S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites Saleh Kaytbaya , Medhat El-Hadekb a Benha University, Department of Mechanical Engineering, Faculty of Engineering, Benha, Egypt University, Department of Production & Mechanical Design, Faculty of Engineering, Port-Said, Egypt b Port-Said IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. Wear resistance and fracture mechanics of WC–Co composites Manufacturing of WC–Co composites using the electroless precipitation method at different sintering temperatures of 1 100, 1 250, 1 350 and 1 500 8C was successfully achieved. The chemical composition of the investigated materials was 90 wt.% WC with 10 wt.% Co, and 80 wt.% WC with 20 wt.% Co. The specific density, densification, and Vickers microhardness measurements were found to increase with increased sintering temperature for both the WC–Co compositions. The composites of tungsten carbide with 10 wt.% Co had a higher specific density and Vickers microhardness measurements than those for the composites of tungsten carbide with 20 wt.% Co. Composites with WC-10 wt.% Co had better wear resistance. The stress– strain and transverse rupture strength increased monotonically with the increase in sintering temperatures, agreeing with the material hardness and wear resistance behavior. Fractographical scanning electron microscopy analysis of the fracture surface demonstrated a rough characteristic conical shape failure in the direction of the maximum shear stress. A proposed mechanism for the formation of the conical fracture surface under compression testing is presented. Keywords: Cobalt tungsten carbide; Fracture mechanics; Wear resistance; Electroless precipitation; Sintering densification 1. Introduction Cemented carbides are made out of metal matrix composite, as the aggregate carbide particles are pressed with metallic binder that serves as the matrix. Throughout this process the binder enters the liquid state and the carbide grains’ higher melting point means they remain in the solid state [1]. Compared to other cemented carbides, tungsten carbide (WC) combines favorable properties such as high hardness and good wettability. Tungsten carbide (WC) is approximately two times stiffer than steel, with a Young’s modulus of approximately 550 GPa and is much denser than steel and titanium [2, 3]. The singular combination of hardness, wear resistance, refractory materials, strength and other desired physical and chemical properties makes cemented carbides candidate materials for many industrial applications. Unlike most other metal powders, cemented carbide powders do not deform during the compacting process. Generally, they cannot be compressed to much above 65 % of the theoretical upper limit for density. Despite this low green density, carbide manufacturers have developed Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 the technology for achieving good dimensional tolerances in the sintered product [4]. WC is used in a number of industrial appliactions such as machining wood tools, in chipless operations such as can tooling, high pressure punches and for wear parts [5]. WC is also reported to be used in the cases where hot hardness and thermal shock resistance are required, such as mining tools, oil and gas drilling tools, metal forming and forestry tools [6]. Cobalt (Co) as a binder metal with WC has been popular due to its excellent compatibility with WC. This is due to the good wettability of the molten metal as WC has high solubility in the Co during liquid phase sintering, leading to good densification, and superior mechanical properties [7]. WC possesses strong ionic interatomic bonding producing desirable material characteristics that exists in several crystalline phases, where all revert to the most stable hexagonal alpha phase at elevated temperatures. Because of WC’s high hardness, excellent dielectric properties, refractoriness and good thermal properties this makes it the material of choice for a wide range of applications [8]. The properties of the composite depend on the variety of techniques used to compact the cemented carbide grade powders. Carbide tools for mining and construction applications are pill pressed in semiautomatic or automatic presses [9]. Metal-cutting inserts are also pill pressed, but may require additional shaping after sintering. Cold isostatic pressing, in which the powder is subjected to equal pressure from all directions followed by green forming, is also a common practice for wear and metal forming tools [10]. Rods and wires can be formed by extrusion. Cemented carbides can be produced by the injection molding process, although this method of consolidation is used less than the aforementioned consolidation processes [11]. Unlike most other metal powders, cemented carbide powders do not deform during the compacting process. Generally, they cannot be pressed above 65 % of the theoretical upper limit for density [10, 11]. Newer-emerged processes such as electroless precipitation with sintering [12, 13] that could increase the green density, achieve good dimensional tolerances and the desired microstructure have been developed. In recent years, WC–Co composites with submicron tungsten carbide grain sizes have been developed [14], as the decrease in the average tungsten carbide grain size involves an improvement in strength [15], hardness and wear resistance [16]. Typical applications include indexable inserts and a wide variety of solid tungsten carbide drilling and milling tools [5]. It has been reported that densification of nanostructured WC–Co powder could be completed in short time intervals if grain growth inhibitors were added [17], because grains 557 IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites grow extremely rapidly during heating. WC–Co composites made from nanostructured powder had better surface crack resistance than samples made from standard sub-micron powders [15]. High purity WC particles were usable in both oxidizing and high temperatures all the way up to 1 925 8C [15, 16]. Fracture analysis of WC–Co composites has been correlated to the strength and mechanical and microstructural parameters [18]. The use of indentations involved in direct measurement of Vickers hardness producing radial cracks as a function of indentation load to measure the fracture toughness has also been reported [19, 20]. However, with more comprehensive analysis, it was found that hardness indentations on the hard metal substrates produce Palmqvist-type cracks [20] and consequently the measurements’ accuracy was questioned. If WC–Co composites could be produced with higher hardness and slightly lower fracture toughness values, they could be used in mining and wear resistant applications with better reliability [20]. When using WC–Co as a coating material, it was found that the wear resistance was superior to that of conventional coatings, specifically under high load conditions [21]. Recently, Stewart et al. [21] studied the abrasive impact wear resistance of WC–Co, where the wear resistance was correlated to the bulk hardness of the material and the cements. Materials with lower binder content and higher hardness exhibited higher wear resistance compared to materials with higher binder content [22]. Interestingly, experimental investigation of phase equilibria shows that coarse grained WC–Co composites have lower wear resistance compared to fine grained materials [23]. The scope of this study was to manufacture WC–Co composites using an electroless precipitation method and different sintering temperatures of 1 100, 1 250, 1 350 and 1 500 8C. This unique technique uses a water-soluble chemical salt as a starting material for the precursor of these composites. The compositions of all investigated WC–Co composites were 90 wt.% WC with 10 wt.% Co, and 80 wt.% WC with 20 wt.% Co. The microstructure was thoroughly examined using high resolution optical investi- gations during the process. The effect of different sintering temperatures on the density, hardness and stress–strain behavior, wear resistance, and fracture mechanics were investigated. 2. Experimental 2.1. Materials preparation The WC particle sizes used in this study were commercially available at 1 – 0.5 lm, apparent density of 3.2 g · cm–3, melting temperature of 2 870 8C, and with purity of 99.5 % [24]. The Co particle sizes used were commercially available at 100 – 200 nm, apparent density of 1.17 g · cm–3, melting temperature of 1 495 8C, and with purity of 99.5 % [24]. The process started with preparing the plating baths that contained the Co salt and a reducing agent. The cobalt was electrolessly deposited from hypophosphite-based solutions, but only from alkaline baths. In addition, complexing agents (generally citrate and ammonium salts) were used. The complexing agents serve the function of simply preventing precipitation of basic salts, as this also affects the deposition rate and properties. The bath pH was usually maintained in the range between 8 and 10 with ammonium hydroxide. Lower deposition rates resulted when the pH was adjusted with sodium hydroxide. Typical electroless Co plating bath compositions are presented in Table 1. Bath 1 was used to produce semi-bright Co deposits containing approximately 4 % phosphorus. Deposition rates were increased with increased bath pH or hypophosphite concentration. The deposition rate in bath 1 was increased from 5.6 lm · h–1 to 10 lm · h–1 at 85 8C by simply adding organosulfur compound of 0.2 g · cm–3 thiourea. Bath 2 contained less citrate than bath 1, which resulted in substantially greater deposition rate, but the resulting deposits had inferior physical properties. Bath 3 was a typical acid electroless Co plating bath, using reducing agents that were incapable of yielding Co deposits from acidic solution. Using the electroless process, WC–Co composites were success- Table 1. The bath composition for electroless cobalt deposition, based on 10 min deposition on palladium activated Mylar. Bath constituents as grams of substances per liter* Cobalt chloride, CoCl2 · 6H2O Cobalt sulfate,CoSo4 · 7H2O Sodium hypophosphite, NaH2PO2 · H2O Dimethylamine Borane, (CH3)2NHBH3 Sodium citrate, Na3C6H5O7 · 2H2O Ammonium chloride, NH4Cl Ammonium sulfate, (NH4)2SO4 Sodium sulfate, CI2H25O4SNa Sodium succinate, C4H4Na2O4 · 6H2O Sodium sulfate, Na2SO4 pH Temperature (8C) Deposition rate (lm · h–1) Bath 1 2 3 30 – 20 – 84.5 50 – – – – 9.5 92 6.8 30 – 20 – 29.6 50 – – – – 9.5 92 15 – 25 – 4 – – – – 25 15 5 70 13 * All chemical solutions were prepared in Port-Said University chemical laboratory. 558 Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites fully manufactured by coating the WC particles with the smaller Co particles and then using the powder metallurgy compaction pressure densification technique. Some of the reported advantages of this method are significant cost reduction, quality improvement of the deposited materials, and elimination of cross contamination [25, 26]. No organic binder or lubricant were used in this study, acetone was used to remove any organic matter, such as grease or oil, and finally sodium hydroxide of 50 % concentration was used for further cleaning the powder and to etch the surfaces. Now after having powders of WC–Co with two different compositions the final step was to prepare standard cylindrical and rectangular specimens for material characterization and testing [10 – 12]. A certain amount of the produced mixture was uniaxially pressed with 30 MPa compaction pressure to form the desired standard sample size. The cylindrical samples were manufactured with a diameter of 10 mm, whereas the rectangular samples were manufactured with a cross-section of 10 mm · 30 mm. After compaction, the green compacts were then heated in a closed vacuum furnace for 3 h at 350 8C to completely remove moisture and volatiles. The capacity of the vacuum furnace used for densification of the powders is 50 kg with a control cabinet consisting of graphite heating elements of temperature up to 1 600 8C and vacuum down to 10–3 bar. The temperature was increased to the desired sintering temperatures of 1 100, 1 250, 1 350, and 1 500 8C in the same furnace under the same conditions, at a rate of 5 K · min–1. At the sintering temperatures, the furnace conditions were fixed for one hour, to ensure homogeneous heat distribution over all the samples. Then the furnace was cooled with the same slow rate of 5 K · min–1until laboratory conditions were reached. tungsten carbide with the nano-scale cobalt particles after the electroless precipitation. Note the uniformity of the shape of the coated particles and the size of the Co particles in the nano-range, as presented with a higher magnification in Fig. 1b. After heat treatment of the WC–Co powders, X-ray diffraction was conducted for phase identification. X-ray diffraction (XRD) provides information regarding the composition of mixtures of crystalline materials, detailed structural information of the atomic arrangements, bond lengths, and bond angles, as well as macroscopic information like crystallite size and relative component concentrations. X-ray diffraction using an HZG4 diffractometer, Bruker D8, with Cu-Ka radiation and secondary monochromator (KV 40) was used. The XRD spectrum shows sharp peaks at 2h varying between 20 – 808, as displayed in Fig. 2 with peaks indicated for both tungsten carbide (WC), and cobalt (Co) that confirm the preserved crystal structure after the electroless process. Consolidation of the WC–Co composites coated powders was carried out using cold compaction, followed by furnace sintering at 1 100, 1 250, 1 350, and 1 500 8C with a holding time of 1 h at each temperature. 2.3. Density and microhardness measurements Significant densification of WC–Co composites was observed to begin at about 1 000 8C and greater than 90 % of the theoretical density can be obtained at 1 300 8C. Lique- 2.2. SEM and XRD measurements Scanning electron microscopy (SEM) investigation was conducted on the WC–Co composites using during the preparation process. Microscopic examination was performed using a Jeol 5400 SEM unit with a Link energy dispersive X-ray spectroscopy (EDS) detector attachment to observe the particle morphology, particle size, particle shape, and agglomeration of particles. Figure 1a shows the fully coated Fig. 2. X-ray diffraction pattern for 2h varying between 20 – 808 with peaks indicated for both tungsten carbide and cobalt. Fig. 1. (a) Tungsten carbide fully coated with nano-scale cobalt particles after electroless precipitation, and (b) higher magnification of the cobalt particles which were in the nanorange. Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 559 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. faction of the Co phase occurs between 1 280 8C and 1 350 8C as a result of the formation of a WC–Co eutectic or eutectic-like phase. Essentially, the part becomes fully dense soon after the liquid phase forms. During densification, the pressed part undergoes a volumetric shrinkage of about 40 to 45 %. It has been reported [27, 28] that complete densification could be accomplished for WC with about 3 to 25 wt.% of Co, whereas full densification was very difficult to achieve for a Co content of less than 3 %. For Co contents greater than 25 %, part shape was difficult to maintain. This was due to the presence of high liquid phase fraction, which allows the part to sag [27, 28]. The densities of the powders sintered at different temperatures were measured according to MPIF Standards 42 [29], using Archimedes’ rule, in which the sintered specimen was weighed in air. The sample was then lowered into water measuring the volume to ascertain density, assuming closed porosity in water. Then the densifications were computed according to: Densification ¼ ðqs qg Þ=ðqt qg Þ ð1Þ where qg, qs, and qt are the density of green compacts, sintered density, and theoretical density, respectively, as shown in Fig. 3. The theoretical density of consolidated WC–Co composites was calculated using the inverse rule of mixture: qWC Co ¼ 1=½1=qWC þ f ð1=qCo 1=qWC Þ ð2Þ where f is the volume fraction of Co and qWC-Co, qWC, qCo were the densities of WC–Co, WC, and Co, respectively, as plotted in Fig. 3. Vickers’ microhardness measurements were carried out for the WC–Co composites at the four sintering temperatures ensuring consistency throughout the material surface and homogeneity. An HM-210/220 model tester with an xy stage unit, and with a digital micrometer head indenter was used for measuring the microhardness at a load of 1 kg for 20 sec time, and the speed of the indenter was 100 lm · sec–1. 2.4. Wear measurements The wear resistance of the WC–Co composites was investigated using a pin-on-disk abrasive wear tester [30]. The wear specimens were 8 mm in diameter and 12 mm in length. Surface preparation was conducted before the wear test, where each specimen was ground with 1 lm alumina suspension. Wear tests were conducted under dry sliding conditions, applied loads of 10 N, and a constant sliding speed of 1.8 m · s–1. The time of wear was 10 min, and the track diameter was 120 mm for all samples. Wear loses were obtained by calculating the weight loss of the specimens before and after the testing using an electronic balance with sensitivity of 0.1 g–3. The samples were cleaned in an acetone bath and dried using hot air before the tests to remove organic substances. 2.5. Mechanical uniaxial compression test Cylindrical samples were prepared with dimensions of 10 mm in diameter and 8 mm in height for the compression test. The tests were carried out at room temperature using an MTS Testing Machine (Model 610) fitted with a 160 kN load cell operating in the displacement control mode. The stress–strain responses of WC–Co composites of different sintering temperatures were measured from uniaxial compression testing performed accordingly to ASTM standard E-9 for metals. The cross-head speed was 1 mm · min–1. SEM investigation was conducted on the fractured WC–Co composites after the compression testing, using the previously described tool, to explain the fracture phenomena. 2.6. Transverse rupture strength measurements Transverse rupture strength (TRS) test is used as an indicator of the quality of the sintered WC–Co composites, as its value will be very inconsistent for high porosity levels. Using three-point bending test the TRS could be used for determining the fracture strength. All WC–Co composite samples were compacted into rectangular specimen dimensions of 30 · 12 · 6.0 mm, at a compaction pressure of 30 MPa. The three point bending test was carried out at room temperature using an MTS Testing Machine (Model 610). 3. Results 3.1. Microhardness results Fig. 3. The specific density, densification, and Vickers microhardness measurements for WC-10wt.%Co, and WC-20wt.%Co composites varying with sintering temperatures (8C). 560 Vickers’ microhardness measurements were made and are plotted in Fig. 3. A minimum of five readings were taken for each case and the average value is plotted. It was observed that the specific density increases with the fraction of WC, this is due to the higher density of WC compared to the density of Co. Hardness increases with the volume fraction of WC, this was also due to the hardness of WC being significantly higher than the hardness of Co. At a given sintering temperature, densification increases when the volume fraction of Co-rich binder increases. The WC10wt.%Co composite had higher specific density and Vickers microhardness measurements than for the WC20wt.%Co composite, whereas the densification of the composites of WC-10wt.%Co was lower than for the comInt. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites posites of WC-20wt.%Co at different sintering temperatures, which reflects the strong bonding between the particles at high sintering temperature. IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. 3.2. Wear resistance The wear measurements for the WC–Co composites under abrasive sliding contact were plotted in terms of the volume losses versus the abrasive content for the composites with WC-10wt.%Co, and WC-20wt.%Co at the four different sintering temperatures (1 100, 1 250, 1 350, and 1 500 8C), as presented in Fig. 4. The wear depends on the time that the tungsten carbide takes to move from its location. Predominantly, this time depends on the adhesion force between the tungsten carbide and the cobalt. The volume loss of WC–Co composites in the case of higher binder content increases in an almost linear manner depending on the increase in the amount of abrasive. Composites with WC-10wt.%Co had more resistance to abrasion and wear, as was evident through the material hardness that was also affected by the level of porosity and microstructure. Figure 5 shows SEM micrographs of the worn surfaces for the two WC–Co composites sintered at temperature (a) of 1 100 8C after the wear test. Note that the pores were clearly visible in the higher magnifications. Therefore, the wear resistance of the WC-10wt.%Co composites was higher compared to WC-20wt.%Co composites. Figure 6 presents SEM photographs of the worn surfaces for the two WC–Co composites sintered at temperature of 1 500 8C after the wear test. 3.3. Mechanical uniaxial compression The tests were repeated with three samples for each experiment. The stress–strain responses for the four different sintering temperatures are presented in Fig. 7. For the two WC–Co composites the stress–strain response was found to increase monotonically with the increase in the sintering temperatures, reflecting material hardness and wear resistance behavior. 3.4. Transverse rupture strength The rupture test was performed using the compression testing machine according to MPIF [13, 31]. The specimens were placed between the platens of the three-point-bending (b) Fig. 4. The volume losses measurements for (a) WC-10wt.%Co, and (b) WC-20wt.%Co composites at the four different sintering temperatures (8C). (a) (b) Fig. 5. SEM images of the worn surface of (a) WC-10wt.%Co composite and with a higher magnification, and (b) WC-10wt.%Co composite and with a higher magnification at the upper right corner, after wear test at a sintering temperature of 1 100 8C. Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 561 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites testing machine and the load was applied at a constant rate of 2.5 mm · min–1, until the specimens fractured. The TRS for the bulk samples was calculated according to: TRS ¼ ð3 P LÞ=ð2 t2 ð3Þ WÞ where TRS is in GPa, P is the fracturing of rupture force in N, L is the distance between the supporting rods as 25.4 mm, t is the thickness of the sample in mm, and W is the width of the sample in mm. The TRS measurements for WC-10wt.%Co and WC-20wt.%Co composites at the four different sintering temperatures are plotted in Fig. 8. IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. 3.5. Fracture surface analysis Tungsten carbide particles coated with cobalt were relatively small, irregularly shaped, and tending to agglomerate. The Co phase is dispersed with many pools or lakes present in the compression fractured samples. The increase in hardness may be attributed to the process of continuous crystallization during plastic deformation. WC–Co composites have been observed to undergo mechanically induced nano-crystallization. Nano-crystal precipitation in WC–Co composites was also observed within vein protrusions on the compression fracture (a) surface and along crack propagation paths, as well as within shear bands resulting from bending [18]. All the WC–Co composites show an apparently classical inclined fracture surface, about 458 to the applied stress axis, which was similar to that encountered for a variety of hard metals [32]. The WC-20wt.%Co composites possessed interesting fracture behavior presented in Fig. 9a. The WC-20wt.%Co composites compression fracture surface had a conical shape with a rough morphology fracture surface. Examination of the fracture surface of the WC–Co composite showed a peculiar anomaly. Deeper investigation of the surface showed a rough conical shape failure in the direction of the maximum shear stress. The formation of the single shear plane was due to the inhomogeneous deformation that leads to localized shear. The conical surface makes an angle of about 458 with the compressive stress axis. The conical surface had contained a layered morphology with one layer showing a well-defined vein pattern reported normally in bulk metallic composites [32] and the other layer with no noticeable characteristic features as shown in Fig. 9a. The layered structure was aligned in the direction of the maximum shear stress. A large number of traces of inhomogeneous shear bands appear in the outer surface of the cylinder as a sign of increased shear activity, as presented in Fig. 9b. (b) Fig. 6. SEM images of the worn surface of (a) WC-10wt.%Co composite and with a higher magnification, and (b) WC-10wt.%Co composite and with a higher magnification at the upper right corner, after wear test at sintering temperature of 1 500 8C. (a) (b) Fig. 7. Stress–strain curves for (a) WC-10wt.%Co, and (b) WC-20wt.%Co composites sintered at four different temperatures (8C). 562 Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites 4. Discussion 4.2. Density and microhardness 4.1. Composite preparation and XRD The WC-10wt.%Co composite had higher specific density and Vickers microhardness measurements than for the WC20wt.%Co composite, whereas the densification for the composites of WC-10wt.%Co was lower than for the composites of WC-20wt.%Co at different sintering temperatures, reflecting the strong bonding between the particles at high sintering temperature. This could be due to the reduction of spaces and pores with the increase in the sintering temperature. Furthermore, with the less Co wt.% in the compositions the resistance to penetration increases and the hardness increases as the dominant behavior of the WC. The homogeneity of the coated particles distributed in the composite was observed. The reason for this homogeneous distribution could be explained using the SEM images. Each WC particle was fully coated with nano-scale Co precipitation, and limited contact between the WC particles occurs, which reduces the debonding zones throughout the materials. IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. Manufacturing of two WC–Co composites using electroless precipitation at different sintering temperatures (1 100, 1 250, 1 350 and 1 500 8C) was successfully achieved. The chemical composition of the investigated materials was 90 wt.% of WC with 10 wt.% of Co, and 80 wt.% of WC with 20 wt.% of Co. SEM shows the fully coated tungsten carbide with the nano cobalt particles after the electroless precipitation. The XRD show sharp peaks at 2h varying between 20 – 808 with peaks indicated for both tungsten carbide (WC), and cobalt (Co) that confirm the preserved crystal structure after the electroless process. 4.3. Wear resistance Fig. 8. Transverse rupture strength measurements for the two WC–Co composites at the four different sintering temperatures (8C), where the error bars are from three measurements of each sample. The wear mechanism of the investigated materials consisted of plastic deformation of the cobalt binder phase and brittle cracking of carbide grains and the carbide network. When the binder content increased the wear increased. Materials with higher amounts of cobalt in the microstructure exhibit higher wear resistance due to the lower bulk hardness. For lower binder content, the brittle cracking of carbide grains was the dominant reason for wear. It was noticed that the number of grooves at the surface formed as a result of wear was lower than of those formed in case of composites sintered at temperature of 1 100 8C, presented previously in Fig. 5. This could be explained by the particulate bonding that increased for the composites sintered at temperature of 1 500 8C and the surface homogeneity that was observed in comparing Figs. 5 and 6. For WC–Co composites of comparable tungsten carbide grain size, hardness and abrasion resistance decreased with increasing the Co content. Additionally, it was evident that with the increase Fig. 9. SEM images of (a) compression fractured WC-10wt.%Co composites, (b) higher magnification at the edge, (c) higher magnification showing melting droplets morphology (d) higher magnification showing the irregularity and rough morphology of the fracture surface. Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 563 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites in the sintering temperature the weight loss decreased in both WC–Co compositions. WC–Co composites with high sintering temperature showed good wettability and less weight loss as the adhesion forces between the particles increases, the presence of the pores between the particles decreases. IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. 4.4. Mechanical compression and rapture strength The stress–strain responses for WC-10wt.%Co composites were higher than for WC-20wt.%Co composites. However, the strain range was somewhat the same for the two types of WC–Co composites. The highest elastic modulus value was found to be for WC-10wt.%Co composites at 1 500 8C sintering temperature as 660 GPa, followed by WC-20wt.%Co composites at 1 500 8C sintering temperature as 423 GPa. The lowest elastic modulus values was found to be for WC10wt.%Co composites at 1 100 8C sintering temperature as 205 GPa, followed by WC-20wt.%Co composites at 1 100 8C sintering temperature as 115 GPa. Approximately 75 % reduction in the elastic modulus between WC–Co composites sintered at 1 500 8C and 1 100 8C was found. During the three point bending test, a specimen would experience plastic deformation, crack initiation, and crack growth before fracture even though the plastic deformation of WC–matrix was usually small and difficult to detect when binder content was very low. Therefore, TRS was defined as the maximum flow stress during the deformation, crack initiation, and crack propagation processes before the specimen fractures. As the effect of porosity was negligible, the crack initiation process would dominate when the hardness was very high and the fracture toughness was very low. Therefore, the higher the hardness, the higher the stress needed for crack initiation and hence the higher the TRS value. For the two WC– Co composites the TRS measurements were found to increase monotonically with the increase in the sintering temperatures, reflecting material hardness. The TRS for WC10wt.%Co composites was higher than for WC-20wt.%Co composites. The highest TRS value was found to be for WC-10wt.%Co composites at 1 500 8C sintering temperature as 2.758 GPa, whereas, the lowest TRS value was found to be for WC-20wt.%Co composites at 1 100 8C sintering temperature as 1.448 GPa. Approximately 25 % reduction in the TRS between WC-10wt.%Co composites and WC-20wt.%Co composites was evident. [34]. The viscosity of the WC–Co composites in the region where fewer nanoparticles exist will be lower. This leads to a less critical shear stress and a more ready plastic deformation. To ensure the reproducibility of the formation of the conical surface, several specimens of the WC–Co composites were tested and the phenomenon was confirmed. The hexagonal tungsten carbide crystals seemed to instigate local atomic rearrangements in regions of high plastic strain. This mechanical deformation explains the formation of the high strength crystalline composites as presented in a schematic diagram of the proposed mechanism for the formation of the conical fracture surface under compression testing shown in Fig. 10. Multishear planes with fracture angles deviating from 458 intersect at the center point (C), the angle difference leads to the development of two stress components parallel to the conical surface and a stress component perpendicular to it, which consecutively plays an important role in the fracture process. The force balance shown in Fig. 10 indicates that the cone surface in part (B) was subjected to a normal tensile stress in addition to the shear stress. As the crack propagates at the fracture surface between the two parts A and B, part A will be forced out leading to the formation of the cone. The difference in the height between the two layers indicates that extensive plastic deformation occurs on the micrometer scale, which was approximately *10 lm, it also reflected the difference in flow properties between each layer. The extensive plastic flow deformation and fracture behavior depends on the nature, size, volume fraction and distribution of the nanoparticles. The formation of mixed fracture mode implies that the conical shape morphology emerges for a critical volume fraction and size of nanoparticles. The size of the nanoparticles relative to the width of the shear bands was critical in understanding the inhomogeneity in the flow deformation and microcrack formation. The nanoparticles could explain the increase in the viscosity of the plastic flow, leading to particle strengthening of the composites. 4.5. Fracture surface morphology Evidence of melting spots at the end of crack path propagation on the top of each cone was noticed, as shown in Fig. 9c. These severe melting spots could be due to the release of the load at the final event of fracture in this limited area. The conical shape fracture surface was linked to the formation of the very fine nanoparticles in the composites. This enhances the homogeneity of the deformation, leading to the formation of multiple shear planes instead of single shear plane normally encountered in WC–Co composites as presented in Fig. 9d. A mechanism for the aggregation of small nanoparticles during deformation to a region with a high density of nanoparticles and another with lower density was presented by Nakamura and Gurland [33]. Moreover, the formation of nanoparticles in shear bands during nanoindentation was reported by Ingelstrom and Nordberg 564 Fig. 10. A schematic diagram of the normal (r) and shear (s) stress representation for the formation mechanism of conical fracture surface under compression testing. Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 S. Kaytbay, M. El-Hadek: Wear resistance and fracture mechanics of WC–Co composites IJMR downloaded from www.hanser-elibrary.com by Harvard on June 7, 2014 For personal use only. 5. Conclusions The effect of different WC compositions at different sintering temperatures on the wear resistance, and fracture mechanical properties were examined. X-ray diffraction examination suggests a preserved crystal structure after the electroless process. The tungsten carbide particles were shown to be fully coated with the cobalt polygon particles. The specific density, densification, and Vickers microhardness measurements were found to increase with the increase in the sintering temperature for both the WC–Co compositions. The WC-10wt.%Co composites were found to have a higher specific density and Vickers microhardness measurements than WC-20wt.%Co composites, whereas the densification for WC-10wt.%Co composites was lower than WC20wt.%Co composites at all the different sintering temperatures. WC-10wt.%Co composites were found to be more resistant to abrasion and wear. This was evident through the material hardness as it was affected by the level of porosity and microstructure for the composites. For the two WC–Co composites the stress–strain response was found to increase monotonically with the increase in the sintering temperatures. The stress–strain response for WC-10wt.%Co composites was higher than WC-20wt.%Co composites. For the two WC–Co composites the TRS measurements were found to increase monotonically with the increase in the sintering temperatures, reflecting material hardness and wear resistance behavior. The TRS for WC-10wt.%Co composites was higher than for WC-20wt.%Co composites. Fractographical SEM analysis of the fracture surface exposed a rough characteristic conical shape failure in the direction of the maximum shear stress. This could be due to the inhomogeneous deformation that leads to localized shear. A proposed mechanism for the formation of the conical fracture surface under compression testing was successfully presented. References [1] A. Gabriel, H. Pastor, D.M. Deo, S. Basu, C.H. Allibert: Sintering 85 (1987) 379. DOI:10.1007/978-1-4613-2851-3_41. [2] W.D. Schubert, H. Neumeister, G. Kinger, B. Lux: Int. J. Refract. Metals. Hard. Mater. 16 (1998) 133. DOI:10.1016/S0263-4368(98)00028-6. [3] G.S. Upadhyaya: Mater. Des. 22 (2001) 483. DOI:10.1016/S0261-3069(01)00007-3. [4] C. Pascal, A. Thomazic, A. Antoni-Zdziobek, J.-M. Chaix: Int. J. Mater. Res. 103 (2012) 296. DOI:10.3139/146.110647. [5] L.J. Prakash: Int. J. Refract. Metals. Hard. Mater. 13 (1995) 257. DOI:10.1016/0263-4368(95)92672-7. [6] E.O. Correa, J.N. Santos, A.N. Klein: Int. J. Mater. Res. 102 (2011) 1369. DOI:10.3139/146.110593. [7] G. Gille, B. Szesny, K. Dreyer, H. van den Berg, J. Schmidt, T. Gestrich, G. Leitner: Int. J. Refract. Metals. Hard. Mater. 20 (2002) 3. DOI:10.1016/S0263-4368(01)00066-X. [8] H.C. Kim, I.K. Jeong, I.J. Shon, I.Y. Ko, J.M. Doh: Int. J. Refract. Metals. Hard. Mater. 25 (2007) 336. DOI:10.1016/j.ijrmhm.2006.09.001. [9] U. Beste, S. Jacobson: Wear 264 (2008) 1129. DOI:10.1016/j.wear.2007.01.030. [10] K. Frisk, A. Markström: Int. J. Mater. Res. 99 (2008) 287. DOI:10.3139/146.101632. [11] F.R.N. Nabarro, B.A. Bilby: Int. J. Mater. Res. 93 (2002) 602. DOI: http://www.ijmr.de/MK020602. [12] M. El-Hadek, S. Kaytbay: Int. J. Mech. Mater. Des. 4 (2008) 63. DOI:10.1007/s10999-008-9058-2. [13] M. El-Hadek, S. Kaytbay: Strain 45 (2009) 506. DOI:10.1111/j.1475-1305.2008.00552.x. Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6 [14] M. El-Hadek, S. Kaytbay: Mater. Manuf. Processes. 28 (2013) 1003. DOI:10.1080/10426914.2012.736662. [15] C.W. Morton, D.J. Wills, K. Stjernberg: Int. J. Refract. Metals. Hard. Mater. 23 (2005) 287. DOI:10.1016/j.ijrmhm.2005.05.011. [16] R.W. Armstrong: Mater. 4 (2011) 1287. DOI:10.3390/ma4071287. [17] K. Choi, N.M. Hwang, D.-Y. Kim : J. Powder. Metall. 43 (2000) 168. DOI:10.1179/003258900665943. [18] L.S. Sigl, H.E. Exner: Metall. Mater. Trans. A. 18 (1987) 1299. DOI:10.1007/BF02647199. [19] J.L. Chermant, F. Osterstock: J. Mater. Sci. 11 (1976) 1939. DOI:10.1007/BF00708272. [20] M.T. Laugier: J. Mater. Sci. Lett. 4 (1985) 207. DOI:10.1007/BF00728078. [21] D.A. Stewart, P.H. Shipway, D.G. McCartney: Wear 225 (1999) 789. DOI:10.1016/S0043-1648(99)00032-0. [22] Y.C. Zhu, C.X. Ding, K. Yukimura, T.D. Xiao, P.R. Strutt: Ceramics. Int. 27 (2001) 669. DOI:10.1016/S0272-8842(01)00016-5. [23] X. Zhang, S. Yang, C. Zhao, Y. Lu, X. Liu, C. Wang: Int. J. Mater. Res. 104 (2013) 836. DOI:10.3139/146.110935. [24] \Sandvik" is an engineering group in tooling, materials technology, mining and construction, http://www.sandvik.com/en/, (Access July 2013). [25] Y. Shacham-Diamand, V.M. Dubin: Microelectron. Eng. 33 (1997) 47. DOI:10.1016/S0167-9317(96)00030-5. [26] S. Shukla, S. Seal, Z. Rahaman, K. Scammon: Mater. Lett. 57 (2002) 151. DOI:10.1016/S0167-577X(02)00722-X. [27] P. Arató, L. Bartha, R. Porat, S. Berger, A. Rosen: Nanostruct. Mater. 10 (1998) 245. DOI:10.1016/S0965-9773(98)00067-1. [28] R. Frykholm, M. Ekroth, B. Jansson, J. Ågren, H.-O. Andrén: Acta Materialia 51 (2003) 1115. DOI:10.1016/S1359-6454(02)00515-3. [29] \MPIF Standard 42. Determination of density of compacted or sintered metal powder products", Standard test methods for metal powders and powder metallurgy products, Metal Powder Industries Federation, Princeton, NJ, USA (2000). [30] S. Guicciardi, C. Melandri, F. Lucchini, G. de Portu: Wear. 252 (2002) 1001. DOI:10.1016/S0043-1648(02)00066-2. [31] \MPIF standard 41: Determination of transverse rupture strength of sintered metal powder test specimen", Standard test methods for metal powders and powder metallurgy products, Metal Powder Industries Federation, Princeton, NJ, USA (1991). [32] B. Roebuck, E.A. Almond: Int. Mater. Rev. 33 (1988) 90. DOI:10.1179/095066088790324094. [33] M. Nakamura, J. Gurland: Metall. Mater. Trans. A 11 (1980) 141. DOI:10.1007/BF02700449. [34] N. Ingelstrom, H. Nordberg: Eng. Fract. Mech. 6 (1974) 597. DOI:10.1016/0013-7944(74)90016-2. (Received October 11, 2013; accepted December 17, 2013; online since April 24, 2014) Correspondence address Prof. Eng. Medhat Awad El-Hadek Department of Production and Mechanical Design Faculty of Engineering Port Fouad 42523 Port-Said Egypt Tel.: +2 0100 827 1778 Fax: +2 066 3400 936 E-mail: melhadek@eng.psu.edu.eg Bibliography DOI 10.3139/146.111069 Int. J. Mater. Res. (formerly Z. Metallkd.) 105 (2014) 6; page 557 – 565 # Carl Hanser Verlag GmbH & Co. KG ISSN 1862-5282 565