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WC-based cermet coatings produced by cold gas dynamic and pulsed gas dynamic spraying processes
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Surface & Coatings Technology 202 (2007) 382 – 390 www.elsevier.com/locate/surfcoat
WC-based cermet coatings produced by cold gas dynamic and pulsed gas dynamic spraying processes M. Yandouzi a,⁎, E. Sansoucy a , L. Ajdelsztajn b , B. Jodoin a a b
University of Ottawa, Ottawa, Ontario, Canada University of California, Davis, California, USA
Received 9 February 2007; accepted in revised form 30 May 2007 Available online 6 June 2007
Abstract Due to their mechanical properties, WC-based cermet coatings are extensively used in wear-resistant applications. These coatings are usually produced using thermal spray processes. However, due to the nature and the environment of these spraying processes, the feedstock powder structure and properties suffer from decomposition, which subsequently degrade the performance of the coatings produced. The cold gas dynamic spraying process appears to be a promising alternative technique to preserve the properties of the feedstock powder during the coating preparation. Although the latter technique can minimize or eliminate the degradation of the sprayed material, the deposition of cermet using this technique is a difficult task. In this study, two types of cermet powders, the nanocrystalline (WC–15Co) and the conventional (WC–10Co4Cr) powders were deposited using the cold gas dynamic spraying and the pulsed gas dynamic spraying processes. The feedstock powders and coatings microstructures were investigated by OM, SEM and XRD, as well as their hardness. The results revealed the possibility of depositing cermet coatings onto aluminum substrates using both processes without any degradation of the carbide phase of the feedstock powder. The cold gas dynamic spraying process experienced difficulty in depositing and building up dense coatings without major defects. The pulsed gas dynamic process produced thick cermet (conventional and nanocrystalline) coatings with low porosity as long as the feedstock powder was preheated above 573 K. © 2007 Elsevier B.V. All rights reserved. Keywords: Cold gas dynamic spraying; Pulsed gas dynamic spraying; Nanocrystalline; WC–15Co; WC–10Co–4Cr; Coatings
1. Introduction WC–Co cemented carbide (or cermet), in the form of sintered components or coatings has been successfully used to provide wear resistance in a wide range of applications, particularly in the heavy machinery sector [1]. The excellent wear resistance exhibited by these cemented carbides is attributed to their unique combination of high hardness and moderate levels of fracture toughness. A large variety of metallic matrix materials (Co, Cr, Ni, etc.) are available to incorporate the WC grains. The best performance in terms of wear resistance is
⁎ Corresponding author. Mechanical Engineering Department, Faculty of Engineering, University of Ottawa, 770 King Edward Avenue, Ottawa, Ontario, Canada K1N-6N5. E-mail address: [email protected] (M. Yandouzi). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.05.095
reached with cobalt [2]. However, the corrosion resistance of this composite is insufficient for many applications. In such cases, preference is given to coatings produced from commercially available WC–CoCr spray powders, in which the chromium addition provides an improvement in the corrosion resistance of the metallic binder phase over that of unalloyed WC–Co [3,4]. Cemented carbides wear properties have been extensively investigated over the years and it has been found that the abrasive wear resistance of sintered cermet is generally improved by a reduction of the binder volume fraction and a decrease of the carbide particle size [5]. Moreover, it has been reported that sintered WC-based nanocrystalline materials show greater wear resistance compared to the conventional (nonnanocrystalline) ones [6]. It is envisioned by the hard metals industry that nanocrystalline cemented carbide coatings could offer new opportunities for achieving a combination of superior hardness and toughness [7].
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Thermally sprayed (TS) cermet coatings, using either plasma spraying or high velocity oxygen fuel (HVOF) processes are widely used as they offer an effective and economical method of conferring wear resistance without compromising other attributes of the component. In particular, HVOF spraying has been extensively utilized to spray cermet coatings and it was shown to be one of the best methods for depositing WC–Co feedstock powders [8]. This was attributed to the higher velocities and lower temperatures achieved by the feedstock particles resulting in reduced decomposition of the WC phase during spraying, as compared to plasma based routes [9]. Consequently, this results in higher quality, improved wear-resistant coatings, with higher retention levels of the WC phase and reduced porosity. However, when compared to sintered cermets, nanocrystalline WC–Co coatings have shown contradictory results. Different studies have observed an inverse trend for the abrasion response of thermally sprayed WC–Co coatings [10]. The powder morphology, type of spraying system as well as the spray parameters have all been shown to affect the coating microstructure and properties [11]. Many studies have revealed that the high temperature characteristics of the traditional thermal spray processes have a negative effect. The WC–Co feedstock powder tends to undergo an unavoidable combination of decarburization, oxidation, reduction by reaction with the H2 (in plasma spraying), and dissolution/reaction between the WC and Co during spraying, which results in the formation of hard and brittle undesired phases (such as W2C, W, η-phases, WO3, etc.) [12,13]. Furthermore, it has been reported that a decrease in WC particle size in the feedstock powder led to an increase in the decomposition of the WC phase because of the increased reactivity of WC phase with decreasing particle size [14]. Therefore, significant amounts of unwanted phases have been reported in nanostructured WC–Co coatings produced using TS processes [15]. It is generally recognized that the promise of nanocrystalline WC–Co cemented carbide coatings will not be realized, unless the particle coalescence and carbide decomposition problems associated with the coating process can be overcome. Over the last decade, significant improvements in the production of nanocrystalline coatings have been obtained by optimization of the TS process parameters in order to maximize the retention of the nanocrystalline size and to minimize the decarburization of the ceramic reinforcement [13,16]. The coating process development has aimed at reducing and controlling the particle temperature while the particle impact velocity was increased. Successfully synthesized near nanostructured WC–18%Co coatings with a low amount of undesired phases were recently achieved and improved hardness and wear resistance were observed [16]. To minimize the degradation of cermet coatings caused by the process environment, it is envisioned that the cold gas dynamic spraying (CGDS) process, an emerging TS process, may be a potential alternative deposition technique not only for conventional cermet but also nanocrystalline materials [17]. In the CGDS process, the feedstock powder is injected in a supersonic gas flow and accelerated above a material dependant critical velocity. Upon impact of the particles on the substrate,
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thin surface films are disrupted providing an intimate contact between the particles and the substrate. The particles then deform plastically and bond to the substrate to form a coating [18]. The process temperature never reaches the softening or melting temperatures of the sprayed material [19]. As a result, the process is referred to as a solid-state process. In addition, both the original chemical composition and the microstructure of the feedstock powder are preserved [20]. The particles are not exposed to temperatures high enough to initiate chemical reactions or grain growth. Several pure metals, alloys, and composite coatings have successfully been produced using the CGDS process [21]. The coatings can be manufactured at high deposition rates and with very little oxidation, low residual stresses, low porosity and good coating–substrate adhesion. However, only a few cermets, such as WC–Co [22], and ceramic [23] coatings produced by the CGDS process have been reported due to the lack of ductile phase in the feedstock powder. Nevertheless, researchers have succeeded in producing nanostructured WC–12–17%Co coatings with low porosity using the CGDS process, and have attributed it to the reasonable powder preheating during the deposition [24]. It has been found that the low particle impact temperature in CGDS is problematic when spraying hard particles due to the difficulty of achieving the extremely high particle velocity required obtaining sufficient plastic deformation to produce a coating. Recently, a new coating process the pulsed gas dynamic spraying (PGDS) has been developed in at the Cold Spray Laboratory at the University of Ottawa [25,26]. It is envisioned that this process could allow the powder feedstock particles to be accelerated to high impact velocities (similar to those found in CGDS) and heated to intermediate temperatures, in a nonreacting gas, thus enabling the production of dense coatings that would exhibit the chemical and microstructural composition of the feedstock particles. It is expected that the intermediate particle impact temperature achieved through this process would be higher than the ones reached in the CGDS process and would lead to a lower required critical velocity, compared to CGDS. The objective of the present work is to produce two types of cemented carbide coatings, nanocrystalline WC–15Co and conventional WC–10Co4Cr, using the CGDS and PGDS processes. The effect of the process and the nature of the feedstock powder on the structure and the phase of the coatings are investigated using scanning electron microscopy (SEM), Xray diffraction (XRD) and hardness measurements. 2. Experimental procedures 2.1. Feedstock materials Two different WC-based cermet materials are used as feedstock powders to produce the coatings. The first is a nanocrystalline WC–Co powder that was produced using the spray drying technique. The second was a conventionally agglomerated and sintered cermet powder (SM5847-Sulzer Metco, USA). The details on the chemical composition and the particles size for both feedstock powders are given in Table 1.
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Table 1 Feedstock powders characteristics Feedstock powder
Composition (wt.%)
Particles size (μm)
Nanocrystalline WC–Co Conventional WC–CoCr
WC85–Co15 WC85–Co10–Cr4
37.1 ± 22.5 μm 32.4 ± 17.1 μm
The particle sizes given in Table 1 were measured using imaging software to analyze a set of SEM images of the powder. 2.2. Coating processes The coatings were produced using two processes, the CGDS process and the newly developed PGDS [26]. The schematic of both techniques is shown in Fig. 1. The CGDS system used in this work was developed at the University of Ottawa Cold Spray Laboratory [27,28]. It includes a spray chamber, a spray gun, a propellant gas heater, and a commercial powder feeder (Praxair Surface Technologies model 1264, Concord, NH, USA). The spray gun consists of a converging–diverging nozzle with an exit diameter of 7.3 mm. In the present work, helium was used as the propellant gas at a gas stagnation pressure and temperature of 1.7 MPa and 823 K, respectively. The coatings were produced on grit-blasted aluminum substrates. The substrate holder is equipped with a two-axis travel displacement system and provides the capability of coating a surface with multiple passes at set overlaps. In the current study, the coatings were produced from a single pass at a constant substrate traversing speed. The PGDS, a new spray process recently developed at the University of Ottawa Cold Spray Laboratory, is illustrated in Fig. 1. The system consists of a long smooth wall pipe of circular cross-section (4.2 mm in diameter in the present study), divided in two compartments separated by a valve (V2). The first compartment of the tube, the shock generator, is initially kept at high pressure (3.0 MPa in the present study). The second
compartment, the spraying gun, is opened at the end, at a lower pressure (atmospheric pressure ∼ 0.1 MPa in the present study). Helium is used as the propellant gas in this study. The initial pressure ratio of the system is set to generate the required driving shock wave in the driven section [25,26]. By rapidly opening/closing the valve V2, compression waves are generated at the valve and coalesce into a shock wave. The shock wave travels along the driven section (gun) and generates a high temperature supersonic gas flow that accelerates and heats the preheated feedstock material towards the substrate. The feedstock powder is dropped down from the feeder by opening the feeder valve (Vf), to let a specific amount of powder flow in the tube, and closing it prior to the passage of the shock wave. The system is equipped with an electrical gas heater, allowing preheating the propellant gas, and an electric feedstock powder heater, to preheat the powder prior to the injection in the system. A detailed description and analysis of the PGDS technique is presented elsewhere [25,26]. In the present work, the feedstock powders were preheated at different temperatures, from room temperature (RT) to 873 K, prior to injection. Coatings were deposited on grit-blasted aluminum substrates. While both techniques accelerate the solid feedstock powder in a high velocity inert gas flow, the PGDS technique, due to the unsteady nature of the flow generated in the spray gun, allows reaching a higher feedstock particle impact temperature [26]. In the PGDS process, the propellant gas does not experience the rapid drop of temperature found during the gas expansion in the CGDS process [29]. 2.3. Analysis techniques Microstructural characterization of the powders and coatings was performed by scanning electron microscopy (Philips XL30, LaB6 Analytical SEM), attached with an Electron Dispersion Spectroscopy (Oxford Instruments, EDS X-ray Microanalysis System) for chemical analysis. The powder samples were
Fig. 1. Schematic of CGDS and PGDS processes.
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examined directly. Prior to microstructural observations, crosssections of the coatings were obtained by resin cold mounting followed by grinding and polishing to a 0.05 μm surface finish. The porosity of the coatings was measured using an optical microscope (Olympus, Metallurgical Microscope) with 500× magnification images and the image analysis software (Clemex, Vision-Lite). The intensity range and thresholds were standardized on reference materials, and fifteen measurements were performed per sample, at various positions within the coating structure. Phase composition (crystallinity of the matrix, stoichiometry of tungsten carbides, formation of mixed phases) was investigated by XRD. The XRD analyses were carried out with an Xray diffractometer (Scintag XDS 2000) equipped with a graphite monochromator using Cu Kα (λ = 0.15406 nm) radiation, 30–80° 2θ range, 0.02° step width, 1 s per step acquisition time. Microhardness measurements were performed on mounted samples polished to a roughness of 0.25 μm using a Vickers microhardness tester (Struers Duramin-2). The indentation measurements were performed using 500 g (HVN0.5) load during 15 s. The measurements were made in the middle part of the coating cross-section of each specimen, parallel to the substrate. To avoid the effect of stress field, the distance between two indentations was kept greater than three times of the indentation diagonal [30,31]. The values presented for each coating is the average of ten measurements.
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3. Results 3.1. Feedstock powder characterization Fig. 2a and b presents respectively low and high magnification back-scattered SEM images of the conventional feedstock powder. This conventional cermet powder, as seen in Fig. 2b, consists of a high volume fraction of the agglomerated carbide micro-particles (from submicron size to 5 μm in diameter). The individual WC particles exhibit a blocky and angular shape, and are embedded in the metallic binder particles (CoCr). The carbide particles are not uniformly distributed inside the feedstock powder particles. The agglomerated particles have a spherical morphology, with an average size of 32.4 ± 17.1 μm. The measured particle size distribution is as follows: 32% from 1–20 μm, 40% from 21–40 μm, 21% from 41–60 μm, and 7% above 61 μm. The powder is highly porous, with large holes within the spherical particle. Fig. 2c and d presents respectively low and high magnification back-scattered SEM images of the nanocrystalline feedstock powder. The feedstock particles are dense and made of small agglomerated WC/Co particles. Although the larger particles exhibit a more spherical-shell morphology, the small ones do not show a specific morphology and look like fragmented shells of different sizes. Unlike the conventional powder, the nanocrystalline one has a broad size distribution,
Fig. 2. Low and high magnification SEM images showing the morphologies of the starting powders, (a, b) conventional WC–10%Co–4%Cr, (c and d) nanocrystalline WC–15%Co.
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Fig. 3. (a) XRD patterns of WC–10%Co–4%Cr cermet powder and (a′, a″) coatings prepared by PGDS process at 573 K and 773 K respectively. (b) XRD patterns of the conventional WC–12Co powder and coating deposited using HVOF process [31].
the feedstock powder. The XRD patterns of all coatings demonstrated the presence of WC phase, with minor secondary phases in the case of the conventional coatings, as initially found in the feedstock powder. Compared to cermet coatings prepared by other TS processes, the use of PGDS process has allowed producing coatings without phase change. As it can be seen in Figs. 3b and 4b, noticeable phase changes were reported when conventional and nanostructured WC–12Co cermet coating that were prepared by the HVOF process [31]. The analysis revealed the formation of W2C and W phases. Furthermore, Figs. 3b and 4b show small diffuse humps at approximately 40°, which is known to be evidence of amorphization of the metallic Co phase. The microstructure of the coatings in the as-polished conditions was investigated by OM and SEM, using the bright-field illumination technique and secondary-electron mode respectively. Fig. 5 (a and b) shows the typical microstructure of the conventional coating produced using the CGDS process. The coating thickness is approximately 200 μm. It is worth noting that the coating was difficult to produce using the CGDS process, with peeling of the coating frequently occurring while spraying. The cross-sectional microstructure shows that the coating is porous and that cracks with different sizes and orientations are observed. Nanocrystalline coatings prepared using the same process and spray conditions are presented in Fig. 5 (c and d). The coatings show similarity with the conventional one with the presence of large cracks.
from few microns fragmented shells to 100 μm in diameter particles. The carbide and the binder materials are combined together such a way that it is hard to distinguish between them at low magnification (Fig. 2c). At high magnification (Fig. 2d), the SEM image reveals the nanostructure features of the carbide particles. The carbides size varies from 50 nm to a submicron size. The observation of the powder has also shown that the nanocrystalline feedstock powder is porous, with the presence of submicron holes within the particles. The XRD patterns of both feedstock powders conventional and nanocrystalline one are shown in Figs. 3a and 4a respectively. The analysis revealed slight differences on the phase composition of the two powders. Although the results show that the two powders are mainly composed of WC, the spectra of the conventional powder indicated a minor presence of other phases such as Co3W3C. (indicated by arrows in Fig. 3a). Note that, in the case of nanocrystalline powder the Co phase was identified while in the conventional one it was very hard to identify the binder phases. 3.2. Coating characterization The XRD analysis of the as-sprayed coatings produced by both processes, CGDS and PGDS, did not show the presence of other phases than those present in the feedstock powders. Indeed, neither the nanocrystalline coating (Fig. 4a′) nor the conventional one (Fig. 3a′ and a″) has exhibited any noticeable phase change during the PGDS process, when compared with
Fig. 4. (a, a′) XRD patterns of WC–15Co nanocrystalline cermet powder and coatings prepared by PGDS process at 673 K respectively. (b) XRD patterns of the nanocrystalline WC–12Co powder and coating deposited using HVOF process [31].
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Fig. 5. OM and SEM images showing the morphology of both the (a, b) conventional and (c, d) nanocrystalline coatings deposited by the CGDS process.
Fig. 6. OM and SEM images showing the morphology of both the (a, b) conventional and (c, d) nanocrystalline coatings deposited by the PGDS process.
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Fig. 7. Vickers microhardness graph showing the HV0.5 measured values for the four different types of coatings and compared to hardness of different cermets using different deposition process reported in recent scientific papers.
Nevertheless, the coatings seem to be denser compared to the ones produced using the conventional feedstock powder. Typical microstructures of PGDS coatings produced using the conventional (a and b) and the nanocrystalline (c and d) cermet powders are shown in Fig. 6. The images are taken from the coatings produced with the feedstock powder preheated to 673 K prior to spraying. In general, the microstructure of the coatings is very dense and exhibits no lamellar or banded microstructure, which is typical of thermally sprayed coatings. Low levels of porosity (less than 5%) were obtained for the conventional coatings, while the nanocrystalline coatings seemed to be denser (levels of porosity less than 2%). The carbide distribution in the conventional coatings is generally less uniform than in the nanocrystalline coatings, as observed in the feedstock powder. The conventional and nanocrystalline coatings did not show the presence of large cracks as it is the case in coatings prepared using the CGDS process. Nevertheless, micro-cracks were observed in all coatings, especially close to the substrate–coating interface. Vickers microhardness (HV0.5) measurements were made in the middle of the coating cross-section of each sample. The measured values are presented in Fig. 7, as well as other hardness values obtained from coatings prepared by other thermal spray processes [33–35] and presented here for comparison purposes. The measurements show that PGDS coatings have higher hardness values (721 ± 95 HV and 899 ± 64 HV), compared to the ones produced by the conventional CGDS process (356 ± 135 HV and 462 ± 92 HV), using the conventional and the nanocrystalline powders respectively. 4. Discussion The cold gas dynamic spraying process has recently shown to be a suitable technique in order to preserve the microstructure of the feedstock powder in the coating [24,36,37]. In com-
parison to other TS processes such as HVOF, where carbide decomposition and amorphization of the binder have been reported [13,16,32], the coatings produced in the present work by both the CGDS and PGDS processes did not exhibit any major degradation of the original carbide phase. Similar results were also obtained in the case of nanostructured WC–12Co prepared by CGDS [24]. This is attributed to the lower deposition temperatures of the feedstock particles compared to those usually found in conventional TS processes and the inert environment of both the CGDS and PGDS processes. Hard materials, such as WC-based composites, are known to be difficult to deposit using the CGDS process as very few successful attempts have been reported [24]. In the present work, when using the CGDS process, neither the conventional nor the nanocrystalline cermet has been forming dense and thick coatings, despite increasing the main gas temperature up to 823 K. Both conventional and nanocrystalline CGDS coatings exhibit a large number of cracks within their microstructures (Fig. 5). This is attributed to the limited impact energy absorption available through plastic deformation of the binder phase (usually present in a small volume fraction for cermet materials). Consequently, in order to absorb the high release of energy during impact and high level of localized stresses in the material, cracks initiate and propagate mainly in the binder phase resulting in a poor coating microstructure. It has been reported that for metal matrix composite (MMC), where the binder volume fraction is much higher than the one in cermets, the presence of crack during deposition can be avoided. In CGDS Al12 wt.%Si/SiC coatings [38] the impact energy is well absorbed by the aluminum matrix. Consequently, higher Co content in WC/Co cermets would be more suited for CGDS but it would likely be detrimental to its wear properties. In CGDS, successful bonding of impacting particles requires localized deformation and adiabatic shear instabilities, which
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occur at high impact velocity, the critical velocity (Vcr). The critical velocity of most metals and alloys for CGDS deposition were reported to be in the range of 500–700 m/s based on theoretical and experimental works [39,40]. The empirical based model [40] has revealed that the critical velocity can be predicted by: Vcr ¼ 667 0:014q þ 0:08ðTm TR Þ þ 10 7ru 0:4ðTi TR Þ Where, Vcr is the critical velocity, ρ is the density of the material, Tm is the melting temperature, TR is the reference temperature (293 K), σu is the yield stress, and Ti is the impact temperature. The above formula indicates that particle preheating temperature of 773 K will decrease the critical particle velocity by 200 m/s (in the case of WC–12Co) [24]. It was shown that it is possible to deposit and to build up thick nanostructured WC– 12Co coating by cold gas dynamic spraying, as long as the powder is preheated to a minimal temperature of 773 K. By using the PGDS process, thick and dense coatings were obtained as long as the feedstock powder was preheated above 573 K (Fig. 6 reports typical coatings obtained when using 673 K). Below a preheating temperature of 573 K, the feedstock powder mostly bounces off the substrate, with very limited adhesion. Consequently, it becomes clear that the particle impact temperature plays a significant role, as predicted by the above formula. Particle preheating will decrease the critical particle velocity since the ductility of the material increases as the material temperature is increased, The microhardness measurements (Fig. 7) show that the PGDS coatings using both types of powder resulted in higher hardness values than the ones produced using the CGDS process. This can be explained by the higher porosity observed in the latter coatings and also the presence of larger number of cracks within these coating as seen in Fig. 6. The hardness of the nanocrystalline coatings produced using the CGDS and PGDS processes was almost two fold the values recorded for the conventional coatings. This observation is in accordance to reported values for nanostructured cermets found in the literature [41]. Usually finer carbide particles result in higher hardness values in bulk samples; however, this trend can be reversed in thermal spray coatings when carbide decomposition occurs. In these cases, the decarburization is increased by the large surface area in the nanostructured samples due to the reactive environment in the thermal spray process [33–35]. Both the CGDS and PGDS coatings the hardness of the nanocrystalline coatings is higher than the conventional ones. This can be rationalized by the fact that the inert spray environment does not promote decarburization in the coatings. Also from Fig. 7 one can observe that the PGDS coatings are harder than the CGDS for both types of powder. This difference is explained by the low porosity in the coating achieved using the PGDS when compared to the CGDS ones. The higher particle impact temperature achieved in the PGDS process allows coating formation at lower particle velocities as a result of a lower critical deposition velocity. It also reduces the impact energy to be absorbed by the binder phase in the form of fracture
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surfaces due to its capacity to absorb the energy through plastic deformation at higher spraying temperatures. 5. Conclusion The production of cermet coatings with conventional and nanocrystalline feedstock powders by CGDS process and PGDS process were carried out in this study. It was found that: – There is no degradation of the phase (phase transformation and/or decarburization of WC) when using either process. – Dense coatings without major defects are difficult to obtain when using the CGDS process. – It is possible to produce coatings with low porosity using the PGDS process if the feedstock powder is preheated above 573 K. – The nanostructure of the feedstock powder is preserved during the CGDS and PGDS depositions. – Hard and dense coatings were obtained when depositing nanocrystalline powder using the PGDS process. Acknowledgements The authors wish to acknowledge the financial support of the Natural Sciences and Engineering Research Council of Canada for this work. The authors express their gratitude to Miss. Renata Zavadil from CANMET laboratory for many helpful technical discussions. References [1] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, J. Mater. Sci. 20 (1985) 1873. [2] H.L. de Villiers Lovelock, J.M. Benson, P.M. Young, J. Therm. Spray Technol. 7 (1998) 97. [3] L.-M. Berger, P. Ettmayer, P. Vuoristo, W. Kunert, J. Therm. Spray Technol. 10 (2) (2001) 311. [4] A. Kasrimi, C. Verdon, Surf. Coat. Technol. 57 (1993) 81. [5] D.G.F. O'Quigley, M.N. James, Int. J. Refract. Met. Hard Mater. 15 (1) (1997) 73. [6] K. Jia, T.E. Fischer, Wear 203 (1997) 310. [7] B.K. Kim, G.H. Ha, D.W. Lee, J. Mater. Process. Technol. 63 (1997) 317. [8] J. Nerz, B. Kushner, A. Rotolico, in: R.M. Yazici (Ed.), Protective Coatings: Processing and Characterization, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 135. [9] R. Schwetzke, H. Kreye, Proceedings of the 15th International Thermal Spray Conference, Nice, France, 1998, p. 187. [10] S. Usmani, S. Sampath, D.L. Houck, D. Lee, Tribol. Trans. 40 (1997) 470. [11] C.J. Li, A. Ohmori, Y. Harada, J. Mater. Sci. 31 (1996) 785. [12] H.L.D. Lovelock, J.M. Benson, P.M. Young, J. Therm. Spray Technol. 7 (1998) 97. [13] B.H. Kear, G. Skandan, R.k. Sadanji, Scr. Mater. 44 (2001) 1703. [14] S. Usmani, S. Sampath, H. Herman, J. Therm. Spray Technol., 7 (1998) 429. [15] D.A. Stewart, P.H. Shipway, D.G. McCartney, Acta Mater. 48 (2000) 1593. [16] J. He, Y. Liu, T.E. Fischer, E.J. Lavernia, Metall. Mater. Trans., A Phys. Metall. Mater. Sci. 33 (1) (2002) 145. [17] T.H. Van Sreenkiste, R.C. McCune, K.J. Barnett, Surf. Coat. Technol. 111 (1999) 62. [18] R.C. Dykhuizen, M.F. Smith, D.L. Gilmore, R.A. Neiser, X. Jiang, S. Sampath, J. Therm. Spray Technol. 8 (1999) 559. [19] A.P. Alkhimov, V.F. Kosarev, A.N. Papyrin, Sov. Phys. Dokl. 35 (1990) 1047.
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[30] S.-H. Leigh, C.-K. Lin, C.C. Berndt, J. Am. Ceram. Soc. 80-8 (1997) 2093. [31] P.H. Shipway, D.G. McCartney, T. Sudaprasert, Wear 259 (2005) 820. [32] J.O. Berghaus, B. Marple, C. Moreau, ITSC 2006 Proceeding, May 15-18, Seattle, Washington, USA, 2006. [33] N. Petsas, M. Vardavoulias, A. Moutsatsou, J.P. Cellis, P. Matteazzi, ITSC 2006 Proceeding, May 15-18, Seattle, Washington, USA, 2006. [34] J.K.N. Murthy, B. Venkataraman, Surf. Coat. Technol. 200 (2006) 2642. [35] D.A. Stewart, P.H. Shipway, D.G. McCartney, Surf. Coat. Technol. 105 (1998) 13. [36] L. Ajdelsztajn, B. Jodoin, P. Richer, E. Sansoucy, E.J. Lavernia, J. Therm. Spray Technol. 15 (4) (2006) 495. [37] L. Ajdelsztajn, B. Jodoin, J.M. Schoenung, Surf. Coat. Technol. 201 (3-4) (2006) 1167. [38] E. Sansoucy, L. Ajdelsztajn, B. Jodoin, ITSC 2007–Global Coating Solutions Proceeding, ASM International, USA, 2007. [39] H. Assadi, F. Gartner, T. Stoltenhoff, H. Kreye, Acta Mater. 51 (2003) 4379. [40] T. Schmidt, F. Gartner, H. Assadi, H. Kreye, Acta Mater. 54 (3) (2006) 729. [41] J. He, J.M. Schoenung, Surf. Coat. Technol. 157 (2002) 72.
Surface & Coatings Technology 202 (2007) 382 – 390 www.elsevier.com/locate/surfcoat
WC-based cermet coatings produced by cold gas dynamic and pulsed gas dynamic spraying processes M. Yandouzi a,⁎, E. Sansoucy a , L. Ajdelsztajn b , B. Jodoin a a b
University of Ottawa, Ottawa, Ontario, Canada University of California, Davis, California, USA
Received 9 February 2007; accepted in revised form 30 May 2007 Available online 6 June 2007
Abstract Due to their mechanical properties, WC-based cermet coatings are extensively used in wear-resistant applications. These coatings are usually produced using thermal spray processes. However, due to the nature and the environment of these spraying processes, the feedstock powder structure and properties suffer from decomposition, which subsequently degrade the performance of the coatings produced. The cold gas dynamic spraying process appears to be a promising alternative technique to preserve the properties of the feedstock powder during the coating preparation. Although the latter technique can minimize or eliminate the degradation of the sprayed material, the deposition of cermet using this technique is a difficult task. In this study, two types of cermet powders, the nanocrystalline (WC–15Co) and the conventional (WC–10Co4Cr) powders were deposited using the cold gas dynamic spraying and the pulsed gas dynamic spraying processes. The feedstock powders and coatings microstructures were investigated by OM, SEM and XRD, as well as their hardness. The results revealed the possibility of depositing cermet coatings onto aluminum substrates using both processes without any degradation of the carbide phase of the feedstock powder. The cold gas dynamic spraying process experienced difficulty in depositing and building up dense coatings without major defects. The pulsed gas dynamic process produced thick cermet (conventional and nanocrystalline) coatings with low porosity as long as the feedstock powder was preheated above 573 K. © 2007 Elsevier B.V. All rights reserved. Keywords: Cold gas dynamic spraying; Pulsed gas dynamic spraying; Nanocrystalline; WC–15Co; WC–10Co–4Cr; Coatings
1. Introduction WC–Co cemented carbide (or cermet), in the form of sintered components or coatings has been successfully used to provide wear resistance in a wide range of applications, particularly in the heavy machinery sector [1]. The excellent wear resistance exhibited by these cemented carbides is attributed to their unique combination of high hardness and moderate levels of fracture toughness. A large variety of metallic matrix materials (Co, Cr, Ni, etc.) are available to incorporate the WC grains. The best performance in terms of wear resistance is
⁎ Corresponding author. Mechanical Engineering Department, Faculty of Engineering, University of Ottawa, 770 King Edward Avenue, Ottawa, Ontario, Canada K1N-6N5. E-mail address: [email protected] (M. Yandouzi). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.05.095
reached with cobalt [2]. However, the corrosion resistance of this composite is insufficient for many applications. In such cases, preference is given to coatings produced from commercially available WC–CoCr spray powders, in which the chromium addition provides an improvement in the corrosion resistance of the metallic binder phase over that of unalloyed WC–Co [3,4]. Cemented carbides wear properties have been extensively investigated over the years and it has been found that the abrasive wear resistance of sintered cermet is generally improved by a reduction of the binder volume fraction and a decrease of the carbide particle size [5]. Moreover, it has been reported that sintered WC-based nanocrystalline materials show greater wear resistance compared to the conventional (nonnanocrystalline) ones [6]. It is envisioned by the hard metals industry that nanocrystalline cemented carbide coatings could offer new opportunities for achieving a combination of superior hardness and toughness [7].
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Thermally sprayed (TS) cermet coatings, using either plasma spraying or high velocity oxygen fuel (HVOF) processes are widely used as they offer an effective and economical method of conferring wear resistance without compromising other attributes of the component. In particular, HVOF spraying has been extensively utilized to spray cermet coatings and it was shown to be one of the best methods for depositing WC–Co feedstock powders [8]. This was attributed to the higher velocities and lower temperatures achieved by the feedstock particles resulting in reduced decomposition of the WC phase during spraying, as compared to plasma based routes [9]. Consequently, this results in higher quality, improved wear-resistant coatings, with higher retention levels of the WC phase and reduced porosity. However, when compared to sintered cermets, nanocrystalline WC–Co coatings have shown contradictory results. Different studies have observed an inverse trend for the abrasion response of thermally sprayed WC–Co coatings [10]. The powder morphology, type of spraying system as well as the spray parameters have all been shown to affect the coating microstructure and properties [11]. Many studies have revealed that the high temperature characteristics of the traditional thermal spray processes have a negative effect. The WC–Co feedstock powder tends to undergo an unavoidable combination of decarburization, oxidation, reduction by reaction with the H2 (in plasma spraying), and dissolution/reaction between the WC and Co during spraying, which results in the formation of hard and brittle undesired phases (such as W2C, W, η-phases, WO3, etc.) [12,13]. Furthermore, it has been reported that a decrease in WC particle size in the feedstock powder led to an increase in the decomposition of the WC phase because of the increased reactivity of WC phase with decreasing particle size [14]. Therefore, significant amounts of unwanted phases have been reported in nanostructured WC–Co coatings produced using TS processes [15]. It is generally recognized that the promise of nanocrystalline WC–Co cemented carbide coatings will not be realized, unless the particle coalescence and carbide decomposition problems associated with the coating process can be overcome. Over the last decade, significant improvements in the production of nanocrystalline coatings have been obtained by optimization of the TS process parameters in order to maximize the retention of the nanocrystalline size and to minimize the decarburization of the ceramic reinforcement [13,16]. The coating process development has aimed at reducing and controlling the particle temperature while the particle impact velocity was increased. Successfully synthesized near nanostructured WC–18%Co coatings with a low amount of undesired phases were recently achieved and improved hardness and wear resistance were observed [16]. To minimize the degradation of cermet coatings caused by the process environment, it is envisioned that the cold gas dynamic spraying (CGDS) process, an emerging TS process, may be a potential alternative deposition technique not only for conventional cermet but also nanocrystalline materials [17]. In the CGDS process, the feedstock powder is injected in a supersonic gas flow and accelerated above a material dependant critical velocity. Upon impact of the particles on the substrate,
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thin surface films are disrupted providing an intimate contact between the particles and the substrate. The particles then deform plastically and bond to the substrate to form a coating [18]. The process temperature never reaches the softening or melting temperatures of the sprayed material [19]. As a result, the process is referred to as a solid-state process. In addition, both the original chemical composition and the microstructure of the feedstock powder are preserved [20]. The particles are not exposed to temperatures high enough to initiate chemical reactions or grain growth. Several pure metals, alloys, and composite coatings have successfully been produced using the CGDS process [21]. The coatings can be manufactured at high deposition rates and with very little oxidation, low residual stresses, low porosity and good coating–substrate adhesion. However, only a few cermets, such as WC–Co [22], and ceramic [23] coatings produced by the CGDS process have been reported due to the lack of ductile phase in the feedstock powder. Nevertheless, researchers have succeeded in producing nanostructured WC–12–17%Co coatings with low porosity using the CGDS process, and have attributed it to the reasonable powder preheating during the deposition [24]. It has been found that the low particle impact temperature in CGDS is problematic when spraying hard particles due to the difficulty of achieving the extremely high particle velocity required obtaining sufficient plastic deformation to produce a coating. Recently, a new coating process the pulsed gas dynamic spraying (PGDS) has been developed in at the Cold Spray Laboratory at the University of Ottawa [25,26]. It is envisioned that this process could allow the powder feedstock particles to be accelerated to high impact velocities (similar to those found in CGDS) and heated to intermediate temperatures, in a nonreacting gas, thus enabling the production of dense coatings that would exhibit the chemical and microstructural composition of the feedstock particles. It is expected that the intermediate particle impact temperature achieved through this process would be higher than the ones reached in the CGDS process and would lead to a lower required critical velocity, compared to CGDS. The objective of the present work is to produce two types of cemented carbide coatings, nanocrystalline WC–15Co and conventional WC–10Co4Cr, using the CGDS and PGDS processes. The effect of the process and the nature of the feedstock powder on the structure and the phase of the coatings are investigated using scanning electron microscopy (SEM), Xray diffraction (XRD) and hardness measurements. 2. Experimental procedures 2.1. Feedstock materials Two different WC-based cermet materials are used as feedstock powders to produce the coatings. The first is a nanocrystalline WC–Co powder that was produced using the spray drying technique. The second was a conventionally agglomerated and sintered cermet powder (SM5847-Sulzer Metco, USA). The details on the chemical composition and the particles size for both feedstock powders are given in Table 1.
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Table 1 Feedstock powders characteristics Feedstock powder
Composition (wt.%)
Particles size (μm)
Nanocrystalline WC–Co Conventional WC–CoCr
WC85–Co15 WC85–Co10–Cr4
37.1 ± 22.5 μm 32.4 ± 17.1 μm
The particle sizes given in Table 1 were measured using imaging software to analyze a set of SEM images of the powder. 2.2. Coating processes The coatings were produced using two processes, the CGDS process and the newly developed PGDS [26]. The schematic of both techniques is shown in Fig. 1. The CGDS system used in this work was developed at the University of Ottawa Cold Spray Laboratory [27,28]. It includes a spray chamber, a spray gun, a propellant gas heater, and a commercial powder feeder (Praxair Surface Technologies model 1264, Concord, NH, USA). The spray gun consists of a converging–diverging nozzle with an exit diameter of 7.3 mm. In the present work, helium was used as the propellant gas at a gas stagnation pressure and temperature of 1.7 MPa and 823 K, respectively. The coatings were produced on grit-blasted aluminum substrates. The substrate holder is equipped with a two-axis travel displacement system and provides the capability of coating a surface with multiple passes at set overlaps. In the current study, the coatings were produced from a single pass at a constant substrate traversing speed. The PGDS, a new spray process recently developed at the University of Ottawa Cold Spray Laboratory, is illustrated in Fig. 1. The system consists of a long smooth wall pipe of circular cross-section (4.2 mm in diameter in the present study), divided in two compartments separated by a valve (V2). The first compartment of the tube, the shock generator, is initially kept at high pressure (3.0 MPa in the present study). The second
compartment, the spraying gun, is opened at the end, at a lower pressure (atmospheric pressure ∼ 0.1 MPa in the present study). Helium is used as the propellant gas in this study. The initial pressure ratio of the system is set to generate the required driving shock wave in the driven section [25,26]. By rapidly opening/closing the valve V2, compression waves are generated at the valve and coalesce into a shock wave. The shock wave travels along the driven section (gun) and generates a high temperature supersonic gas flow that accelerates and heats the preheated feedstock material towards the substrate. The feedstock powder is dropped down from the feeder by opening the feeder valve (Vf), to let a specific amount of powder flow in the tube, and closing it prior to the passage of the shock wave. The system is equipped with an electrical gas heater, allowing preheating the propellant gas, and an electric feedstock powder heater, to preheat the powder prior to the injection in the system. A detailed description and analysis of the PGDS technique is presented elsewhere [25,26]. In the present work, the feedstock powders were preheated at different temperatures, from room temperature (RT) to 873 K, prior to injection. Coatings were deposited on grit-blasted aluminum substrates. While both techniques accelerate the solid feedstock powder in a high velocity inert gas flow, the PGDS technique, due to the unsteady nature of the flow generated in the spray gun, allows reaching a higher feedstock particle impact temperature [26]. In the PGDS process, the propellant gas does not experience the rapid drop of temperature found during the gas expansion in the CGDS process [29]. 2.3. Analysis techniques Microstructural characterization of the powders and coatings was performed by scanning electron microscopy (Philips XL30, LaB6 Analytical SEM), attached with an Electron Dispersion Spectroscopy (Oxford Instruments, EDS X-ray Microanalysis System) for chemical analysis. The powder samples were
Fig. 1. Schematic of CGDS and PGDS processes.
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examined directly. Prior to microstructural observations, crosssections of the coatings were obtained by resin cold mounting followed by grinding and polishing to a 0.05 μm surface finish. The porosity of the coatings was measured using an optical microscope (Olympus, Metallurgical Microscope) with 500× magnification images and the image analysis software (Clemex, Vision-Lite). The intensity range and thresholds were standardized on reference materials, and fifteen measurements were performed per sample, at various positions within the coating structure. Phase composition (crystallinity of the matrix, stoichiometry of tungsten carbides, formation of mixed phases) was investigated by XRD. The XRD analyses were carried out with an Xray diffractometer (Scintag XDS 2000) equipped with a graphite monochromator using Cu Kα (λ = 0.15406 nm) radiation, 30–80° 2θ range, 0.02° step width, 1 s per step acquisition time. Microhardness measurements were performed on mounted samples polished to a roughness of 0.25 μm using a Vickers microhardness tester (Struers Duramin-2). The indentation measurements were performed using 500 g (HVN0.5) load during 15 s. The measurements were made in the middle part of the coating cross-section of each specimen, parallel to the substrate. To avoid the effect of stress field, the distance between two indentations was kept greater than three times of the indentation diagonal [30,31]. The values presented for each coating is the average of ten measurements.
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3. Results 3.1. Feedstock powder characterization Fig. 2a and b presents respectively low and high magnification back-scattered SEM images of the conventional feedstock powder. This conventional cermet powder, as seen in Fig. 2b, consists of a high volume fraction of the agglomerated carbide micro-particles (from submicron size to 5 μm in diameter). The individual WC particles exhibit a blocky and angular shape, and are embedded in the metallic binder particles (CoCr). The carbide particles are not uniformly distributed inside the feedstock powder particles. The agglomerated particles have a spherical morphology, with an average size of 32.4 ± 17.1 μm. The measured particle size distribution is as follows: 32% from 1–20 μm, 40% from 21–40 μm, 21% from 41–60 μm, and 7% above 61 μm. The powder is highly porous, with large holes within the spherical particle. Fig. 2c and d presents respectively low and high magnification back-scattered SEM images of the nanocrystalline feedstock powder. The feedstock particles are dense and made of small agglomerated WC/Co particles. Although the larger particles exhibit a more spherical-shell morphology, the small ones do not show a specific morphology and look like fragmented shells of different sizes. Unlike the conventional powder, the nanocrystalline one has a broad size distribution,
Fig. 2. Low and high magnification SEM images showing the morphologies of the starting powders, (a, b) conventional WC–10%Co–4%Cr, (c and d) nanocrystalline WC–15%Co.
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Fig. 3. (a) XRD patterns of WC–10%Co–4%Cr cermet powder and (a′, a″) coatings prepared by PGDS process at 573 K and 773 K respectively. (b) XRD patterns of the conventional WC–12Co powder and coating deposited using HVOF process [31].
the feedstock powder. The XRD patterns of all coatings demonstrated the presence of WC phase, with minor secondary phases in the case of the conventional coatings, as initially found in the feedstock powder. Compared to cermet coatings prepared by other TS processes, the use of PGDS process has allowed producing coatings without phase change. As it can be seen in Figs. 3b and 4b, noticeable phase changes were reported when conventional and nanostructured WC–12Co cermet coating that were prepared by the HVOF process [31]. The analysis revealed the formation of W2C and W phases. Furthermore, Figs. 3b and 4b show small diffuse humps at approximately 40°, which is known to be evidence of amorphization of the metallic Co phase. The microstructure of the coatings in the as-polished conditions was investigated by OM and SEM, using the bright-field illumination technique and secondary-electron mode respectively. Fig. 5 (a and b) shows the typical microstructure of the conventional coating produced using the CGDS process. The coating thickness is approximately 200 μm. It is worth noting that the coating was difficult to produce using the CGDS process, with peeling of the coating frequently occurring while spraying. The cross-sectional microstructure shows that the coating is porous and that cracks with different sizes and orientations are observed. Nanocrystalline coatings prepared using the same process and spray conditions are presented in Fig. 5 (c and d). The coatings show similarity with the conventional one with the presence of large cracks.
from few microns fragmented shells to 100 μm in diameter particles. The carbide and the binder materials are combined together such a way that it is hard to distinguish between them at low magnification (Fig. 2c). At high magnification (Fig. 2d), the SEM image reveals the nanostructure features of the carbide particles. The carbides size varies from 50 nm to a submicron size. The observation of the powder has also shown that the nanocrystalline feedstock powder is porous, with the presence of submicron holes within the particles. The XRD patterns of both feedstock powders conventional and nanocrystalline one are shown in Figs. 3a and 4a respectively. The analysis revealed slight differences on the phase composition of the two powders. Although the results show that the two powders are mainly composed of WC, the spectra of the conventional powder indicated a minor presence of other phases such as Co3W3C. (indicated by arrows in Fig. 3a). Note that, in the case of nanocrystalline powder the Co phase was identified while in the conventional one it was very hard to identify the binder phases. 3.2. Coating characterization The XRD analysis of the as-sprayed coatings produced by both processes, CGDS and PGDS, did not show the presence of other phases than those present in the feedstock powders. Indeed, neither the nanocrystalline coating (Fig. 4a′) nor the conventional one (Fig. 3a′ and a″) has exhibited any noticeable phase change during the PGDS process, when compared with
Fig. 4. (a, a′) XRD patterns of WC–15Co nanocrystalline cermet powder and coatings prepared by PGDS process at 673 K respectively. (b) XRD patterns of the nanocrystalline WC–12Co powder and coating deposited using HVOF process [31].
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Fig. 5. OM and SEM images showing the morphology of both the (a, b) conventional and (c, d) nanocrystalline coatings deposited by the CGDS process.
Fig. 6. OM and SEM images showing the morphology of both the (a, b) conventional and (c, d) nanocrystalline coatings deposited by the PGDS process.
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Fig. 7. Vickers microhardness graph showing the HV0.5 measured values for the four different types of coatings and compared to hardness of different cermets using different deposition process reported in recent scientific papers.
Nevertheless, the coatings seem to be denser compared to the ones produced using the conventional feedstock powder. Typical microstructures of PGDS coatings produced using the conventional (a and b) and the nanocrystalline (c and d) cermet powders are shown in Fig. 6. The images are taken from the coatings produced with the feedstock powder preheated to 673 K prior to spraying. In general, the microstructure of the coatings is very dense and exhibits no lamellar or banded microstructure, which is typical of thermally sprayed coatings. Low levels of porosity (less than 5%) were obtained for the conventional coatings, while the nanocrystalline coatings seemed to be denser (levels of porosity less than 2%). The carbide distribution in the conventional coatings is generally less uniform than in the nanocrystalline coatings, as observed in the feedstock powder. The conventional and nanocrystalline coatings did not show the presence of large cracks as it is the case in coatings prepared using the CGDS process. Nevertheless, micro-cracks were observed in all coatings, especially close to the substrate–coating interface. Vickers microhardness (HV0.5) measurements were made in the middle of the coating cross-section of each sample. The measured values are presented in Fig. 7, as well as other hardness values obtained from coatings prepared by other thermal spray processes [33–35] and presented here for comparison purposes. The measurements show that PGDS coatings have higher hardness values (721 ± 95 HV and 899 ± 64 HV), compared to the ones produced by the conventional CGDS process (356 ± 135 HV and 462 ± 92 HV), using the conventional and the nanocrystalline powders respectively. 4. Discussion The cold gas dynamic spraying process has recently shown to be a suitable technique in order to preserve the microstructure of the feedstock powder in the coating [24,36,37]. In com-
parison to other TS processes such as HVOF, where carbide decomposition and amorphization of the binder have been reported [13,16,32], the coatings produced in the present work by both the CGDS and PGDS processes did not exhibit any major degradation of the original carbide phase. Similar results were also obtained in the case of nanostructured WC–12Co prepared by CGDS [24]. This is attributed to the lower deposition temperatures of the feedstock particles compared to those usually found in conventional TS processes and the inert environment of both the CGDS and PGDS processes. Hard materials, such as WC-based composites, are known to be difficult to deposit using the CGDS process as very few successful attempts have been reported [24]. In the present work, when using the CGDS process, neither the conventional nor the nanocrystalline cermet has been forming dense and thick coatings, despite increasing the main gas temperature up to 823 K. Both conventional and nanocrystalline CGDS coatings exhibit a large number of cracks within their microstructures (Fig. 5). This is attributed to the limited impact energy absorption available through plastic deformation of the binder phase (usually present in a small volume fraction for cermet materials). Consequently, in order to absorb the high release of energy during impact and high level of localized stresses in the material, cracks initiate and propagate mainly in the binder phase resulting in a poor coating microstructure. It has been reported that for metal matrix composite (MMC), where the binder volume fraction is much higher than the one in cermets, the presence of crack during deposition can be avoided. In CGDS Al12 wt.%Si/SiC coatings [38] the impact energy is well absorbed by the aluminum matrix. Consequently, higher Co content in WC/Co cermets would be more suited for CGDS but it would likely be detrimental to its wear properties. In CGDS, successful bonding of impacting particles requires localized deformation and adiabatic shear instabilities, which
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occur at high impact velocity, the critical velocity (Vcr). The critical velocity of most metals and alloys for CGDS deposition were reported to be in the range of 500–700 m/s based on theoretical and experimental works [39,40]. The empirical based model [40] has revealed that the critical velocity can be predicted by: Vcr ¼ 667 0:014q þ 0:08ðTm TR Þ þ 10 7ru 0:4ðTi TR Þ Where, Vcr is the critical velocity, ρ is the density of the material, Tm is the melting temperature, TR is the reference temperature (293 K), σu is the yield stress, and Ti is the impact temperature. The above formula indicates that particle preheating temperature of 773 K will decrease the critical particle velocity by 200 m/s (in the case of WC–12Co) [24]. It was shown that it is possible to deposit and to build up thick nanostructured WC– 12Co coating by cold gas dynamic spraying, as long as the powder is preheated to a minimal temperature of 773 K. By using the PGDS process, thick and dense coatings were obtained as long as the feedstock powder was preheated above 573 K (Fig. 6 reports typical coatings obtained when using 673 K). Below a preheating temperature of 573 K, the feedstock powder mostly bounces off the substrate, with very limited adhesion. Consequently, it becomes clear that the particle impact temperature plays a significant role, as predicted by the above formula. Particle preheating will decrease the critical particle velocity since the ductility of the material increases as the material temperature is increased, The microhardness measurements (Fig. 7) show that the PGDS coatings using both types of powder resulted in higher hardness values than the ones produced using the CGDS process. This can be explained by the higher porosity observed in the latter coatings and also the presence of larger number of cracks within these coating as seen in Fig. 6. The hardness of the nanocrystalline coatings produced using the CGDS and PGDS processes was almost two fold the values recorded for the conventional coatings. This observation is in accordance to reported values for nanostructured cermets found in the literature [41]. Usually finer carbide particles result in higher hardness values in bulk samples; however, this trend can be reversed in thermal spray coatings when carbide decomposition occurs. In these cases, the decarburization is increased by the large surface area in the nanostructured samples due to the reactive environment in the thermal spray process [33–35]. Both the CGDS and PGDS coatings the hardness of the nanocrystalline coatings is higher than the conventional ones. This can be rationalized by the fact that the inert spray environment does not promote decarburization in the coatings. Also from Fig. 7 one can observe that the PGDS coatings are harder than the CGDS for both types of powder. This difference is explained by the low porosity in the coating achieved using the PGDS when compared to the CGDS ones. The higher particle impact temperature achieved in the PGDS process allows coating formation at lower particle velocities as a result of a lower critical deposition velocity. It also reduces the impact energy to be absorbed by the binder phase in the form of fracture
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surfaces due to its capacity to absorb the energy through plastic deformation at higher spraying temperatures. 5. Conclusion The production of cermet coatings with conventional and nanocrystalline feedstock powders by CGDS process and PGDS process were carried out in this study. It was found that: – There is no degradation of the phase (phase transformation and/or decarburization of WC) when using either process. – Dense coatings without major defects are difficult to obtain when using the CGDS process. – It is possible to produce coatings with low porosity using the PGDS process if the feedstock powder is preheated above 573 K. – The nanostructure of the feedstock powder is preserved during the CGDS and PGDS depositions. – Hard and dense coatings were obtained when depositing nanocrystalline powder using the PGDS process. Acknowledgements The authors wish to acknowledge the financial support of the Natural Sciences and Engineering Research Council of Canada for this work. The authors express their gratitude to Miss. Renata Zavadil from CANMET laboratory for many helpful technical discussions. References [1] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, J. Mater. Sci. 20 (1985) 1873. [2] H.L. de Villiers Lovelock, J.M. Benson, P.M. Young, J. Therm. Spray Technol. 7 (1998) 97. [3] L.-M. Berger, P. Ettmayer, P. Vuoristo, W. Kunert, J. Therm. Spray Technol. 10 (2) (2001) 311. [4] A. Kasrimi, C. Verdon, Surf. Coat. Technol. 57 (1993) 81. [5] D.G.F. O'Quigley, M.N. James, Int. J. Refract. Met. Hard Mater. 15 (1) (1997) 73. [6] K. Jia, T.E. Fischer, Wear 203 (1997) 310. [7] B.K. Kim, G.H. Ha, D.W. Lee, J. Mater. Process. Technol. 63 (1997) 317. [8] J. Nerz, B. Kushner, A. Rotolico, in: R.M. Yazici (Ed.), Protective Coatings: Processing and Characterization, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 135. [9] R. Schwetzke, H. Kreye, Proceedings of the 15th International Thermal Spray Conference, Nice, France, 1998, p. 187. [10] S. Usmani, S. Sampath, D.L. Houck, D. Lee, Tribol. Trans. 40 (1997) 470. [11] C.J. Li, A. Ohmori, Y. Harada, J. Mater. Sci. 31 (1996) 785. [12] H.L.D. Lovelock, J.M. Benson, P.M. Young, J. Therm. Spray Technol. 7 (1998) 97. [13] B.H. Kear, G. Skandan, R.k. Sadanji, Scr. Mater. 44 (2001) 1703. [14] S. Usmani, S. Sampath, H. Herman, J. Therm. Spray Technol., 7 (1998) 429. [15] D.A. Stewart, P.H. Shipway, D.G. McCartney, Acta Mater. 48 (2000) 1593. [16] J. He, Y. Liu, T.E. Fischer, E.J. Lavernia, Metall. Mater. Trans., A Phys. Metall. Mater. Sci. 33 (1) (2002) 145. [17] T.H. Van Sreenkiste, R.C. McCune, K.J. Barnett, Surf. Coat. Technol. 111 (1999) 62. [18] R.C. Dykhuizen, M.F. Smith, D.L. Gilmore, R.A. Neiser, X. Jiang, S. Sampath, J. Therm. Spray Technol. 8 (1999) 559. [19] A.P. Alkhimov, V.F. Kosarev, A.N. Papyrin, Sov. Phys. Dokl. 35 (1990) 1047.
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