Effect of WC size on interface fracture toughness of WC–Co HVOF sprayed coatings

Surface & Coatings Technology 201 (2006) 619 – 627 www.elsevier.com/locate/surfcoat

Effect of WC size on interface fracture toughness of WC–Co HVOF sprayed coatings M. Watanabe a,⁎, A. Owada a,b , S. Kuroda a , Y. Gotoh b a

b

Thermal Spray Group, MEL, National Institute for Materials Science, Ibaraki, Japan Department of Materials Science and Technology, Science University of Tokyo, Chiba, Japan Received 11 July 2005; accepted in revised form 8 December 2005 Available online 25 January 2006

Abstract The interface fracture toughness of high velocity oxygen fuel (HVOF) sprayed coatings on carbon steel made of various types of WC–12 wt.% Co powders with different WC particle sizes of 0.2 to 7.0 μm was evaluated by the pre-notched four-point bending test to clarify the size effect of WC particles and to explore the superior adhesion mechanisms of this coating system. The correlation between the splat microstructure and the toughness variation was investigated by observing a cross-section of WC–Co splats around the interface using the focused-ion-beam (FIB) technique. The interface fracture toughness of WC–12 wt.% Co coating/carbon steel under a mixed Mode I/Mode II loading condition increases from 600 to 1800 J/m2 as the WC particle size increases. Compaction of the underlying microstructure, the intrusion of WC particles into the substrate, and the volume fraction of the metallic binder phase in the coating are key factors in achieving excellent adhesion of this coating system. © 2005 Elsevier B.V. All rights reserved. PACS: 62.20.Mk Keywords: High velocity oxy-fuel; WC–Co; Fracture toughness; Focused ion beam

1. Introduction Adhesive strength is one of the most important factors in thermal spray coating since it directly relates to the durability of the coating [1,2]. Several sprayed coatings used in thermal spray industries are well known to have superior adhesion strength on metal substrates. The family of tungsten carbide– cobalt (WC–Co) cermet coatings is one example. The adhesion strength of a coating is typically evaluated by tensile adhesion tests, such as the ASTM C633-79 [3]. A polymer-based adhesive is used in these test methods to attach a jig on the coating surface; hence, the measurable adhesion strength is limited by the strength of the adhesive, which is usually in the range of 60 to 80 MPa. Attempts to evaluate the adhesion strength of WC–Co coatings sprayed by the High Velocity Oxy Fuel (HVOF) process have ended with failure in the adhesives before debonding between the coating and substrate [4,5], and the critical loads could not be determined. The only information ⁎ Corresponding author. Tel.: +81 29 859 2469. E-mail address: [email protected] (M. Watanabe). 0257-8972/$ - see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.12.019

obtained by these experimental assessments is that this coating system has excellent adhesion strength. Staia et al. studied the effect of substrate roughness induced by grit blasting on the adhesion strength of WC–17% Co HVOF coatings [6]. They applied two test methods: the Vickers indentation method and the tensile adhesion test, and experimentally demonstrated the substrate roughness effect. While they could measure the adhesion strength by applying the tensile adhesion test (ASTM C633-79), the adhesion strength was considerably low, in a range of 32.7 to 52.5 MPa, as the HVOF-sprayed WC–Co coatings. They used considerably bulky starting powders, with particle sizes of 325 ± 5 μm, which is one order of magnitude larger than the widely used powders with 30 to 50 μm diameters. Evidently, the temperature and the velocity of the molten particles under the spraying conditions were not adequate to enable satisfactory bonding between the coatings and substrate. There have been several studies to investigate the mechanisms of the strong adhesion of the WC–Co coating system [4,5,7–10]. Nutting et al. investigated the interface microstructure between an HVOF-sprayed WC–Co coating and low

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Table 1 Spraying conditions using the High Velocity Oxygen Fuel (HVOF) gas spray process

Table 2 Classification of the starting powders of WC–12 wt.% Co with different WC particle-size distributions

Barrel length (mm) Spray distance (mm) Fuel (l/min) Oxygen (l/min) Powder feed gas

Type

WC size (μm)

A B C D E F

0.2 1.0∼1.5 2.0∼2.5 3.0∼4.0 5.0∼7.0 1.0∼5.0

102 380 0.38 944 Nitrogen

carbon steel by transmission electron microscopy (TEM) [8]. They found that there were no deleterious oxides or intermetallics, such as iron oxides, at the interface that suggested direct bonding between the W–Co alloy phases and the metal substrate. In addition, no evidence of the transfer of Co and W into the steel substrate was observed. Li et al. conducted a series of detailed studies of carbide coatings that focused on the effects of carbide particle types and sizes on the coating microstructures and the adhesion strength [4] as well as the wear resistance [10] by comparing various coatings produced by HVOF, such as WC–Co, WC–Ni–Fe, and W– Ni. It was found that the WC–Co particle is in a solid–liquid two-phase state during the spraying process with WC in the solid state and the metallic matrix in the liquid state, and that this two-phase state during spraying is a necessary condition to achieve superior adhesion [4]. They also determined that the carbide particle sizes and shapes significantly influence the wear behavior of cermet coatings [10]. The adhesion strength of all WC–Co and W–Ni coatings could not be determined quantitatively due to failure in the adhesives used in those experiments. However, they suggested the important concept that the adhesion strength of HVOF cermet coatings are increased with an increase in the density of the solid materials in two-phase particles, such as WC, due to the increased dynamic impact pressure caused by the impact of the solid particles. Experimental proof remains unavailable due to the lack of quantitative data. The present study was undertaken to evaluate the adhesion strength of WC–Co coatings and to explore the mechanisms of superior adhesion in terms of the coating microstructures. Spray powders with a wide range of carbide particle sizes were selected, and we focused on the effects of carbide particle sizes on adhesion and the splat–substrate interface structure to highlight the effects of the dynamic impact energy of carbide particles. We examined the interface fracture toughness using the pre-notched four-point bending test. Cross sections of WC– Co splats deposited on a steel plate were prepared and observed by a focused ion beam (FIB) and a scanning electron microscope (SEM) to investigate the features of carbide particles inside a splat as well as the microstructure of the splat–steel interface.

(A∼E: Company A, F: Company B).

steel (0.45% C) substrates. Six commercially available WC– Co powders were sprayed under the conditions listed in Table 1. These powders had similar size distributions of 15 to 45 μm with the same composition, but differed in the size distributions of the WC particles contained in one WC–12% Co powder particle. Table 2 presents a classification of those powders. Backscattered electron (BSE) images of type A (about 0.2 μm WC) and type D (5 to 7 μm WC) powders are provided in Fig. 1 as examples. The substantial size differences of the WC particles (white regions) between the two can be clearly seen. Type F was obtained from a different source than powders A to E. It contains particles with much wider carbide size distributions, 1 to 5 μm, in the powder. A substrate surface was mirror-finished by mechanical polishing before spraying to ensure the same deposition

(a)

5µm

(b)

WC

Co

2. Experiments 2.1. Spraying conditions and powders HVOF spraying equipment (JP5000, Praxair Technology Inc., USA) was used to spray WC–12 wt.% Co on carbon

5µm

Fig. 1. BSE images of WC–12% Co powders. (a) Type A, containing 0.2 μm size WC particles. (b) Type D, with 5 to 7 μm WC particles.

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conditions among the different powders and to correlate the microstructure of the splats with the coatings. The temperature of a substrate significantly affects the splat formation and adhesion of the coating [11–13]. Intentional substrate heating and cooling were not applied in the present study. Two types of samples were prepared for all powders, one for the splat microstructure observations by FIB and SEM and the other for interface toughness measurements by the pre-notched fourpoint bending test, as described in the next section.

Splat (Top view)

621

Splat

Substrate surface

2.2. Evaluation of interface fracture toughness The interface fracture toughness was measured by the prenotched four-point bending test to determine the effect of WC particle sizes. There is no need to use any adhesives in this test method, such as those required for the commonly used tensile adhesion test (ASTM C633-79) and the pin test. Therefore, it is possible to apply this method without the limitations imposed by adhesive strength. Fig. 2a and b provide a schematic of the pre-notched fourpoint bending test and an image of the introduced pre-notch. The bending moment increases increasing load P, and a precrack initiates from the notch (mechanically facilitated) and propagates vertically to the interface. The crack deflects into the interface and propagates along it when the interface is sufficiently weaker than the substrate fracture toughness. Another advantage of this test design is that the strain energy release rate for the interface crack is independent of the crack length when crack length a is within a range of a / l ≤ 1.5, where l is the distance between the inner and outer load lines [14–17]. This steady-state behavior reduces experimental difficulties due

Sectioning by FIB FIB sectioned region

Stage rotation

Substrate surface

Splat

Cross section to be observed

Fig. 3. Schematics of the splat sectioning process by focused ion beam (FIB) equipment.

(a)

P/2

to crack length measurements. The strain energy release rate Γ for the plane strain condition can be expressed as [15]

P/2 h2

Substrate

Crack, 2a

C¼

h1

Coating

l

Notch


 P2 l 2 ð1−m22 Þ 1 k − 8b2 E2 I2 IC

ð1Þ

k ¼ E2 ð1−m21 Þ=E1 ð1−m22 Þ

ð2Þ

I2 ¼ h32 =12

ð3Þ

(b) Substrate

and Interface

Ic ¼ h31 =12 þ kh32 =12 þ kh1 h2 ðh1 þ h2 Þ2 =4ðh1 þ kh2 Þ WC-Co coating

Notch

200µm

Fig. 2. (a) Schematic of the four-point bending test configuration. (b) Highly magnified image of a notch tip in the coating.

ð4Þ

where Ei is Young's modulus, νi is Poisson's ratio, hi is the thickness, and b is the width of the sample. Subscripts 1 and 2 refer to the coating layer and the substrate. The Young's modulus of coating E1 was evaluated by the resonant technique following the ASTM C623-92 [18] and the values were about 290∼340 GPa. The Poisson ratio ν1 = 0.3 for the coatings, E2 = 210 GPa and ν2 = 0.28 for the substrates were adopted in all cases. The dimensions of the test sample substrate were

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80 × 10 × 5t mm, and the thickness of the coating was about 800 μm. A 100 μm wide pre-notch was introduced with a diamond notch blade, leaving a 100 to 200 μm thick coating from the tip of the notch to the interface, as indicated in Fig. 2b. A sample was placed and carefully aligned between two pairs of loading pins with the machined notch at the tensile side. The bending and fracture behavior was monitored by a CCD camera during the experiment to observe the crack initiation and the propagation behavior. The load was applied under a constant load point displacement rate of 0.1 mm/min. The tests were conducted twice for each type of powder.

(a)

Coating

Substrate 100µm

2.3. Splat microstructure observations by focused ion beam (FIB)

(b) Mount

Cross-sectional microstructures of the splats were characterized by the FIB system. These splats were deposited on a mirror-finished substrate surface under the same spray conditions used for the coatings. A schematic of the procedure is provided in Fig. 3. A splat was sectioned with a Ga ion beam along an arbitrary line containing carbide particles. The sample stage was rotated by 60° after sectioning, and the microstructure in the cross section was observed.

Substrate

Remaining coating layer

3. Results and discussion

5µm

3.1. Interface fracture toughness

Fig. 5. (a) Interface crack formation during loading. (b) Cross section of the coating after a test, indicating failure at the interface and in the coating.

Fig. 4 depicts the typical load displacement curve of the four-point bending test. The displacement was converted from the cross-head displacement rate, and thus the load curve increased nonlinearly at the beginning due to compliance of the loading system. The load then began to increase linearly. Crack propagation and spallation of coatings occurred instantaneously after a small load drop, shown in the figure inset, occurred a few times corresponding to the vertical crack growth to the interface and the deflection

along the interface (Fig. 5a). Observations after testing revealed that a fracture occurred either at the interface or in the WC–Co coating just above the interface in all cases (Fig. 5b). The interface fracture toughness, Γc, was obtained by substituting the load, Pc, at the final failure into Eq. (1). Fig. 6 indicates the interface fracture toughness for six different coatings as a function of the carbide particle size in the starting powders. The interface fracture toughness increased as the carbide sizes in the original powders increased from Interface Fracture Toughness (J/m2)

5.0 Crack propagation and spallation of coatings

Load (kN)

4.0

3.0

2.0

1.0

0

0

0.2

0.4

0.6

0.8

1.0

2000 F E

1500 C

D

1000

A

500

0

0

B

2

4

6

8

WC particle size (µm)

Cross head displacement (mm) Fig. 4. An example of a load–displacement curve.

Fig. 6. Evaluated interface fracture toughness as a function of the WC particle size in the starting powders.

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about 600 J/m2 (type A) to 1800 J/m2 (type E), except for type F, which had wider carbide size distributions and the greatest interface fracture toughness. The interface fracture toughness of conventional SUS316 HVOF coating on steel substrates is reportedly 200 to 300 J/m2 for the tensile fracture mode (mode I) [19]. While the difference in the failure mode must be considered, the interface fracture toughness of the WC–Co coatings/steel substrates was remarkable. Type F exhibited slightly more toughness than

type E, even though the carbide size was smaller on average. It is also notable that the two data points for the same powder are very close. This indicates the superb reproducibility of the present test method compared to the conventional tensile adhesion test, which normally yields much greater scatter in the strength value. The results of the splat microstructure observations are discussed in the following sections to clarify the microstructural differences among all the samples, focusing on the effects of WC size.

Type A (WC: 0.2μm) (a)

623

Type E (WC: 5~7μm) (b)

10μm

10μm

(d)

(c)

WC-Co splat

Cross section by FIB

Substrate Surface 10μm

10μm

(e)

(f)

WC Particle Alloy phase

3 μm

Substrate

3 μm

Fig. 7. FIB images of type A and type E splats in top view (a and b), in a cross-section of the splats (c and d), and high magnification images of the microstructure around the interface (e and f).

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Type F

3.2. Splat microstructures Fig. 7 contains images of the disc-shaped type A and type E splats (Fig. 7a and b) and cross sections along the dotted lines (Fig. 7c to f). These images were obtained by FIB equipment. The microstructures of the steel substrate, such as ferrite and pearlite surrounding a splat and under a splat, can be seen due to sputtering of the surface layer with a Ga ion beam during FIB observation. EDX analysis revealed that the splats consisted of W–Co metal binder phases and the carbides. XRD analysis of the coatings showed the crystalline peaks of WC, W2C, and W (Fig. 8). There is a clear tendency of the increase of W2C peak as the carbide size decreases, which agrees with a previous study by Stewart et al. [20]. It is considered that the larger surface to volume ratio of the smaller carbides promotes more rapid dissolution and decarburization [20]. The difference of decarburization degree causes the variation of binder phase compositions and may have considerable effects on the mechanical properties of coatings. Nutting et al. investigated the microstructure evolution of a low-carbon steel substrate under a WC–12% Co splat by TEM [8]. They reported that the microstructure of a substrate under a coating of WC–Co/low carbon steel displays gradual microstructural changes over a 100 μm depth due to heat transfer during deposition. The images in Fig. 7 depict various microstructures on sectioned surfaces; there is no noticeable difference between the substrate microstructure and the asreceived substrate. The disparity may be due to the heating history difference between a splat sprayed with only one path and thick coating deposition with multiple paths. Nutting et al. also reported that no deleterious phases, such as iron oxides, existed at the interfaces. Greater magnification images of the interface regions (Fig. 7e and f) suggest that the interface between the splat and the substrate bonds tightly without any intermediate layer. The interface microstructure and the substrate phase evolution during thick coating deposition remains of interest.

WC

Normalized Intensity

1.0

0.0 30

WC

WC

W2C

35

W2C W2C W

40 Difraction Angle (2θ)

45

Type F Type E Type D Type C Type B Type A 50

Fig. 8. XRD scans from all coatings in the 2θ range 30∼50°; the increase of W2C crystalline peak is apparent as the reduction of the carbide size.

5μm Fig. 9. Cross section of a splat made of type F powder.

Splats of types A and E consist of the W–Co alloy phase (brighter region) with WC particles (dark gray region). The type E splat appears to contain a greater volume fraction of the alloy phase. The interfaces of both splats display concave peripheries, implying substrate deformation by the high kinetic energy of the molten particles upon impact. The deepest penetration depth around the center is 1 to 2 μm. The maximum thickness of the splats was about 4 μm for both, but the type A splat had more uniform thickness over the splat. In contrast, a cross section of the type E splat indicates a nonuniform feature due to the existence of large particles. The type A splat consists of a large number of uniformly dispersed WC particles about 100 nm in size and the matrix phase, while type E contains much larger WC particles of about 3 μm diameter, but only six carbide particles appear in the cross section. These carbide sizes are a little smaller than those in the original powders but still correspond well to the original sizes. The smaller WC particle at the interface in the type E splat was compressed by the larger WC particle just above it and intruded into the substrate (Fig. 7f). The deformation and the compression of the substrate microstructure under the smaller carbide particles can be observed, indicating the high impact pressure of carbide particles. In contrast, no such carbides intruded into the substrate in type A and thus the interface is much smoother than that of type E. Fig. 9 depicts the splat microstructures of type F, which contains carbides with a wide size distribution of 1 to 5 μm in the original powder. The characteristic feature of this splat is the uniformly spread thin alloy phase (white region) and the sharp carbide particles pinning the splat onto the substrate (marked with dotted circles in the figure). This powder appears to have melted well during deposition compared with other powders. The lower viscosity of the matrix liquid phase upon impact may have assisted the intrusion of carbide particles into the substrate at the interface. In addition to the presence of the WC particles pinning the splat, the bonding area between the W–Co alloy phase and the substrate surface is much wider than in type A.

M. Watanabe et al. / Surface & Coatings Technology 201 (2006) 619–627

3.3. Coating microstructures Fig. 10 depicts the coating microstructure around the failure path on a cross section of type F after the test. Delamination of the coating layer occurred just above this region. The intrusions of large WC particles into the substrate could be clearly detected, consistent with the FIB images. There was no evidence of crack formations at the interface between the WC particles and the substrate, implying superior bonding strength between the two. Moreover, compressive residual stress could be generated in the radial direction along the WC particles peripheral in the substrate due to forced intrusion of the particles and strain hardening of the substrate. This effect is similar to the shot peening technique, which is a common surface-treatment process used to improve the fatigue life of machine parts by throwing many small hard balls with a specific velocity at a target surface. The introduction of a compressive residual field into the surface layer improves fatigue life of the materials [21,22]. The same mechanism functions for WC–Co

(a)

Fracture plane

Crack

Voids and Crack

Voids

WC Interface

Crack Substrate

3μm

(b)

Broken WC-1

Crack

WC

W-Co alloy phase

Broken WC-1'

Voids

500nm

Fig. 10. Coating microstructure around the fracture path after the four-point bending test.

25 A

Bare substrate area fraction on fracture surface (%)

Type E and F coatings exhibited much greater interface fracture toughness than that of type A, as described in the previous section. A comparison of the splat microstructures revealed the following factors: (i) the intrusion of the carbides into the substrates; (ii) the compression of the microstructure under the impacted surface; (iii) rougher interfaces; and (iv) the large area fraction of the bonding between the W–Co alloy phase and the substrate surface. These factors are thought to be responsible for the substantial increase of the interface fracture toughness.

625

20

15 B

10 C

F

0

E

D

5

0

1

2

3

4

5

6

7

WC particle size (µm) Fig. 11. Bare substrate surface area fraction on the fractured surface as a function of the volume fraction of the alloy phase in the coating.

HVOF coatings, since many hard WC particles impact the substrate surface during deposition with a 500 to 1500 m/s velocity. The introduced compressive residual stress will constrain deformation and debonding of the WC particles from the substrate, preventing crack propagation at the interface. Dynamic numerical analyses will be required to clarify this effect on the interface fracture toughness, including the thermal and residual stress induced during the deposition process. Formations of sharp cracks can be seen in the W–Co alloy phases and at the interface between the alloy phase and the substrate in Fig. 10a, particularly in the area where the volume fraction of WC particles is relatively lower than in other places. This clearly indicates that the regions surrounded by fewer carbide particles are not compressed well during the depositions, resulting in weak bonding or a locally porous microstructure. Compression of the underlying microstructure by the impact of the carbides must enhance the bonding, whether it is at the interface or in the coating. Fig. 10b depicts another region where a crack passed through the WC particles. Many voids were formed around the broken WC particles, indicating that the W–Co alloy phases supported the load by plastic deformation after corruption of the WC particles. This failure process is the same behavior that is observed in sintered bulk WC–Co cermets. The binder phase (Co) networks and their load sustainability due to plastic deformation around a crack tip are considered to be the reasons for the superior toughness of bulk WC–Co [23–31]. This implies that the regions in Fig. 10b had satisfactory bonding between the WC and the surrounding matrix and a level of fracture resistance similar to bulk materials. The area fraction of the bare substrate surface on the fractured surface of the substrate side was calculated by an image analysis of backscattered electron images of the fractured surfaces and is plotted in Fig. 11. The bare substrate surface on the fractured surface indicates that failure occurred exactly at

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Fracture toughness, KIC (MPa m1/2)

30 J. L. Chermant et. al. [29] M. Nakamura et. al. [24] V.D. Kristic et. al. [23] Present expriment results, Keq

25

20

15

10

5

0

0.1

0.2

0.3

0.4

0.5

0.6

Volume fraction of binder phase, Vb Fig. 12. Equivalent fracture toughness Keq as a function of the binder phase volume fraction with reference data of various bulk WC–Co [23,24,29].

the interface because no deflection of cracking into the steel substrate would occur. It must be noted that the maximum fraction of the interface failure was only 20% with the finest WC size. A strong tendency for additional failure to occur in the coating as the carbide size increases is recognized. These facts ensure that the fracture toughness at the interface is significantly greater than the toughness of the coating itself, and the former increases more strongly with the carbide size. Therefore, failure of this coating system would be determined by the fracture toughness of the coating around the interface. The mixed mode phase angle in the steady-state region of the present experiment, φ = tan−1(KII / KI), was estimated to be around 40° for h1 / h2 = 0.16 according to the literature [15], where KI and KII are the stress intensity factors for modes I and II. The equivalent critical stress intensity factor Keq ¼ ffi pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 þ K 2 can be estimated based on the following relation KIC IIC between the critical energy release rate and the critical stress intensity factor for an interface crack [32,33]:   1 v1 þ 1 v2 þ 1 2 CC ¼ þ ð5Þ Keq l1 l2 16cosh2 ðepÞ e ¼ ð1=2pÞln

1−b ; v ¼ 3−4mi ; li ¼ Ei =2ð1 þ mi Þ 1þb i

ð6Þ

where β is the Dunders' parameter. The estimated Keq is plotted for a binder volume ratio Vb with the KIC of various bulk WC– Co (containing 7∼20 wt.% Co) in Fig. 12 [23,24,29]. The alloy phase area fraction in the coatings obtained by image analysis was used as Vb (Table 3). The mode pffiffiffiffi I fracture toughness of the bulk WC–Co is 10 to 20 MPa m for Vb b 0.4 and increases as the volume fraction increases. The Keq values derivedpfrom ffiffiffiffi the present experiment results were 13 to 23 MPa m for 0.25 b V b b 0.6 and

corresponded well with the value expected from the bulk data for a higher Vb. The fracture toughness increased as the volume fraction of the alloy phase in the deposited coating increased. One of the reasons of high volume fraction of the alloy phase in the coatings is considered to be rebounding and loss of larger WC particles during deposition due to higher kinetic energy. This is evident from the fact that the coatings made of the larger WC particle powders, such as types D and E, contain more alloy phases in the microstructure. Type F coating also contains a very high volume fraction of alloy phases (54%) and exhibited superior interface fracture toughness, even though the average WC particle size was medium. The sprayed particles of type F appeared to be melted well, as indicated in Fig. 9, and thus the splat becomes wider and thinner, increasing the probability of rebounding of the large WC particles. In addition, decarburization of WC during deposition can cause the increase of the binder phase volume fraction. Since smaller carbide particles tend to be dissolved and decarburized more [20] and the Type F powders contain not only large carbides but also considerable amount of small carbides, these small carbides can be decarburized, and thus the coatings fabricated from type F could result in higher amount of binder phases. The alloy phase in Fig. 9is apparently compressed well and bonds with the surrounding microstructure enough to cause plastic deformation for crack propagations, resulting in superior fracture toughness around the interface. The implication is that coating with a greater volume ratio of alloy phases will yield greater fracture toughness if the sprayed coating is dense enough; however, sufficient densification cannot be achieved without impact and compaction by solid particles. The high kinetic energy of a larger WC particle also simultaneously induces the other possible toughening mechanisms described in previous sections, such as the intrusion of WC particles into the substrate at the interface. Therefore, the interface fracture toughness dramatically increases as the WC size increases. Also, it should be noted that the mechanical property variations depending on both the composition of binder phases and the ratio of WC to W2C might have considerable effects. The contribution of each factor to improvement of the interface fracture toughness is still unclear and further study is required. The discussions above are not confined to WC–Co coatings, as Li et al. [4] suggested. The addition of a secondary phase with high hardness and a high melting point to the original powders will improve the adhesion strength of the coating. These composite coatings could be useful as bond coatings. For example, coatings that consist of nano- or submicron-order microstructures are expected to have superior properties, such as higher wear resistance than conventional Table 3 Volume fraction of the W–Co alloy phase in WC–Co coatings Type Powder Coatings

A 0.26

B

C

0.37

0:19 ~ 0:2 0.36 0.42

D

E

F

0.52

.054

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coatings. However, the adhesion strength of those coatings sprayed by the HVOF process will decrease significantly, as observed in type A coating in the present study. Application of bond coating with larger particles will be one option to improve adhesion and to control the total performance of the coating system. The substrates in the current study were polished and mirror-finished before spraying in order to correlate with splat microstructure observations. The effect of surface conditions must be considered when interpreting the interface fracture toughness curve. The interface becomes much rougher as the carbide particles become larger, as described in the previous section. In addition, the deposition efficiency decreases as the particle sizes increase. For example, type A is twice as efficient as type E under the current spray conditions. Therefore, the greater the carbide size, the longer the substrate surface must be blasted, which constitutes a selfpeening effect. Staia et al. reported the effect of the surface roughness on the adhesion strength of WC–17% Co by controlling the substrate surface by varying the carrier gas pressure during grit-blasting. The adhesion strength increased by about 60% in their experiment due to the increase of the surface roughness [6]. Although detailed mechanisms are still unclear, their study implies that the roughness of a substrate surface has large effects on local impact phenomena due to various reasons. For example, the impacts onto grit-blasted surface have wide range of the impact angles with sprayed particles and also consist of various geometrical situation such as particle-flat plane, -crest, -valley, and so on. These factors will also affect the carbide intrusions observed in the present FIB observations depending on spray conditions. Therefore, the interfacial fracture toughness for the WC particle size studied in this paper could vary depending on the surface condition before spraying, but the similar trend would be observed. A comparison of the interface fracture toughness of coatings on blasted substrates with the present results is underway and will provide useful assessment for practical situations. 4. Conclusions Splats and coatings of six types of WC-12 wt.% Co with various WC particle sizes were deposited by HVOF spraying. The interface fracture toughness was evaluated by the prenotched four-point bending test, and the microstructure of the splat around the interface was examined by SEM and FIB crosssectioning. The WC particle size in the original powders was found to significantly affect the interface fracture toughness, which generally increases as the WC particle size in the original powder increases. The key factors in realizing the excellent adhesion of this coating system is considered to be compression and densification of surrounding microstructure by the impact of WC particles, intrusion of the particles into the substrate, and plastic deformation of W–Co alloy phases around a crack tip, which functions as the energy dissipation mechanism for crack propagation in the coating.

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Acknowledgements This investigation was sponsored by the National Institute for Materials Science under grant, “High performance coatings for severe environments,” No. 161319 (2004–2005). We thank to M. Komatsu and N. Kakeya for supporting fabrication of coatings. M.W. also thanks to the Special Coordination Funds for Promoting Science and Technology from the Ministry of Education, Culture, Sports, Science and Technology of the Japanese Government for its financial support. References [1] H.D. Steffens, B. Wielage, J. Drozak, Surf. Coat. Technol. 45 (1991) 229. [2] S.D. Brown, Thin Solid Films 119 (1984) 127. [3] ASTM 633-79, American Society of Testing Materials, Philadelphia, 1979. [4] C.-J. Li, Y.-Y. Wang, T. Wu, G.-C. Ji, Ohmori, Surf. Coat. Technol. 145 (2001) 113. [5] C.-J. Li, H. Li, Thermal Spray Meeting the Challenges of the 21st Century, ASM International, Materials Park, OH, 1988, p. 723. [6] M.H. Staia, E. Ramos, A. Carrasquero, A. Roman, J. Lesage, D. Chicot, G. Mesmacque, Thins Solid Films 377–378 (2000) 657. [7] M.M. Lima, C. Godoy, J.C. Avelar-Batista, P.J. Modenesi, Mater. Sci. Eng., A 357 (2003) 337. [8] J. Nutting, J.M. Guilemany, Z. Dong, Mater. Sci. Technol. 11 (1995) 961. [9] J.M. Guilemany, J.M. Paco, J. Nutting, J.R. Miguel, Metal. Mater. Trans. A 30A (1999) 1913. [10] C.-J. Li, Y.-Y. Wang, G.-J. Yang, A. Ohmori, K.A. Khor, Mater. Sci. Technol. 20 (2004) 1087. [11] M. Fukumoto, Y. Huang, J. Therm. Spray Technol. 8 (1999) 427. [12] M. Fukumoto, E. Nishioka, T. Nishiyama, Surf. Coat. Technol. 161 (2002) 103. [13] M. Fukumoto, E. Nishioka, T. Matsubara, Surf. Coat. Technol. 121 (1999) 131. [14] Z. Suo, J.W. Hutchinson, Mater. Sci. Eng., A 107 (1989) 135. [15] P.G. Charalambides, J. Lund, A.G. Evans, R.M. McMeeking, J. Appl. Mech. 111 (1989) 77. [16] Q. Ma, H. Fujimoto, P. Flinn, V. Jain, F. Adibi-Rizi, F. Moghadam, H. Dauskardt, Mater. Res. Soc. Symp. Proc. 391 (1995) 91. [17] Q. Ma, C. Pan, H. Fujimoto, B. Triplett, P. Coon, C. Chiang, Mater. Res. Soc. Symp. Proc. 436 (1997) 379. [18] ASTM 623-92, American Society of Testing Materials, Philadelphia, 1992. [19] S. Kuroda, National Institute for Materials Science, unpublished work. [20] D.A. Stewart, P.H. Shipway, D.G. McCartney, Acta. Mater. 48 (2000) 1593. [21] J.K. Li, Y. Mei, W. Dui, Mater. Sci. Eng., A 147 (1991) 167. [22] S.A. Meguid, G. Shagal, J.C. Stranart, K.M. Liew, L.S. Ong, J. Eng. Mater. Technol. 127 (2005) 170. [23] V.D. Krstic, M. Komac, Philos. Mag., A. 51 (1985) 191. [24] M. Nakamura, J. Gurland, Metall. Tans. 11A (1980) 141. [25] A.G. Evans, A. Heuer, D.L. Porter, in: D.M.R. Taplin (Ed.), 4th Int. Conf. on Fracture, Pergamon Press, Oxford, 1977, p. 529. [26] L.S. Sigl, H. Exner, Metall. Trans. 18A (1987) 1299. [27] L.S. Sigl, H.F. Fischmeister, Acta. Mater. 36 (1988) 887. [28] M.J. Murray, Proc. R. Soc. Lond., A 356 (1977) 438. [29] J.L. Chermant, F. Osterstock, J. Mater. Sci. 11 (1976) 1939. [30] R.K. Viswanadham, T.S. Sun, E.F. Drake, J.A. Peck, J. Mater. Sci. 16 (1981) 1029. [31] J.R. Pickens, J. Gurland, Mater. Sci. Eng. 33 (1978) 135. [32] B.M. Malyshev, R.L. Salganik, Int. J. Fract. Mech. 1 (1965) 114. [33] Y. Torres, V.K. Sarin, M. Anglada, L. Llanes, Scripta. Mater. 52 (2005) 1087.