Abrasion wear, thermal shock and impact resistance of WC-cemented carbides produced by PECS and LPS

Abrasion wear, thermal shock and impact resistance of WC-cemented carbides produced by PECS and LPS

Int. Journal of Refractory Metals and Hard Materials 49 (2015) 133–142 Contents lists available at ScienceDirect Int. Journal of Refractory Metals a...

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Int. Journal of Refractory Metals and Hard Materials 49 (2015) 133–142

Contents lists available at ScienceDirect

Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM

Abrasion wear, thermal shock and impact resistance of WC-cemented carbides produced by PECS and LPS☆ R.M. Genga a,b,⁎, G. Akdogan b,c, C. Polese d, J.C. Garrett a,b, L.A. Cornish a,b a

School of Chemical and Metallurgical Engineering, University of the Witwatersrand, South Africa DST/NRF Centre of Excellence in Strong Materials, University of the Witwatersrand, South Africa Department of Process Engineering, University of Stellenbosch, South Africa d School of Mechanical, Industrial and Aeronautical Engineering, University of the Witwatersrand, South Africa b c

a r t i c l e

i n f o

Article history: Received 4 May 2014 Received in revised form 22 July 2014 Accepted 29 July 2014 Available online 2 August 2014 Keywords: Pulse electric current sintering Grain growth inhibition Nickel Binder content Mechanical properties

a b s t r a c t The effects of rapid sintering by pulse electric current sintering (PECS), variation in starting WC size (0.1–0.8 μm), Ni as a Co binder substitute and TiC, NbC and Mo2C additions on the microstructure, abrasion wear, thermal shock and impact resistance of WC–Co and WC–Ni alloys were studied. Abrasion wear tests were done using a ball-ondisk tribometer, with 100Cr6 steel and silicon nitride balls. Use of PECS gave finer microstructures with poorly distributed binder pools than similar liquid phase sintered (LPS) samples, although large angular WC grains of up to 1 μm occurred in the nano (0.1 μm) and ultrafine (0.4 μm) grades. Addition of 5 wt.% NbC to WC–10Co (wt.%) had negligible effect on the WC grain size, while 5 wt.% Mo2C to WC–6.25TiC–9.3Ni significantly improved WC grain growth inhibition, leading to increased hardness (from ~13 to N21 GPa) and reduced wear rate (from 2.73 × 10−4 to b8.0 × 10−5 mm3·N−1·m−1), compared to the LPS WC–9.3Ni. Thermal shock and impact resistance were measured using a thermal imaging camera and force gauges while testing the samples as cutting tools on Ti–6Al–4V under aggressive interrupted milling conditions. The LPS samples had better thermal shock and impact resistance, due to their higher fracture toughness (N 12.5 MPa·m0.5) and B3B transverse rupture strength (N 2100 MPa), as a result of the larger and better binder pool distribution. © 2014 Elsevier Ltd. All rights reserved.

1. Introduction Cemented carbides based on WC and a ductile metal matrix (Co, Ni and/or Fe) are used in a wide range oftribological applications such as cutting, mining, metal forming, drilling and turning [1], because of their good combinations of mechanical properties such as hardness, toughness, wear resistance and strength [1,2]. These properties depend mainly on the WC grain size, type and amount of binder [3], and thus to improve on the operation life of these cemented carbides, successful manipulation of these factors is required. The hardness and abrasion wear resistance of WC cemented carbides produced by powder metallurgy increase with decreased WC grain size from the micron to the nanometric scale, and Co binder strengthening from increased W dissolution [1,4]. Additionally, reduced hardness with increased binder content is less in cemented carbides with finer WC grains than with micron-sized WC [5]. Fine grain sizes have high sintering rates because of the large interfacial surface area between the WC particles and the binder [3,6]. Compared to micron-sized cemented carbides, sintering of nanosized powders has an additional ☆ This paper belongs to the Special Issue on the Science of Hard Materials (ICSHM10). ⁎ Corresponding author. E-mail address: [email protected] (R.M. Genga).

http://dx.doi.org/10.1016/j.ijrmhm.2014.07.031 0263-4368/© 2014 Elsevier Ltd. All rights reserved.

challenge of retaining the nanoscaled grain size on achieving full density, because of rapid grain growth during liquid phase sintering (LPS) [3], which increases with increased sintering temperature and time. Grain growth can be reduced by pulse electric current sintering (PECS), which allows high degrees of densification at low temperatures within a short sintering time and lower sintering temperatures [6–8], and limiting continuous Ostwald ripening [8]. Inhibition of grain growth can also be achieved through the introduction of transition metal carbides such as TiC, Cr3C2, Mo2C and NbC [9,10], which alter the interfacial energies, interfering with the Ostwald ripening and inhibiting grain growth [4]. Furthermore, TiC and NbC improve the chemical stability, preventing diffusion wear during machining [4,8,11]. High temperature hardness can also be improved through NbC, TiC, Mo2C [4] and Cr3C2[12] additions, forming complex carbides which prevent deformation of the cutting edge by WC grain boundary slip [4]. Cobalt is the most used metal binder because of its very good solubility of WC at low and high temperatures, and the good wettability of WC [13]. Additionally, Co has very good communition characteristics, giving better distribution during milling, and subsequent homogenous distribution in the microstructure [4]. However, because of unstable market prices due to short supply, health hazards for plant maintenance personnel and limited corrosion resistance, Ni has been studied as a possible alternative [1,13]. Nickel retains its ductile face

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centered cubic (fcc) structure at all temperatures [13,14]. Nickel also has better resistance to corrosion and oxidation [15], and improves hot hardness and resistance to thermal cracking [1,10]. Additionally, its more ductile austenitic structure prevents loss of WC grains by extruding outwards to replace the eroded binder between the WC grains, improving the wear resistance [13]. Nickel has higher solubility and better wetting of TiC than Co [11], improving the resistance to diffusion wear, particularly for the machining of steel [11]. Generally, LPS WC–Ni cemented carbides have lower hardness than WC–Co cemented carbides [4], due to poor wettability of WC [16], higher plasticity and the absence of any phase transition [13,14]. Increased hardness can be achieved by PECS and addition of grain growth inhibitors, such as TiC and Mo2C [4], as well as by optimizing the binder content [4], to increase the hardness without significantly reducing the fracture toughness. The addition of Mo or Mo2C also improves the wettability of WC by Ni [4,17], as well as of TiC by Ni [11], improving the mechanical properties [4,11,17]. In this study, WC–Co and WC–Ni cemented carbides with NbC, TiC and Mo2C additions were prepared by PECS, and their mechanical properties were compared to LPS WC–Co and WC–Ni cemented carbides under aggressive tribological conditions.

Table 2 Sintering conditions. Sample (wt.%)

Temperature (°C)

WC–0.5Cr3C2–10Co

200 °C/min to 1000 100 °C/min to 1220 1220 °C for 5 min 200 °C/min to 1050 100 °C/min to 1280 1280 °C for 5 min 200 °C/min to 1050 100 °C/min to 1260 1260 °C for 5 min 200 °C/min to 1050 100 °C/min to 1280 1280 °C for 5 min 200 °C/min to 1100 100 °C/min to 1300 1300 °C for 5 min 200 °C/min to 1100 100 °C/min to 1320 1320 °C for 5 min 200 °C/min to 1200 100 °C/min to 1380 1380 °C for 5 min

WC–0.5Cr3C2–9.3Ni

WC–0.5Cr3C2–5NbC–10Co

WC–6.25TiC–9.3Ni

WC–1 Mo2C–6.25TiC–9.3Ni

WC–5Mo2C–6.25TiC–9.3Ni

WC–5Mo2C–6.25TiC–7Ni

Pressure (MPa) °C °C °C °C °C °C °C °C °C °C °C °C °C °C

16 30 50 16 30 50 16 30 60 16 40 70 16 40 70 16 40 70 16 40 70

2. Experimental procedure 2.1. Materials The starting powders and their characteristics are given in Table 1. The powders were wet milled in 99% pure ethanol for 10 h in a steel container with WC milling balls. They were then dried using a rota evaporator at 64 °C and 80 rpm for 1 h. 2.2. Sintering Milled composite powders were consolidated in a spark plasma sintering (HP D5, FCT Systeme, Germany) furnace. The powders were poured into cylindrical graphite dies with inner and outer diameters of 20.9 mm and 40 mm respectively, and 48 mm in height. The composite powder assemblies were heated in a vacuum (2 Pa) in two steps as shown in Table 2, for example WC–0.5Cr3C2–10Co (wt.%) powders were first heated to 1000 °C at a rate of 200 °C/min and subsequently to 1220 °C at a heating rate of 100 °C/min, and the temperature was held at 1220 °C for 5 min during sintering. A cooling rate of 200 °C/min was given to all samples. The applied pressure was adjusted within 30 s from 16 MPa to 30 MPa at 1000 °C, and from 30 MPa to 50 MPa at 1220 °C. The pressure was then held constant at 50 MPa throughout the rapid sintering cycle. Horizontal and vertical graphite papers were used to separate the powders from the die and punch setup. Hexagonal boron nitride was placed on the graphite paper to prevent carbon diffusion from the graphite paper to the powders during sintering. The graphite die was wrapped in a carbon cloth to minimize the heat loss from the die surface. The temperature was controlled by an optical pyrometer focused on a central borehole on

Table 1 Specifications of starting powders. Materials

Particle size (μm)

Crystal structure

Purity (wt.%)

Source

WC doped with 0.5 wt.% Cr3C2 WC

0.8

Hexagonal

N99.00

H.C. Starck, Germany

0.1

Hexagonal

N99.95

Co Ni

0.9 0.2

hcp fcc

N99.80 N99.80

TiC NbC Mo2C

1.5 1.2 1.7

Cubic Cubic Cubic

N99.00 N99.00 N99.00

Dong Yang (HK) Int'l Group, China OMG Americas, USA Dong Yang (HK) Int'l Group, China Treibacher, Austria Treibacher, Austria Treibacher, Austria

the upper punch, 1 mm above the top surface of the sample, to give an accurate estimation of the sample temperature [8]. Consolidation of the composite powder assemblies in the axial direction was monitored by following the position of the plunger. Different sintering profiles, depending on the powder compositions, were used to achieve good densification, as shown in Table 2. All the powders were sintered for a dwell time of 5 min. Liquid phase sintering (HIP, Ultra Temp, USA) of WC–0.5Cr3C2–10Co (wt.%) and WC–9.3Ni (wt.%) samples was done by heating the compositions in a vacuum (0.04 MPa) at an initial heating rate of 2.4 °C/min up to 1200 °C. At 1200 °C, cobalt loss protection (CLP) was carried out by the addition of argon gas at a pressure of 0.37 MPa. The CLP was done using a heating rate 3.5 °C/min up to 1430 °C. The temperature was held constant for 75 min, and for the last 20 min, hot isostatic pressing (HIP) was done at 4.4 MPa to eliminate all the surface porosity [4]. The furnace was then water cooled at a rate of 3.5 °C/min. 2.3. Characterization and mechanical testing Archimedes' principle was used to determine the density of the sintered samples (ED224S, Sartorius, Germany). Microstructures of the cemented carbides were examined by scanning electron microscopy (SEM) in the backscattered electron and secondary electron modes (JSM-7500F, JEOL, Japan), with 15 kV accelerating voltage. Both point and area analyses were done to identify phases with an energy dispersive X-ray spectroscopy (EDX) system (Inca Penta FETx3, Oxford Instruments, UK). Image analysis was done on SEM micrographs using SIS-pro and ImageJ software, and calculations were done using MathConnex 2000 with MathCad 2000 Professional to derive the mean grain size. Vickers hardness (HV30) was measured on polished specimens after standard metallographic preparation, using a load of 30 N (VHT 003 MTA, Vickers Limited, United Kingdom), calculating an average from five indentations at different regions on each sample. The criteria for the accurate derivation of fracture toughness (K1C) using Shetty's equation were satisfied [18]: 1.25 ≤ c / a ≤ 2.25 and 0.25 ≤ I / a ≤ 2.5, where c is the crack length from the center of indentation to the crack tip, a is the half diagonal length of indentation and I is the difference between c and a. The transverse rupture strength (TRS) was determined by the ball-on-three-ball (B3B) method (5500R Universal tester, Instron, USA). Resistance to abrasion during sliding wear was investigated using a ball-on-disk tribometer, applying a 10 N force, for 300 m at a sliding speed of 0.21 m/s, using 100Cr6 steel and silicon nitride balls as the

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abrasive materials, without lubricant. The sample worn volume (ΔVS) was measured using the wear track profile from laser confocal 3D topography imaging, imagJ software and the JISR1613 wear formula (Eq. (1)) [19], then used to calculate the wear rate (KS) using Eq. (2)[20].

Kistler Multicomponent Force Link, Switzerland), while testing the samples as cutting tool inserts on Ti–6Al–4V under aggressive interrupted face milling conditions. The cutting speeds were varied between 40 and 75 m/min and the depth of cuts between 1 and 2 mm (Table 3). The sample width was 40 mm and the radial depth of cut was 45 mm, ensuring that the inserts would disengage from the workpiece for rapid cooling every revolution. Due to the low thermal conductivity of Ti–6Al–4V, about 80% of the heat generated was conducted by the insert [21], producing high temperatures at the insert cutting edge which would rapidly decrease when the insert disengaged from the workpiece. The thermal camera took 30 thermal readings per second, facilitating measurement of the temperature change of the inserts. The forces during milling were measured by Kistler force gauges attached to the workpiece clamping vice.

ΔV S ¼ ðπRðS1 þ S2 þ S3 þ S4 ÞÞ=2

ð1Þ

3. Results

K S ¼ ΔV S =F  s:

ð2Þ

3.1. Density

Table 3 Milling conditions for the WC cemented carbide insets. Milling test variables

CNC parameters

Axial depth of cut (mm)

Cutting speed (m/min)

Feed per tooth (mm/tooth)

Spindle speed (rpm)

Feed rate (mm/min)

1 1 2

40 60 75

0.10 0.10 0.10

283 424 530

28.3 42.4 53.0

Thermal shock and impact resistance were measured using a thermal imaging camera (D5000, Nikon, Japan) and force gauges (9366CC0,

All the WC–Co and WC–Ni cemented carbides produced by PECS had similar relative densities to those produced by LPS (N99%) (Table 4), but

Table 4 Densification behavior of cemented carbides. Sample (wt.%)

Abbreviation

Sintering temperature (°C)

Densification (%) (n)

Densification (%) (u)

Densification (%) (s)

WC–0.5Cr3C2–10Co WC–9.3Ni WC–0.5Cr3C2–10Co WC–9.3Ni WC–6.25TiC–9.3Ni WC–1 Mo2C–6.25TiC–9.3Ni WC–0.5Cr3C2–5NbC–10Co WC–5 Mo2C–6.25TiC–9.3Ni WC–5Mo2C–6.25TiC–7Ni

10Co (LPS) 9.3Ni (LPS) 10Co 9.3Ni 0M 1M 5N 5M 5M–7Ni

1430 1430 1220 1280 1280 1290 1260 1330 1380

– 99.65 – 99.52 99.14 99.36 – 99.27 99.14

– – – 99.24 99.27 99.25 – 99.42 99.45

99.71 – 99.17 99.21 99.35 99.23 99.33 99.61 –

± 0.14 ± 0.31 ± 0.57 ± 0.35 ± 0.13 ± 0.10

± 0.17 ± 0.05 ± 0.07 ± 0.23 ± 0.15

± 0.20 ± ± ± ± ± ±

0.51 0.08 0.10 0.31 0.14 0.12

(s) 0.8 μm starting size, (u) 0.4 μm starting size, (n) 0.1 μm starting size.

Fig. 1. SEM-BSE images of (a) LPS WC–0.5Cr3C2–10Co (wt.%), WC (light) and Co (dark); and WC–9.3Ni (wt.%) sintered by (b) LPS and (c) PECS, both produced from nanoWC powders, WC (light) and Ni (dark).

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Table 5 Variation in grain size of sintered compositions. Sample abbreviation

10Co (LPS) 9.3Ni (LPS) 9.3Ni 5N 5M

WC sintered grain size (μm) WC (s)

WC (u)

WC (n)

2.15 ± 0.11 – 0.95 ± 0.11 0.82 ± 0.05 –

– – 1.12 ± 0.23 – 0.54 ± 0.03

– 3.85 ± 0.36 1.56 ± 0.09 – 0.62 ± 0.06

(s) 0.8 μm starting size, (u) 0.4 μm starting size, (n) 0.1 μm starting size.

NbC Co

The microstructure of the WC–0.5Cr3C2–5NbC–10Co (wt.%) (Fig. 2) showed Co mainly distributed between the WC grains and rarely between the NbC grains. There was negligible WC grain growth (Table 5), since the starting size was ~ 0.8 μm (Table 1). Micrographs of WC–9.3Ni, with 6.25 wt.% TiC and 5 wt.% Mo2C additions are shown in Fig. 3. Titanium carbide was mainly found at the WC/(Ni) interfaces as layers of TiC atoms by high resolution (HR) STEM electron energy loss spectroscopy (EELS) mapping (Fig. 4). Molybdenum carbide was identified by SEM-EDX area analysis, although exact locations were difficult to discern in the fine microstructures. However, Mo was confirmed to be between the WC grains by HAADF STEM mapping (Fig. 5), where areas of overlap between W and Mo were observed. The overlap between W and Mo was confirmed to be WC grains under Mo2C grains by bright field (BF) STEM (Fig. 6). Additions of Mo2C to WC–6.25TiC–9.3Ni (wt.%) significantly improved the grain growth inhibition, with 5 wt.% giving the finest WC grains in both the nano and ultrafine grades (Fig. 3 and Table 5). The WC–5Mo2C–6.25TiC– 9.3Ni (wt.%) sample produced from ultrafine WC powder had finer mean grain size than the similar nano. 3.3. Mechanical properties

Fig. 2. SEM-BSE image of WC–0.5Cr3C2–5NbC–10Co (wt.%): WC (light), NbC (medium) and Co (dark).

higher sintering temperatures and pressures were required to achieve good densities for the WC–Ni samples than the WC–Co samples. Addition of 5 wt.% NbC to WC–0.5Cr3C2–10Co (wt.%) needed an increased sintering temperature and pressure (Tables 2 and 4) to achieve good densification. Similarly, 6.25 wt.% TiC and 1–5 wt.% Mo2C additions to WC–9.3Ni (wt.%.) necessitated higher temperatures and pressures to achieve good densification. The sintering temperature increased with increased amount of Mo2C (Tables 2 and 4).

3.3.1. Sliding wear The 10Co (LPS) and 9.3Ni (LPS) samples had higher sample wear rates (SWRs) than all the PECS samples, for both the 100Cr6 steel and Si3N4 abrasive balls (Table 6 and Fig. 7), while the 5M–7Ni-u sample had the lowest SWR (Table 6). Apart from the 9.3Ni (LPS) sample, generally the SWR reduced with increased sample hardness (Table 6), for both the 100Cr6 steel and Si3N4 abrasive balls. However, the coefficient of friction (μ) reduced with increased sample hardness for the 100Cr6 steel balls, and increased with sample hardness for the Si3N4 balls (Table 6). The μ due to the 100Cr6 steel balls also decreased with increased ball wear rate (BWR) (Fig. 8). This was confirmed by the iron oxide debris on the wear tracks of the harder 5N-s and 5M–7Ni-u samples (Fig. 9), which was not present on the wear track of the 10Co (LPS) sample. The SEM-SE images of the 10Co-LPS sample wear track from the Si3N4 ball showed depressions and fragmentation of WC grains (Fig. 10). The wear track for 5N-s sample mainly had scratches and fragmentation of WC grains, while the 5M–7Ni-u sample had fewer scratches than 5N-s sample and no fragmentation of WC grains.

3.2. Microstructure The micrographs of the liquid phase sintered WC–0.5Cr3C2–10Co (wt.%) and WC–9.3Ni (wt.%) (nanograde) (Fig. 1(a) and (b)) had larger WC grains (Table 5) and thicker, more homogeneously distributed binder pools than PECS nanograde WC–9.3Ni (wt.%) sample. Although the PECS sample had finer WC grains, the mean grain size was N 1 μm (Table 5), indicating that rapid grain growth took place, even though a short sintering dwell time and lower sintering temperature were used.

3.3.2. Thermal shock and impact resistance analysis Optical images of the 10Co (LPS) and 5N-s insert cutting edges after milling at a cutting speed (vc) of 40 m/min and a depth of cut (ap) of 1 mm had negligible flank wear rates (FWR), but the 5M–7Ni-u insert had a FWR of 143.6 μm/min (Table 7) (only one test was due for the FWR, so no errors were reported.). This was confirmed by the SEM micrographs of the 5M–7Ni-u insert (Fig. 11) which showed the fractured surface of the cutting edge. High magnification of the exposed

Fig. 3. SEM-BSE images of WC–5Mo2C–6.25TiC–9.3Ni (wt.%): WC (light), Ni (medium) and TiC (dark), produced from: (a) nano, and (b) ultrafine WC powders.

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and Av. FR of N 550 °C/s and N700 N, as well as a significant increase in FWRs (Table 7). The 5M–7Ni-u insert's cutting edge suffered a severe failure, having the highest flank wear (VB) of 707.5 μm (Fig. 13(c)), the 10Co (LPS) insert had a much lower VB (Fig. 13(a)). The 10Co (LPS) fracture surface was smaller (Fig. 14) than in the 5M–7Ni-u insert, even though a higher vc (75 m/min) and longer ap (2 mm) were used. Disconnected cracks parallel to the insert cutting edge were observed (indicated by the arrows in Fig. 14(b)). The 10Co (LPS) insert experienced slightly higher resultant forces and had a slightly higher FWR than the 9.3Ni (LPS) insert (Table 7). 4. Discussion

Fig. 4. HR-STEM image of the WC–Ni interface of WC–5Mo2C–6.25TiC–9.3Ni (wt.%), with the inset showing a STEM-EELS elemental map of (left to right) W (orange), C (red), Ti (green) and Ni (blue). Note: Since the journal is printed in black and white only the readers are advised to visit the internet version of this paper where color is retained for a better study of the element distribution.

fracture surface indicated fragmentation and pull-out of WC grains and interconnected micro-voids (~1 μm diameter) between them (indicated by arrows in Fig. 11). The 5M–7Ni-u insert had the highest FWR, even though all inserts had similar thermal variations per second (ΔT/s), and it experienced lower average resultant force (Av. FR) than the 10Co (LPS) insert (Table 7). Fig. 12 shows that after ~100 s, a significant increase in resultant force occurred rapidly (inflection) for the 5M–7Ni-u insert only. The 5M–7Ni-u and 5N-s inserts had points of rapid increase and decrease of resultant force (spikes) approximately every 30 s, which was not seen for the 10Co (LPS) insert (Fig. 12). Increased vc from 40 to 60 m/min, while maintaining the ap, resulted in increased FWR, ΔT/s and average resultant force (Av. FR) in all inserts (Table 7); the 10Co–LPS and 9.3Ni–LPS inserts had the lowest FWRs, followed by the 5N-s insert, the 5M–7Ni-u had the highest FWRs. Further increase of the vc to 75 m/min and ap to 2 mm resulted in ΔT/s

Although both the WC–Co and WC–Ni cemented carbides produced by PECS achieved good densification (N99%) (Table 4), higher sintering temperatures and pressures were required for the WC–Ni samples than the WC–Co samples. This was because of the lower solubility of WC in Ni [4], as well as the poor wetting of WC by Ni, both of which increase with sintering temperature [4,22]. An addition of 5 wt.% NbC to WC– 0.5Cr3C2–10Co (wt.%) needed higher sintering temperature and pressure to achieve good densification, because NbC inhibits WC–Co densification during PECS [8]. Additions of TiC and Mo2C inhibit densification of WC–Ni cemented carbide during PECS [23], hence higher sintering pressures and temperatures are required. Good densification was achieved with increased sintering temperature since the solubility of WC in Ni and wetting of WC by Ni, as well as Mo solubility in Ni increases with sintering temperature [4,22]. This is evident in Figs. 1(c) and 3(a), where additions of Mo2C produced smaller and more evenly distributed Ni binder pools [23]. A higher sintering temperature (1380 °C) was used for the WC–5Mo2C–6.25TiC–7Ni (wt.%) (5M–7Ni) samples than for WC–5Mo2C–6.25TiC–9.3Ni (wt.%) (5M) samples (1330 °C), because of the reduced volume fraction of the lower melting point Ni binder (from 9.3 to 7 wt.%) [4,17]. The larger WC grains in the LPS samples than the PECS samples (Table 5) originated from the higher sintering temperature (1430 °C) and longer sintering time (75 min) which allowed continuous Ostwald ripening [3]. These larger grains meant lower hardness and poor abrasion resistance during sliding wear (Table 6). Abrasion resistance increases with decreased grain size [4], and amount and size of binder

Fig. 5. HAADF-STEM EDX mapping images of WC–6.25TiC–5Mo2C–9.3Ni (wt.%), showing C (purple), Mo (blue), W (green), Ti (yellow) and Ni (red). Note: Since the journal is printed in black and white only the readers are advised to visit the internet version of this paper where color is retained for a better study of the element distribution.

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Fig. 6. (a) HAADF-STEM EDX mapping image of WC–5Mo2C–6.25TiC–9.3Ni (wt.%), showing overlap between W and Mo, and (b) BF-STEM image, showing a WC grain (dark) under Mo2C (light). Note: Since the journal is printed in black and white only the readers are advised to visit the internet version of this paper where color is retained for a better study of the element distribution.

pools [24]. The 9.3Ni (LPS) sample had larger WC grains than the 10Co (LPS) sample, because it was produced from nanostarting powders which had a higher driving force for grain growth from their higher surface energies, leading to much WC grain growth [3,6]. The smaller grains in the 10Co (LPS) sample than the 9.3Ni (LPS) samples was also due to grain growth inhibition by Cr3C2[10]. However, the 9.3Ni (LPS) sample had slightly better sample wear rates (SWRs) than the 10Co (LPS) from both the 100Cr6 steel and Si3N4 balls, even though it had lower hardness. This was because of the slightly lower binder amount [24] and the Ni binder [13]. Nickel reduces loss of WC grains by extruding outwards to replace the eroded binder between the WC grains, since it has a stable ductile fcc structure [13]. The depressions on the wear track of the 10Co (LPS) sample were from the pressure at the contact between the ball and the sample from the applied load (Hertzian pressure), which caused yielding of the Co binder below the wear track, and WC was shifted below the wear surface [25]. The formation of depressions increases with binder pool size and binder amount [25]. The fragmentation of WC grains on the 10Co (LPS) wear was from abrasive plowing of the exposed WC grains [26] by the harder Si3N4 ball (~ 15.7 GPa) [27], through gross plastic deformation [26]. The thicker binder pools were associated with larger WC grains, because binder thickness decreases with decreased WC grain size [4]. The more homogenous distribution of the binders compared to the PECS sample was due to the formation of the binder liquid phase during LPS which enhanced WC solubility [4], as well as the capillarity action of the liquid phase in the pores during the secondary rearrangement stage of sintering [4]. The thicker binder pool size and good binder distribution between WC

grains promoted significantly higher K1C[28] and TRS [29] (Table 6) of both the LPS samples than all PECS samples. The lower TRS in the PECS samples was also due to retained stresses, because of the rapid cooling rates compared to LPS [30]. The 10Co (LPS) and 9.3Ni (LPS) inserts had lower flank wear rates (FWRs) than the PECS inserts, irrespective of vc or ap, due to good thermal shock and impact resistance [31,32]. The good thermal shock and impact resistance were associated with higher TRS and K1C of the LPS inserts than the PECS inserts [4,33]. The FWR increased with vc because of the increased impact (from 5 to 7 collisions per second with the workpiece) and higher thermal variations per second (Table 7) [32]. The 9.3Ni (LPS) sample had a slightly lower TRS than the 10Co (LPS) sample, because the large WC grains which acted as internal defects that prematurely initiated the fractures [34]. The 10Co (LPS) insert had cracks parallel to the cutting edge from the increased mechanical impact of the high cutting speed of 75 m/min (9 collisions with the workpiece per second) [33]. Although the PECS 9.3Ni sample had finer WC grains than the 9.3Ni (LPS) sample (both produced from nanoWC starting powders), its mean grain size was N 1 μm (Table 5), indicating that rapid grain growth took place, even though a short sintering dwell time and lower sintering temperature was used. This rapid grain growth occurred during the heating stage of sintering by coalescence of neighboring grains [3]. Grain growth due to coalescence is thermally activated, and increases with sintering temperatures [3]. Thus, the grain growth of nanopowders during LPS took place in two stages: an initial rapid grain growth by coalescence during heating, and a later stage during the isothermal

Table 6 Sliding wear and mechanical property results. Sample abbreviation

10Co (LPS) 9.3Ni (LPS) 9.3Ni-n 9.3Ni-u 9.3Ni-s 0M-n 0M-u 0M-s 5N-s 5M-s 5M–7Ni-n 5M–7Ni-u

Wear from 100Cr6 steel balls

Wear from Si3N4 balls

Sample wear rate (× 10−5 mm3/N·m)

Av. coefficient of friction, μ (× 10−1)

Sample wear rate (× 10−5 mm3/N·m)

Av. coefficient of friction, μ (× 10−1)

Vickers hardness (GPa)

K1C (MPa·m1/2)

Transverse rupture strength (GPa)

30.210 26.827 18.155 17.343 16.317 12.253 12.071 12.565 11.126 9.823 2.918 2.830

0.88 0.85 0.83 0.82 0.84 0.80 0.82 0.80 0.81 0.77 0.75 0.76

33.812 ± 0.011 30.417 ± 0.010 25.165 ± 0.008 – – – – – 17.332 ± 0.006 – 5.441 ± 0.005 5.314 ± 0.003

0.62 0.63 0.66 – – – – – 0.66 – 0.69 0.71

14.07 12.64 15.54 15.89 15.98 16.17 16.21 16.15 17.47 18.88 19.96 20.33

13.61 14.63 10.35 11.87 11.54 10.52 11.18 11.03 12.51 11.01 10.83 10.27

2286 2175 1785 1817 1830 1718 1741 1718 1954 1658 1496 1467

± ± ± ± ± ± ± ± ± ± ± ±

0.010 0.009 0.006 0.004 0.005 0.006 0.004 0.004 0.005 0.007 0.005 0.007

± ± ± ± ± ± ± ± ± ± ± ±

0.02 0.01 0.02 0.03 0.02 0.03 0.02 0.01 0.03 0.02 0.03 0.02

(s) 0.8 μm starting size, (u) 0.4 μm starting size, (n) 0.1 μm starting size.

Mechanical properties

± 0.02 ± 0.01 ± 0.03

± 0.01 ± 0.02 ± 0.01

± ± ± ± ± ± ± ± ± ± ± ±

0.39 0.92 0.43 0.17 0.20 0.25 0.19 0.18 0.47 0.49 0.31 0.49

± ± ± ± ± ± ± ± ± ± ± ±

0.52 1.54 0.44 0.39 0.24 0.31 0.56 0.40 0.22 0.40 0.32 0.27

± ± ± ± ± ± ± ± ± ± ± ±

62 71 82 82 97 43 79 43 72 85 52 71

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139

Iron oxide debris

Fig. 9. SEM-BSE image of the wear tracks of the 5M–7Ni-u sample from 100Cr6 abrasive ball, showing iron oxide debris.

Fig. 7. Effect of hardness on the wear rates from 100Cr6 steel (steel) and Si3N4 balls.

hold (Ostwald ripening). The slightly smaller sintered mean grain size of the 9.3Ni sample produced from submicron starting powders than the similar compositions produced from nano and ultrafine starting WC powders (Table 5) was due to the higher surface energies in the smaller nano and ultrafine starting powders, giving a higher driving force for WC grains to coalesce and grow [3,6]. The finer grain size of all the PECS 9.3Ni grades improved the hardness and abrasion resistance during sliding wear (Table 6) over the 9.3Ni (LPS) sample, which also had poorer binder distribution than the 9.3Ni (1430) sample (Fig. 1), giving lower K1C[23] and TRS [28,29] (Table 6). The microstructure of the WC–0.5Cr3C2–5NbC–10Co (wt.%) (5N-s) (Fig. 2) shows Co mainly distributed between the WC grains and rarely between the NbC grains, this could be due to the different solubility of NbC and WC in Co, as well as wetting behavior with Co [8]. Compared to NbC, the solubility of WC in Co and wetting of WC by Co are much higher [8], leading to the Co binder beingpreferably located between the WC grains. Niobium was not found in the Co binder either by SEM EDX and HAADF STEM EDX mapping analyses, although according to the Co–Nb phase diagram [35], under the sintering temperature (1280 °C) Nb solubility of up to 8 wt.% in the Co is expected. This could be due to the fact that PECS is a rapid process occurring out of equilibrium [36], while the phases present in the phase diagram are formed under equilibrium conditions. The 5N-s sample had finer

Fig. 8. Effect of the samples' Vickers hardness on the wear rate of the steel balls.

mean grain size than all the 9.3Ni grades (Table 5) because of the submicron WC starting size, which had a lower driving force for grain growth than the nanopowders [3,7] and the Cr3C2 and NbC grain growth inhibitors [10]. This resulted in higher hardness and abrasion wear resistance of the 5N-s sample than 9.3Ni grades (Table 6). The 5N-s sample had the highest K1C and TRS for all the PECS WC–Ni samples, from the better wetting of WC by Co, facilitating better binder distribution [4,23], as well as the slightly higher binder proportion [4]. The good K1C and TRS meant better thermal shock and impact resistance than the WC– 5Mo2C–6.25TiC–Ni (wt.%) (5 M–7Ni-u) insert. However, the 5N-s sample had lower K1C and TRS than the 10Co (LPS) sample, because of the better binder distribution during LPS and higher sintering temperatures that increased the amount of W and C in the (Co), increasing the amount of the tougher fcc allotrope [4]. Addition of 6.25 wt.% TiC to WC–9.3Ni (wt.%) improved hardness and abrasion resistance due to TiC's higher hardness than WC (by ~5 GPa) [37], and its WC grain growth inhibition effect [10]. Layers of TiC atoms were observed at the WC/(Ni) interface (Fig. 4), which may have interfered with the WC/(Ni) interfacial energies, thereby inhibiting solution re-precipitation (Ostwald ripening) [10]. Up to 5 wt.% Mo2C to WC–6.25TiC–9.3Ni (wt.%) significantly reduced the grain size, particularly in the nano and ultrafine grades (Table 5), because Mo2C delayed the aggregation of the fine WC grains, thus preventing coalescence [3]. The ultrafine 5M sample had finer mean grain size than the nano5M sample, because the higher surface energies in the smaller nanostarting powders had a higher driving force for WC grain coalescence [3]. The fine WC grains in the 5M samples gave significantly higher hardness than all the other samples, and the reduction in amount of Ni from 9.3 to 7 wt.% (5M–7Ni) further increased the hardness to N 20 GPa (Table 6), which significantly improved the abrasion

Fig. 10. SEM-SE image of the wear track of 10Co (LPS) from Si3N4 abrasive ball.

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Table 7 Insert properties with different milling conditions; Av. FR is the average resultant force, ΔT/s is the average thermal variation per second and FWR is the flank wear rate. Milling conditions

Properties

Ave. FR (N) ΔT/s (°C/s) FWR (μm/min) Ave. FR (N) ΔT/s (°C/s) FWR (μm/min) Ave. FR (N) ΔT/s (°C/s) FWR (μm/min)

vc = 40 m/min, ap = 1 mm vc = 60 m/min, ap = 1 mm vc = 75 m/min, ap = 2 mm

Cutting tool inserts 10Co (LPS)

9.3Ni (LPS)

5N-s

5M–7Ni-u

354.4 ± 27.7 396.2 ± 46.4 Negligible 341.4 ± 20.8 489.2 ± 34.4 58.4 752.1 ± 44.2 603.4 ± 26.8 147.3

– – – 363.9 504.6 54.9 720.3 596.5 119.8

279.6 ± 13.1 358.1 ± 26.0 Negligible 358.1 ± 32.5 494.3 ± 29.9 126.7 749.6 ± 24.7 610.0 ± 19.2 293.5

301.2 411.4 143.6 435.2 518.5 277.7 776.8 617.6 849.4

(a)

± 24.3 ± 37.2 ± 18.2 ± 31.7

± 19.1 ± 24.2 ± 30.4 ± 18.7 ± 21.3 ± 22.5

(b)

Fig. 11. SEM-SE images of the fractured cutting edge of the 5M–7Ni-u cutting tool insert after a 40 m/min vc and 1 mm ap, showing: (a) Ti–6Al–4V debris (dark gray), and (b) interconnected micro-voids (arrows).

wear resistance (lowest SWRs) (Fig. 7). The 5M–7Ni samples gave the highest 100Cr6 steel ball wear rates (Fig. 8), due to significantly higher hardness than the steel balls (hardness ~7 GPa [38]). The wear rate of the 100Cr6 balls increased with sample hardness, and there was an iron oxide debris (tribolayer) on the wear tracks of the 5N-s and 5M– 7Ni-u samples (Fig. 9), the latter having more of the tribolayer [39]. The coefficient of friction, μ, with the 100Cr6 steel balls reduced with increased sample hardness (Table 6), because the iron oxide tribolayer acted as a lubricant [38]. The 5N-s and 5M–7Ni-u samples had higher

Fig. 12. Variation of resultant forces with time for 10Co (LPS), 5N–Cr and 5M–7Ni-u cutting tool inserts from a vc of 40 m/min. Note: Since the journal is printed in black and white only the readers are advised to visit the internet version of this paper where color is retained for a better study of the element distribution.

hardness than the Si3Ni4 balls, so acted as abrasive bodies, wearing the balls [39]. This increased the friction due to three-body abrasion from the break-up of asperities from the balls and the cemented carbides [40], leading to scratches on the wear track and so μ increased with increased sample hardness (Table 6). The 5M–7Ni-u wear track had fewer scratches than the 5N-s sample, and no WC grain fragmentation because of its higher hardness, which improved the wear resistance. The 5M–7Ni-u sample had the lowest K1C and TRS, due to the lower binder content and retained stress from the rapid cooling from 1380 °C sintering temperature [29]. These lowered the thermal shock and impact resistance, giving the highest FWRs independent of the vc or ap (Table 7). The interconnected micro-voids on the fracture surface of the 5M–7Ni-u insert suggested ductile fracture [41]. Resistance to ductile fracture increases with increased binder proportion [41,42], explaining the lower FWRs of the 5N-s and 10Co (LPS) inserts with a higher binder proportion (10 wt.%). Fracture in fine grained (b2 μm) WC cemented carbides occurs in two stages, interfacial decohesion of the WC grains (stage 1), followed by extensive plastic deformation and rapture of the binder ligaments (ductile fracture) (stage 2) [42]. Fracture in WC cemented carbides is governed by the second stage fracture (ductile fracture) [43], and is dependent on the critical strain energy release rate (G1C) [42]. The resistance to ductile fracture increases with increased G1C[42]. The G1C increases with increased area fraction of the binder and binder mean free path, both of which increase with increased binder proportion [4,42]. The inflection after ~100 s in the 5M–7Ni-u insert (Fig. 12) was due to flank wear which blunted the cutting edge, increasing the cutting force because of increased shear area [44]. This inflection was used as an indication of cutting edge failure. The 10Co (LPS) insert experienced higher resultant cutting forces than 5N-s and 5M–7Ni-u (before the inflection point) inserts (Table 7 and Fig. 12), because of its lower hardness which allowed deformation of the cutting surface (blunting), increasing the cutting forces [44]. The 5M–7Ni-u and 5N-s inserts had rapid increases and decreases

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(b)

(a)

Flank wear (VB) = 122.4 µm

Flank wear (VB) = 264.6 µm

141

(c)

Flank wear (VB) = 707.5 µm

Fig. 13. Flank wear on (a) 10Co (LPS), (b) 5N–Cr and (c) 5M–7Ni-u cutting tool insets from a 75 m/min cutting speed (vc) and 2 mm depth of cut (ap).

of resultant forces (“spikes”) every ~30 s. The “spikes” were the points at which the inserts would re-engage the workpiece after a complete length of cut (i.e. points of maximum impact). The 10Co (LPS) insert did not show spikes because of its higher TRS. The 5N-s insert experienced the lowest resultant forces, leading to the lower thermal variation per second (Table 7), because of its hard cutting edge, good TRS and K1C (Table 6). The sample wear rate reduced with increased sample hardness, showing that the abrasion resistance during sliding wear increased with WC grain size refinement and reduction of the binder amount, explaining the significantly higher abrasion wear resistance of the PECS 5M–7Ni samples compared to the LPS samples. However, the LPS inserts had much lower flank wear rates than the PECS inserts because of their significantly higher K1C and TRS that improved their thermal shock and impact resistance during the milling test. Hence, LPS samples were better against thermal shock and impact, whereas PECS samples were better against abrasion. 5. Conclusions Through pulse electric current sintering, different starting particle sizes, different binders and binder amounts and the use of Cr3C2, NbC,

(a)

TiC and Mo2C additions, WC-cemented carbides with good mechanical properties were achieved. Pulse electric current sintering resulted in finer WC grains and poorlydistributed binder pools than liquid phase sintering, resulting in higher hardness and improved abrasion wear resistance, but lower K1C and TRS. The lower TRS was also attributed to the retained stresses from the rapid cooling after PECS. The 10Co (LPS) sample had a higher TRS than the 9.3Ni (LPS) sample, due to the larger WC grains which acted as defects. The LPS samples had the highest thermal shock and impact resistance, shown by their higher TRS and K1C. The 5N-s sample had a good hardness, because of PECS and addition of NbC and Cr3C2 grain growth inhibitors, as well as good fracture toughness and TRS, because of a high binder amount and good wetting of WC by Co, giving good combinations of abrasion wear, thermal shock and impact resistance. Although PECS was used, the WC–9.3Ni (wt.%) samples produced from nano and ultrafine WC starting powders had mean grain sizes N1 μm, attributed to rapid grain growth from grain coalescence. Addition of 6.25 TiC (wt.%) slightly improved the hardness, and a further 5 Mo2C (wt.%) significantly reduced the grain growth, improving the hardness and abrasion wear resistance. Reduction of Ni binder from 9.3 wt.% to 7 wt.% increased both hardness and abrasion wear resistance, but reduced the K1C and TRS, leading to poor thermal shock and impact resistance.

(b)

Fig. 14. SEM-SE images of the fractured cutting edge of the 10Co (LPS) cutting tool insert after 75 m/min vc and 1 mm ap, showing cracks parallel to the cutting edge (arrows).

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