co –C phase diagram (see Appendix A), this precipitation could also be induced by carbon loss. This is supported by the fact that the carbon concentration in the matrix was measured to be low near the island boundaries (Section 3.2.2). The metallic cobalt dendrites can only grow if the composition of the matrix is on the Co rich side of the eutectic. This corresponds to an atomic ratio W/Co = 0.16. Even if the EDS analyses of the matrix near the cobalt dendrites (W/Co 50.3, Section 3.1.4) are slightly higher than that value, they remain low as compared to the mean W/Co ratio in the binder (2.0 and 0.9 for coatings 1 and 2, respectively). Fig. 8a,b indicate that only few data points correspond to such low concentrations. They are related to a limited WC dissolution and

probably correspond to areas located at the center of the largest powder grains. This explains the scarcity of metallic cobalt.

4.4. Comparison of coatings 1 and 2 The above results exhibit significant differences between the two coatings. They are illustrated on Fig. 7a,b and in Table 2, Table 3. It can be seen that different spraying conditions result in different element concentrations and phase mass fractions strongly correlated to each other. In particular, some measured quantities follow the same trend: as the carbon loss increases (45% for coating 1 as compared to 19% for coating 2, Table 3), the mass fractions of W2C (25 and 21% for coatings 1 and 2, respectively) and W precipitates (15 and 6% for coatings 1 and 2, respectively) increase, as well as the amount of the nano-crystalline binder (I= 42 and 73% for coatings 1 and 2, respectively); the mean W concentration in the binder increases (mean value of W/Co are 2.0 and 0.9 for coatings 1 and 2, respectively) and the distribution of this element becomes more heterogeneous (compare Fig. 8a,b). The extent of the transformations that occur during spraying can be related to the fuel gas which is used. Indeed, the thermal conduction of a molecular gas increases at its dissociation temperature [35]. Since the latter temperature is lower for H2 (lower dissociation energy) as compared to C3H8, the resulting flame has a

C. Verdon et al. / Materials Science and Engineering A246 (1998) 11–24

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5. Conclusions The detailed features and chemical composition of the different phases which constitute HVOF sprayed WC–Co coatings have been quantitatively determined at a nanometric scale. These observations allow us to propose realistic physico-chemical processes which account for the microstructural changes in the material during spraying. In particular the proposed mechanism for the net carbon loss is the oxidation of carbon at the surface of the particles, following partial dissolution of WC grains in the liquid matrix. These phenomena account for the appearance of sub-carbide W2C, tungsten grains, and only small amount of metallic cobalt in a nano-crystalline matrix under cooling. The matrix contains cobalt, tungsten and carbon in proportions which fluctuate through the coating, in relation with the fluctuation of temperature and decarburization levels experienced by the impacting particles. The resistance of the coating microstructure with respect to slurry erosion is the subject of the second part of this article.

Acknowledgements Fig. 11. Schematics of microstructure formation during the spraying process. (a) A particle during flight. (b) Resulting islands in the coatings. The brighter color of the gray scale gradation corresponds to a lower W concentration in the liquid (a) and in the matrix (b), respectively.

higher thermal conduction at the same temperature. This allows for a higher heat transfer to the powder grains which induces a greater dissolution of WC and facilitate the loss of carbon according to the mechanisms previously described. Therefore, the differences which are observed between coatings 1 and 2, can be understood in the light of the transformation mechanisms which are thought to operate during HVOF spraying and that we have already described.

Fig. A1. Simulated isothermal sections of the W–Co–C phase diagram at (a) 1600°C and (b) 2000°C.

This work was supported by The Comission d’Encouragement a` la Recherche Scientifique presently ISFC, the Fond National Suisse de la Recherche Scientifique, and Sulzer AG. The authors are indebted to CIME-EPFL for providing electron microscopy facilities. Thanks are due to Sulzer Metco AG for manufacturing the coatings.

Appendix A. Simulation of the W–Co–C ternary diagram Owing to the very high heating and cooling rates occurring during thermal spraying, the transformations which take place in the material of the powder grain can not be considered as processes occurring at thermal equilibrium. However, the knowledge of the system equilibrium state as a function of the temperature gives some insight on the trend that the transformations will follow. This motivated the simulation of the W–Co–C diagram at high temperatures which is not currently available. The phase diagram was simulated using a program commercialized by the firm Thermodata [36]. This program uses a hill-climbing minimization procedure of Gibbs energy [37] to calculate phase equilibrium at given temperature and composition. Besides the pure elements C, Co and W and the liquid, the compounds WC, W2C, Co3W, Co7W6, M12C and M23C were considered for the calculation.

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Fig. A1a,b present isothermal sections of the W– Co–C diagram calculated at 1600 and 2000°C, respectively. Combined with the published section at 1425°C [30], they help to picture the evolution of the equilibrium state as temperature rises.

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