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WC – (Cu: AISI304) composites processed from high energy ball milled powders
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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WC – (Cu: AISI304) composites processed from high energy ball milled powders J.P. Cardoso, J. Puga, A.M. Ferro Rocha, C.M. Fernandes, A.M.R. Senos
T
⁎
Department of Materials and Ceramics Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal
A R T I C LE I N FO
A B S T R A C T
Keywords: Nanopowders Nanocomposites Cemented carbides Copper: Stainless steel alloys Microstructure Mechanical properties
Alternative binders to cobalt, based on stainless steel (SS, AISI304) and copper were investigated for tungsten carbide (WC) based cemented carbides. The binder content was fixed at 12 wt%, and the Cu:SS ratio varied in proportions of 0:1, 1:5, 1:2, 1:1, 1:0. High energy ball milling was applied to ensure high homogenization, nanometric particle size and mechanical alloying of binder elements in the powders' mixtures. To assess an adequate sintering route, wettability testing and constant heating rate dilatometry in vacuum were performed. The composites were analyzed in terms of their structural, microstructural and mechanical characteristics. The poor wettability of melted Cu on WC surfaces was increased by alloying it with SS and highly dense compacts could be successfully attained at reduced vacuum sintering temperatures with binders having a Cu:SS ratio equal to or lower than 1:2. The microstructures show secondary phases and significant grain coarsening during sintering, whereas the average grain size was kept in the nanometric range. The composites that attained almost full densification present high hardness, comparable to that of nanometric WC-12Co cemented carbides processed by similar routes, but lower toughness values.
1. Introduction Tungsten carbide (WC) “hardmetals” are well-known materials used in wear and cutting applications, due to their excellent properties, such as high hardness and good toughness values [1]. The increasing demands for alternative binders to Co, which are greener, have lower health risks and with improved properties, motivates research in this topic. Tungsten carbide based composites are on the leading edge of some of the most demanding materials in the industrial world, being used to produce valves for flux control, in oil and gas prospection components, and in forming technologies for the pharmaceutical industry, amongst others [2,3]. Since the beginning of the twentieth century, the development of these composites has drastically improved the production efficiency in various sectors, namely in cutting tools and wear components, encouraging a research wave leading to the development of the hardmetal of today [4]. Given the evolution in the past decades, a broad range of hardmetal grades can be produced, mainly in the WC-Co system, to attain the desired properties for specific mechanical applications, and suitable lifetimes in service. With variable compositions in terms of binder content and WC grain size, very different mechanical properties can be achieved. The variation of grain size and binder content allow the composite to range in properties from high values of hardness and low ⁎
Corresponding author. E-mail address: [email protected] (A.M.R. Senos).
https://doi.org/10.1016/j.ijrmhm.2019.104990 Received 13 February 2019; Received in revised form 11 June 2019 Available online 18 June 2019 0263-4368/ © 2019 Elsevier Ltd. All rights reserved.
fracture toughness, to high fracture toughness and, consequently, lower values of hardness [5]. Being the standard cemented carbide system, WC-Co has been extensively studied and characterized. In the search for alternative binders, new possibilities arose, with opportunities to develop special grades for new applications in the hardmetal world. With the goal of a satisfactory substitution of Co in the processing and properties, and improvement of some poor characteristics such as corrosion and creep resistance, new prototypes have been created and successfully processed [6,7]. The motivation for the replacement of cobalt also partly resides in health concerns associated with its processing, and economic factors from cobalt's rising market value, particularly in an era where it is very much sought by the battery industry, due to the boom in electric car production [8]. Extensive work was developed by Fernandes et al. [9–11] using stainless steel (SS) as a binder in WC based cemented carbides. The good wettability of SS AISI304L on tungsten carbide makes it a suitable candidate, where liquid phase sintering mechanisms are dominant for densification [12,13]. Similar mechanical properties - namely hardness, toughness and transverse rupture strength - were achieved with WC-SS, comparable with WC-Co compositions [6]. Moreover, the total binder replacement of Co by SS slightly improved wear and corrosion resistance in the WC-SS composites [14]. The balance of carbon to
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
J.P. Cardoso, et al.
prevent η-phase (i.e. M6C) formation is difficult, due to the narrow two phase (fcc + WC) region in the phase diagrams of WC-Fe alloys [15], but the formation of η-phase can be hindered by a fine tuning of the quantity of carbon added [6]. Copper based binders may present an interesting option providing good thermal conductivity and lower sintering temperatures, for instance in wear tools for oil prospection [16]. However, Cu presents low wettability on WC [12], and W has an extremely low solubility in Cu (< 1 at.% at 1500 °C) [17]. Research has been carried out on the WCCu system through various processing techniques, with a highlight being the work developed by Shinoda et al. [7], where good results could only be attained through spark plasma sintering (SPS). A possibility to overcome the low Cu wettability on W/WC surfaces is alloying it with other metals showing improved wetting behavior on those surfaces, as shown for the case of Cu alloyed with Co [12]. This way, processing through conventional techniques (i.e. vacuum sintering and hot isostatic pressing) might be feasible, posing an interesting alternative for the industry. In this work, we are investigating the possibility of using Cu - stainless steel AISI304L (SS) alloys to process WC cemented carbides. It is expected that the good wettability of stainless steel on WC surfaces [13] may enable the successful use of SS alloys with other elements of lower wettability, such as copper. Additionally, high energy ball milling (HEBM), a powder metallurgy processing technique with inexpensive investment, simplicity of use and potential for large-scale production [18,19], was used to achieve more homogeneous and finer powder blends, leading to high quality nanocrystalline/submicrometer hardmetal grades with increased hardness.
The WC powder, Fig. 1 a), shows agglomerates of smaller particles with angular sides, characteristic of the hexagonal crystal system, whereas Figs. 1 b) and c), for the Cu and SS powders, respectively, present near spherical agglomerated particles, differing in particle size (Table 1). The WC composites were prepared from the above powders, with a fixed binder content of 12 wt%, while varying the ratio between SS and Cu (from no SS, to no Cu). The proportionated powders were milled in a high energy ball mill (HEBM) Fritsch Pulverisette 6, using a AISI303 steel vessel and WC-6wt%Co balls, in argon atmosphere to prevent powder and grinding media oxidation. The HEBM parameters (rotation speed, ball-to-powder weight ratio and time) were optimized in previous work [18] and no Co contamination was detected by EDS in the milled powders. Optimized HEBM parameters of 350 rpm, with a ball-powder ratio (BPR) of 20:1 were applied in all the compositions. The milling time was only slightly increased from 8 to 10 h for the richer Cu compositions, in order to obtain composite powders with near equivalent particle size, ~235 ± 25 nm and crystallite size ~11 ± 1 nm, as shown in Table 2. The thermal behavior of some compositions was assessed by dilatometric analysis, in a homemade vertical graphite dilatometer, under low vacuum pressure (1 Pa), at a CHR of 10 °C/min to 1450 °C. The samples were vacuum sintered (VS) in a vertical tubular furnace, SUPER KANT Termolab, controlled by a Eurotherm PID controller. The furnace was coupled to a Criolab vacuum system comprising a primary vacuum Alcatel A120 pump and a 917 Pirani pressure sensor. The temperature was controlled with a type R thermocouple (Pt/Pt 13% Rh), which was in contact with the graphite crucible. Temperature and pressure during sintering were recorded by a two channel Kipp & Zonen recorder. The sintering parameters (sintering temperature, heating and cooling rates, vacuum pressure, holding time) were defined from the dilatometric results presented in this work and previous studies on WCSS [13, 18]. After the first sintering step, the samples were submitted to hot isostatic pressing (HIP), in an industrial equipment. All the sintering parameters are presented in Table 3. The sintered samples were characterized in terms of structure and microstructure. The crystallographic characterization was performed by X-ray diffraction (XRD) with a Panalytical X'Pert Pro Diffractometer using Cu Kα1 (λCu = 0.154056 nm). The XRD patterns were recorded in the range 10–80°, with a 0.02° step and time per step of 3 s. The crystallographic phases were determined using the High Score Plus software [20], and quantified through Rietveld refinement in the full range of acquisition. The patterns are presented in the range 30–65°, which is representative of the phases found. The WC crystallite size was determined using the Scherrer equation [21]:
2. Experimental procedure The first assessment involved evaluating the wettability of copper alloyed with stainless steel (SS) AISI304L on WC. In previous work, the wettability of Cu on WC and WC-Co substrates was investigated [12]. For this work, a high temperature wetting experiment was performed by the sessile drop method, using copper on carbon-optimized WC-SS surfaces, i.e. without η-phase, expecting that at high temperature the liquid Cu and SS would alloy at the surface. The experiment was performed in a Termolab horizontal furnace with an Eurotherm PID controller. A Cu piece, cut into a parallelepiped geometry (0.5 × 0.5 × 0.2 cm), was placed on a dense substrate of WC-10wt%SS. The system was put on a graphite support inside an alumina tube furnace, and the contact angle was measured from photographs captured every 10 s, using a Canon 1100D camera coupled with an 18–55 mm EF lens. The photographs were captured from 1080 to 1210 °C, at a constant heating rate (CHR) of 7 °C/min in a vacuum atmosphere (16 Pa) to observe the evolution of the wettability of copper on the substrate. The sample was later observed in cross-section to ascertain the interdiffusion of iron and copper, from and to the substrate. To prepare the composite powder, powders of tungsten carbide (WC), stainless steel (SS) AIS304L and high purity copper were used. The chemical and physical characteristics of the precursor powders, as given by the suppliers, are summarized in Table 1. The morphologies of the raw powders were observed in a scanning electron microscope (SEM), Hitachi SU-70, and are presented in Fig. 1.
DScherrer = K x λ/(β x cosθ)
(1)
where DScherrer is the weighted apparent crystallite size, K is a dimensionless shape factor with a value close to unity, λ is the X-ray wavelength (i.e. 0.154056 nm for CuKα1), β is the full width at height maximum (FWHM) and θ is the Bragg angle, θ=35.62° corresponding to WC (100). For microstructural observation, the sintered samples and the wettability test sample (in cross-section) were hot mounted in resin using a Struers LaboPress-3. The samples were ground and polished,
Table 1 Chemical composition and physical properties of the raw powders as provided by the suppliers. Characteristics/suppliers
WC W
Chemical composition (wt.%) Average particle size (μm) Density (g/cm3) Melting temperature (°C) Suppliers
Cu C
93.87 6.13 1 15.6 2785 H. C. Starck, HSCT-Germany
SS AISI 304L Fe
99.9 9 8.9 1083 Sandvik Osprey, Ltd.
2
Cr
71.15 18.80 5 7.9 1400–1455 Alfa Aesar Gmbh & Co. KG
Ni
C
10.00
0.03
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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Fig. 1. SEM micrographs of precursor powders: a) WC; b) Cu and c) SS.
Table 2 Compositions, respective milling parameters and milled powder characteristics. Sample designation
Binder (wt% ratio Cu:SS)
Rotation speed (RPM)
Ball-powder ratio (BPR)
Time (h)
Average particle size (nm)
Crystallite size (nm)
WC-12Cu WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS
1: 1: 1: 1: 0:
350
20:1
10
257 241 211 223 234
12 12 10 11 12
0 1 2 5 1
8
firstly with SiC paper (down to 18.3 μm finish) and afterwards with diamond paste (down to ¼ μm finish). After removing from the resin, the sintered samples were chemically etched using Murakami's solution. Etching was performed for 2 s to reveal the η-phase and afterwards for 2 min to reveal grain boundaries. The samples were observed by SEM, Hitachi SU-70, and the chemical analysis performed with a coupled energy dispersive X-ray spectroscope (EDS), Bruker – Quantax 400. The cross-sectioned wetting sample was not chemically etched but cleaned ultrasonically for 30 min before the EDX analysis. For samples that densify above 95% of relative density, average grain size, hardness and Palmqvist toughness were also assessed. Average grain size was estimated from grain area analysis in SEM micrographs, using the software ImageJ, and the Vickers hardness, HV, using a Zwick/Roell ZHU equipment, where 8–10 measurements were made; the hardness was calculated according to the following equation:
± ± ± ± ±
1 1 1 1 1
3.1. Wettability test The wettability results performed with Cu on a WC-SS substrate, with increasing temperature up to 1210 °C, are presented in Figs. 2 and 3. The wettability test was performed to assess the behavior of copper on a dense WC-SS substrate, with WC and ferritic/martensitic Fe rich phases [24]. It is thought that, at temperatures beyond Cu melting point (1083 °C, Table 1), Cu would dissolve Fe and other metallic elements, such as Ni or Cr, from the substrate binder, forming an alloy (as expected from the phase diagrams Cr-Cu-Fe and Cu-Fe-Ni [25,26]), with increased wettability on the WC surfaces. In Fig. 3, the wetting angle measured from the photographs captured during the experiment (as shown in Fig. 2), using ImageJ, was plotted against the increasing temperature. The wetting behavior observed in Figs. 2 and 3 validates the hypothesis previously presented, where Cu is expected to dissolve elements from the SS alloy with increasing temperatures; at 1080 °C, melting occurred and the measured wetting angle was 67°; the angle dropped sharply to ~40° at 1120 °C; from that point onward, the wetting angle decreased below that which observed by Vera et al. [12], where copper obtained the minimum angle of 25° over a pure WC film, whereas in the present case a minimum of 15° was obtained at 1180 °C. The local copper concentration was considerably higher than that of SS in the compact, so it is expected that copper would dissolve SS elements, enabling gradually lower wetting angles, as evidenced by the experimental data. Interestingly, the final ~15° wetting angle achieved for Cu on the WC-SS substrate was very close to that of ~13° for Cu on a WC-Co substrate [12]. The sample from the wettability experiment was cross-sectioned
where P is the applied load (30 kgf for Vickers HV30), and d1 and d2 are the indentation diagonals (mm). The diagonals were measured using a Zeiss Jenaphot 2000 optical microscope at 200× magnification. The method presented by Roebuck et al. [22] was followed. Palmqvist toughness was determined from Eq. (3), observing the values of hardness and crack length (which resulted from the indentation hardness) [23]:
WK = A ∗ (HV)^0.5 x (WG)^0.5
25 25 25 25 25
3. Results and discussion
(2)
HV = 1.8544 x P/[(d1+d2 )/2]^2
± ± ± ± ±
(3)
where WK is the Palmqvist fracture toughness (MPa.m1/2), A is a numerical constant, A = 0.0028, HV the measured hardness (HV30) and WG the load applied divided by the sum of the crack lengths.
Table 3 Densities of the different compositions before and after the sintering cycles. Sample designation
ρgreen (%)
Vacuum Sintering (°C/min)
ρsint. (%)
Weight loss (%)
HIP (°C/min/MPa)
ρsint. (%)
WC-12Cu WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS
65 60 59 61 59
1300/30 1350/30 1350/30 1380/30 1450/30
78 91 94 96 99
4.1 3.6 3.4 2.8 3.0
1400/60/60
79 94 98 99 100
ρgreen- relative green density; ρsint- relative sintered density. 3
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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a)
b)
1080 °C
10 mm
d)
1122 °C
10 mm
10 mm
e)
1143 °C
c)
1101 °C
f)
1164 °C
10 mm
1185 °C
10 mm
10 mm
Fig. 2. Photographs taken during the wetting experiment of Cu on WC-SS, from 1080 °C to 1185 °C, at CHR of 7 °C/min.
elements, it seems possible to densify WC composites with Cu-SS binders by conventional liquid phase sintering, which was investigated in the present work. 3.2. Sintering behavior The shrinkage and shrinkage rate curves from the dilatometric analysis of WC-12Cu, WC-6SS6Cu and WC-12SS are presented in Fig. 5. The dilatometric curves were quite different, as expected from the variations in the binder composition. This led to differences in the melting temperature, wetting behavior and thermal reactions. Nonetheless, shrinkage started in all compositions at the same temperature, ~600 °C, and the densification rate was very similar at the initial stage, increasing gradually with rising temperature. The initial stage was attributed to diffusion during solid state sintering, and proceeds until the appearance of a liquid phase, where the shrinkage rate suddenly increased by the faster matter transport mechanisms in liquid phase assisted sintering [13,27]. This second stage started at ~810 °C for WC-12Cu, and ~1320 °C for WC-12SS, while for WC-6Cu-6SS there is a short acceleration of the shrinkage ~810 °C, as in WC-Cu, followed by a decrease in shrinkage rate, which created a small peak in Fig. 5b. At ~940 °C, a significant increase in the shrinkage rate occurred. In fact, liquid phases are expected to form at lower temperatures than the melting temperature of the individual components due to alloying and a lower temperature liquidus in the WC–Cu system compared to the WCSS systems is in agreement with the melting of Cu at temperatures significantly lower than SS (Table 1). Playing an important role as well is the initial particle size: agglomerates had an average size of 240 nm, but crystallite size was around 11 nm (Table 2). With such small grain sizes, melting may occur at lower temperatures (i.e. lower energy) than in submicrometer/micrometer powders [28]. However, in the case of WC-Cu, due to the poor wettability of Cu on the WC surfaces, the liquid phase mechanisms are not very effective at producing shrinkage, and only ~4% linear shrinkage was observed in the second stage. After that, the shrinkage proceeded at a lower rate, which was attributed to slow diffusion processes with coarsening [29]. It should be noted that, when the binder consisted of Cu and SS, in WC-6Cu-6SS, as discussed above for the wettability experiments, a Cu-rich liquid phase started to form at ~810 °C (dissolving elements of the SS binder), but only at higher temperature, after the reaction and melting of the binder elements, were the liquid phase sintering mechanisms fully activated at ~1050 °C. This reaction behavior between Cu and metallic elements was also responsible for the multiple peaks, occurring within an enlarged temperature range in the shrinkage rate curve. Furthermore, the better
Fig. 3. Variation of the wetting angle of copper on WC-SS substrate, measured from the photographs using ImageJ, at CHR of 7 °C/min.
Fig. 4. BSE image and EDX line profile analysis, showing the diffusion of the binder (e.g. Fe, Cr) to the interface and the copper to the substrate.
and observed on the SEM and an EDX line profile was obtained to determine the concentration of the different elements across the interface, as shown in Fig. 4. As stated in the hypothesis, there was evident interdiffusion of copper and the binder elements, to and from the substrate respectively, and formation of an interface alloy containing iron, copper and chromium. Taking the improved wettability results for Cu alloyed with SS 4
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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Fig. 6. Diffractograms, fit and difference (in black) for WC-12Cu, WC-6SS-6Cu and WC-12SS, represented in the range 30–65°.
poor wettability, was reduced by the addition of stainless steel, and increased densification was obtained for the samples with decreasing Cu:SS binder ratios. The compositions with binders having Cu:SS ratios equal to, or lower than, 1:2 achieved almost full densification (Table 3). It is important to note the significant amounts of weight loss, 3–4 wt %, in Table 3, due to the nanometric size of the starting powders which facilitated the binder volatilization. The weight loss was similar in all compositions.
Fig. 5. Dilatometric curves for WC-12Cu, WC-6SS-6Cu, and WC-12SS: a) shrinkage percentage; b) shrinkage rate.
wettability of the Cu/SS rich liquid phase enabled a high shrinkage in this stage, ~9% (Fig. 5a). On the other hand, the liquid phase assisted mechanisms in the shrinkage curve of WC-12SS were delayed up to higher temperatures, due to the nature of the binder, and to the high decarburization observed in these nanocomposite powders [18]. This stage extended above 1450 °C, and a larger shrinkage than for the other compositions was obtained, in accordance with the excellent wetting behavior of the SS binder [13]. Based on the dilatometric results and previous studies [13,18], the sintering cycles were established, considering a range of temperatures between 1300–1450 °C, starting with 1300 °C for the WC-12Cu composition, and increasing the maximum temperature with the increase of SS content in the binder composition (Table 3). A heating/cooling rate of 10 °C/min and a holding time of 30 min were used for all the compositions. A second hot isostatic pressing (HIP) densification cycle was applied, to increase the densification as indicated in Table 3. The densities, before and after vacuum sintering (VS), followed by HIP, and weight losses (due to metal volatilization and decarburization) are also presented in Table 3. Even though the densities prior to sintering, Table 3, were higher for WC-12Cu, likely due to the softer nature of copper, greater densification was not obtained. WC-12Cu demonstrated the poor wettability of copper in the system. The low densification through VS (i.e. 78%) gave a highly open porous compact, making HIP ineffective, since the applied pressure is only efficient on the closed porosity. This effect, i.e.
3.3. Composite characterization The WC-12Cu, WC-6SS-6Cu and WC-12SS sintered compacts were mechanically reduced to a fine powder to assess their phase composition. In Fig. 6 the respective diffractograms are shown, and Table 4 presents their compositions for the crystalline phases. These were quantified by Rietveld refinement, whose agreement factors, Rexp, Rwp, X2, indicated good fits. The diffractogram for WC-12Cu presents three main phases, WC, Cu0.4W0.6 and Cu. The Cu0.4W0.6 formed by the reaction between Cu and W, from decarburized WC. The formation of this compound, together with the observed binder volatilization, gave a net reduction of the free Cu, up to ~2 wt%. For WC-6SS-6Cu, there was no formation of Cu0.4W0.6, but, instead, the carbide M6C (η-phase) was formed, having W and Fe as main constituents (Fe3W3C). Free Cu was not detected and ~3 wt% of austenitic/martensite iron rich phase is estimated. The last XRD spectrum, for WC-12SS, showed only two phases, WC and η-phase (~48 wt%), without clearly discernible peaks related to the iron rich phase. The formation of such a large amount of stable η-phase was predicted in the closest WC-Fe phase diagram [15], where significant decarburization occurred, consuming a large amount of the iron-rich phase elements, [13]. The diffractograms of the remaining 5
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
J.P. Cardoso, et al.
Table 4 Rietveld crystalline phase quantification after sintering, agreement factors, Rexp, Rp and Rwp, and goodness of fit X2. Sample designation
Cu (%)
Cu0.4W0.6 (%)
γ (%)
η-phase (Fe3W3C) (%)
WC (%)
Rexp
Rp
Rwp
X2
WC-12Cu WC-6SS-6Cu WC-12SS
2±1 – –
42 ± 1 – –
– 3±1 –
– 19 ± 1 48 ± 1
56 ± 1 78 ± 1 52 ± 1
9.81 5.45 5.16
12.67 7.82 6.16
15.81 10.60 7.95
1.61 1.95 1.54
Fig. 7. SEM micrographs of sintered samples: a) WC-12Cu; b) WC-8SS-4Cu; c) WC-10SS-2Cu; d) WC-12SS. The phases are identified in b) by the contrast: WC is light gray, η-phase is medium gray, and iron-rich phase is dark gray.
Fig. 8. SEM micrographs of the composites: a) WC-6SS-6Cu, secondary electrons (SE) images and EDX map; b) WC-8SS-4Cu, SE images and EDX map.
angular grains due to its hexagonal structure, η-phase, and an iron-rich phase, though the latter appeared in small amounts, due to the formation of the more stable η-phase. The microstructure and EDX maps of the WC-6SS-6Cu and WC-8SS4Cu samples are shown in Fig. 8. The EDX map allows the identification of the different elements and how they are distributed, with iron and
compositions are not shown since they presented similar trends. Fig. 7 shows typical microstructures of the composites. WC-12Cu had a poorly densified microstructure. The compact was highly porous, with WC grains remaining in the same range of the initial particle size (~150 nm). WC-10SS-2Cu and WC-12SS had microstructures with less porosity, and three distinguishable phases: WC with the characteristic 6
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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Table 5 Sintering technique and parameters, microstructural and mechanical properties for WC-(Cu, Fe, Ni, Cr) and WC-12Co composites. Sample designation
Sintering technique and parameter
WC grain size: range, average (nm)
Relative density (%)
Vickers Hardness (HV30)
WK (MPa.m1/2)
Ref
WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS WC-12Co WC-12Co
See Table 3 See Table 3 See Table 3 See Table 3 LPS, 30 min@1400°C HIP, 120 MPa 30 min@1100 °C
50–5000, Gavg=638 50–4000, Gavg=235 50–3000, Gavg=161 50–4000, Gavg=151 40–250, Gavg=780 40–80, Gavg=241 – 265
94 98 99 100 99 100
– 1456 ± 80 1588 ± 80 1850 ± 50 1381 1800
– 8.1 ± 1.5 7.2 ± 1.5 8.2 ± 0.6 17.3 10.7
– – – – [30] [31]
LPS – liquid phase sintering; HIP – Hot Isostatic Pressing; Gavg- average grain size.
system [24], and this will be further investigated in forthcoming studies. 4. Conclusions The wettability of melted Cu on WC was enhanced by alloying it with stainless steel (SS, AISI304), and a systematic study of the development of WC-(Cu:SS) composites processed from high energy ball milled powders was undertaken. Binder compositions with varying Cu:SS ratios were studied, with a fixed total binder amount of 12 wt%. Dilatometric experiments showed that the shrinkage curves were very dependent on the binder composition, and despite the fact that the liquid phase appeared at increasing temperatures with higher SS contents, larger shrinkage was obtained for SS-rich binders. Accordingly, highly dense composites with binders having a Cu:SS ratio equal to, or lower than, 1:2 could be attained using vacuum sintering, followed by hot isostatic pressing. The microstructure of the densified composites depended on the binder compositions, and significant amounts of secondary phases formed during sintering. By increasing the amount of SS in the binder composition, higher contents of η-phase and lower amounts of the remaining ductile phase occurred. Densification was accompanied by grain coarsening, resulting in wide grain size distributions, with some abnormal grains in the micrometric range, but maintaining the average grain size in the nanometric range. Along with its nanometric structure and low content of remaining ductile phase, high hardness values from 1450 to 1800 HV30, with modest values of 7–8 MPa m1/2 for the Palmqvist toughness, were determined for the composites which achieved almost full densification (Cu:SS ≤1:2).
Fig. 9. Vickers hardness, HV30, and fracture toughness, Wk, for composites with different Cu:SS ratios.
copper being distributed around the WC grains, and the formation of ηphase between the grains. The darkest regions appeared richer in copper, whereas the medium dark gray areas were rich only in Fe and W (since Cu cannot form η-phase with either of the other elements). A good homogenization in terms of the Fe, Cr binder elements was attained, though quite a large WC grain size distribution can be seen. The grain size ranges and estimated average values, from the SEM micrographs, are presented in Table 5. By increasing the Cu content in the Cu-SS binders, despite the poor densification, the grain growth was slightly increased; the average grain size, Gavg, varied from 151 nm for WC-SS, up to 638 nm in WC-6SS-6Cu composites. The largest Gavg was observed for WC-6SS-6Cu, despite the poor densification of this sample (94%, even after HIP). In addition, all the microstructures in Figs. 7 and 8, with exception of the porous WC-Cu, had large, abnormal grains in the micrometer range. High grain growth rates with abnormal grain growth (AGG) are frequently observed in nanometric and ultrafine grain powders compacts [30]. The mechanical properties of hardness and toughness, measured for the compacts that densified above 95%, are presented in Table 5 and Fig. 9. The Vickers hardness varied from 1456 to 1850 HV30, when reducing the ratio of Cu:SS in the binder from 1:2 to 0. Such high hardness values were comparable with those obtained for nanometric WC-12 wt% Co compacts processed by similar routes [31,32]. The negative influence of the copper content on the composite hardness was evident, since copper lowered the hardness, when directly compared to any steel. The smaller grain size observed in WC-12SS might have increased the hardness. On the other hand, the effect of the Cu amount was not as clear in terms of Palmqvist fracture toughness, and very similar values were obtained. These values were also significantly lower than WC-12 wt% Co [32], which is not surprising considering the low amount of ductile phase in the composites (< 3 wt%). The fracture toughness may be improved by optimized processing, adding carbon to avoid the formation of the mixed carbide phase, as shown in the WC-SS
Acknowledgements C.M. Fernandes gratefully acknowledges the financial support from the program COMPETE and by national funds through FCT under the grant SFRH/BPD/43402/2008. This work was financed by Portugal 2020 through European Regional Development Fund (ERDF) in the frame of Operational Competitiveness and Internationalization Programme (POCI) and in the scope of the project HardCorr (POCI-010247-FEDER-003382). This work was developed within the scope of the project CICECO-Aveiro Institute of Materials, POCI-01-0145-FEDER007679 (FCT Ref. UID /CTM /50011/2013), financed by national funds through the FCT/MEC and when appropriate co-financed by FEDER under the PT2020 Partnership Agreement. We thank R. C. Pullar for assistance with the English editing of this article. References [1] Z.Z. Fang, M.C. Koopman, H. Wang, 1.04 - cemented tungsten carbide hardmetal-an introduction, in: V.K. Sarin Materials (Ed.), Comprehensive Hard, Elsevier, Oxford, 2014, pp. 123–137. [2] U. Beste, T. Hartzell, H. Engqvist, N. Axén, Surface damage on cemented carbide rock-drill buttons, Wear 249 (3) (May 2001) 324–329. [3] X. Ren, H. Miao, Z. Peng, A review of cemented carbides for rock drilling: an old but still tough challenge in geo-engineering, Int. J. Refract. Met. Hard Mater. 39 (Jul. 2013) 61–77. [4] G.E. Spriggs, A history of fine grained hardmetal, Int. J. Refract. Met. Hard Mater.
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1018–1024. [19] R.M. Raihanuzzaman, Z. Xie, S.J. Hong, R. Ghomashchi, Powder refinement, consolidation and mechanical properties of cemented carbides – an overview, Powder Technol. 261 (Jul. 2014) 1–13. [20] T. Degen, M. Sadki, E. Bron, U. König, G. Nénert, The HighScore suite, Powder Diffraction, vol. 29, 2014, pp. 13–18. [21] B.D. Cullity, Elements of X ray diffraction, Addison-Wesley Publishing Company, Inc, 1956. [22] B. Roebuck, E.A. Almond, Deformation and fracture processes and the physical metallurgy of WC–Co hardmetals, Int. Mater. Rev. 33 (1) (Jan. 1988) 90–112. [23] ISO28079:9001, Hardmetals – Palmqvist Toughness Test, (2015). [24] A.M.R. Senos, C.M. Fernandes, E. Soares, F.J. Oliveira, J. Sacramento, Optimizing the mechanical properties of WC-AISI304, World PM2016 - Alternative Binders, Germany, 2016. [25] T. Velikanova, M. Turchanin, Chromium–copper–iron, Ternary Alloy Systems, Springer, 2008, pp. 88–128. [26] N. Lebrun, P. Perrot, M. Turchanin, A. Serbruyns, M. Materials Science International Team, Copper – iron – nickel, in: G. Effenberg, S. Ilyenko (Eds.), Ternary Alloy Systems, Springer, Berlin Heidelberg, 2008, pp. 473–508. [27] A.G.P. da Silva, W.D. Schubert, B. Lux, The role of the binder phase in the WC-Co sintering, Mater. Res. 4 (2) (2001) 59–62. [28] G.L. Allen, R.A. Bayles, W.W. Gile, W.A. Jesser, Small particle melting of pure metals, Thin Solid Films 144 (2) (Nov. 1986) 297–308. [29] S.-J.L. Kang, 6 - normal grain growth and second-phase particles, Sintering, Butterworth-Heinemann, Oxford, 2005, pp. 91–96. [30] T. Li, Q. Li, L. Lu, J.Y.H. Fuh, P.C. Yu, Abnormal grain growth of WC with small amount of cobalt, Philos. Mag. 87 (36) (Dec. 2007) 5657–5671. [31] V. Bonache, M.D. Salvador, A. Fernández, A. Borrell, Fabrication of full density near-nanostructured cemented carbides by combination of VC/Cr3C2 addition and consolidation by SPS and HIP technologies, Int. J. Refract. Met. Hard Mater. 29 (2) (Mar. 2011) 202–208. [32] D. Sivaprahasam, S.B. Chandrasekar, R. Sundaresan, Microstructure and mechanical properties of nanocrystalline WC–12Co consolidated by spark plasma sintering, Int. J. Refract. Met. Hard Mater. 25 (2) (Mar. 2007) 144–152.
13 (5) (Jan. 1995) 241–255. [5] R.W. Armstrong, The hardness and strength properties of WC-Co composites, Materials 4 (7) (Jul. 2011) 1287–1308. [6] C.M. Fernandes, et al., WC-stainless steel hardmetals, Int. J. Refract. Met. Hard Mater. 72 (2018) 21–26. [7] Y. Shinoda, Y. Yanagisawa, T. Akatsu, F. Wakai, H. Fujii, Development of creepresistant tungsten carbide copper cemented carbide, Mater. Trans. 50 (6) (2009) 1250–1254. [8] J.-R. Riba, C. López-Torres, L. Romeral, A. Garcia, Rare-earth-free propulsion motors for electric vehicles: a technology review, Renew. Sust. Energ. Rev. 57 (May 2016) 367–379. [9] C.M. Fernandes, A.M.R. Senos, M.T. Vieira, Particle surface properties of stainless steel-coated tungsten carbide powders, Powder Technol. 164 (3) (Jun. 2006) 124–129. [10] C.M. Fernandes, A.M.R. Senos, M.T. Vieira, J.V. Fernandes, Composites from WC powders sputter-deposited with iron rich binders, Ceram. Int. 35 (4) (2009) 1617–1623. [11] C.M. Fernandes, et al., Mechanical characterization of WC–10 wt% AISI 304 cemented carbides, Mater. Sci. Eng. A 618 (Nov. 2014) 629–636. [12] V.L. Silva, C.M. Fernandes, A.M.R. Senos, Copper wettability on tungsten carbide surfaces, Ceram. Int. 42 (1) (Jan. 2016) 1191–1196. [13] C.M. Fernandes, F.J. Oliveira, A.M.R. Senos, Reactive sintering and microstructure development of tungsten carbide-AISI 304 stainless steel cemented carbides, Mater. Chem. Phys. 193 (Jun. 2017) 348–355. [14] A.B. Oliveira, et al., Corrosion behaviour of WC-10% AISI 304 cemented carbides, Corros. Sci. 100 (Nov. 2015) 322–331. [15] C.M. Fernandes, A.M.R. Senos, Cemented carbide phase diagrams: a review, Int. J. Refract. Met. Hard Mater. 29 (4) (Jul. 2011) 405–418. [16] C. Zhang, G. Luo, J. Zhang, Y. Dai, Q. Shen, L. Zhang, Synthesis and thermal conductivity improvement of W-Cu composites modified with WC interfacial layer, Mater. Des. 127 (Aug. 2017) 233–242. [17] T.B. Massalski, H. Okamoto, P.R. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, 1990. [18] C.M. Fernandes, J. Puga, A.M.R. Senos, Nanometric WC-12 wt% AISI 304 powders obtained by high energy ball milling, Advanced Powder Technology 30 (5) (2019)
8
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International Journal of Refractory Metals & Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
WC – (Cu: AISI304) composites processed from high energy ball milled powders J.P. Cardoso, J. Puga, A.M. Ferro Rocha, C.M. Fernandes, A.M.R. Senos
T
⁎
Department of Materials and Ceramics Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal
A R T I C LE I N FO
A B S T R A C T
Keywords: Nanopowders Nanocomposites Cemented carbides Copper: Stainless steel alloys Microstructure Mechanical properties
Alternative binders to cobalt, based on stainless steel (SS, AISI304) and copper were investigated for tungsten carbide (WC) based cemented carbides. The binder content was fixed at 12 wt%, and the Cu:SS ratio varied in proportions of 0:1, 1:5, 1:2, 1:1, 1:0. High energy ball milling was applied to ensure high homogenization, nanometric particle size and mechanical alloying of binder elements in the powders' mixtures. To assess an adequate sintering route, wettability testing and constant heating rate dilatometry in vacuum were performed. The composites were analyzed in terms of their structural, microstructural and mechanical characteristics. The poor wettability of melted Cu on WC surfaces was increased by alloying it with SS and highly dense compacts could be successfully attained at reduced vacuum sintering temperatures with binders having a Cu:SS ratio equal to or lower than 1:2. The microstructures show secondary phases and significant grain coarsening during sintering, whereas the average grain size was kept in the nanometric range. The composites that attained almost full densification present high hardness, comparable to that of nanometric WC-12Co cemented carbides processed by similar routes, but lower toughness values.
1. Introduction Tungsten carbide (WC) “hardmetals” are well-known materials used in wear and cutting applications, due to their excellent properties, such as high hardness and good toughness values [1]. The increasing demands for alternative binders to Co, which are greener, have lower health risks and with improved properties, motivates research in this topic. Tungsten carbide based composites are on the leading edge of some of the most demanding materials in the industrial world, being used to produce valves for flux control, in oil and gas prospection components, and in forming technologies for the pharmaceutical industry, amongst others [2,3]. Since the beginning of the twentieth century, the development of these composites has drastically improved the production efficiency in various sectors, namely in cutting tools and wear components, encouraging a research wave leading to the development of the hardmetal of today [4]. Given the evolution in the past decades, a broad range of hardmetal grades can be produced, mainly in the WC-Co system, to attain the desired properties for specific mechanical applications, and suitable lifetimes in service. With variable compositions in terms of binder content and WC grain size, very different mechanical properties can be achieved. The variation of grain size and binder content allow the composite to range in properties from high values of hardness and low ⁎
Corresponding author. E-mail address: [email protected] (A.M.R. Senos).
https://doi.org/10.1016/j.ijrmhm.2019.104990 Received 13 February 2019; Received in revised form 11 June 2019 Available online 18 June 2019 0263-4368/ © 2019 Elsevier Ltd. All rights reserved.
fracture toughness, to high fracture toughness and, consequently, lower values of hardness [5]. Being the standard cemented carbide system, WC-Co has been extensively studied and characterized. In the search for alternative binders, new possibilities arose, with opportunities to develop special grades for new applications in the hardmetal world. With the goal of a satisfactory substitution of Co in the processing and properties, and improvement of some poor characteristics such as corrosion and creep resistance, new prototypes have been created and successfully processed [6,7]. The motivation for the replacement of cobalt also partly resides in health concerns associated with its processing, and economic factors from cobalt's rising market value, particularly in an era where it is very much sought by the battery industry, due to the boom in electric car production [8]. Extensive work was developed by Fernandes et al. [9–11] using stainless steel (SS) as a binder in WC based cemented carbides. The good wettability of SS AISI304L on tungsten carbide makes it a suitable candidate, where liquid phase sintering mechanisms are dominant for densification [12,13]. Similar mechanical properties - namely hardness, toughness and transverse rupture strength - were achieved with WC-SS, comparable with WC-Co compositions [6]. Moreover, the total binder replacement of Co by SS slightly improved wear and corrosion resistance in the WC-SS composites [14]. The balance of carbon to
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prevent η-phase (i.e. M6C) formation is difficult, due to the narrow two phase (fcc + WC) region in the phase diagrams of WC-Fe alloys [15], but the formation of η-phase can be hindered by a fine tuning of the quantity of carbon added [6]. Copper based binders may present an interesting option providing good thermal conductivity and lower sintering temperatures, for instance in wear tools for oil prospection [16]. However, Cu presents low wettability on WC [12], and W has an extremely low solubility in Cu (< 1 at.% at 1500 °C) [17]. Research has been carried out on the WCCu system through various processing techniques, with a highlight being the work developed by Shinoda et al. [7], where good results could only be attained through spark plasma sintering (SPS). A possibility to overcome the low Cu wettability on W/WC surfaces is alloying it with other metals showing improved wetting behavior on those surfaces, as shown for the case of Cu alloyed with Co [12]. This way, processing through conventional techniques (i.e. vacuum sintering and hot isostatic pressing) might be feasible, posing an interesting alternative for the industry. In this work, we are investigating the possibility of using Cu - stainless steel AISI304L (SS) alloys to process WC cemented carbides. It is expected that the good wettability of stainless steel on WC surfaces [13] may enable the successful use of SS alloys with other elements of lower wettability, such as copper. Additionally, high energy ball milling (HEBM), a powder metallurgy processing technique with inexpensive investment, simplicity of use and potential for large-scale production [18,19], was used to achieve more homogeneous and finer powder blends, leading to high quality nanocrystalline/submicrometer hardmetal grades with increased hardness.
The WC powder, Fig. 1 a), shows agglomerates of smaller particles with angular sides, characteristic of the hexagonal crystal system, whereas Figs. 1 b) and c), for the Cu and SS powders, respectively, present near spherical agglomerated particles, differing in particle size (Table 1). The WC composites were prepared from the above powders, with a fixed binder content of 12 wt%, while varying the ratio between SS and Cu (from no SS, to no Cu). The proportionated powders were milled in a high energy ball mill (HEBM) Fritsch Pulverisette 6, using a AISI303 steel vessel and WC-6wt%Co balls, in argon atmosphere to prevent powder and grinding media oxidation. The HEBM parameters (rotation speed, ball-to-powder weight ratio and time) were optimized in previous work [18] and no Co contamination was detected by EDS in the milled powders. Optimized HEBM parameters of 350 rpm, with a ball-powder ratio (BPR) of 20:1 were applied in all the compositions. The milling time was only slightly increased from 8 to 10 h for the richer Cu compositions, in order to obtain composite powders with near equivalent particle size, ~235 ± 25 nm and crystallite size ~11 ± 1 nm, as shown in Table 2. The thermal behavior of some compositions was assessed by dilatometric analysis, in a homemade vertical graphite dilatometer, under low vacuum pressure (1 Pa), at a CHR of 10 °C/min to 1450 °C. The samples were vacuum sintered (VS) in a vertical tubular furnace, SUPER KANT Termolab, controlled by a Eurotherm PID controller. The furnace was coupled to a Criolab vacuum system comprising a primary vacuum Alcatel A120 pump and a 917 Pirani pressure sensor. The temperature was controlled with a type R thermocouple (Pt/Pt 13% Rh), which was in contact with the graphite crucible. Temperature and pressure during sintering were recorded by a two channel Kipp & Zonen recorder. The sintering parameters (sintering temperature, heating and cooling rates, vacuum pressure, holding time) were defined from the dilatometric results presented in this work and previous studies on WCSS [13, 18]. After the first sintering step, the samples were submitted to hot isostatic pressing (HIP), in an industrial equipment. All the sintering parameters are presented in Table 3. The sintered samples were characterized in terms of structure and microstructure. The crystallographic characterization was performed by X-ray diffraction (XRD) with a Panalytical X'Pert Pro Diffractometer using Cu Kα1 (λCu = 0.154056 nm). The XRD patterns were recorded in the range 10–80°, with a 0.02° step and time per step of 3 s. The crystallographic phases were determined using the High Score Plus software [20], and quantified through Rietveld refinement in the full range of acquisition. The patterns are presented in the range 30–65°, which is representative of the phases found. The WC crystallite size was determined using the Scherrer equation [21]:
2. Experimental procedure The first assessment involved evaluating the wettability of copper alloyed with stainless steel (SS) AISI304L on WC. In previous work, the wettability of Cu on WC and WC-Co substrates was investigated [12]. For this work, a high temperature wetting experiment was performed by the sessile drop method, using copper on carbon-optimized WC-SS surfaces, i.e. without η-phase, expecting that at high temperature the liquid Cu and SS would alloy at the surface. The experiment was performed in a Termolab horizontal furnace with an Eurotherm PID controller. A Cu piece, cut into a parallelepiped geometry (0.5 × 0.5 × 0.2 cm), was placed on a dense substrate of WC-10wt%SS. The system was put on a graphite support inside an alumina tube furnace, and the contact angle was measured from photographs captured every 10 s, using a Canon 1100D camera coupled with an 18–55 mm EF lens. The photographs were captured from 1080 to 1210 °C, at a constant heating rate (CHR) of 7 °C/min in a vacuum atmosphere (16 Pa) to observe the evolution of the wettability of copper on the substrate. The sample was later observed in cross-section to ascertain the interdiffusion of iron and copper, from and to the substrate. To prepare the composite powder, powders of tungsten carbide (WC), stainless steel (SS) AIS304L and high purity copper were used. The chemical and physical characteristics of the precursor powders, as given by the suppliers, are summarized in Table 1. The morphologies of the raw powders were observed in a scanning electron microscope (SEM), Hitachi SU-70, and are presented in Fig. 1.
DScherrer = K x λ/(β x cosθ)
(1)
where DScherrer is the weighted apparent crystallite size, K is a dimensionless shape factor with a value close to unity, λ is the X-ray wavelength (i.e. 0.154056 nm for CuKα1), β is the full width at height maximum (FWHM) and θ is the Bragg angle, θ=35.62° corresponding to WC (100). For microstructural observation, the sintered samples and the wettability test sample (in cross-section) were hot mounted in resin using a Struers LaboPress-3. The samples were ground and polished,
Table 1 Chemical composition and physical properties of the raw powders as provided by the suppliers. Characteristics/suppliers
WC W
Chemical composition (wt.%) Average particle size (μm) Density (g/cm3) Melting temperature (°C) Suppliers
Cu C
93.87 6.13 1 15.6 2785 H. C. Starck, HSCT-Germany
SS AISI 304L Fe
99.9 9 8.9 1083 Sandvik Osprey, Ltd.
2
Cr
71.15 18.80 5 7.9 1400–1455 Alfa Aesar Gmbh & Co. KG
Ni
C
10.00
0.03
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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Fig. 1. SEM micrographs of precursor powders: a) WC; b) Cu and c) SS.
Table 2 Compositions, respective milling parameters and milled powder characteristics. Sample designation
Binder (wt% ratio Cu:SS)
Rotation speed (RPM)
Ball-powder ratio (BPR)
Time (h)
Average particle size (nm)
Crystallite size (nm)
WC-12Cu WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS
1: 1: 1: 1: 0:
350
20:1
10
257 241 211 223 234
12 12 10 11 12
0 1 2 5 1
8
firstly with SiC paper (down to 18.3 μm finish) and afterwards with diamond paste (down to ¼ μm finish). After removing from the resin, the sintered samples were chemically etched using Murakami's solution. Etching was performed for 2 s to reveal the η-phase and afterwards for 2 min to reveal grain boundaries. The samples were observed by SEM, Hitachi SU-70, and the chemical analysis performed with a coupled energy dispersive X-ray spectroscope (EDS), Bruker – Quantax 400. The cross-sectioned wetting sample was not chemically etched but cleaned ultrasonically for 30 min before the EDX analysis. For samples that densify above 95% of relative density, average grain size, hardness and Palmqvist toughness were also assessed. Average grain size was estimated from grain area analysis in SEM micrographs, using the software ImageJ, and the Vickers hardness, HV, using a Zwick/Roell ZHU equipment, where 8–10 measurements were made; the hardness was calculated according to the following equation:
± ± ± ± ±
1 1 1 1 1
3.1. Wettability test The wettability results performed with Cu on a WC-SS substrate, with increasing temperature up to 1210 °C, are presented in Figs. 2 and 3. The wettability test was performed to assess the behavior of copper on a dense WC-SS substrate, with WC and ferritic/martensitic Fe rich phases [24]. It is thought that, at temperatures beyond Cu melting point (1083 °C, Table 1), Cu would dissolve Fe and other metallic elements, such as Ni or Cr, from the substrate binder, forming an alloy (as expected from the phase diagrams Cr-Cu-Fe and Cu-Fe-Ni [25,26]), with increased wettability on the WC surfaces. In Fig. 3, the wetting angle measured from the photographs captured during the experiment (as shown in Fig. 2), using ImageJ, was plotted against the increasing temperature. The wetting behavior observed in Figs. 2 and 3 validates the hypothesis previously presented, where Cu is expected to dissolve elements from the SS alloy with increasing temperatures; at 1080 °C, melting occurred and the measured wetting angle was 67°; the angle dropped sharply to ~40° at 1120 °C; from that point onward, the wetting angle decreased below that which observed by Vera et al. [12], where copper obtained the minimum angle of 25° over a pure WC film, whereas in the present case a minimum of 15° was obtained at 1180 °C. The local copper concentration was considerably higher than that of SS in the compact, so it is expected that copper would dissolve SS elements, enabling gradually lower wetting angles, as evidenced by the experimental data. Interestingly, the final ~15° wetting angle achieved for Cu on the WC-SS substrate was very close to that of ~13° for Cu on a WC-Co substrate [12]. The sample from the wettability experiment was cross-sectioned
where P is the applied load (30 kgf for Vickers HV30), and d1 and d2 are the indentation diagonals (mm). The diagonals were measured using a Zeiss Jenaphot 2000 optical microscope at 200× magnification. The method presented by Roebuck et al. [22] was followed. Palmqvist toughness was determined from Eq. (3), observing the values of hardness and crack length (which resulted from the indentation hardness) [23]:
WK = A ∗ (HV)^0.5 x (WG)^0.5
25 25 25 25 25
3. Results and discussion
(2)
HV = 1.8544 x P/[(d1+d2 )/2]^2
± ± ± ± ±
(3)
where WK is the Palmqvist fracture toughness (MPa.m1/2), A is a numerical constant, A = 0.0028, HV the measured hardness (HV30) and WG the load applied divided by the sum of the crack lengths.
Table 3 Densities of the different compositions before and after the sintering cycles. Sample designation
ρgreen (%)
Vacuum Sintering (°C/min)
ρsint. (%)
Weight loss (%)
HIP (°C/min/MPa)
ρsint. (%)
WC-12Cu WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS
65 60 59 61 59
1300/30 1350/30 1350/30 1380/30 1450/30
78 91 94 96 99
4.1 3.6 3.4 2.8 3.0
1400/60/60
79 94 98 99 100
ρgreen- relative green density; ρsint- relative sintered density. 3
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
J.P. Cardoso, et al.
a)
b)
1080 °C
10 mm
d)
1122 °C
10 mm
10 mm
e)
1143 °C
c)
1101 °C
f)
1164 °C
10 mm
1185 °C
10 mm
10 mm
Fig. 2. Photographs taken during the wetting experiment of Cu on WC-SS, from 1080 °C to 1185 °C, at CHR of 7 °C/min.
elements, it seems possible to densify WC composites with Cu-SS binders by conventional liquid phase sintering, which was investigated in the present work. 3.2. Sintering behavior The shrinkage and shrinkage rate curves from the dilatometric analysis of WC-12Cu, WC-6SS6Cu and WC-12SS are presented in Fig. 5. The dilatometric curves were quite different, as expected from the variations in the binder composition. This led to differences in the melting temperature, wetting behavior and thermal reactions. Nonetheless, shrinkage started in all compositions at the same temperature, ~600 °C, and the densification rate was very similar at the initial stage, increasing gradually with rising temperature. The initial stage was attributed to diffusion during solid state sintering, and proceeds until the appearance of a liquid phase, where the shrinkage rate suddenly increased by the faster matter transport mechanisms in liquid phase assisted sintering [13,27]. This second stage started at ~810 °C for WC-12Cu, and ~1320 °C for WC-12SS, while for WC-6Cu-6SS there is a short acceleration of the shrinkage ~810 °C, as in WC-Cu, followed by a decrease in shrinkage rate, which created a small peak in Fig. 5b. At ~940 °C, a significant increase in the shrinkage rate occurred. In fact, liquid phases are expected to form at lower temperatures than the melting temperature of the individual components due to alloying and a lower temperature liquidus in the WC–Cu system compared to the WCSS systems is in agreement with the melting of Cu at temperatures significantly lower than SS (Table 1). Playing an important role as well is the initial particle size: agglomerates had an average size of 240 nm, but crystallite size was around 11 nm (Table 2). With such small grain sizes, melting may occur at lower temperatures (i.e. lower energy) than in submicrometer/micrometer powders [28]. However, in the case of WC-Cu, due to the poor wettability of Cu on the WC surfaces, the liquid phase mechanisms are not very effective at producing shrinkage, and only ~4% linear shrinkage was observed in the second stage. After that, the shrinkage proceeded at a lower rate, which was attributed to slow diffusion processes with coarsening [29]. It should be noted that, when the binder consisted of Cu and SS, in WC-6Cu-6SS, as discussed above for the wettability experiments, a Cu-rich liquid phase started to form at ~810 °C (dissolving elements of the SS binder), but only at higher temperature, after the reaction and melting of the binder elements, were the liquid phase sintering mechanisms fully activated at ~1050 °C. This reaction behavior between Cu and metallic elements was also responsible for the multiple peaks, occurring within an enlarged temperature range in the shrinkage rate curve. Furthermore, the better
Fig. 3. Variation of the wetting angle of copper on WC-SS substrate, measured from the photographs using ImageJ, at CHR of 7 °C/min.
Fig. 4. BSE image and EDX line profile analysis, showing the diffusion of the binder (e.g. Fe, Cr) to the interface and the copper to the substrate.
and observed on the SEM and an EDX line profile was obtained to determine the concentration of the different elements across the interface, as shown in Fig. 4. As stated in the hypothesis, there was evident interdiffusion of copper and the binder elements, to and from the substrate respectively, and formation of an interface alloy containing iron, copper and chromium. Taking the improved wettability results for Cu alloyed with SS 4
International Journal of Refractory Metals & Hard Materials 84 (2019) 104990
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Fig. 6. Diffractograms, fit and difference (in black) for WC-12Cu, WC-6SS-6Cu and WC-12SS, represented in the range 30–65°.
poor wettability, was reduced by the addition of stainless steel, and increased densification was obtained for the samples with decreasing Cu:SS binder ratios. The compositions with binders having Cu:SS ratios equal to, or lower than, 1:2 achieved almost full densification (Table 3). It is important to note the significant amounts of weight loss, 3–4 wt %, in Table 3, due to the nanometric size of the starting powders which facilitated the binder volatilization. The weight loss was similar in all compositions.
Fig. 5. Dilatometric curves for WC-12Cu, WC-6SS-6Cu, and WC-12SS: a) shrinkage percentage; b) shrinkage rate.
wettability of the Cu/SS rich liquid phase enabled a high shrinkage in this stage, ~9% (Fig. 5a). On the other hand, the liquid phase assisted mechanisms in the shrinkage curve of WC-12SS were delayed up to higher temperatures, due to the nature of the binder, and to the high decarburization observed in these nanocomposite powders [18]. This stage extended above 1450 °C, and a larger shrinkage than for the other compositions was obtained, in accordance with the excellent wetting behavior of the SS binder [13]. Based on the dilatometric results and previous studies [13,18], the sintering cycles were established, considering a range of temperatures between 1300–1450 °C, starting with 1300 °C for the WC-12Cu composition, and increasing the maximum temperature with the increase of SS content in the binder composition (Table 3). A heating/cooling rate of 10 °C/min and a holding time of 30 min were used for all the compositions. A second hot isostatic pressing (HIP) densification cycle was applied, to increase the densification as indicated in Table 3. The densities, before and after vacuum sintering (VS), followed by HIP, and weight losses (due to metal volatilization and decarburization) are also presented in Table 3. Even though the densities prior to sintering, Table 3, were higher for WC-12Cu, likely due to the softer nature of copper, greater densification was not obtained. WC-12Cu demonstrated the poor wettability of copper in the system. The low densification through VS (i.e. 78%) gave a highly open porous compact, making HIP ineffective, since the applied pressure is only efficient on the closed porosity. This effect, i.e.
3.3. Composite characterization The WC-12Cu, WC-6SS-6Cu and WC-12SS sintered compacts were mechanically reduced to a fine powder to assess their phase composition. In Fig. 6 the respective diffractograms are shown, and Table 4 presents their compositions for the crystalline phases. These were quantified by Rietveld refinement, whose agreement factors, Rexp, Rwp, X2, indicated good fits. The diffractogram for WC-12Cu presents three main phases, WC, Cu0.4W0.6 and Cu. The Cu0.4W0.6 formed by the reaction between Cu and W, from decarburized WC. The formation of this compound, together with the observed binder volatilization, gave a net reduction of the free Cu, up to ~2 wt%. For WC-6SS-6Cu, there was no formation of Cu0.4W0.6, but, instead, the carbide M6C (η-phase) was formed, having W and Fe as main constituents (Fe3W3C). Free Cu was not detected and ~3 wt% of austenitic/martensite iron rich phase is estimated. The last XRD spectrum, for WC-12SS, showed only two phases, WC and η-phase (~48 wt%), without clearly discernible peaks related to the iron rich phase. The formation of such a large amount of stable η-phase was predicted in the closest WC-Fe phase diagram [15], where significant decarburization occurred, consuming a large amount of the iron-rich phase elements, [13]. The diffractograms of the remaining 5
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Table 4 Rietveld crystalline phase quantification after sintering, agreement factors, Rexp, Rp and Rwp, and goodness of fit X2. Sample designation
Cu (%)
Cu0.4W0.6 (%)
γ (%)
η-phase (Fe3W3C) (%)
WC (%)
Rexp
Rp
Rwp
X2
WC-12Cu WC-6SS-6Cu WC-12SS
2±1 – –
42 ± 1 – –
– 3±1 –
– 19 ± 1 48 ± 1
56 ± 1 78 ± 1 52 ± 1
9.81 5.45 5.16
12.67 7.82 6.16
15.81 10.60 7.95
1.61 1.95 1.54
Fig. 7. SEM micrographs of sintered samples: a) WC-12Cu; b) WC-8SS-4Cu; c) WC-10SS-2Cu; d) WC-12SS. The phases are identified in b) by the contrast: WC is light gray, η-phase is medium gray, and iron-rich phase is dark gray.
Fig. 8. SEM micrographs of the composites: a) WC-6SS-6Cu, secondary electrons (SE) images and EDX map; b) WC-8SS-4Cu, SE images and EDX map.
angular grains due to its hexagonal structure, η-phase, and an iron-rich phase, though the latter appeared in small amounts, due to the formation of the more stable η-phase. The microstructure and EDX maps of the WC-6SS-6Cu and WC-8SS4Cu samples are shown in Fig. 8. The EDX map allows the identification of the different elements and how they are distributed, with iron and
compositions are not shown since they presented similar trends. Fig. 7 shows typical microstructures of the composites. WC-12Cu had a poorly densified microstructure. The compact was highly porous, with WC grains remaining in the same range of the initial particle size (~150 nm). WC-10SS-2Cu and WC-12SS had microstructures with less porosity, and three distinguishable phases: WC with the characteristic 6
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Table 5 Sintering technique and parameters, microstructural and mechanical properties for WC-(Cu, Fe, Ni, Cr) and WC-12Co composites. Sample designation
Sintering technique and parameter
WC grain size: range, average (nm)
Relative density (%)
Vickers Hardness (HV30)
WK (MPa.m1/2)
Ref
WC-6SS-6Cu WC-8SS-4Cu WC-10SS-2Cu WC-12SS WC-12Co WC-12Co
See Table 3 See Table 3 See Table 3 See Table 3 LPS, 30 min@1400°C HIP, 120 MPa 30 min@1100 °C
50–5000, Gavg=638 50–4000, Gavg=235 50–3000, Gavg=161 50–4000, Gavg=151 40–250, Gavg=780 40–80, Gavg=241 – 265
94 98 99 100 99 100
– 1456 ± 80 1588 ± 80 1850 ± 50 1381 1800
– 8.1 ± 1.5 7.2 ± 1.5 8.2 ± 0.6 17.3 10.7
– – – – [30] [31]
LPS – liquid phase sintering; HIP – Hot Isostatic Pressing; Gavg- average grain size.
system [24], and this will be further investigated in forthcoming studies. 4. Conclusions The wettability of melted Cu on WC was enhanced by alloying it with stainless steel (SS, AISI304), and a systematic study of the development of WC-(Cu:SS) composites processed from high energy ball milled powders was undertaken. Binder compositions with varying Cu:SS ratios were studied, with a fixed total binder amount of 12 wt%. Dilatometric experiments showed that the shrinkage curves were very dependent on the binder composition, and despite the fact that the liquid phase appeared at increasing temperatures with higher SS contents, larger shrinkage was obtained for SS-rich binders. Accordingly, highly dense composites with binders having a Cu:SS ratio equal to, or lower than, 1:2 could be attained using vacuum sintering, followed by hot isostatic pressing. The microstructure of the densified composites depended on the binder compositions, and significant amounts of secondary phases formed during sintering. By increasing the amount of SS in the binder composition, higher contents of η-phase and lower amounts of the remaining ductile phase occurred. Densification was accompanied by grain coarsening, resulting in wide grain size distributions, with some abnormal grains in the micrometric range, but maintaining the average grain size in the nanometric range. Along with its nanometric structure and low content of remaining ductile phase, high hardness values from 1450 to 1800 HV30, with modest values of 7–8 MPa m1/2 for the Palmqvist toughness, were determined for the composites which achieved almost full densification (Cu:SS ≤1:2).
Fig. 9. Vickers hardness, HV30, and fracture toughness, Wk, for composites with different Cu:SS ratios.
copper being distributed around the WC grains, and the formation of ηphase between the grains. The darkest regions appeared richer in copper, whereas the medium dark gray areas were rich only in Fe and W (since Cu cannot form η-phase with either of the other elements). A good homogenization in terms of the Fe, Cr binder elements was attained, though quite a large WC grain size distribution can be seen. The grain size ranges and estimated average values, from the SEM micrographs, are presented in Table 5. By increasing the Cu content in the Cu-SS binders, despite the poor densification, the grain growth was slightly increased; the average grain size, Gavg, varied from 151 nm for WC-SS, up to 638 nm in WC-6SS-6Cu composites. The largest Gavg was observed for WC-6SS-6Cu, despite the poor densification of this sample (94%, even after HIP). In addition, all the microstructures in Figs. 7 and 8, with exception of the porous WC-Cu, had large, abnormal grains in the micrometer range. High grain growth rates with abnormal grain growth (AGG) are frequently observed in nanometric and ultrafine grain powders compacts [30]. The mechanical properties of hardness and toughness, measured for the compacts that densified above 95%, are presented in Table 5 and Fig. 9. The Vickers hardness varied from 1456 to 1850 HV30, when reducing the ratio of Cu:SS in the binder from 1:2 to 0. Such high hardness values were comparable with those obtained for nanometric WC-12 wt% Co compacts processed by similar routes [31,32]. The negative influence of the copper content on the composite hardness was evident, since copper lowered the hardness, when directly compared to any steel. The smaller grain size observed in WC-12SS might have increased the hardness. On the other hand, the effect of the Cu amount was not as clear in terms of Palmqvist fracture toughness, and very similar values were obtained. These values were also significantly lower than WC-12 wt% Co [32], which is not surprising considering the low amount of ductile phase in the composites (< 3 wt%). The fracture toughness may be improved by optimized processing, adding carbon to avoid the formation of the mixed carbide phase, as shown in the WC-SS
Acknowledgements C.M. Fernandes gratefully acknowledges the financial support from the program COMPETE and by national funds through FCT under the grant SFRH/BPD/43402/2008. This work was financed by Portugal 2020 through European Regional Development Fund (ERDF) in the frame of Operational Competitiveness and Internationalization Programme (POCI) and in the scope of the project HardCorr (POCI-010247-FEDER-003382). This work was developed within the scope of the project CICECO-Aveiro Institute of Materials, POCI-01-0145-FEDER007679 (FCT Ref. UID /CTM /50011/2013), financed by national funds through the FCT/MEC and when appropriate co-financed by FEDER under the PT2020 Partnership Agreement. We thank R. C. Pullar for assistance with the English editing of this article. References [1] Z.Z. Fang, M.C. Koopman, H. Wang, 1.04 - cemented tungsten carbide hardmetal-an introduction, in: V.K. Sarin Materials (Ed.), Comprehensive Hard, Elsevier, Oxford, 2014, pp. 123–137. [2] U. Beste, T. Hartzell, H. Engqvist, N. Axén, Surface damage on cemented carbide rock-drill buttons, Wear 249 (3) (May 2001) 324–329. [3] X. Ren, H. Miao, Z. Peng, A review of cemented carbides for rock drilling: an old but still tough challenge in geo-engineering, Int. J. Refract. Met. Hard Mater. 39 (Jul. 2013) 61–77. [4] G.E. Spriggs, A history of fine grained hardmetal, Int. J. Refract. Met. Hard Mater.
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