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Microstructural and mechanical properties of advanced HVOF-sprayed WC-based cermet coatings
Surface & Coatings Technology 286 (2016) 95–102
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Microstructural and mechanical properties of advanced HVOF-sprayed WC-based cermet coatings S.M. Nahvi a,⁎, M. Jafari b a b
Steel Institute, Isfahan University of Technology, Isfahan 84156-83111, Iran Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
a r t i c l e
i n f o
Article history: Received 27 July 2015 Revised 20 October 2015 Accepted in revised form 9 December 2015 Available online 10 December 2015 Keywords: HVOF WC–NiMoCrFeCo WC–FeCrAl WC–Co Mechanical properties
a b s t r a c t The aim of this research was to investigate the microstructural and mechanical properties of WC–FeCrAl and WC–NiMoCrFeCo coatings deposited by high velocity oxygen fuel (HVOF) spraying. Microstructural characterizations of feedstock powders and coatings were carried out using X-ray diffractometry (XRD) and scanning electron microscopy (SEM) equipped with energy dispersive X-ray (EDX) analysis. Image analysis techniques were utilized to measure the porosity level of the coatings, the volume fraction of different phases and the carbide grain size in the powders and coatings. Microhardness and indentation fracture toughness measurements were executed to evaluate the mechanical properties of the coatings. For comparison, the same experiments were performed on conventional WC–Co coating deposited by similar HVOF spraying parameters. WC–NiMoCrFeCo coating showed the maximum W2C/WC peak ratio of 40.42% in comparison to WC–FeCrAl (12.48%) and WC– Co (9.14%) coatings indicating a larger extent of W2C phase precipitated during solidification of WC–NiMoCrFeCo coating. The highest porosity level of 5.1 vol% was observed in the case of WC–FeCrAl coating due to the lower temperature of the powder particles causing the FeCrAl matrix not to be fully melted. The microhardness of WC–FeCrAl coating was found to be 1498 HV0.3 indicating higher value as compared to WC–Co and WC–NiMoCrFeCo coatings with 1305 and 1254 HV0.3, respectively. Moreover, a mean fracture toughness of 5.9 MPam1/2 was obtained for WC–Co, which was substantially greater than that for both WC–FeCrAl and WC–NiMoCrFeCo coatings with 3.1 and 2.8 MPam1/2, respectively. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Most engineering materials used for applications in which abrasive wear resistance is a major requirement such as tool steels, white cast irons, cobalt-based alloys and metallic matrix composites are multiphase materials consisting of a metallic matrix reinforced by a dispersion of hard particles [1]. The reason for the success of this type of material in tribological applications can be explained in a simplified form by stating that the toughness of the matrix together with the hardness of the reinforcement particles provides optimal wear resistance. Due to their excellent combination of hardness, fracture toughness and wear resistance, ceramic–metal (cermet) composites are extensively used to produce wear resistant parts such as cutting, drilling and machining tools [2–4]. Among different types of cermet composites, WC-based cermets have frequently been used in industry to improve the wear resistance of machine parts [2,5–7]. The WC-based cermets combine the hard WC phase and a ductile metallic binder phase, normally cobalt, in different proportions to produce materials with a wide range of properties [8]. Other metallic or alloyed binders such as nickel ⁎ Corresponding author. E-mail address: [email protected] (S.M. Nahvi).
http://dx.doi.org/10.1016/j.surfcoat.2015.12.016 0257-8972/© 2015 Elsevier B.V. All rights reserved.
are used to increase the corrosion resistance of WC cermets [9]. The characteristic high hardness and fracture toughness of sintered WC cermets have made them ideal materials for abrasive wear resistant components in a variety of industrial applications. As an alternative technique to the bulk material fabrication, surface engineering aims to develop wear resistant coatings onto different substrates to benefit from their physical, chemical, mechanical and thermal properties. In this context, WC-based thermal spray coatings, mostly deposited by high velocity oxy-fuel (HVOF) spraying, are promising candidates for wear protection due to their excellent combination of high hardness and fracture toughness, low friction and chemical inertness [10–14]. The properties and performance of WC cermet coatings are attributed to a complex function of size, shape and distribution of carbides, composition and content of the metallic matrix, and also microstructural evolution during HVOF spraying. The investigation of the causal relationship between deposition process parameters, microstructure and wear resistance has shown that to achieve the optimal performance, the coating should have large extent of retained WC particles finely dispersed within the metallic matrix [15–19]. This depends essentially on the level of WC decarburization during HVOF leading to the formation of non-WC phases such as W2C, amorphous/nanocrystalline Co–W–C phase and complex carbides in
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Table 1 The details of the commercially produced powders. Powder
Manufacturer
Commercial designation
Powder type
Nominal size (μm)
WC–C (83–17 wt.%) WC–F (85–15 wt.%) WC–N (85–15 wt.%)
Sulzer Metco
Woka 3202
Agglomerated and sintered
−45 + 15
H.C.Starck
Amperit 618
Agglomerated and sintered
−45 + 15
H.C.Starck
Amperit 529
Agglomerated and sintered
−45 + 15
the coating microstructure [20–23]. Due to their high brittleness, these non-WC phases deteriorate the wear performance of the cermet coatings by decreasing the fracture toughness [24–27]. The extent to which WC decarburization takes place is a function of powder characteristics, flame temperature and particles velocity. In recent decades, the processing, properties and applications of the WC–Co coatings with different cobalt contents have been extensively studied. There is also a considerable work on the characterization and properties of thermally sprayed WC cermets with other binders such as CoCr [28,29] and Ni [30–33]. However, there is little information in the literature about the new WC cermet coatings with complicated alloyed binders, e.g. nickel-based and iron-based alloys. In this work, the microstructural and mechanical properties of WC–NiMoCrFeCo and WC–FeCrAl coatings, deposited using a Top Gun HVOF system, are investigated and compared with those of the conventional WC–Co coating. 2. Experimental 2.1. Materials In this study, three different WC-based cermet powders including WC–NiMoCrFeCo (denoted as WC–N), WC–FeCrAl (denoted as WC–F) and WC–Co (denoted as WC–C) were used as feedstock materials. The average WC grain size in WC–N, WC–F and WC–C powders were 0.7, 0.5 and 1.0, respectively. All powders were agglomerated and sintered spheroids with diameters in the range of 15–45 μm. Details of the powders provided by two manufacturers (H.C. Starck, Laufenburg, Germany and Sulzer Metco, Hattersheim, Germany) are given in Table 1. The measured chemical composition of the feedstock powders is given in Table 2. The particle size distribution of the powders was measured by a Malvern Mastersizer S (Malvern Instruments Ltd., Worcestershire, UK) laser particle size analyzer. The substrates used for coating deposition were plain-carbon steel (0.12% C, 0.7% Mn) sheets with hardness of 246 HV0.3 and a dimension of 59 × 25 × 3 mm. The substrates were cleaned and grit blasted with ~250 μm brown alumina just before the coating process in order to degrease and roughen the surface.
2.3. Microstructural characterization The phase composition of the powders and coatings was identified by X-ray diffraction (XRD) (Siemens D500 diffractometer, 40 kV, 25 mA) utilizing a monochromatic Cu Kα (λ = 0.15406 nm) radiation. The diffraction data were collected over a 2θ range of 30°–80° with a step size of 0.010° and 4 s dwell time per step. Microstructural examination of feedstock powders and as-sprayed coatings was carried out using a scanning electron microscope (SEM) (Philips XL30, FEI Ltd.) equipped with energy dispersive X-ray (EDX) analysis. All SEM investigations were performed at an accelerating voltage of 20 kV in both secondary electron (SE) and back-scattered electron (BSE) modes. Image analysis (IA) was performed on SEM/BSE images at a magnification of 2500× obtained from polished cross-sections of the coatings in order to measure the porosity level of the coatings. Image analysis software (ImageJ 1.41) was employed to identify and measure porosity. Ten images were recorded to calculate the mean pore volume fraction. The volume fraction of phases and the carbide grain size in the powders and coatings were estimated by the method of line analysis from BSE micrographs at magnifications in the range of 5000–10000×. Chemical analysis was executed on as-sprayed coatings by LSM Ltd. (London and Scandinavian Metallurgical Co. Limited, South Yorkshire, UK). Oxygen and carbon contents were determined by XRF-HSS (quantitative), while the other elements were determined by X-Ray Fluorescence (XRF) technique using a XRF-Uniquant (semiquantitative).
2.4. Mechanical properties evaluation Vickers microhardness of as-sprayed coatings was measured using a LECO M-400 microhardness tester under a load of 300 gf for dwell time of 15 s. The mean value of 10 indents taken along the mid-plane of the cross-section parallel to the coating/substrate interface was quoted as the coatings hardness. The fracture toughness of coatings was determined by an indentation method. Vickers indentation measurements were performed on the metallographically prepared cross-sections of the coating under a
2.2. HVOF spraying The feedstock powders were sprayed onto the substrates using a Praxair/UTP Top-Gun HVOF spray system with parameters listed in Table 3. Hydrogen and nitrogen were employed as the fuel and carrier gases, respectively, and the samples were cooled with compressed air jets during spraying.
Table 2 The measured compositions of the feedstock powders. Designation
WC–C WC–F WC–N
Composition (wt%) W
Ni
Mo
Cr
Fe
Co
Al
C
O
77.98 79.05 79.97
– – 8.47
– – 2.24
– 3.40 2.15
0.04 10.79 0.84
16.82 – 0.62
– 1.02 -
5.16 5.58 5.65
– 0.16 0.06
Table 3 Spray parameters employed for coating depositions. Spray parameter −1
O2 flow rate (l min ) Fuel gas (H)2 flow rate (l min−1) Carrier gas (N2) flow rate (l min−1) Spray distance (mm) Number of pass Length of pass (mm) Carousel diameter (mm) Substrate velocity (m s−1) Gun transverse speed (mm s−1) Coating time (s) Consumption of powder (g) Coating thickness (μm) Powder feed rate (g min−1)
WC–C
WC–F
WC–N
240 640 17 250 40 77 280 1 5 674 710 445 63
240 640 17 250 40 77 280 1 5 729 635 436 52
240 640 17 250 51 76 280 1 5 924 555 260 36
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load of 5 kgf. The length of the cracks parallel to the substrate/coating interface was measured from optical micrographs at a magnification of 400× using the image analysis software (ImageJ 1.41). For each coating, at least 35 indentations were conducted. The fracture toughness (Kc) of the coatings was calculated according to the Evans and Wilshaw model [34]: K C ¼ 0:079
P a3=2
log
4:5a c
ð1Þ
where P is the applied indentation load (N), a is the indentation half diagonal (m), and c is the crack length from the center of the indent (m). The recommended c/a ratio for valid use of this equation is 0.6 ≤ c/a b 4.5. Fig. 1. XRD patterns of (a) WC–C, (b) WC–F and (c) WC–N powders.
Fig. 2. SEM images from morphology and cross-section of (a,d) WC–C, (b,e) WC–F and (c,f) WC–N powders.
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3. Results and discussion 3.1. Characterization of powders The XRD patterns of feedstock powders (Fig. 1) show that the main peaks correspond to the WC phase for all powders. In addition, different crystal phases related to the binder of each powder are detected. These phases include Co (fcc) binder for WC–C powder, AlFe and Fe3W3C for WC–F powder with FeCrAl alloy binder, and NiCrFe and Mo for the WC–N powder with Ni-based Hastelloy binder. These patterns verify that in the powders there are no extra carbide phases to a level which is detectable by the XRD system. SEM images from morphology and cross-section of different feedstock powders are shown in Fig. 2. As illustrated in Figs. 2(a–c), the agglomerated and sintered particles for all powders have a spherical morphology and are highly porous, with large holes within the spherical particles. Fig. 3 shows the result of particle size distribution analysis for the powders measured using the laser diffractometry technique. The median sizes of particles (d50%) and the size distribution ranges (d5%–d95%) are also included in Fig. 3. It can be seen that there are small differences in the size range and median size of particles among the powders. These results indicate similarity of powders in morphology and particle size although they have different binder phase and/or carbide grain size. Figs. 2(d–f) show the cross-sectional BSE images of powder particles at high magnification. Two phases with different contrasts are visible: the WC particles with bright contrast and the metallic binder phases with dark contrast. It is also obvious that the carbide grains are completely surrounded by metallic binders. The blocky shape of WC grains in WC–C powder is clear (Fig. 2d), while in the case of WC–F and WC–N powders, WC grains with more rounded morphology can be observed (Fig. 2(e,f)); also, it is apparent that WC–C contains larger WC grains as compared to WC–F and WC–N powders. The grain size and volume fraction of WC in the powders, measured based on the cross-sectional BSE images utilizing the line analysis method, are given in Table 4. 3.2. Characterization of as-sprayed coatings 3.2.1. Microstructure Table 5 shows the chemical composition of as-sprayed coatings. It is clear that WC decarburization has occurred upon HVOF spraying for all materials designations. The minimum carbon loss is obtained for WC–F (16%), while WC–C and WC–N coatings undergo more severe
Table 4 Microstructural properties of WC–C, WC–F and WC–N powders and coatings. Designation
State
Carbide phase (vol.%)
WC size (μm)
Mean free path (μm)
(W2C/WC) ratio
WC–C
Powder Coating Powder Coating Powder Coating
67 55 77 58 73 59
1.0 0.8 0.5 0.4 0.7 0.6
– 0.68 – 0.31 – 0.4
– 9.14 – 12.48 – 40.42
WC–F WC–N
decarburization of 30% and 36%, respectively. In addition, there has been no significant oxygen pick-up for WC–C coating whereas a more pronounced oxygen pick-up is observed during spraying WC–N (0.14%) and WC–F (0.46%) coatings. Fig. 4 shows the comparative XRD patterns of the as-sprayed coatings. XRD peaks corresponding to the W2C and W phases formed during deposition are detected for all coatings but at different intensities. WC–C coating consists of lower W2C phase than the WC–F one, while the highest level of W2C is observed in case of WC–N coating. As tabulated in Table 4, the calculated W2C/WC peak ratio of WC–N (40.42%) is significantly higher than that of WC–F (12.48%) and WC–C (9.14%) coatings, signifying that the greatest fraction of W2C phase is formed in WC–N coating as compared to other coatings. There is also a broad diffraction halo between 2θ values of approximately 37 and 47° for each coating indicating the presence of an amorphous phase in the deposits. Although this broad halo exists for all coatings, it is narrower and less significant for WC–F and WC–N in contrast to WC–C coating. Besides, no crystalline peaks related to the binder phases can be observed on the XRD patterns of the coatings. Decomposition of WC phase during HVOF spraying can be described in the following stages: (a) Melting of binder phase: When the powder particles are exposed to the hot gas jet (with a temperature of ~2000 K), the temperature of particles is increased until the metallic binder phase reaches its melting point while the WC particles, with melting point of 3143 K, mostly remain in solid state. (b) Dissolution of WC in the binder phase: At this stage, WC begins to be dissolved in the liquid binder phase enriching the metallic matrix in W and C. As a result, the binder phase of sprayed coatings shows a wide range of compositions depending on the temperature the powder particles reach. (c) Decarburization: Due to the high temperature involved in the spraying process, oxygen diffuses quickly throughout the liquid phase and reacts with the dissolved C to form CO2. Carbon will be removed from the melt either by reaction with oxygen at the melt/gas interface or through oxygen diffusion into the rim of the molten particles, leading to CO formation. (d) Solidification: Final stage of decarburization is rapid solidification. It may occur during the particles flight either when they are near the substrate or when the particles impact on the substrate. When the temperature decreases, new phases precipitate due to the decreasing solubility of W in the binder phase. As the C has been partially removed due to the oxidation reactions in the previous stages, it is only possible to form new phases with a lower amount of carbon (W2C, W) and also nanocrystalline/ Table 5 Chemical composition of as-sprayed coatings for all material designations. Designation Composition (wt.%) W
Fig. 3. Size distribution analysis of the feedstock powders.
WC–C WC–F WC–N
Ni
Mo
Cr
Fe
Co
Al
C
O
Carbon loss (%)
79.20 – – – 0.04 17.09 – 3.61 0.06 30 79.44 – – 3.42 10.80 – 1.03 4.69 0.62 16 81.58 8.66 2.28 2.20 0.86 0.61 – 3.61 0.2 36
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Fig. 4. XRD patterns of (a) WC–C, (b) WC–F and (c) WC–N coatings.
amorphous phase of the binder containing the remained W and C. Carbon loss brings the melt composition much closer to the W2C and/or W phase field.
In addition, W2C phase can be formed through oxidation of WC particles located on the surface of starting powder because they are directly exposed to oxidizing HVOF flame [14]. Fig. 5 demonstrates the cross-sectional BSE images from microstructure of the sprayed coatings. These images exhibit a typical splat-like microstructure associated with thermal spraying with dark and bright contrast matrix layers corresponding to the regions of lower and higher mean atomic number, respectively. In the darker areas of the matrix, WC particles have retained their angular morphology representing a negligible WC dissolution into the matrix. In the brighter regions, WC
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particles have a more rounded appearance and are often partially or fully enclosed by an irregular-shaped W2C phase with brighter contrast indicating that WC particles act as efficient nucleation sites for W2C precipitation and growth [35,36]. These features associate with the part of cermet particles in which a larger extent of WC dissolution and decarburization has occurred due to the higher local temperature [14,24]. Therefore, the coating structure is made up of retained WC grains, which remain in solid state during spraying, new precipitated phases (WC, W2C and W) depending on the local composition of binder along with the a nanocrystalline/amorphous binder phase containing dissolved W and C. It is evident that bright shells surrounding WC grains are much more prevalent for WC–N coating (Fig. 5c,d), in agreement with XRD spectrum of this coating (Fig. 4c) representing the maximum W2C/WC peak ratio. Since the HVOF spray parameters for all coatings were nearly identical (see Table 3), the size of the powder particles becomes a key factor determining the level of decarburization; that is, the smaller the particle size, the higher the particles temperature during HVOF spraying [37]. Based on this explanation, it is proposed that the higher level of WC decarburization and W2C formation in WC–N coating arises from its smallest powder particle size (~32 μm), which causes the in-flight particles to experience higher temperatures during spraying leading to the greater WC dissolution into the binder phase. Another mechanism for larger extent of WC decarburization of WC–N coating in comparison to WC–C is attributed to the W and C solubility in different metallic binders. When the temperature is decreased upon solidification, new phases are formed due to decreasing the solubility of W and C in the liquid binder phase. Since the solubility of W and C in Ni is lower than that in Co binder, the new carbide phases including W2C can be more easily crystallized during solidification of WC–N coating. This is in good agreement with findings of Shaw et al. [38], who reported the larger extent of W2C and W formation for WC–NiAl coating than that for WC–C coating because of lower solubility of W and C in NiAl binder. In contrast, higher solubility of W and C in the Co binder after cooling results in the
Fig. 5. The cross-sectional BSE images from microstructure of (a) WC–C, (b) WC–F and (c,d) WC–N coatings.
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Fig. 6. Image analysis performed on the cross-sectional BSE images of (a) WC–C, (b) WC–F and (c) WC–N coatings.
formation of binder phase rich in W and C in a solution state; this causes the precipitation of new carbide phases to be insignificant, and also stimulates the amorphous and nanocrystalline Co–W–C phase formation. The average diameter of WC grains and the volume fraction of the binder were measured on the coatings cross-sections by line analysis technique. The mean free path of the binder (λ) was also measured from the BSE images using the following equation [39]: λ¼
ð1− f Þ NL
ð2Þ
where NL is the number of non-continuous grains intersected on a metallographic plane by a line of unit length and f is the volume fraction of the dispersed phase. For each coating, the results of 5 measurements were quoted to obtain an average value. Table 4 presents the results of the mean volume fraction of carbide phases, grain size of WC, the mean free path of the binders, measured by line analysis method, for different coatings. The results indicate that an almost similar carbide phase volume fraction in the range of 55–59% exists in all coatings. The WC–C coating exhibits the higher WC grain size (~ 0.8 μm) and mean free path (0.68 μm) in comparison to WC–F (WC grain size: ~0.4 μm, mean free path: 0.31 μm) and WC–N (WC grain size: ~ 0.6 μm, mean free path: 0.4 μm) coatings. The porosity percentage of the coatings was measured using image analysis of cross-sectional BSE images. Fig. 6 shows the typical results of image analysis performed on the coatings cross-section to determine the porosity level. Table 6 presents the volume fraction of the porosity along with the mean pore size of the coatings.
According to Table 6, the WC–F coating reveals the maximum porosity of 5.1 vol.%, while WC–N and WC–C coatings show lower porosity percentages of 2.2 and 1.8 vol.%, respectively. The minimum porosity of WC–C coating is due to a good wetting property of Co binder for WC grains compared to those of FeCrAl and NiMoCrFeCo binders [40–42]. The micrograph corresponding to WC–C coating (Fig. 5c) also shows evidences of the greater degree of particle melting and flow after impact on the substrate. On the other hand, the highest level of porosity occurring in the Fe coating can be interpreted by the following mechanism. Based on data given in Table 5, the lowest carbon loss corresponds to WC–F coating. This implies the lowest extent of W and C dissolution into the binder due to lower temperature of the in-flight particles. Under such a circumstance, the binder has no chance to be fully melted, and thereby the porosity significantly increases for WC–F coating. This is in consistence with XRD pattern of WC–F coating which indicates a lower extent of W2C compared to WC–N coating and also a narrower halo (representing only a minor content of amorphous phase in the deposits) in contrast to WC–C coatings.
Table 6 Porosity percentage and mechanical properties of WC–C, WC–F and WC–N coatings. Coating
Porosity (vol.%)
Mean pore size (μm)
Microhardness (HV0.3)
KIC (MPam1/2)
WC–C WC–F WC–N
1.8 5.1 2.2
0.36 0.25 0.30
1305 ± 71 1498 ± 82 1254 ± 38
5.9 ± 0.13 3.1 ± 0.23 2.8 ± 0.27
Fig. 7. Cumulative percentiles corresponding to the indentation fracture toughness of WC–C, WC–F and WC–N coatings.
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3.2.2. Mechanical properties The microhardness and fracture toughness of WC–C, WC–F and WC– N coatings are summarized in Table 6. The WC–F coating exhibits the maximum microhardness of 1498 HV0.3, while WC–C and WC–N coatings reveal the lower values of 1305 and 1254 HV0.3, respectively. The hardness of thermally sprayed WC cermet coatings depends on the volume fractions of new phases (e.g., W2C and W) and retained WC phase, hardness of binder phase and microstructural properties of the coatings such as porosity, mean free path of binder and WC grain size. The WC–N and WC–F coatings, which have almost identical volume fractions of carbide phases comprising both the retained WC and newly precipitated W2C (see Table 4), show a large difference in their hardness values. This can be explained by the lower carbon loss and W2C/WC ratio of WC–F coating resulting in the higher fraction of retained WC phase in comparison to WC–N coating. The important role of the retained WC phase in enhancing the hardness of thermally sprayed WC cermet coatings is well known and, thus, many efforts have been made to control the decomposition of the WC phase during the spraying process [43,44]. Usmani et al. [45] showed that increasing W2C content in microstructure of HVOFsprayed WC cermet coatings yield to a decrease in hardness. The formation of the harder and more brittle W2C phase surrounding the WC grains decreases their cohesion to the coating matrix consequently deteriorating the mechanical properties [46,47]. Fig. 7 shows the plots of the cumulative distribution of the fracture toughness of the coatings. The results clearly demonstrate the higher fracture toughness of WC–C coating with mean value of 5.9 MPam1/2 in comparison to both WC–F and WC–N coatings with 3.1 and 2.8 MPam1/2, respectively. In HVOF coatings, the cracks mostly propagate parallel to the substrate because of the weak inter-splat interfaces. As for HVOF-sprayed WC cermet coatings, the brittle regions including the amorphous metal matrix, embrittled through WC dissolution [46], and W2C phase are preferential paths for crack propagation [48]; also, the high level of porosity and weak cohesion between matrix-carbide interfaces facilitate the crack propagation, thereby decreasing the fracture toughness of the coatings [49]. Considering the above-mentioned factors, it can be concluded that the higher fracture toughness of WC–C coating results from the minor porosity, which provides a greater WC-matrix cohesion and stronger inter-splat interface, and the lower W2C/WC ratio of this coating with respect to WC–F and WC–N coatings. 4. Conclusions In this study, the microstructural and mechanical properties of advanced HVOF-sprayed WC–N and WC–F coatings were evaluated and compared to those of the conventional WC–C coating. The following conclusions can be drawn: (1) Based on carbon content analysis, a minimum carbon loss of 16% was obtained for WC–F, while WC–C and WC–N coatings suffered from more severe decarburization of 30% and 36%, respectively. (2) WC–N coating showed the highest W2C/WC peak ratio of 40.42% in comparison to WC–F (12.48%) and WC–C (9.14%) coatings representing the higher amounts of W2C phase precipitated during solidification of WC–N coating. (3) The WC–F coating revealed a maximum porosity of 5.1 vol.% due to the lower temperature of particles causing the FeCrAl matrix not to be fully melted. This was confirmed by chemical analysis results indicating the lowest extent of W and C dissolution into the FeCrAl binder because of the lower temperature of the inflight WC–F particles. (4) A maximum microhardness of 1498 HV0.3 was measured for WC–F coating, whereas WC–C and WC–N coatings revealed lower values of 1305 and 1254 HV0.3.
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(5) The results showed a mean fracture toughness of 5.9 MPam1/2 for WC–C which was substantially greater than that for both WC–F and WC–N coatings, respectively with 3.1 and 2.8 MPam1/2. The higher fracture toughness of WC–C coating results from the minor porosity and stronger inter-splat interface together with the lower W2C/WC ratio of this coating in comparison to WC–F and WC–N coatings.
References [1] R. Colaço, R. Vilar, A model for the abrasive wear of metallic matrix particlereinforced materials, Wear 254 (7–8) (2003) 625–634. [2] C.J. Li, G.J. Yang, Relationships between feedstock structure, particle parameter, coating deposition, microstructure and properties for thermally sprayed conventional and nanostructured WC–Co, Int. J. Refract. Met. Hard Mater. 39 (2013) 2–17. [3] G.Q. Shao, X.L. Duan, J.R. Xie, Sintering of nanocrystalline of WC–Co composite powder, Rev. Adv. Mater. Sci. 5 (2003) 281–286. [4] H.E. Exner, J. Gurland, A review of parameters influencing some of mechanical properties of tungsten carbide–cobalt alloys, Powder Metall. 13 (1970) 13–63. [5] W. Żórawski, The microstructure and tribological properties of liquid-fuel HVOF sprayed nanostructured WC–12Co coatings, Surf. Coat. Technol. 220 (2013) 276–281. [6] W.H. Gu, Y.S. Jeong, K. Kim, J.C. Kim, S.H. Son, S. Kim, Thermal oxidation behavior of WC–Co hard metal machining tool tip scraps, J. Mater. Process. Technol. 212 (2012) 1250–1256. [7] M.I. Dvornik, A.V. Zaytsev, Research of surfaces and interfaces increasing during planetary ball milling of nanostructured tungsten carbide/cobalt powder, Int. J. Refract. Met. Hard Mater. 36 (2013) 271–277. [8] J. Gurland, P. Bardzil, Relation of strength, composition, and grain size of sintered WC–Co alloys, Trans. Am. Inst. Min. Metall. Pet. Eng. 203 (2) (1955) 311–315. [9] M. Aristizabal, J.M. Sanchez, N. Rodriguez, F. Ibarreta, R. Martinez, Comparison of the oxidation behaviour of WC–Co and WC–Ni–Co–Cr cemented carbides, Corros. Sci. 53 (2011) 2754–2760. [10] R.J.K. Wood, Tribology of thermal sprayed WC–Co coatings, Int. J. Refract. Met. Hard Mater. 28 (2010) 82–94. [11] P. Chivavibul, M. Watanabe, S. Kuroda, K. Shinoda, Effects of carbide size and Co content on the microstructure and mechanical properties of HVOF-sprayed WC–Co coatings, Surf. Coat. Technol. 202 (2007) 509–521. [12] Z.G. Ban, L.L. Shaw, Characterization of thermal sprayed nanostructured WC–Co coatings derived from nanocrystalline WC–18 wt.%Co powders, J. Therm. Spray Technol. 12 (1) (2003) 112–119. [13] G. Barbezat, A.R. Nicol, A. Sickinger, Abrasion, erosion and scuffing resistance of carbide and oxide ceramic thermal sprayed coatings for different applications, Wear 162-164 (1993) 529–537. [14] J.M. Guilemany, M. De Paco, J. Nutting, J.R. Miguel, Characterization of the W2C phase formed during the high velocity oxygen fuel spraying of a WC + 12 Pct Co powder, Metall. Mater. Trans. A 30A (1999) 1913–1921. [15] A. Mateen, G.C. Saha, T.I. Khan, F.A. Khalid, Tribological behaviour of HVOF sprayed near-nanostructured and microstructured WC–17 wt.%Co coatings, Surf. Coat. Technol. 206 (2011) 1077–1084. [16] M. Aristizabal, N. Rodriguez, F. Ibarreta, R. Martinez, J.M. Sanchez, Liquid phase sintering and oxidation resistance of WC–Ni–Co–Cr cemented carbides, Int. J. Refract. Met. Hard Mater. 28 (2010) 516–522. [17] N. Birks, G.H. Meier, F.S. Pettit, Introduction to the High-Temperature Oxidation of Metals, second ed. Cambridge University Press, 2006. [18] J.E. Nerz, B.A. Kushner, A.J. Rotolico, Microstructural Evaluation of Tungsten CarbideCobalt Coatings. In 4th Nationaal Thermal Spray Conf. ASM, Pittsburgh, Pennsylvania, USA, 1992. [19] G.C. Saha, T.I. Khan, The corrosion and wear performance of microcrystalline WC10Co-4Cr and near-nanocrystalline WC-17Co high velocity oxy-fuel sprayed coatings on steel substrate, Metal. Mater. Trans. A 41A (2010) 3000–3009. [20] P. Chivavibul, M. Watanabe, S. Kuroda, J. Kawakita, M. Komatsu, K. Sato, J. Kitamura, Effects of particle strength of feedstock powders on properties of warm-sprayed WC–Co coatings, J. Therm. Spray Technol. 20 (5) (2011) 1098–1109. [21] P. Suresh Babu, B. Basu, G. Sundararajan, Processing–structure–property correlation and decarburization phenomenon in detonation sprayed WC–12Co coatings, Acta Mater. 56 (2008) 5012–5026. [22] Q. Yang, T. Senda, A. Ohmori, Effect of carbide size on microstructure and sliding wear behavior of HVOF-sprayed WC–12%Co coatings, Wear 254 (2003) 23–34. [23] C.J. Li, A. Ohmori, Y. Harada, Effect of powder structure on the structure of thermally sprayed WC–Co coatings, J. Mater. Sci. 31 (1996) 785–794. [24] J.M. Guilemany, S. Dosta, J. Nin, R. Miguel, Study of the properties of WC–Co nanostructured coatings sprayed by high-velocity oxyfuel, J. Therm. Spray Technol. 14 (2005) 405–413. [25] D.A. Stewart, P.H. Shipway, D.G. McCartney, Abrasive wear behaviour of conventional and nanocomposite HVOF-sprayed WC–Co coatings, Wear 789 (1999) 225–229. [26] Q. Yang, T. Senda, A. Hirose, Sliding wear behavior of WC–12%Co coatings at elevated temperatures, Surf. Coat. Technol. 200 (2006) 4208–4212. [27] Y. Ishikawa, J. Kawakita, S. Sawa, T. Itsukaichi, Y. Sakamoto, M. Takaya, S. Kuroda, Evaluation of corrosion and wear resistance of hard cermet coatings sprayed by using an improved HVOF process, J. Therm. Spray Technol. 14 (3) (2005) 384–390.
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[28] S.L. Liu, X.P. Zheng, G.Q. Geng, Influence of nano-WC–12Co powder addition in WC– 10Co–4Cr AC-HVAF sprayed coatings on wear and erosion behaviour, Wear 269 (2010) 362–367. [29] L. Thakur, N. Arora, R. Jayaganthan, R. Sood, An investigation on erosion behavior of HVOF sprayed WC–CoCr coatings, Appl. Surf. Sci. 258 (2011) 1225–1234. [30] J.M. Guilemany, J. Nutting, J.R. Miguel, Z. Dong, Microstructure characterization of WC–Ni coatings obtained by HVOF thermal spraying, Scr. Metall. Mater. 33 (1995) 55–61. [31] L.M. Berger, S. Saaro, T. Naumann, M. Wiener, V. Weihnacht, S. Thiele, J. Suchánek, Microstructure and properties of HVOF-sprayed chromium alloyed WC–Co and WC–Ni coatings, Surf. Coat. Technol. 202 (2008) 4417–4421. [32] E. Celik, O. Culha, B. Uyulgan, N.F. Ak Azem, I. Ozdemir, A. Turk, Assessment of microstructural and mechanical properties of HVOF sprayed WC-based cermet coatings for a roller cylinder, Surf. Coat. Technol. 200 (2006) 4320–4328. [33] V.V. Sobolev, J.M. Guilemany, J.R. Miguel, J.A. Calero, Influence of thermal processes on coating formation during high velocity oxy-fuel (HVOF) spraying of WC–Ni powder particles, Surf. Coat. Technol. 82 (1996) 121–129. [34] A.G. Evans, T.R. Wilshaw, Quasi-static solid particle damage in brittle solids: observations analysis and implications, Acta Metall. 24 (1976) 939–956. [35] J. He, M. Ice, S. Dallek, E.J. Lavernia, Synthesis of nanostructured WC–12 Pct Co coating using mechanical milling and high velocity oxygen fuel thermal spraying, Metall. Mater. Trans. A 31A (2000) 541–553. [36] H. Chen, G. Gou, M. Tu, Y. Liu, Research on the friction and wear behavior at elevated temperature of plasma-sprayed nanostructured WC–Co coatings, J. Mater. Eng. Perform. 19 (1) (2010) 1–6. [37] L. Zhao, M. Maurer, F. Fischer, R. Dicks, E. Lugscheider, Influence of spray parameters on the particle in-flight properties and the properties of HVOF coating of WC–CoCr, Wear 257 (1–2) (2004) 41–46. [38] K.G. Shaw, M.F. Gruninger, W.J. Jarosinski, High Temperature Intermetallic Binders for HVOF Carbides, in Thermal Spray Industrial Applications, 1994.
[39] R.L. Fullman, Measurement of particle sizes in opaque bodies, Trans. Am. Inst. Min. Metall. Pet. Eng. 197 (3) (1953) 447–452. [40] H.J. Scussel, Friction and wear of cemented carbides, in: S.D. Henry (Ed.), ASM Handbook, ASM International 1992, pp. 795–800. [41] H.L. De Villiers Lovelock, Powder/processing/structure relationships in WC–Co thermal spray coatings: a review of the published literature, J. Therm. Spray Technol. 7 (3) (1998) 357–373. [42] H.J. Scussel, Friction and wear of cemented carbides, ASM Handbook, ASM International, 1992. [43] G. Di Girolamo, L. Pilloni, G. Pulci, F. Marra, Tribological characterization of WC–Co plasma sprayed coatings, J. Am. Ceram. Soc. 92 (5) (2009) 1118–1124. [44] E. Sánchez, E. Bannier, M. Salvador, V. Bonache, J. García, J. Morgiel, J. Grzonka, Microstructure and wear behavior of conventional and nanostructured plasmasprayed WC–Co coatings, J. Therm. Spray Technol. 19 (5) (2009) 964–974. [45] S. Usmani, S. Sampath, D.L. Houck, D. Lee, Effect of carbide grain size on the sliding and abrasive wear behavior of thermally sprayed WC–Co coatings, Tribol. T. 40 (3) (1997) 470–478. [46] T. Sudaprasert, P.H. Shipway, D.G. McCartney, Sliding wear behaviour of HVOF sprayed WC–Co coatings deposited with both gas-fuelled and liquid-fuelled systems, Wear 255 (2003) 943–949. [47] J.A. Picas, M. Punset, M.T. Baile, E. Martin, A. Forn, Properties of WC–CoCr based coatings deposited by different HVOF thermal spray processes, Plasma Process. Polym. 6 (2009) S948–S953. [48] B.R. Marple, R.S. Lima, Process temperature/velocity-hardness-wear relationships for high-velocity oxyfuel-sprayed nanostructured and conventional cermet coatings, J. Therm. Spray Technol. 14 (2005) 67–76. [49] M. Rodriguez, J. Klisans, L. Bavaresco, A. Scagni, F. Arenas, Wear resistance of HVOF sprayed carbide coatings, Thermal Spray: New Surfaces for a New Millennium, 2001.
Contents lists available at ScienceDirect
Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
Microstructural and mechanical properties of advanced HVOF-sprayed WC-based cermet coatings S.M. Nahvi a,⁎, M. Jafari b a b
Steel Institute, Isfahan University of Technology, Isfahan 84156-83111, Iran Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
a r t i c l e
i n f o
Article history: Received 27 July 2015 Revised 20 October 2015 Accepted in revised form 9 December 2015 Available online 10 December 2015 Keywords: HVOF WC–NiMoCrFeCo WC–FeCrAl WC–Co Mechanical properties
a b s t r a c t The aim of this research was to investigate the microstructural and mechanical properties of WC–FeCrAl and WC–NiMoCrFeCo coatings deposited by high velocity oxygen fuel (HVOF) spraying. Microstructural characterizations of feedstock powders and coatings were carried out using X-ray diffractometry (XRD) and scanning electron microscopy (SEM) equipped with energy dispersive X-ray (EDX) analysis. Image analysis techniques were utilized to measure the porosity level of the coatings, the volume fraction of different phases and the carbide grain size in the powders and coatings. Microhardness and indentation fracture toughness measurements were executed to evaluate the mechanical properties of the coatings. For comparison, the same experiments were performed on conventional WC–Co coating deposited by similar HVOF spraying parameters. WC–NiMoCrFeCo coating showed the maximum W2C/WC peak ratio of 40.42% in comparison to WC–FeCrAl (12.48%) and WC– Co (9.14%) coatings indicating a larger extent of W2C phase precipitated during solidification of WC–NiMoCrFeCo coating. The highest porosity level of 5.1 vol% was observed in the case of WC–FeCrAl coating due to the lower temperature of the powder particles causing the FeCrAl matrix not to be fully melted. The microhardness of WC–FeCrAl coating was found to be 1498 HV0.3 indicating higher value as compared to WC–Co and WC–NiMoCrFeCo coatings with 1305 and 1254 HV0.3, respectively. Moreover, a mean fracture toughness of 5.9 MPam1/2 was obtained for WC–Co, which was substantially greater than that for both WC–FeCrAl and WC–NiMoCrFeCo coatings with 3.1 and 2.8 MPam1/2, respectively. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Most engineering materials used for applications in which abrasive wear resistance is a major requirement such as tool steels, white cast irons, cobalt-based alloys and metallic matrix composites are multiphase materials consisting of a metallic matrix reinforced by a dispersion of hard particles [1]. The reason for the success of this type of material in tribological applications can be explained in a simplified form by stating that the toughness of the matrix together with the hardness of the reinforcement particles provides optimal wear resistance. Due to their excellent combination of hardness, fracture toughness and wear resistance, ceramic–metal (cermet) composites are extensively used to produce wear resistant parts such as cutting, drilling and machining tools [2–4]. Among different types of cermet composites, WC-based cermets have frequently been used in industry to improve the wear resistance of machine parts [2,5–7]. The WC-based cermets combine the hard WC phase and a ductile metallic binder phase, normally cobalt, in different proportions to produce materials with a wide range of properties [8]. Other metallic or alloyed binders such as nickel ⁎ Corresponding author. E-mail address: [email protected] (S.M. Nahvi).
http://dx.doi.org/10.1016/j.surfcoat.2015.12.016 0257-8972/© 2015 Elsevier B.V. All rights reserved.
are used to increase the corrosion resistance of WC cermets [9]. The characteristic high hardness and fracture toughness of sintered WC cermets have made them ideal materials for abrasive wear resistant components in a variety of industrial applications. As an alternative technique to the bulk material fabrication, surface engineering aims to develop wear resistant coatings onto different substrates to benefit from their physical, chemical, mechanical and thermal properties. In this context, WC-based thermal spray coatings, mostly deposited by high velocity oxy-fuel (HVOF) spraying, are promising candidates for wear protection due to their excellent combination of high hardness and fracture toughness, low friction and chemical inertness [10–14]. The properties and performance of WC cermet coatings are attributed to a complex function of size, shape and distribution of carbides, composition and content of the metallic matrix, and also microstructural evolution during HVOF spraying. The investigation of the causal relationship between deposition process parameters, microstructure and wear resistance has shown that to achieve the optimal performance, the coating should have large extent of retained WC particles finely dispersed within the metallic matrix [15–19]. This depends essentially on the level of WC decarburization during HVOF leading to the formation of non-WC phases such as W2C, amorphous/nanocrystalline Co–W–C phase and complex carbides in
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Table 1 The details of the commercially produced powders. Powder
Manufacturer
Commercial designation
Powder type
Nominal size (μm)
WC–C (83–17 wt.%) WC–F (85–15 wt.%) WC–N (85–15 wt.%)
Sulzer Metco
Woka 3202
Agglomerated and sintered
−45 + 15
H.C.Starck
Amperit 618
Agglomerated and sintered
−45 + 15
H.C.Starck
Amperit 529
Agglomerated and sintered
−45 + 15
the coating microstructure [20–23]. Due to their high brittleness, these non-WC phases deteriorate the wear performance of the cermet coatings by decreasing the fracture toughness [24–27]. The extent to which WC decarburization takes place is a function of powder characteristics, flame temperature and particles velocity. In recent decades, the processing, properties and applications of the WC–Co coatings with different cobalt contents have been extensively studied. There is also a considerable work on the characterization and properties of thermally sprayed WC cermets with other binders such as CoCr [28,29] and Ni [30–33]. However, there is little information in the literature about the new WC cermet coatings with complicated alloyed binders, e.g. nickel-based and iron-based alloys. In this work, the microstructural and mechanical properties of WC–NiMoCrFeCo and WC–FeCrAl coatings, deposited using a Top Gun HVOF system, are investigated and compared with those of the conventional WC–Co coating. 2. Experimental 2.1. Materials In this study, three different WC-based cermet powders including WC–NiMoCrFeCo (denoted as WC–N), WC–FeCrAl (denoted as WC–F) and WC–Co (denoted as WC–C) were used as feedstock materials. The average WC grain size in WC–N, WC–F and WC–C powders were 0.7, 0.5 and 1.0, respectively. All powders were agglomerated and sintered spheroids with diameters in the range of 15–45 μm. Details of the powders provided by two manufacturers (H.C. Starck, Laufenburg, Germany and Sulzer Metco, Hattersheim, Germany) are given in Table 1. The measured chemical composition of the feedstock powders is given in Table 2. The particle size distribution of the powders was measured by a Malvern Mastersizer S (Malvern Instruments Ltd., Worcestershire, UK) laser particle size analyzer. The substrates used for coating deposition were plain-carbon steel (0.12% C, 0.7% Mn) sheets with hardness of 246 HV0.3 and a dimension of 59 × 25 × 3 mm. The substrates were cleaned and grit blasted with ~250 μm brown alumina just before the coating process in order to degrease and roughen the surface.
2.3. Microstructural characterization The phase composition of the powders and coatings was identified by X-ray diffraction (XRD) (Siemens D500 diffractometer, 40 kV, 25 mA) utilizing a monochromatic Cu Kα (λ = 0.15406 nm) radiation. The diffraction data were collected over a 2θ range of 30°–80° with a step size of 0.010° and 4 s dwell time per step. Microstructural examination of feedstock powders and as-sprayed coatings was carried out using a scanning electron microscope (SEM) (Philips XL30, FEI Ltd.) equipped with energy dispersive X-ray (EDX) analysis. All SEM investigations were performed at an accelerating voltage of 20 kV in both secondary electron (SE) and back-scattered electron (BSE) modes. Image analysis (IA) was performed on SEM/BSE images at a magnification of 2500× obtained from polished cross-sections of the coatings in order to measure the porosity level of the coatings. Image analysis software (ImageJ 1.41) was employed to identify and measure porosity. Ten images were recorded to calculate the mean pore volume fraction. The volume fraction of phases and the carbide grain size in the powders and coatings were estimated by the method of line analysis from BSE micrographs at magnifications in the range of 5000–10000×. Chemical analysis was executed on as-sprayed coatings by LSM Ltd. (London and Scandinavian Metallurgical Co. Limited, South Yorkshire, UK). Oxygen and carbon contents were determined by XRF-HSS (quantitative), while the other elements were determined by X-Ray Fluorescence (XRF) technique using a XRF-Uniquant (semiquantitative).
2.4. Mechanical properties evaluation Vickers microhardness of as-sprayed coatings was measured using a LECO M-400 microhardness tester under a load of 300 gf for dwell time of 15 s. The mean value of 10 indents taken along the mid-plane of the cross-section parallel to the coating/substrate interface was quoted as the coatings hardness. The fracture toughness of coatings was determined by an indentation method. Vickers indentation measurements were performed on the metallographically prepared cross-sections of the coating under a
2.2. HVOF spraying The feedstock powders were sprayed onto the substrates using a Praxair/UTP Top-Gun HVOF spray system with parameters listed in Table 3. Hydrogen and nitrogen were employed as the fuel and carrier gases, respectively, and the samples were cooled with compressed air jets during spraying.
Table 2 The measured compositions of the feedstock powders. Designation
WC–C WC–F WC–N
Composition (wt%) W
Ni
Mo
Cr
Fe
Co
Al
C
O
77.98 79.05 79.97
– – 8.47
– – 2.24
– 3.40 2.15
0.04 10.79 0.84
16.82 – 0.62
– 1.02 -
5.16 5.58 5.65
– 0.16 0.06
Table 3 Spray parameters employed for coating depositions. Spray parameter −1
O2 flow rate (l min ) Fuel gas (H)2 flow rate (l min−1) Carrier gas (N2) flow rate (l min−1) Spray distance (mm) Number of pass Length of pass (mm) Carousel diameter (mm) Substrate velocity (m s−1) Gun transverse speed (mm s−1) Coating time (s) Consumption of powder (g) Coating thickness (μm) Powder feed rate (g min−1)
WC–C
WC–F
WC–N
240 640 17 250 40 77 280 1 5 674 710 445 63
240 640 17 250 40 77 280 1 5 729 635 436 52
240 640 17 250 51 76 280 1 5 924 555 260 36
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load of 5 kgf. The length of the cracks parallel to the substrate/coating interface was measured from optical micrographs at a magnification of 400× using the image analysis software (ImageJ 1.41). For each coating, at least 35 indentations were conducted. The fracture toughness (Kc) of the coatings was calculated according to the Evans and Wilshaw model [34]: K C ¼ 0:079
P a3=2
log
4:5a c
ð1Þ
where P is the applied indentation load (N), a is the indentation half diagonal (m), and c is the crack length from the center of the indent (m). The recommended c/a ratio for valid use of this equation is 0.6 ≤ c/a b 4.5. Fig. 1. XRD patterns of (a) WC–C, (b) WC–F and (c) WC–N powders.
Fig. 2. SEM images from morphology and cross-section of (a,d) WC–C, (b,e) WC–F and (c,f) WC–N powders.
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3. Results and discussion 3.1. Characterization of powders The XRD patterns of feedstock powders (Fig. 1) show that the main peaks correspond to the WC phase for all powders. In addition, different crystal phases related to the binder of each powder are detected. These phases include Co (fcc) binder for WC–C powder, AlFe and Fe3W3C for WC–F powder with FeCrAl alloy binder, and NiCrFe and Mo for the WC–N powder with Ni-based Hastelloy binder. These patterns verify that in the powders there are no extra carbide phases to a level which is detectable by the XRD system. SEM images from morphology and cross-section of different feedstock powders are shown in Fig. 2. As illustrated in Figs. 2(a–c), the agglomerated and sintered particles for all powders have a spherical morphology and are highly porous, with large holes within the spherical particles. Fig. 3 shows the result of particle size distribution analysis for the powders measured using the laser diffractometry technique. The median sizes of particles (d50%) and the size distribution ranges (d5%–d95%) are also included in Fig. 3. It can be seen that there are small differences in the size range and median size of particles among the powders. These results indicate similarity of powders in morphology and particle size although they have different binder phase and/or carbide grain size. Figs. 2(d–f) show the cross-sectional BSE images of powder particles at high magnification. Two phases with different contrasts are visible: the WC particles with bright contrast and the metallic binder phases with dark contrast. It is also obvious that the carbide grains are completely surrounded by metallic binders. The blocky shape of WC grains in WC–C powder is clear (Fig. 2d), while in the case of WC–F and WC–N powders, WC grains with more rounded morphology can be observed (Fig. 2(e,f)); also, it is apparent that WC–C contains larger WC grains as compared to WC–F and WC–N powders. The grain size and volume fraction of WC in the powders, measured based on the cross-sectional BSE images utilizing the line analysis method, are given in Table 4. 3.2. Characterization of as-sprayed coatings 3.2.1. Microstructure Table 5 shows the chemical composition of as-sprayed coatings. It is clear that WC decarburization has occurred upon HVOF spraying for all materials designations. The minimum carbon loss is obtained for WC–F (16%), while WC–C and WC–N coatings undergo more severe
Table 4 Microstructural properties of WC–C, WC–F and WC–N powders and coatings. Designation
State
Carbide phase (vol.%)
WC size (μm)
Mean free path (μm)
(W2C/WC) ratio
WC–C
Powder Coating Powder Coating Powder Coating
67 55 77 58 73 59
1.0 0.8 0.5 0.4 0.7 0.6
– 0.68 – 0.31 – 0.4
– 9.14 – 12.48 – 40.42
WC–F WC–N
decarburization of 30% and 36%, respectively. In addition, there has been no significant oxygen pick-up for WC–C coating whereas a more pronounced oxygen pick-up is observed during spraying WC–N (0.14%) and WC–F (0.46%) coatings. Fig. 4 shows the comparative XRD patterns of the as-sprayed coatings. XRD peaks corresponding to the W2C and W phases formed during deposition are detected for all coatings but at different intensities. WC–C coating consists of lower W2C phase than the WC–F one, while the highest level of W2C is observed in case of WC–N coating. As tabulated in Table 4, the calculated W2C/WC peak ratio of WC–N (40.42%) is significantly higher than that of WC–F (12.48%) and WC–C (9.14%) coatings, signifying that the greatest fraction of W2C phase is formed in WC–N coating as compared to other coatings. There is also a broad diffraction halo between 2θ values of approximately 37 and 47° for each coating indicating the presence of an amorphous phase in the deposits. Although this broad halo exists for all coatings, it is narrower and less significant for WC–F and WC–N in contrast to WC–C coating. Besides, no crystalline peaks related to the binder phases can be observed on the XRD patterns of the coatings. Decomposition of WC phase during HVOF spraying can be described in the following stages: (a) Melting of binder phase: When the powder particles are exposed to the hot gas jet (with a temperature of ~2000 K), the temperature of particles is increased until the metallic binder phase reaches its melting point while the WC particles, with melting point of 3143 K, mostly remain in solid state. (b) Dissolution of WC in the binder phase: At this stage, WC begins to be dissolved in the liquid binder phase enriching the metallic matrix in W and C. As a result, the binder phase of sprayed coatings shows a wide range of compositions depending on the temperature the powder particles reach. (c) Decarburization: Due to the high temperature involved in the spraying process, oxygen diffuses quickly throughout the liquid phase and reacts with the dissolved C to form CO2. Carbon will be removed from the melt either by reaction with oxygen at the melt/gas interface or through oxygen diffusion into the rim of the molten particles, leading to CO formation. (d) Solidification: Final stage of decarburization is rapid solidification. It may occur during the particles flight either when they are near the substrate or when the particles impact on the substrate. When the temperature decreases, new phases precipitate due to the decreasing solubility of W in the binder phase. As the C has been partially removed due to the oxidation reactions in the previous stages, it is only possible to form new phases with a lower amount of carbon (W2C, W) and also nanocrystalline/ Table 5 Chemical composition of as-sprayed coatings for all material designations. Designation Composition (wt.%) W
Fig. 3. Size distribution analysis of the feedstock powders.
WC–C WC–F WC–N
Ni
Mo
Cr
Fe
Co
Al
C
O
Carbon loss (%)
79.20 – – – 0.04 17.09 – 3.61 0.06 30 79.44 – – 3.42 10.80 – 1.03 4.69 0.62 16 81.58 8.66 2.28 2.20 0.86 0.61 – 3.61 0.2 36
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Fig. 4. XRD patterns of (a) WC–C, (b) WC–F and (c) WC–N coatings.
amorphous phase of the binder containing the remained W and C. Carbon loss brings the melt composition much closer to the W2C and/or W phase field.
In addition, W2C phase can be formed through oxidation of WC particles located on the surface of starting powder because they are directly exposed to oxidizing HVOF flame [14]. Fig. 5 demonstrates the cross-sectional BSE images from microstructure of the sprayed coatings. These images exhibit a typical splat-like microstructure associated with thermal spraying with dark and bright contrast matrix layers corresponding to the regions of lower and higher mean atomic number, respectively. In the darker areas of the matrix, WC particles have retained their angular morphology representing a negligible WC dissolution into the matrix. In the brighter regions, WC
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particles have a more rounded appearance and are often partially or fully enclosed by an irregular-shaped W2C phase with brighter contrast indicating that WC particles act as efficient nucleation sites for W2C precipitation and growth [35,36]. These features associate with the part of cermet particles in which a larger extent of WC dissolution and decarburization has occurred due to the higher local temperature [14,24]. Therefore, the coating structure is made up of retained WC grains, which remain in solid state during spraying, new precipitated phases (WC, W2C and W) depending on the local composition of binder along with the a nanocrystalline/amorphous binder phase containing dissolved W and C. It is evident that bright shells surrounding WC grains are much more prevalent for WC–N coating (Fig. 5c,d), in agreement with XRD spectrum of this coating (Fig. 4c) representing the maximum W2C/WC peak ratio. Since the HVOF spray parameters for all coatings were nearly identical (see Table 3), the size of the powder particles becomes a key factor determining the level of decarburization; that is, the smaller the particle size, the higher the particles temperature during HVOF spraying [37]. Based on this explanation, it is proposed that the higher level of WC decarburization and W2C formation in WC–N coating arises from its smallest powder particle size (~32 μm), which causes the in-flight particles to experience higher temperatures during spraying leading to the greater WC dissolution into the binder phase. Another mechanism for larger extent of WC decarburization of WC–N coating in comparison to WC–C is attributed to the W and C solubility in different metallic binders. When the temperature is decreased upon solidification, new phases are formed due to decreasing the solubility of W and C in the liquid binder phase. Since the solubility of W and C in Ni is lower than that in Co binder, the new carbide phases including W2C can be more easily crystallized during solidification of WC–N coating. This is in good agreement with findings of Shaw et al. [38], who reported the larger extent of W2C and W formation for WC–NiAl coating than that for WC–C coating because of lower solubility of W and C in NiAl binder. In contrast, higher solubility of W and C in the Co binder after cooling results in the
Fig. 5. The cross-sectional BSE images from microstructure of (a) WC–C, (b) WC–F and (c,d) WC–N coatings.
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Fig. 6. Image analysis performed on the cross-sectional BSE images of (a) WC–C, (b) WC–F and (c) WC–N coatings.
formation of binder phase rich in W and C in a solution state; this causes the precipitation of new carbide phases to be insignificant, and also stimulates the amorphous and nanocrystalline Co–W–C phase formation. The average diameter of WC grains and the volume fraction of the binder were measured on the coatings cross-sections by line analysis technique. The mean free path of the binder (λ) was also measured from the BSE images using the following equation [39]: λ¼
ð1− f Þ NL
ð2Þ
where NL is the number of non-continuous grains intersected on a metallographic plane by a line of unit length and f is the volume fraction of the dispersed phase. For each coating, the results of 5 measurements were quoted to obtain an average value. Table 4 presents the results of the mean volume fraction of carbide phases, grain size of WC, the mean free path of the binders, measured by line analysis method, for different coatings. The results indicate that an almost similar carbide phase volume fraction in the range of 55–59% exists in all coatings. The WC–C coating exhibits the higher WC grain size (~ 0.8 μm) and mean free path (0.68 μm) in comparison to WC–F (WC grain size: ~0.4 μm, mean free path: 0.31 μm) and WC–N (WC grain size: ~ 0.6 μm, mean free path: 0.4 μm) coatings. The porosity percentage of the coatings was measured using image analysis of cross-sectional BSE images. Fig. 6 shows the typical results of image analysis performed on the coatings cross-section to determine the porosity level. Table 6 presents the volume fraction of the porosity along with the mean pore size of the coatings.
According to Table 6, the WC–F coating reveals the maximum porosity of 5.1 vol.%, while WC–N and WC–C coatings show lower porosity percentages of 2.2 and 1.8 vol.%, respectively. The minimum porosity of WC–C coating is due to a good wetting property of Co binder for WC grains compared to those of FeCrAl and NiMoCrFeCo binders [40–42]. The micrograph corresponding to WC–C coating (Fig. 5c) also shows evidences of the greater degree of particle melting and flow after impact on the substrate. On the other hand, the highest level of porosity occurring in the Fe coating can be interpreted by the following mechanism. Based on data given in Table 5, the lowest carbon loss corresponds to WC–F coating. This implies the lowest extent of W and C dissolution into the binder due to lower temperature of the in-flight particles. Under such a circumstance, the binder has no chance to be fully melted, and thereby the porosity significantly increases for WC–F coating. This is in consistence with XRD pattern of WC–F coating which indicates a lower extent of W2C compared to WC–N coating and also a narrower halo (representing only a minor content of amorphous phase in the deposits) in contrast to WC–C coatings.
Table 6 Porosity percentage and mechanical properties of WC–C, WC–F and WC–N coatings. Coating
Porosity (vol.%)
Mean pore size (μm)
Microhardness (HV0.3)
KIC (MPam1/2)
WC–C WC–F WC–N
1.8 5.1 2.2
0.36 0.25 0.30
1305 ± 71 1498 ± 82 1254 ± 38
5.9 ± 0.13 3.1 ± 0.23 2.8 ± 0.27
Fig. 7. Cumulative percentiles corresponding to the indentation fracture toughness of WC–C, WC–F and WC–N coatings.
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3.2.2. Mechanical properties The microhardness and fracture toughness of WC–C, WC–F and WC– N coatings are summarized in Table 6. The WC–F coating exhibits the maximum microhardness of 1498 HV0.3, while WC–C and WC–N coatings reveal the lower values of 1305 and 1254 HV0.3, respectively. The hardness of thermally sprayed WC cermet coatings depends on the volume fractions of new phases (e.g., W2C and W) and retained WC phase, hardness of binder phase and microstructural properties of the coatings such as porosity, mean free path of binder and WC grain size. The WC–N and WC–F coatings, which have almost identical volume fractions of carbide phases comprising both the retained WC and newly precipitated W2C (see Table 4), show a large difference in their hardness values. This can be explained by the lower carbon loss and W2C/WC ratio of WC–F coating resulting in the higher fraction of retained WC phase in comparison to WC–N coating. The important role of the retained WC phase in enhancing the hardness of thermally sprayed WC cermet coatings is well known and, thus, many efforts have been made to control the decomposition of the WC phase during the spraying process [43,44]. Usmani et al. [45] showed that increasing W2C content in microstructure of HVOFsprayed WC cermet coatings yield to a decrease in hardness. The formation of the harder and more brittle W2C phase surrounding the WC grains decreases their cohesion to the coating matrix consequently deteriorating the mechanical properties [46,47]. Fig. 7 shows the plots of the cumulative distribution of the fracture toughness of the coatings. The results clearly demonstrate the higher fracture toughness of WC–C coating with mean value of 5.9 MPam1/2 in comparison to both WC–F and WC–N coatings with 3.1 and 2.8 MPam1/2, respectively. In HVOF coatings, the cracks mostly propagate parallel to the substrate because of the weak inter-splat interfaces. As for HVOF-sprayed WC cermet coatings, the brittle regions including the amorphous metal matrix, embrittled through WC dissolution [46], and W2C phase are preferential paths for crack propagation [48]; also, the high level of porosity and weak cohesion between matrix-carbide interfaces facilitate the crack propagation, thereby decreasing the fracture toughness of the coatings [49]. Considering the above-mentioned factors, it can be concluded that the higher fracture toughness of WC–C coating results from the minor porosity, which provides a greater WC-matrix cohesion and stronger inter-splat interface, and the lower W2C/WC ratio of this coating with respect to WC–F and WC–N coatings. 4. Conclusions In this study, the microstructural and mechanical properties of advanced HVOF-sprayed WC–N and WC–F coatings were evaluated and compared to those of the conventional WC–C coating. The following conclusions can be drawn: (1) Based on carbon content analysis, a minimum carbon loss of 16% was obtained for WC–F, while WC–C and WC–N coatings suffered from more severe decarburization of 30% and 36%, respectively. (2) WC–N coating showed the highest W2C/WC peak ratio of 40.42% in comparison to WC–F (12.48%) and WC–C (9.14%) coatings representing the higher amounts of W2C phase precipitated during solidification of WC–N coating. (3) The WC–F coating revealed a maximum porosity of 5.1 vol.% due to the lower temperature of particles causing the FeCrAl matrix not to be fully melted. This was confirmed by chemical analysis results indicating the lowest extent of W and C dissolution into the FeCrAl binder because of the lower temperature of the inflight WC–F particles. (4) A maximum microhardness of 1498 HV0.3 was measured for WC–F coating, whereas WC–C and WC–N coatings revealed lower values of 1305 and 1254 HV0.3.
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(5) The results showed a mean fracture toughness of 5.9 MPam1/2 for WC–C which was substantially greater than that for both WC–F and WC–N coatings, respectively with 3.1 and 2.8 MPam1/2. The higher fracture toughness of WC–C coating results from the minor porosity and stronger inter-splat interface together with the lower W2C/WC ratio of this coating in comparison to WC–F and WC–N coatings.
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