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Microstructure and mechanical properties of WC Ni–Si based cemented carbides developed by powder metallurgy
Int. Journal of Refractory Metals and Hard Materials 28 (2010) 572–575
Contents lists available at ScienceDirect
Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Microstructure and mechanical properties of WC Ni–Si based cemented carbides developed by powder metallurgy E.O. Correa a,⁎, J.N. Santos a, A.N. Klein b a b
Universidade Federal de Itajuba, Instituto de Engenharia Mecânica, Av. BPS, 1303, Pinheirinho, Itajubá, Minas Gerais 37500-903, Brazil Universidade Federal de Santa Catarina, Departamento de Engenharia Mecânica, CP 476, Trindade, Florianópolis, Santa Catarina 88040-970, Brazil
a r t i c l e
i n f o
Article history: Received 12 March 2010 Accepted 8 April 2010 Keywords: WC–Ni–Si cemented carbides Microstructure Mechanical properties
a b s t r a c t In this paper the influence of the Ni binder metal and silicon as an additional alloying element on the microstructure and mechanical properties of WC-based cemented carbides processed by conventional powder metallurgy was studied. Microstructural examinations of specimens indicated the presence of a very low and even distributed porosity and the presence of islands of metal binder in the microstructure of the cemented carbides. Furthermore, despite the addition of silicon and carbon in the cemented carbides, it was not observed the presence of small fractions of undissolved SiC and free graphite nodules in their microstructure. Vickers hardness and Flexural strength tests indicated that the cemented carbide WC–Ni–Si with 10 wt.% of binder presented bulk hardness similar to the conventional WC–Co cemented carbides and superior flexure strength and fracture toughness. © 2010 Elsevier Ltd. All rights reserved.
1. Introduction Cemented carbides are widely used in metal cutting, molds, mineral and petroleum industries and, more recently, as catalyst electrodes in fuel cells and coatings for aerospace components, due to their excellent properties such as high hardness, high hot-hardness, high wear/corrosion resistance and low thermal expansion coefficient [1,2]. Due to their technological importance, these composites of WC have been subject to a great deal of investigation in order to optimize the compositions and the processing route leading to the highest mechanical properties and a reduction of the manufacturing cost [3]. Most cemented carbides of tungsten utilize cobalt as the metal binder due to its excellent wetting, adhesion and adequate mechanical properties [4]. However, three main reasons exist to substitute it as the metal binder in these composites: (1) primary cobalt is a high-priced commodity, (2) cobalt released by the wear and corrosion of components is largely responsible for occupational diseases of plant maintenance personnel and (3) low corrosion resistance of the WC–Co cermet [2,5]. Because of the reasons mentioned above, considerable research efforts have been aimed at finding a satisfactory alternative binder phase [6,7]. It has been shown that by using Ni instead of Co, besides solving the two first problems mentioned, the corrosion and oxidation resistance of the resulting cemented carbide is improved; however, the hardness is somewhat lower than those of WC–Co [8].
⁎ Corresponding author. Tel.: +55 35 3629 1298; fax: +55 35 3629 1148. E-mail address: [email protected] (E.O. Correa). 0263-4368/$ – see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2010.04.003
In order to overcome this deficiency, strengthening of the Ni–W–C alloys by using solid solution techniques have been studied once nickel is capable to dissolve substantial alloying additions without losing the advantage of long-range order [9]. The phase diagrams indicate that alloying elements such as Si, Al, Mn, Cr, Nb and Fe can produce the strengthening of the nickel by solid solution. In the present study, silicon was chosen as the strengthener element because it can be dissolved in nickel in higher amounts (approximately 5 wt.%). Also, a significant difference in the radii of these elements (ratio is approximately 1.13), should result in the strengthening of nickel by silicon. Furthermore, silicon acts as a melting point depressant when dissolved in nickel, which allows the decreasing of the sintering temperature. The purpose of this study is therefore to evaluate the influence of the Ni binder metal strengthened with Si on the microstructure and mechanical properties of WC-based cemented carbide processed by powder metallurgy. 2. Experimental procedure Mixtures of tungsten carbide powder (an average particle size of 2.5 µm), nickel oxide powder (an average particle size of 4 µm), pure carbon black powders (an average particle size of 2 µm) and 99% pure silicon carbide (an average particle size of 2 µm) were mixed in heptane with a ball-to-powder weight ratio and a powder-to-heptane weight ratio of about 2:1. The milling was performed at a horizontal rotation velocity of 50 rpm for 80 h. After milling, due to the presence of nickel oxide in the mixtures, an oxygen reducing treatment in flowing dry hydrogen was carried out at 750 °C for 1 h. After that, it was added in the mixture 1.5 wt.% of paraffin wax dissolved in
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Fig. 2. Micrograph of the polished surface of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal (94.1% Ni + 4.1% Si + 1.8% C) sintered under high vacuum at 1460 °C during 1 h.
Fig. 1. Thermal cycle of sintering.
heptane to improve its compaction. To form samples of five different binder contents and of size 8 × 10 × 50 mm, appropriate quantity of mixed powder was weighed out and poured into a stainless steel die and then cold pressed at 130 MPa for 3 min. Pre-sintering operation was carried out in a pure hydrogen atmosphere according to the thermal cycle shown in Fig. 1 (a). After that, pre-sintered samples of high density were machined into rectangular beams with nominal dimension of 5 mm × 6 mm × 20 mm to assess the flexural rupture resistance of cemented carbides studied. Transient liquid phase (TLP) sintering was carried out in a high vacuum atmosphere (2 to 6 × 10−5 bar) according to the thermal cycle shown in Fig. 1 (b) to complete the metallurgical bonds between powder particles. Flexural strength testing was done in three-point bending with rollers span of 15 mm. The samples dimensions (5mm×6.25mm ×19 mm) were in accordance with ASTM 8865 standard and the testing procedure was in accordance with the ASTM B-406 standard. Additionally, sintered samples were cut with a diamond saw, embedded in resin, grounded and polished with diamond suspensions. Bulk Vickers hardness measurements of samples were carried out on a hardness tester using an indenting load of 30kgf, as recommended by ISO 3878. Fracture toughness was determined by means of the palmqvist indentation cracking test using an indenting load of 100kg. Optical and SEM imaging were performed on polished and etched surfaces (Murakami's reagent; ∼2 min) to characterize the microstructure and to analyze the fracture region. 3. Results and discussion Table 1 presents the set of WC–Ni–Si cemented carbides with five different binder contents developed in this study and their significant attributes. The chemical composition of the binder metal shown in table is related to the powder mixture composition before sintering. The final Ni binder composition is obtained only after sintering by the inter-diffusion of the mixture components. It is worthwhile mention-
ing that the binder phase solubilizes a certain amount of W and C from the WC besides silicon added in the mixture in form of SiC.
3.1. Microstructural characterization Figs. 2 and 3 show the polished surface of the cemented carbide WC–10(Ni–4.1%Si), which presented a better correlation between hardness and flexural strength. Microstructures quite similar to those ones are also observed in cemented carbides with cobalt as the binder alloy. It can be seen from the Fig. 2 that there was no significant presence of pores in the microstructure. The relative density of the cemented carbides, obtained by the Arquimedes method, varied between 98.3 and 98.8% of the theoretical density. Additionally, the pores were small and uniformly distributed in the matrix. However, a lower uniformity of the binder distribution is detected in the microstructure in comparison with that observed in the cobaltbased counterpart, showing an accumulation of the binder phase in some well dispersed regions (see Fig. 3). This may be attributed to the lower wetting and WC dissolution rate of the Ni binder when compared with Co binder; which leads to a deficient spreading of the Ni liquid phase between the grains as it dissolves WC. In addition, the lower chemical homogeneity in the prepared compacts, as a result of insufficient mixing conditions also may contribute significantly for this uneven distribution of the binder metal [10,11]. However, it is worthwhile mentioning that, differently from the porosity, the presence of these islands of binder phase does not produce a deleterious effect on the mechanical properties of the cemented carbide.
Table 1 Chemical composition and properties of the cemented carbides. Alloy
1 2 3 4 5
WC (wt.%)
94 92 90 88 86
Binder (wt.%)
6 8 10 12 ∼14
Mechanical properties
Binder composition (wt.%) Ni
Si
C
Hardness HV30
Flexural strength (N/mm2)
94.1 94.1 94.1 94.1 94.1
4.1 4.1 4.1 4.1 4.1
1.8 1.8 1.8 1.8 1.8
1510 ± 10 1436 ± 26 1362 ± 30 1220 ± 20 1190 ± 25
1700 ± 162 1890 ± 222 2018 ± 208 2784 ± 279 2610 ± 197
Obs. (1) the hardness values shown in the Table 1 are the average of 10 measurements; and (2) the flexural strength values shown in the table are the average of five flexural test results.
Fig. 3. Micrograph of the polished surface of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal (94.1% Ni + 4.1% Si + 1.8% C) showing the presence of islands of binder metal in the microstructure.
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cemented carbides during flexural tests was controlled by the mechanical properties of the binder phase and by the degree of “squeletum formation” or contiguity of the WC grains. In fact, this observations is in a good agreement with findings of several authors [15–17] who proposed that as the rigid continuous “squeletum” of WC particles is submitted to plastic deformation, it appears high stresses in the binder phase. As a result, the rupture of the binder layers occurs between the WC particles due to the rapid nucleation and propagation of microcracks in the WC–WC interface of the squeletum [15–17]. Investigations of the fracture surface also showed that there was not an excessive grain growth of WC particles for the temperature and time of sintering adopted. 3.2. Mechanical characterization Fig. 4. Micrograph of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal after etching. Etching: Murakami reagent.
Fig. 4 shows the micrograph of the cemented carbide WC–Ni–Si after etching. It can be observed that the regions where there are no accumulation of binder phase presented a shorter mean free path (0.6 to 0.75 µm) and a finer dispersion, similar to that observed in cemented carbides WC–10Co. Energy-dispersive spectrometry (EDS) microanalyses showed that the chemical composition of the Ni binder phase was approximately 89.3 wt.% Ni, 6.6 wt.% W, 3.90 wt.% Si and traces of carbon. From the EDS microanalyses results obtained so far, together with information found in the literature [2,3,7,12] is reasonable to suggest that the Si, W and C elements were in solid solution in the Ni phase. Additionally, the amount of W dissolved in the binder metal (approximately 6 wt.%) is higher than that found by Fiedler et al. [13] for pure Ni binder metal (approximately 4.6 wt.%), indicating that the silicon contributed to the higher dissolution of W in the Ni binder metal. In fact, during the microstructural examinations, it was not detected the presence of undissolved SiC particles as well as the presence of small free graphite nodules in the microstructure. In addition, the absence of these phases in the microstructure indicates that there was a satisfactory sinterability of the cemented carbide for the Si and C contents and sintering parameters adopted. The micrograph of the fracture surface of the WC–Ni–Si cemented carbide with 10 wt.% of binder is presented in Fig. 5. Qualitative analysis of this SEM micrograph indicates that the fracture mode through the binder phase, characterized by deep dimples and the fracture mode along the carbide–carbide interface, identified by a smooth carbide surface with no visible plastic, were predominant [14]. Also, these analysis suggest that the fracture resistance of these
The results presented in Table 1 show that, for a WC grain size essentially constant in 2.5 µm, the hardness and flexural strength of the developed cemented carbides is highly dependent of the binder metal content. As can be seen, the hardness decreased continuously with increasing binder phase content whereas, the flexural strength increased. According to the literature [7], this is because part of the hard WC phase is replaced by the ductile nickel binder as the binder content is increased. It can also be observed from the results that the cemented carbide with 10 wt.% of binder alloy presented a relatively better compromise between hardness and flexural strength. Fig. 6 shows the hardness and flexural strength of the cemented carbide WC–Ni–Si with 10 wt.% of metal binder. For comparison, data for WC–Co and WC–Ni materials with comparable binder content and WC particle size from the literature are also shown [18]. As can be observed from the figure, the hardness of cemented carbide WC–Ni–Si is similar to comparable WC–Co cemented carbide and quite superior to comparable WC–Ni cemented carbide. This result demonstrates the effective action of silicon as strengthener by solution in Ni leading to an overall improvement of the mechanical properties of the material; mainly when it is compared with the WC–Ni alloy. According to the literature [19], the most widely accepted mechanism of solid solution strengthening of the nickel binder by silicon arises from the elastic interaction between the strain field of the solute (Si) and that of dislocations. The strain field of the solute is attributed to the atomic size misfit between the solute and solvent. It is worthwhile mentioning that the hardness value of the WC–Ni–Si alloy slightly lower than that of the WC–Co may be attributed to the fact of the more ductile nickel binder presents a higher amount of plastic micro-deformations than the cobalt binder during the cooling. As a result, a considerable relief of residual stress (originated from the different values of linear dilatation coefficients of the hard phase and
Fig. 5. Fracture surface of WC–Ni–Si with 10 wt.% of binder metal. Fracture modes: carbide–carbide interface –C/C and fracture along the binder metal –B.
Fig. 6. Hardness and flexural strength of the WC–Ni–Si cemented carbide with 10 wt.% of binder metal and WC particle size of 2.5 µm.
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In terms of industrial applications, the combination of hardness and toughness obtained for the cemented carbides WC–Ni–Si qualify them to be used to produce several forming molds and perforation tools. 4. Conclusions Based on the analyses and results obtained so far, it can be drawn the following conclusions:
Fig. 7. Data of Flexural strength versus hardness of WC–Ni–Si cemented carbides studied and WC–Co cemented carbides with WC particle size of 2.5 µm and binder metal content variable.
binder) of the alloy occurs together with the decreasing of the hardness [20]. An alternative explanation to the lower hardness of the WC–Ni–Si cemented carbide is that the Ni based binder usually is present in the mixture in higher volume fraction than that of cobalt once its density is lower. Concerning to the flexural strength of the cemented carbide WC–Ni–Si, it can be seen that it is quite superior to that of both WC–Co and WC–Ni materials. This result suggests that the addition of approximately 4.1 wt.% of silicon to strength the Ni binder by solid solution improved the hardness of the materials without sacrificing their ductility. In order to point out the good correlation between the properties of the developed cemented carbides, it is provided in Fig. 7, a comparison of the data of hardness versus flexural strength of the cemented carbides WC–Ni–Si obtained in this work with those of WC–Co material with comparable binder contents and WC particle size reported by literature [18]. It can be seen from the data shown in Fig. 6 that the cemented carbide WC–Ni–Si exhibits a higher flexural strength than that of conventional WC–Co for hardness values quite similar. In Table 2, it is provided a comparison of the mean fracture toughness value obtained for the cemented carbide WC–Ni–Si with binder content of 10 wt.% with those fracture toughness values of the WC–Co and WC–Ni materials with comparable binder content and WC particle size from the literature [14,21,22]. As expected, the cemented carbides with tough nickel binder presented higher fracture toughness than that of WC–Co alloy. However, as can also be seen from the results, even though there is the presence of silicon as a strengthener, the cemented carbide WC– Ni–Si exhibited fracture toughness quite similar to that of the WC–Ni cemented carbide. The results shown in Table 2 and Fig. 2 confirm that silicon produce an improvement of the hardness of the Ni binder without sacrificing considerably its toughness.
Table 2 Comparison of fracture toughness of WC–Ni–Si sintered in this study with previously reported values. Ref.
Binder content (wt.%)
WC grain Size (µm)
Vickers hardness
KIC (MPa m1/2)
[20,21] [22] This work
10Ni 10Co 9.59Ni–0.41Si
2.1–3.6 3.0 2.5
∼1032 ∼1200 ∼1362
14.7 ± 0.4 12.0 ± 0.8 13.8 ± 0.4
1) The solution strengthening of the nickel binder by silicon improved significantly the mechanical properties of the cemented carbides WC–Ni–Si in comparison with similar cemented carbides WC–Ni. 2) The cemented carbide developed presented flexural strength and fracture toughness higher than those observed in similar WC–Co cemented carbides without sacrificing significantly the hardness. This indicates that these materials may be more suitable for applications where the better balance between toughness and hardness is required. References [1] Zhong M, Liu W, Zhang Y, Zhu X. Formation of WC/Ni hard alloy coating by laser cladding of W/C/Ni pure element powder blend. Int J Refract Met Hard Mater 2006;24:453–60. [2] Shon Kim HC, IJ Yoon JK, Doh JM, Munir ZA. Rapid sintering of ultrafine WC–Ni cermets. Int J Refract Met Hard Mater 2006;24:427–31. [3] Fernandes CM, Senos AMR, Vieira MT, Antunes JM. Mechanical characterization of composites prepared from WC powders coated with Ni rich binders. Int J Refract Met Hard Mater 2008;26:491–8. [4] Jiang G, Zhuang H, Li W. Parameters investigation during simultaneous synthesis and densification WC–Ni composites by field-activated combustion. Mater Sci Eng A 2003;360:377–84. [5] Ohriner EK, Wada T, Whelan EP, Ocken H. The chemistry and structure of wearresistant, iron-base hardfacing alloys. Metall Trans A 1991;22:983–91. [6] Gille G, Bredthauer J, Gries B, Mende B, Heinrich W. Advanced and new grades of WC and binder powder – their properties and application. Int J Refract Met Hard Mater 2000;18:87–102. [7] Taheri-Nassaj E, Mirhosseini SH. An in situ WC–Ni composite fabricated by the SHS method. J Mater Process Technol 2003;142:422–6. [8] Almond EA, Roebuck B. Indentification of optimum binder phase compositions for improved WC hard metals. Mat Sci Eng A 1988;105:237–48. [9] Ohiai S, Oya Y, Suzuki T. Alloying behaviour of Ni3Al, Ni3Ga and Ni3G. Acta Metall 1984;32(2):289–98. [10] Lin C, Kny E, Yuan G, Djuricic B. Microstructure and properties of ultrafine WC– 0.6VC–10Co hardmetals densified by pressure-assisted critical liquid phase sintering. J Alloy Compd 2004;383:98–102. [11] Lisovskii AF, Vishneskii AS, Sirota KI. An investigation of sintered WC–Ni hard alloys with an uneven distribution of the binder metal. Powder Metall Mater, Parts Coat 1977;16(5):345–8. [12] Penrice TW. Some Characteristics of the binder phase in cemented carbides. Int J Refract Met Hard Mater 1997;15:113–21. [13] Flieder ML, Stadelmaier HH. The ternary system Nickel-tungsten-carbon. Z Metallkunde 1975;67:402–4. [14] Shatov AV, Ponomarev SS, Firstov SA. Fracture of WC–Ni cemented carbides with different shape of WC crystals. Int J Refr Met Hard Mater 2008;26:68–76. [15] Bengtsson B, Johannesson T, Lindau L. The effect of milling on the sintering of a WC + 6% Co powder. Planseeberichte fur Pulvermetallurgie 1974;22:285–91 In German. [16] Hara A, Nishikawa T, Yazu S. The observation of the fracture path in WC–Co cemented carbides using a newly developed replica method with na electron microscope. Planseeberichte fur Pulvermetallugie 1970;18:28–42 In german. [17] Gathwohl G, Warren R. The effect of cobalt content fo the microstructure of liquid phase sintered TaC-Co alloys. Mat Sci Eng A 1974;14:55–65. [18] Schatt W. Powder Metallurgy. Sintered and composite Materials. (Pulvermetallurgie. Sinter-und Verbundwerstaffe)Leipzig: VEB Deutscher Verlag fur Grundstoffindustrie; 1979. in German. [19] Chen K, Zhao LR. Application of bond border in solid solution strengthening of nickel. J Mater Sci Lett 2003;22:603–5. [20] Exner HE, Gurland J. The effect of small plastic deformations on the strength and hardness of a WC–Co alloy. J Mater 1970;5:75–85. [21] Wittiman B, Schubert WD, Lux B. WC grain growth and grain growth inhibition in nickel and iron binder hardmetals. Int J Refract Met Had Mater 2002;20:51–60. [22] Scieszka SF. Wear transition as a means of fracture toughness evaluation of hardmetals. Tribol Lett 2001;11(3–4):185–94.
Contents lists available at ScienceDirect
Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Microstructure and mechanical properties of WC Ni–Si based cemented carbides developed by powder metallurgy E.O. Correa a,⁎, J.N. Santos a, A.N. Klein b a b
Universidade Federal de Itajuba, Instituto de Engenharia Mecânica, Av. BPS, 1303, Pinheirinho, Itajubá, Minas Gerais 37500-903, Brazil Universidade Federal de Santa Catarina, Departamento de Engenharia Mecânica, CP 476, Trindade, Florianópolis, Santa Catarina 88040-970, Brazil
a r t i c l e
i n f o
Article history: Received 12 March 2010 Accepted 8 April 2010 Keywords: WC–Ni–Si cemented carbides Microstructure Mechanical properties
a b s t r a c t In this paper the influence of the Ni binder metal and silicon as an additional alloying element on the microstructure and mechanical properties of WC-based cemented carbides processed by conventional powder metallurgy was studied. Microstructural examinations of specimens indicated the presence of a very low and even distributed porosity and the presence of islands of metal binder in the microstructure of the cemented carbides. Furthermore, despite the addition of silicon and carbon in the cemented carbides, it was not observed the presence of small fractions of undissolved SiC and free graphite nodules in their microstructure. Vickers hardness and Flexural strength tests indicated that the cemented carbide WC–Ni–Si with 10 wt.% of binder presented bulk hardness similar to the conventional WC–Co cemented carbides and superior flexure strength and fracture toughness. © 2010 Elsevier Ltd. All rights reserved.
1. Introduction Cemented carbides are widely used in metal cutting, molds, mineral and petroleum industries and, more recently, as catalyst electrodes in fuel cells and coatings for aerospace components, due to their excellent properties such as high hardness, high hot-hardness, high wear/corrosion resistance and low thermal expansion coefficient [1,2]. Due to their technological importance, these composites of WC have been subject to a great deal of investigation in order to optimize the compositions and the processing route leading to the highest mechanical properties and a reduction of the manufacturing cost [3]. Most cemented carbides of tungsten utilize cobalt as the metal binder due to its excellent wetting, adhesion and adequate mechanical properties [4]. However, three main reasons exist to substitute it as the metal binder in these composites: (1) primary cobalt is a high-priced commodity, (2) cobalt released by the wear and corrosion of components is largely responsible for occupational diseases of plant maintenance personnel and (3) low corrosion resistance of the WC–Co cermet [2,5]. Because of the reasons mentioned above, considerable research efforts have been aimed at finding a satisfactory alternative binder phase [6,7]. It has been shown that by using Ni instead of Co, besides solving the two first problems mentioned, the corrosion and oxidation resistance of the resulting cemented carbide is improved; however, the hardness is somewhat lower than those of WC–Co [8].
⁎ Corresponding author. Tel.: +55 35 3629 1298; fax: +55 35 3629 1148. E-mail address: [email protected] (E.O. Correa). 0263-4368/$ – see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2010.04.003
In order to overcome this deficiency, strengthening of the Ni–W–C alloys by using solid solution techniques have been studied once nickel is capable to dissolve substantial alloying additions without losing the advantage of long-range order [9]. The phase diagrams indicate that alloying elements such as Si, Al, Mn, Cr, Nb and Fe can produce the strengthening of the nickel by solid solution. In the present study, silicon was chosen as the strengthener element because it can be dissolved in nickel in higher amounts (approximately 5 wt.%). Also, a significant difference in the radii of these elements (ratio is approximately 1.13), should result in the strengthening of nickel by silicon. Furthermore, silicon acts as a melting point depressant when dissolved in nickel, which allows the decreasing of the sintering temperature. The purpose of this study is therefore to evaluate the influence of the Ni binder metal strengthened with Si on the microstructure and mechanical properties of WC-based cemented carbide processed by powder metallurgy. 2. Experimental procedure Mixtures of tungsten carbide powder (an average particle size of 2.5 µm), nickel oxide powder (an average particle size of 4 µm), pure carbon black powders (an average particle size of 2 µm) and 99% pure silicon carbide (an average particle size of 2 µm) were mixed in heptane with a ball-to-powder weight ratio and a powder-to-heptane weight ratio of about 2:1. The milling was performed at a horizontal rotation velocity of 50 rpm for 80 h. After milling, due to the presence of nickel oxide in the mixtures, an oxygen reducing treatment in flowing dry hydrogen was carried out at 750 °C for 1 h. After that, it was added in the mixture 1.5 wt.% of paraffin wax dissolved in
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Fig. 2. Micrograph of the polished surface of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal (94.1% Ni + 4.1% Si + 1.8% C) sintered under high vacuum at 1460 °C during 1 h.
Fig. 1. Thermal cycle of sintering.
heptane to improve its compaction. To form samples of five different binder contents and of size 8 × 10 × 50 mm, appropriate quantity of mixed powder was weighed out and poured into a stainless steel die and then cold pressed at 130 MPa for 3 min. Pre-sintering operation was carried out in a pure hydrogen atmosphere according to the thermal cycle shown in Fig. 1 (a). After that, pre-sintered samples of high density were machined into rectangular beams with nominal dimension of 5 mm × 6 mm × 20 mm to assess the flexural rupture resistance of cemented carbides studied. Transient liquid phase (TLP) sintering was carried out in a high vacuum atmosphere (2 to 6 × 10−5 bar) according to the thermal cycle shown in Fig. 1 (b) to complete the metallurgical bonds between powder particles. Flexural strength testing was done in three-point bending with rollers span of 15 mm. The samples dimensions (5mm×6.25mm ×19 mm) were in accordance with ASTM 8865 standard and the testing procedure was in accordance with the ASTM B-406 standard. Additionally, sintered samples were cut with a diamond saw, embedded in resin, grounded and polished with diamond suspensions. Bulk Vickers hardness measurements of samples were carried out on a hardness tester using an indenting load of 30kgf, as recommended by ISO 3878. Fracture toughness was determined by means of the palmqvist indentation cracking test using an indenting load of 100kg. Optical and SEM imaging were performed on polished and etched surfaces (Murakami's reagent; ∼2 min) to characterize the microstructure and to analyze the fracture region. 3. Results and discussion Table 1 presents the set of WC–Ni–Si cemented carbides with five different binder contents developed in this study and their significant attributes. The chemical composition of the binder metal shown in table is related to the powder mixture composition before sintering. The final Ni binder composition is obtained only after sintering by the inter-diffusion of the mixture components. It is worthwhile mention-
ing that the binder phase solubilizes a certain amount of W and C from the WC besides silicon added in the mixture in form of SiC.
3.1. Microstructural characterization Figs. 2 and 3 show the polished surface of the cemented carbide WC–10(Ni–4.1%Si), which presented a better correlation between hardness and flexural strength. Microstructures quite similar to those ones are also observed in cemented carbides with cobalt as the binder alloy. It can be seen from the Fig. 2 that there was no significant presence of pores in the microstructure. The relative density of the cemented carbides, obtained by the Arquimedes method, varied between 98.3 and 98.8% of the theoretical density. Additionally, the pores were small and uniformly distributed in the matrix. However, a lower uniformity of the binder distribution is detected in the microstructure in comparison with that observed in the cobaltbased counterpart, showing an accumulation of the binder phase in some well dispersed regions (see Fig. 3). This may be attributed to the lower wetting and WC dissolution rate of the Ni binder when compared with Co binder; which leads to a deficient spreading of the Ni liquid phase between the grains as it dissolves WC. In addition, the lower chemical homogeneity in the prepared compacts, as a result of insufficient mixing conditions also may contribute significantly for this uneven distribution of the binder metal [10,11]. However, it is worthwhile mentioning that, differently from the porosity, the presence of these islands of binder phase does not produce a deleterious effect on the mechanical properties of the cemented carbide.
Table 1 Chemical composition and properties of the cemented carbides. Alloy
1 2 3 4 5
WC (wt.%)
94 92 90 88 86
Binder (wt.%)
6 8 10 12 ∼14
Mechanical properties
Binder composition (wt.%) Ni
Si
C
Hardness HV30
Flexural strength (N/mm2)
94.1 94.1 94.1 94.1 94.1
4.1 4.1 4.1 4.1 4.1
1.8 1.8 1.8 1.8 1.8
1510 ± 10 1436 ± 26 1362 ± 30 1220 ± 20 1190 ± 25
1700 ± 162 1890 ± 222 2018 ± 208 2784 ± 279 2610 ± 197
Obs. (1) the hardness values shown in the Table 1 are the average of 10 measurements; and (2) the flexural strength values shown in the table are the average of five flexural test results.
Fig. 3. Micrograph of the polished surface of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal (94.1% Ni + 4.1% Si + 1.8% C) showing the presence of islands of binder metal in the microstructure.
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cemented carbides during flexural tests was controlled by the mechanical properties of the binder phase and by the degree of “squeletum formation” or contiguity of the WC grains. In fact, this observations is in a good agreement with findings of several authors [15–17] who proposed that as the rigid continuous “squeletum” of WC particles is submitted to plastic deformation, it appears high stresses in the binder phase. As a result, the rupture of the binder layers occurs between the WC particles due to the rapid nucleation and propagation of microcracks in the WC–WC interface of the squeletum [15–17]. Investigations of the fracture surface also showed that there was not an excessive grain growth of WC particles for the temperature and time of sintering adopted. 3.2. Mechanical characterization Fig. 4. Micrograph of the cemented carbide WC–Ni–Si with 10 wt.% of binder metal after etching. Etching: Murakami reagent.
Fig. 4 shows the micrograph of the cemented carbide WC–Ni–Si after etching. It can be observed that the regions where there are no accumulation of binder phase presented a shorter mean free path (0.6 to 0.75 µm) and a finer dispersion, similar to that observed in cemented carbides WC–10Co. Energy-dispersive spectrometry (EDS) microanalyses showed that the chemical composition of the Ni binder phase was approximately 89.3 wt.% Ni, 6.6 wt.% W, 3.90 wt.% Si and traces of carbon. From the EDS microanalyses results obtained so far, together with information found in the literature [2,3,7,12] is reasonable to suggest that the Si, W and C elements were in solid solution in the Ni phase. Additionally, the amount of W dissolved in the binder metal (approximately 6 wt.%) is higher than that found by Fiedler et al. [13] for pure Ni binder metal (approximately 4.6 wt.%), indicating that the silicon contributed to the higher dissolution of W in the Ni binder metal. In fact, during the microstructural examinations, it was not detected the presence of undissolved SiC particles as well as the presence of small free graphite nodules in the microstructure. In addition, the absence of these phases in the microstructure indicates that there was a satisfactory sinterability of the cemented carbide for the Si and C contents and sintering parameters adopted. The micrograph of the fracture surface of the WC–Ni–Si cemented carbide with 10 wt.% of binder is presented in Fig. 5. Qualitative analysis of this SEM micrograph indicates that the fracture mode through the binder phase, characterized by deep dimples and the fracture mode along the carbide–carbide interface, identified by a smooth carbide surface with no visible plastic, were predominant [14]. Also, these analysis suggest that the fracture resistance of these
The results presented in Table 1 show that, for a WC grain size essentially constant in 2.5 µm, the hardness and flexural strength of the developed cemented carbides is highly dependent of the binder metal content. As can be seen, the hardness decreased continuously with increasing binder phase content whereas, the flexural strength increased. According to the literature [7], this is because part of the hard WC phase is replaced by the ductile nickel binder as the binder content is increased. It can also be observed from the results that the cemented carbide with 10 wt.% of binder alloy presented a relatively better compromise between hardness and flexural strength. Fig. 6 shows the hardness and flexural strength of the cemented carbide WC–Ni–Si with 10 wt.% of metal binder. For comparison, data for WC–Co and WC–Ni materials with comparable binder content and WC particle size from the literature are also shown [18]. As can be observed from the figure, the hardness of cemented carbide WC–Ni–Si is similar to comparable WC–Co cemented carbide and quite superior to comparable WC–Ni cemented carbide. This result demonstrates the effective action of silicon as strengthener by solution in Ni leading to an overall improvement of the mechanical properties of the material; mainly when it is compared with the WC–Ni alloy. According to the literature [19], the most widely accepted mechanism of solid solution strengthening of the nickel binder by silicon arises from the elastic interaction between the strain field of the solute (Si) and that of dislocations. The strain field of the solute is attributed to the atomic size misfit between the solute and solvent. It is worthwhile mentioning that the hardness value of the WC–Ni–Si alloy slightly lower than that of the WC–Co may be attributed to the fact of the more ductile nickel binder presents a higher amount of plastic micro-deformations than the cobalt binder during the cooling. As a result, a considerable relief of residual stress (originated from the different values of linear dilatation coefficients of the hard phase and
Fig. 5. Fracture surface of WC–Ni–Si with 10 wt.% of binder metal. Fracture modes: carbide–carbide interface –C/C and fracture along the binder metal –B.
Fig. 6. Hardness and flexural strength of the WC–Ni–Si cemented carbide with 10 wt.% of binder metal and WC particle size of 2.5 µm.
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In terms of industrial applications, the combination of hardness and toughness obtained for the cemented carbides WC–Ni–Si qualify them to be used to produce several forming molds and perforation tools. 4. Conclusions Based on the analyses and results obtained so far, it can be drawn the following conclusions:
Fig. 7. Data of Flexural strength versus hardness of WC–Ni–Si cemented carbides studied and WC–Co cemented carbides with WC particle size of 2.5 µm and binder metal content variable.
binder) of the alloy occurs together with the decreasing of the hardness [20]. An alternative explanation to the lower hardness of the WC–Ni–Si cemented carbide is that the Ni based binder usually is present in the mixture in higher volume fraction than that of cobalt once its density is lower. Concerning to the flexural strength of the cemented carbide WC–Ni–Si, it can be seen that it is quite superior to that of both WC–Co and WC–Ni materials. This result suggests that the addition of approximately 4.1 wt.% of silicon to strength the Ni binder by solid solution improved the hardness of the materials without sacrificing their ductility. In order to point out the good correlation between the properties of the developed cemented carbides, it is provided in Fig. 7, a comparison of the data of hardness versus flexural strength of the cemented carbides WC–Ni–Si obtained in this work with those of WC–Co material with comparable binder contents and WC particle size reported by literature [18]. It can be seen from the data shown in Fig. 6 that the cemented carbide WC–Ni–Si exhibits a higher flexural strength than that of conventional WC–Co for hardness values quite similar. In Table 2, it is provided a comparison of the mean fracture toughness value obtained for the cemented carbide WC–Ni–Si with binder content of 10 wt.% with those fracture toughness values of the WC–Co and WC–Ni materials with comparable binder content and WC particle size from the literature [14,21,22]. As expected, the cemented carbides with tough nickel binder presented higher fracture toughness than that of WC–Co alloy. However, as can also be seen from the results, even though there is the presence of silicon as a strengthener, the cemented carbide WC– Ni–Si exhibited fracture toughness quite similar to that of the WC–Ni cemented carbide. The results shown in Table 2 and Fig. 2 confirm that silicon produce an improvement of the hardness of the Ni binder without sacrificing considerably its toughness.
Table 2 Comparison of fracture toughness of WC–Ni–Si sintered in this study with previously reported values. Ref.
Binder content (wt.%)
WC grain Size (µm)
Vickers hardness
KIC (MPa m1/2)
[20,21] [22] This work
10Ni 10Co 9.59Ni–0.41Si
2.1–3.6 3.0 2.5
∼1032 ∼1200 ∼1362
14.7 ± 0.4 12.0 ± 0.8 13.8 ± 0.4
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