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TaC doped ultrafine WC-Co tool material by spark plasma sintering
Journal Pre-proof Preparation and properties of the VC/Cr3C2/TaC doped ultrafine WC-Co tool material by spark plasma sintering Boxiang Wang, Zhenhua Wang, Zengbin Yin, Juntang Yuan, Jiheng Jia PII:
S0925-8388(19)33844-7
DOI:
https://doi.org/10.1016/j.jallcom.2019.152598
Reference:
JALCOM 152598
To appear in:
Journal of Alloys and Compounds
Received Date: 20 May 2019 Revised Date:
10 September 2019
Accepted Date: 6 October 2019
Please cite this article as: B. Wang, Z. Wang, Z. Yin, J. Yuan, J. Jia, Preparation and properties of the VC/Cr3C2/TaC doped ultrafine WC-Co tool material by spark plasma sintering, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.152598. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Preparation and properties of the VC/Cr3C2/TaC doped ultrafine WC-Co tool material by spark plasma sintering Boxiang Wang, Zhenhua Wang*, Zengbin Yin, Juntang Yuan*, Jiheng Jia School of Mechanical Engineering, Nanjing University of Science and Technology, Nanjing 210094, China * Corresponding author. E-mail address: [email protected] (Zhenhua Wang); [email protected] (Juntang Yuan)
Abstract This work is focused on the preparation and properties of ultrafine WC-Co cemented carbides. Samples were prepared by spark plasma sintering technology using WC and Co nanopowders and different inhibitors such as VC, Cr3C2 and TaC. The influence of inhibitors on the microstructure and mechanical properties of cemented carbides was investigated by using separate addition of VC/Cr3C2/TaC (0, 0.4 and 0.8 wt.%) and pairwise addition (VC-Cr3C2, VC-TaC and Cr3C2-TaC). The results showed that with the addition of VC/Cr3C2/TaC, the average WC grain size decreased and VC led to smaller grain size than Cr3C2 and TaC. WC grain morphologies of cemented carbides with VC were triangular prisms with multi-step structure, whereas tended to be round and blunt in cemented carbides with Cr3C2/TaC. Moreover, the addition of inhibitors caused hardness to increase and fracture toughness to decrease in various degrees. Transverse rupture strength (TRS) was affected by changes in grain size and WC/Co interfacial energy caused by inhibitors. The sample with 0.4 wt.% VC-0.4 wt.% Cr3C2 sintered at 1250 °C had comprehensive mechanical properties. Its relative density, average grain size, hardness, fracture toughness and TRS were 98.9%, 200 nm, 2110 kg/mm2, 10.4 MPa·m1/2 and 1990 MPa, respectively. Keywords: WC-Co; Grain growth inhibitor; Grain morphology; Mechanical properties; Spark plasma sintering
1. Introduction WC-Co cemented carbides are typical tool materials and have been widely used for manufacturing various cutting tools due to their high hardness and wear resistance [1]. To meet increasing demands in the area of high speed and precision machining, further improvement of hardness and wear resistance is highly required for WC-Co cemented carbides [2,3]. For this purpose, a lot of efforts have been made to produce nanoscale WC-Co materials [4]. For example, it is possible to sinter cemented carbides with nanoscale microstructures using a small amount of carbide powders as grain growth inhibitors (GGI). Comparing to cemented carbides with micron-sized grains, the hardness and wear resistance of nanoscale cemented carbides are greatly improved due to their finer WC grains. It has been commonly recognized that adding inhibitors in initial powder mixtures is indispensable to suppress grain growth. So far, almost all GGI reported in publications are transition metal carbides, such as VC, Cr3C2, TiC, TaC, NbC and Mo2C, among which VC and Cr3C2 have best inhibition effect [5-10]. Moreover, inhibition mechanisms have been studied by many publications. GGI can alter WC interfacial
energy,
reduce
solubility
of
WC
in
binder
phases
and
interfere
dissolution-reprecipitation effect [7,9,11,12]. Complexions generated by inhibitors effectively hinder migration of grain boundaries and suppressed grain growth [3]. With the addition of VC, V is segregated along WC/Co interfaces during sintering process, and high density hemisphere (W, V) Cx adjacent to (1 01 0) planes inhibits the grain growth [13,14]. Lay et al. [15] indicated that (V, W) Cx layers are formed in the early stage of sintering process. However, Kawakami et al. [16] thought that (V, W) Cx layers are formed in the cooling stage rather than sintering stage. During the conventional liquid phase sintering, WC grain size increases dramatically when
cemented carbides hold for a few minutes at liquidus temperature [17-20]. Even with pressure assistance, conventional liquid phase sintering is not suitable for producing nanocrystalline microstructures. Therefore, it is feasible to sinter WC nanopowders containing GGI at under-eutectic temperature using very fast heating rate and short sintering time [11,21-24]. Spark plasma sintering is a rapid sintering technology which is beneficial for suppressing grain growth. Some studies showed that samples produced by spark plasma sintering have finer grain structures than the samples produced by hot isostatic pressing [12,25]. It is well known that grain size decreases with the addition of GGI, resulting in hardness of cemented carbides to increase generally. However, the changes in fracture toughness and TRS are uncertain. The effect of GGI on the comprehensive mechanical properties of cemented carbides has not been sufficiently understood yet. In this work, ultrafine WC-8Co-(VC/Cr3C2/TaC) cemented carbides were prepared by spark plasma sintering technology. The influence of GGI on the microstructure and mechanical properties of cemented carbides were studied in detail.
2. Experimental 2.1 Materials preparation WC powder with average grain size of 60 nm and Co powder with average grain size of 500 nm were used as raw materials (purity: 99.99%, ChaoWei Nanotechnology Co., Ltd, Shanghai). VC, Cr3C2 and TaC (purity: 99.99%, ChaoWei Nanotechnology Co., Ltd, Shanghai) were added as GGI. SEM images of as-received powders are shown in Fig. 1. As-received powders were mixed in glass container with pure ethanol. Then, the mixture was mixed by ultrasonic vibration and mechanical stirring for two hours in an ultrasonic platform equipped with stirrer (DSA200-JY1-9.0L, China). The powder mixture was dried in a vacuum
drying oven (DZF, China) at 80°C and sieved with a 100 mesh sieve. Then, the powder mixture was sintered in a spark plasma sintering furnace (LABOX-650F, Japan) to fabricate samples with dimensions of 15×15×5 mm, with the heating rate of 100 °C/min, sintering temperature of 1250, 1300 and 1350 °C, constant pressure of 70 MPa and holding time of 5 min. After that, samples were cooled to the room temperature in the furnace. Composition ratios and symbol of samples in present work are given in Table 1.
Fig. 1. SEM images of as-received raw powders, (a) WC, (b) Co, (c) VC, (d) Cr3C2 and (e) TaC. Table 1 Composition ratios and symbol of cemented carbides (wt.%). Cemented carbides
Symbol
WC
Co
VC
Cr3C2
TaC
WC-8Co
A1 (S1)
Bal.
8
0
0
0
WC-0.4VC-8Co
— (S2)
Bal.
8
0.4
0
0
WC-0.4Cr3C2-8Co
— (S3)
Bal.
8
0
0.4
0
WC-0.4TaC-8Co
— (S4)
Bal.
8
0
0
0.4
WC-0.8VC-8Co
A2 (S5)
Bal.
8
0.8
0
0
WC-0.8Cr3C2-8Co
A3 (S6)
Bal.
8
0
0.8
0
WC-0.8TaC-8Co
A4 (S7)
Bal.
8
0
0
0.8
WC-0.4VC-0.4Cr3C2-8Co
A5 (S8)
Bal.
8
0.4
0.4
0
WC-0.4VC-0.4TaC-8Co
A6 (S9)
Bal.
8
0.4
0
0.4
WC-0.4Cr3C2-0.4TaC -8Co
— (S10)
Bal.
8
0
0.4
0.4
2.2 Characterization Density of sintered samples was measured by Archimedes principle. TRS was measured on samples with dimensions of 15×5×3 mm using a three-point bending instrument (UTM5105-G, China). Span and loading velocity was 12 mm and 0.1 mm/min, respectively. Vickers hardness was measured using a hardness tester (MN-9631-130-C, China) with load of 30 kg and loading duration of 15 s. Fracture toughness was calculated based on the length of crack generated by indentation using the following equation [26]. KIC = 0.0028 · (HV·P/L)1/2
(1)
where KIC is fracture toughness (MPa·m1/2); HV is Vickers hardness (N/mm2), P is load (N); and L is total length of cracks (mm). At least five measurements were taken and average value was obtained. Phase in sintered samples was analyzed by X–ray diffraction (XRD, D8 Advance, Germany) with copper Kα radiation. A scanning electron microscope (SEM, Quant 250FEG, USA) was used for observing the microstructure on polished surfaces of sintered samples. Based on SEM images of polished surfaces, size distribution and average size of WC grains were calculated by measuring more than 400 grains using the image processing software (Image-Pro Plus 6.0). Element distribution in the microstructure was investigated using a SEM-energy dispersive spectrometer (SEM-EDS) analysis system. WC grains in sintered samples were extracted by placing samples in hydrochloric acid to remove Co phases, and then the WC grain morphology was observed by SEM.
3. Results and discussion 3.1 Microstructure and phase constitution analysis Fig. 2 shows SEM (BSE) images of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at 1250 °C. Bright areas represent WC phases and black areas are Co phases. It was found that Co phases were distributed between WC grains. There were a few Co pools and pores with irregular shapes and size of several hundred nanometers. Because Co segregation and insufficient wettability were typical characteristics of solid phase sintering, Co phases were not sufficiently diffused in spaces formed by WC skeleton [21,22], which led to formation of Co pools and pores. Lack of Co in the region near Co pool led to increase of WC grain contiguity [7]. In these regions, contact of WC grains promoted grain growth. This phenomenon was prominent in sample A1. As shown in Fig. 2b-f, although Co rich phases did not completely disappear with the addition of VC/Cr3C2/TaC, the microstructure became fine and grain shape transformed from faceted shape to blunt shape. In addition, from Fig. 2b-d, it was found that the addition of VC led to finer WC grains than addition of Cr3C2/TaC. Still Cr3C2 and TaC had important effect on grain refinement despite the existence of coarse grains.
Fig. 2. SEM (BSE) images showing microstructures of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5, (f) A6 and (g) EDS line analysis from point A to C in A1.
As the sintering temperature increased, high thermal energy was generated and dissolution-reprecipitation effect was promoted, which enhanced driving force for diffusion of Co phases [17]. Therefore, the increasing sintering temperature resulted in the densified microstructure, as shown in Fig. 3. It was apparent that WC grains in sample A1 sintered at 1350 °C almost transformed to faceted shape compared to that sintered at 1250 °C. Compared Fig. 2 with Fig. 3, limited but significant grain growth was clearly observed at 1350 °C regardless of chemical composition. However, obvious inhibition effects were still observed in samples with GGI as compared to sample A1, as shown in Fig. 3. Moreover, no abnormal grain growth was
observed in any of samples with the increase in temperature. Microstructures of samples sintered at 1300 °C were not shown because it was similar to that of samples sintered at 1350 °C.
Fig. 3. SEM (BSE) images showing microstructures of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1350 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5 and (f) A6.
Fig. 4 shows XRD patterns of composite powders and sintered samples. No obvious differences were observed in phases formed before and after sintering. There were only WC and Co phases in all samples. No graphite or η phase was observed in any of sintered samples. Diffraction peaks of VC, Cr3C2 and TaC were not detected due to the low content, which was also mentioned in many studies [5,23,27]. Moreover, as there are not different in phases formed at different sintering temperatures, XRD patterns of samples sintered at 1300 and 1350 °C are not given.
Fig. 4. XRD patterns of composite powders and sintered samples, (a) composite powders and (b) samples sintered at 1250 °C.
Fig. 5 shows size distribution and average size of WC grains in WC-8Co-(VC/Cr3C2/TaC) cemented carbides. It can be seen that size distribution became narrow and average grain size decreased with the addition of GGI. As shown in Fig. 5a, the grain size of sample A1 ranged from 100 to 900 nm and average grain size was 280 nm at sintering temperature of 1250 °C. By adding 0.8 wt.% VC/Cr3C2/TaC, the average grain size of samples A2, A3 and A4 reduced to 170, 230 and 250 nm, respectively. Compared to sample A1, the average grain size of sample A2 reduced by 40%, which was similar to results of some studies [5,12]. In these studies, 0.8 wt.% VC was added during the conventional liquid phase sintering. Therefore, inhibition effect of VC was not limited to liquid phase sintering stage. Even if at under-eutectic temperature, VC still had significant effect on suppressing grain growth. Moreover, as half of VC was replaced by Cr3C2/TaC, the average grain size of samples A5 and A6 increased compared to sample A2. Average grain size of sample A5 with 0.4 wt.% VC-0.4 wt.% Cr3C2 was 200 nm. Sample A6 with 0.4 wt.% VC-0.4 wt.% TaC had average grain size of 220 nm. This result showed that inhibition effect of VC was better than that of Cr3C2 and TaC.
Fig. 5. Cumulative size distribution and average size of WC grains in WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at (a) 1250 °C, (b) 1300 °C and (c) 1350 °C.
From Fig. 5b and c, it was found that proportion of coarse grains increased due to the grain growth with the increasing sintering temperature, resulting in the increasing average grain size. From 1250 to 1350 °C, the average grain size of samples increased by about 40-70%. At 1350 °C, the average grain size of sample A2 increased from 170 to 260 nm, but was still smaller than that of sample A1 sintered at 1250 °C. During the liquid phase sintering, the inhibition effect of GGI greatly depends on limiting dissolution-reprecipitation mechanism [12,18]. GGI can be dispersed
at WC surfaces to reduce surface energy of WC and solubility of WC in Co phases [16,28]. Meanwhile, GGI are dissolved in binder phase and further reduces solubility of WC [7,10]. During the spark plasma sintering, dissolution-reprecipitation effect are limited by low sintering temperature (1250 °C). However, VC still has significant effect on inhibiting grain growth, which is likely related to formation of V rich layers. Thin (V, W) Cx films are formed at WC/Co interfaces, which limits diffusion phenomenon involved in grain growth. 3.2 Effect of GGI on the WC grain morphology WC grain morphologies of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at 1250 °C are shown in Fig. 6. From Fig. 6a, it was found that WC grains of sample A1 were truncated triangular prisms, which is the equilibrium morphology of WC grains [29]. It is well known that WC has two (0 0 0 1) basal planes and three (1 01 0) prismatic planes. (0 0 0 1) and (1 01 0) planes were seen as triangular shapes and rectangular shapes, respectively. As shown in Fig. 6b, WC grains of sample A2 had triangular prismatic shapes with multi-steps based on basal planes. Multi-steps were composed of (0 0 0 1) and (1 01 0) planes parallel to basal and prismatic planes, respectively.
Fig. 6. SEM images showing WC grain morphologies of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5 and (f) A6.
The growth process of WC grains is summarized in Fig. 7. Fig. 7a shows growth process of WC grains in cemented carbides without GGI. Coalescence of WC particles occurs in the early sintering stage of nano-sized WC-Co powder [20]. After coalescence, WC grains exhibit round shapes and irregular shapes at low sintering temperature. As the temperature increases, diffusion mobility of atoms enhances as well as W and C diffuse through Co films at surfaces of round grains, resulting in faceted shapes of grains. Truncated triangular prisms are formed from faceting
of round grains driven by surface energy minimization [19].
Fig. 7. Schematic drawing showing growth mechanisms of WC grains, (a) WC-Co and (b) WC-Co-VC.
The existence of multi-steps indicates that grain growth is mainly controlled by 2D nucleation mechanism [30]. Generally, W and C atoms diffuse from the edges of (0 0 0 1) planes towards center and reach the center to form a nucleus. Then, the nucleus grows rapidly until it reaches edges. However, the addition of VC causes (V, W) Cx layers to form at (0 0 0 1) and (1 01 0) planes of WC grains [13-16]. (V, W) Cx layers enhance energy barrier of 2D nucleation in the edge region of (0 0 0 1) planes and hinder diffusion of W and C atoms. The new nucleus can appear at the center of (0 0 0 1) planes before growth completion of preceding nucleus, and then the existing nucleuses grow simultaneously on (0 0 0 1) planes [19]. Therefore, WC grains are stacked along [0 0 0 1] direction to form multi-steps based on (0 0 0 1) basal planes. Grain growth along [0 0 0 1] direction is effectively inhibited because (W, V) Cx layers are thicker on (0 0 0 1) planes than on (1 01 0) planes [3]. As Cr has high solubility in Co phases and WC dissolves in TaC to form (Ta, W) C solid solution, segregation layers with Cr and Ta are not formed at WC/Co interfaces. The multi-step
structure was not observed in samples A3 and A4, as shown in Fig. 6c and d. Compared to sample A1, grain morphologies of samples A3 and A4 tended to become round and blunt. From Fig. 6e and f, it was found that WC grains with multi-step structure in samples A5 and A6 reduced significantly when half of VC was replaced by Cr3C2/TaC. 3.3 Mechanical properties 3.3.1 Effect of sintering temperature on the mechanical properties Fig. 8 shows relative density and mechanical properties of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at different sintering temperatures. From Fig. 8a, it was found that almost all samples exhibited good relative density of above 98.5% at 1250 °C. Sample A1 had higher relative density of 99.5%. For samples A2, A3 and A4, the relative density was 99.4, 98.9, and 98.4%, respectively. With the addition of VC/Cr3C2/TaC, the relative density presented a slight variation. This variation was related to the different solubility of GGI in Co phases and the limited diffusion of Co phases at low sintering temperature [9,21,22]. As the sintering temperature increased, relative density of sample A1 was not improved but that of samples with GGI slightly increased. For samples with GGI, the GGI dispersing in binder phases reduced the solubility of WC in binder phases and the driving force of WC particles. As the temperature increased to 1300 and 1350 °C, the solubility of WC in binder phases increased. Meanwhile, the higher temperature provided a better driving force for the binder phases, causing the liquid phases to diffuse along the capillary formed by WC grains. Therefore, the dissolution-reprecipitation and rearrangement of carbides as well as the enhancing liquid phase fluidity caused the relative density to increase slightly.
100.00
(a)
(b)
2200
99.75
2150
99.50
2100
99.25
2050
99.00
2000
98.75
1950 A1 A4
98.50 98.25
1250
A2 A5
1300
A1 A4
1850 1250
1350
A1 A4
1300
1350
Temperature (°C)
12.0 11.5
A3 A6
1900
A3 A6
Temperature (°C)
(c)
A2 A5
A2 A5
A3 A6
(d)
2000
A1 A4
A2 A5
A3 A6
1900
11.0 1800 10.5 1700 10.0 1600 9.5 1500 9.0 1400 8.5
1250
1300
Temperature (°C)
1350
1250
1300
1350
Temperature (°C)
Fig. 8. (a) relative density, (b) hardness, (c) fracture toughness and (d) transverse rupture strength of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at different temperatures.
From Fig. 8b and c, it was clear that all samples sintered at 1250 °C achieved maximum hardness and fracture toughness. From 1250 to 1350 °C, the hardness presented a slow downward trend (reduced by 5-11%), while the fracture toughness exhibited few changes. Densification and grain size of cemented carbides have significant effect on their hardness [32]. Generally, the hardness increases as the relative density increases while decreases as the WC grain size increases (Hall-Petch relation). With the increasing temperature, the change in relative density was slight (Fig. 8a), while the WC grain size increased significantly (Fig. 5). As a result, the grain growth caused by the increasing temperature is the primary reason of the decrease in hardness. From Fig. 8d, it can be seen that the TRS of most samples reduced gradually with the increasing temperature, but that of sample A4 presented a deviation behavior. At 1250 °C, most
samples reached good densification which effectively decreased strain points and improved the TRS. From 1250 to 1350 °C, the fracture paths along WC/Co and WC/WC interfaces decreased due to the grain coarsening, resulting in the reduction of fracture energy. As a result, the TRS decreased gradually. However, an acceptable densification that was beneficial for the TRS was not achieved for sample A4 sintered at 1250 °C. The increasing temperature was effective on eliminating pores, causing the TRS of sample A4 to increase. 3.3.2 Effect of GGI content on the mechanical properties From analysis in Sec. 3.3.1, it is clear that good mechanical properties are reached at sintering temperature of 1250 °C. Therefore, the influence of GGI content on the mechanical properties of cemented carbides sintered at 1250 °C was further studied. Fig. 9 shows hardness and fracture toughness of samples sintered at 1250 °C. From Fig. 9a, it was found that the hardness increased gradually (1970→2080→2180 kg/mm2) with the increasing VC content (S1→S2→S5: 0→0.4→0.8 wt.%). By adding 0.4 wt.% Cr3C2/TaC, the hardness of samples S3 and S4 increased to 2020 and 2030 kg/mm2, respectively. However, as Cr3C2/TaC content increased to 0.8 wt.%, the hardness of samples S6 and S7 were not further improved. Moreover, as half of VC was replaced by Cr3C2/TaC, the hardness of samples S8 and S9 tended to decrease compared to sample S5. Because sample with VC had smaller grain size than that with Cr3C2/TaC, the sample with VC had higher hardness. It is well known that there is an antagonistic correlation between fracture toughness and hardness and improvement of hardness is generally accompanied by decrease of fracture toughness [8]. As shown in Fig. 9b, the fracture toughness decreased in various degrees with the addition of GGI. Sample S1 had higher fracture toughness (11.5 MPa·m1/2) compared to other
samples. However, sample S5 achieved minimum fracture toughness of 9.4 MPa·m1/2. Compared to sample S5, the fracture toughness of sample S8 and S9 increased because Cr3C2/TaC replaced half of VC. WC grain size has an important influence on fracture toughness of WC-Co when Co content keeps constant. When cracks propagate to WC grains with sub-micron and micron sizes, crack can pass directly through WC grains, which leads to the formation of transgranular fracture. As the intragranular strength is higher than intergranular strength, the transgranular fracture prevents effectively crack propagation, which is conducive to fracture toughness. However, with the decreasing grain size, cracks can propagate easily along grain boundaries and the transgranular fracture occurs rarely. In addition, the crack propagation morphologies can explain well for changes in fracture toughness, as shown in Fig. 10. From sample S1, it was clear that crack paths changed significantly during the crack propagation. The existence of micron-sized grains had a deflection effect on crack propagation [32]. Moreover, crack bridging increased fracture surface areas and enhanced energy consumption. As a result, crack deflection and crack bridging enhanced the resistance against crack propagation and improved fracture toughness. As shown in Fig. 10b-d, with the decrease of grain size, cracks extended mainly along grain boundary, which limited the contribution of crack deflection to fracture toughness. Therefore, in this study, sample without GGI had the higher fracture toughness due to its relatively larger grain size, while the fracture toughness decreased due to the decreasing grain size with the addition of GGI.
Fracture toughness (MPa·m½)
Vickers hardness (kg/mm2)
Fig. 9. (a) hardness and (b) fracture toughness of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co.
Fig. 10. SEM (BSE) images showing crack propagation morphologies of cemented carbides sintered at 1250 °C, (a) S1: WC-8Co, (b) S2: WC-0.4VC-8Co, (c) S5: WC-0.8VC-8Co and (d) S8: WC-0.4VC-0.4Cr3C2-8Co.
The influence of GGI on the TRS of samples is shown in Fig. 11. It can be seen that the TRS first increased and then decreased (1670→1810→1690 MPa) with the increasing VC content (S1 →S2→S5: 0→0.4→0.8 wt.%). As the WC grain size decreases, the volume fraction of total boundary areas between WC and Co increases. The increase of fracture path through WC/Co interfaces can bring about numerous fracture energies and enhance the TRS of cemented carbides
[33]. Therefore, if there are no pores, inclusions and abnormally large grains in the microstructure, the TRS is usually in inverse ratio to average grain size. However, compared to sample S1, the TRS of sample S5 did not increase significantly despite significant grain refinement. This phenomenon is likely related to VC additions. Generally, the addition of GGI can reduce WC/Co interface coherency, and low fraction of WC/Co interfacial coherency is detrimental to the TRS [4,34]. On the one hand, samples with VC has high stability of (V, W) Cx layers on WC surfaces and low fraction of WC/Co interface coherency [3], which decreases significantly WC/Co interfacial energy and results in the low separation energy. On the other hand, multi-step structure of WC grains causes stress concentration and increases fracture sensitivity [35], which makes it easy for cemented carbides to fracture along WC/Co interfaces. Therefore, the grain refinement, low fraction of WC/Co interface coherency and fracture sensitivity affected comprehensively the TRS of samples with VC. These factors worked together to cause the TRS to first increase and then decrease with the increasing VC content. The TRS of samples S8 and S9 increased by 17 and 5% compared to sample S5, respectively. As half of VC was replaced by Cr3C2/TaC, the V segregation and multi-step structure decreased, which was beneficial to improve the TRS. In addition, the TRS first decreased and then increased (1670→1540→1830 MPa) with the increasing Cr3C2 content (S1→S3→S6: 0→0.4→0.8 wt.%). Cr3C2 has weakest effect on reducing WC/Co interface coherency compared to other GGI [3]. As a result, the TRS of sample S6 increased by 13% compared to sample S1, which was a comprehensive outcome of grain refinement and large fraction of WC/Co interface coherency.
Transverse rupture strength (MPa) Fig. 11. Transverse rupture strength of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co.
Fig. 12 shows fracture surfaces of samples sintered at 1250 °C. For sample S1, transgranular fracture of coarsened WC grains and intergranular fracture of carbides were observed clearly, as shown in Fig. 12a. With the decrease of grain size due to the GGI additions, fracture surfaces presented that intergranular fracture became predominant while transgranular fracture decreased, as shown in Fig. 12b-f. In addition, some pores were observed on the fracture surface of sample S7, as shown in Fig. 12d. Pores are main fracture sources of cemented carbides which can cause stress concentration and produce cracks during deformation [31]. Therefore, the TRS of sample S7 was 1440 MPa and decreased by 17% compared to sample S1.
Fig. 12. SEM images showing fracture surfaces of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) S1: WC-8Co, (b) S5: WC-0.8VC-8Co, (c) S6: WC-0.8Cr3C2-8Co, (d) S7: WC-0.8TaC-8Co, (e) S8: WC-0.4VC-0.4Cr3C2-8Co and (f) S9: WC-0.4VC-0.4TaC-8Co.
Fig. 13 shows the relationship between mechanical properties of samples. It was found that distribution between hardness and fracture toughness exhibited an inverse relationship. Samples with mixed GGI were in the intermediary position (S8, S9 and S10). High hardness were achieved by using VC or mixed GGI with VC (S2, S5, S8 and S9). The addition of TaC neither provided high fracture toughness nor significantly improved hardness (S4 and S7).
Comprehensive mechanical properties were achieved by adding 0.4 wt.% VC-0.4 wt.% Cr3C2, with the hardness of 2110 kg/mm2, fracture toughness of 10.4 MPa·m1/2 and TRS of 1990 MPa. Properties of some WC-Co cemented carbides prepared by different GGI additives and sintering technologies are given in Table 2. It was clear that the cemented carbide prepared by spark plasma sintering in this work has a good comprehensive hardness and fracture toughness compared to other research results.
Fig. 13. Relationship between hardness, fracture toughness and transverse rupture strength of samples sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co. Table 2 Properties and sintering technologies of some WC-Co cemented carbides. GGI additives (wt.%)
Co content (wt.%)
Grain size (nm)
Hardness (kg/mm2)
KIC (MPa·m1/2)
Sintering technology
0.9NbC [36]
12
320
1631
10.8
Vacuum sintering
0.5VC [6]
12
550
1490
10.9
Spark plasma sintering
1Cr3C2 [12]
12
214
1870
10.5
Hot isostatic pressing
0.26VC-0.45Cr3C2 [37]
6
210
2041
9.0
Liquid phase sintering
0.5VC-0.5Cr3C2 [21]
12
190
1925
10.2
Spark plasma sintering
0.6VC-0.3Cr3C2 [24]
10
–
2070
8.2
Hot isostatic pressing
4. Conclusions (1) Average WC grain size of cemented carbides decreased from 280 to 170, 230 and 250 nm with the addition of 0.8 wt.% VC/Cr3C2/TaC. The effect of VC on inhibiting grain growth was better than that of Cr3C2 and TaC. With the addition of VC, WC grain morphologies transformed from equilibrium to triangular prismatic shapes with multi-step structure because (V, W) Cx layers were formed at WC surfaces. However, grain morphologies became round and blunt with the addition of Cr3C2/TaC. (2) As sintering temperature increased from 1250 to 1350 °C, the hardness and TRS decreased gradually due to grain growth, while fracture toughness had few changes. At 1250 °C, with the addition of VC/Cr3C2/TaC, the smaller the grain size of sintered samples, the higher the hardness of sintered samples, whereas fracture toughness is in an inverse ratio to hardness. Grain size and WC/Co interfacial energy comprehensively affected the TRS of cemented carbides. (3) Excellent comprehensive mechanical properties can be achieved by using mixed GGI to balance and optimize the relationship between grain refinement and WC/Co interfacial energy.
Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 51775280 and No. 51675285) and the Jiangsu Provincial Six Talent Peaks Project (2016-HKHT-019).
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Highlights: • The effect of VC on inhibiting grain growth is better than that of Cr3C2 and TaC. • WC grains of alloys with VC are triangular prisms with multi-step structure. • VC/Cr3C2/TaC causes hardness to increase and fracture toughness to decrease. • Grain size and interfacial energy comprehensively affect transverse rupture strength.
S0925-8388(19)33844-7
DOI:
https://doi.org/10.1016/j.jallcom.2019.152598
Reference:
JALCOM 152598
To appear in:
Journal of Alloys and Compounds
Received Date: 20 May 2019 Revised Date:
10 September 2019
Accepted Date: 6 October 2019
Please cite this article as: B. Wang, Z. Wang, Z. Yin, J. Yuan, J. Jia, Preparation and properties of the VC/Cr3C2/TaC doped ultrafine WC-Co tool material by spark plasma sintering, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.152598. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Preparation and properties of the VC/Cr3C2/TaC doped ultrafine WC-Co tool material by spark plasma sintering Boxiang Wang, Zhenhua Wang*, Zengbin Yin, Juntang Yuan*, Jiheng Jia School of Mechanical Engineering, Nanjing University of Science and Technology, Nanjing 210094, China * Corresponding author. E-mail address: [email protected] (Zhenhua Wang); [email protected] (Juntang Yuan)
Abstract This work is focused on the preparation and properties of ultrafine WC-Co cemented carbides. Samples were prepared by spark plasma sintering technology using WC and Co nanopowders and different inhibitors such as VC, Cr3C2 and TaC. The influence of inhibitors on the microstructure and mechanical properties of cemented carbides was investigated by using separate addition of VC/Cr3C2/TaC (0, 0.4 and 0.8 wt.%) and pairwise addition (VC-Cr3C2, VC-TaC and Cr3C2-TaC). The results showed that with the addition of VC/Cr3C2/TaC, the average WC grain size decreased and VC led to smaller grain size than Cr3C2 and TaC. WC grain morphologies of cemented carbides with VC were triangular prisms with multi-step structure, whereas tended to be round and blunt in cemented carbides with Cr3C2/TaC. Moreover, the addition of inhibitors caused hardness to increase and fracture toughness to decrease in various degrees. Transverse rupture strength (TRS) was affected by changes in grain size and WC/Co interfacial energy caused by inhibitors. The sample with 0.4 wt.% VC-0.4 wt.% Cr3C2 sintered at 1250 °C had comprehensive mechanical properties. Its relative density, average grain size, hardness, fracture toughness and TRS were 98.9%, 200 nm, 2110 kg/mm2, 10.4 MPa·m1/2 and 1990 MPa, respectively. Keywords: WC-Co; Grain growth inhibitor; Grain morphology; Mechanical properties; Spark plasma sintering
1. Introduction WC-Co cemented carbides are typical tool materials and have been widely used for manufacturing various cutting tools due to their high hardness and wear resistance [1]. To meet increasing demands in the area of high speed and precision machining, further improvement of hardness and wear resistance is highly required for WC-Co cemented carbides [2,3]. For this purpose, a lot of efforts have been made to produce nanoscale WC-Co materials [4]. For example, it is possible to sinter cemented carbides with nanoscale microstructures using a small amount of carbide powders as grain growth inhibitors (GGI). Comparing to cemented carbides with micron-sized grains, the hardness and wear resistance of nanoscale cemented carbides are greatly improved due to their finer WC grains. It has been commonly recognized that adding inhibitors in initial powder mixtures is indispensable to suppress grain growth. So far, almost all GGI reported in publications are transition metal carbides, such as VC, Cr3C2, TiC, TaC, NbC and Mo2C, among which VC and Cr3C2 have best inhibition effect [5-10]. Moreover, inhibition mechanisms have been studied by many publications. GGI can alter WC interfacial
energy,
reduce
solubility
of
WC
in
binder
phases
and
interfere
dissolution-reprecipitation effect [7,9,11,12]. Complexions generated by inhibitors effectively hinder migration of grain boundaries and suppressed grain growth [3]. With the addition of VC, V is segregated along WC/Co interfaces during sintering process, and high density hemisphere (W, V) Cx adjacent to (1 01 0) planes inhibits the grain growth [13,14]. Lay et al. [15] indicated that (V, W) Cx layers are formed in the early stage of sintering process. However, Kawakami et al. [16] thought that (V, W) Cx layers are formed in the cooling stage rather than sintering stage. During the conventional liquid phase sintering, WC grain size increases dramatically when
cemented carbides hold for a few minutes at liquidus temperature [17-20]. Even with pressure assistance, conventional liquid phase sintering is not suitable for producing nanocrystalline microstructures. Therefore, it is feasible to sinter WC nanopowders containing GGI at under-eutectic temperature using very fast heating rate and short sintering time [11,21-24]. Spark plasma sintering is a rapid sintering technology which is beneficial for suppressing grain growth. Some studies showed that samples produced by spark plasma sintering have finer grain structures than the samples produced by hot isostatic pressing [12,25]. It is well known that grain size decreases with the addition of GGI, resulting in hardness of cemented carbides to increase generally. However, the changes in fracture toughness and TRS are uncertain. The effect of GGI on the comprehensive mechanical properties of cemented carbides has not been sufficiently understood yet. In this work, ultrafine WC-8Co-(VC/Cr3C2/TaC) cemented carbides were prepared by spark plasma sintering technology. The influence of GGI on the microstructure and mechanical properties of cemented carbides were studied in detail.
2. Experimental 2.1 Materials preparation WC powder with average grain size of 60 nm and Co powder with average grain size of 500 nm were used as raw materials (purity: 99.99%, ChaoWei Nanotechnology Co., Ltd, Shanghai). VC, Cr3C2 and TaC (purity: 99.99%, ChaoWei Nanotechnology Co., Ltd, Shanghai) were added as GGI. SEM images of as-received powders are shown in Fig. 1. As-received powders were mixed in glass container with pure ethanol. Then, the mixture was mixed by ultrasonic vibration and mechanical stirring for two hours in an ultrasonic platform equipped with stirrer (DSA200-JY1-9.0L, China). The powder mixture was dried in a vacuum
drying oven (DZF, China) at 80°C and sieved with a 100 mesh sieve. Then, the powder mixture was sintered in a spark plasma sintering furnace (LABOX-650F, Japan) to fabricate samples with dimensions of 15×15×5 mm, with the heating rate of 100 °C/min, sintering temperature of 1250, 1300 and 1350 °C, constant pressure of 70 MPa and holding time of 5 min. After that, samples were cooled to the room temperature in the furnace. Composition ratios and symbol of samples in present work are given in Table 1.
Fig. 1. SEM images of as-received raw powders, (a) WC, (b) Co, (c) VC, (d) Cr3C2 and (e) TaC. Table 1 Composition ratios and symbol of cemented carbides (wt.%). Cemented carbides
Symbol
WC
Co
VC
Cr3C2
TaC
WC-8Co
A1 (S1)
Bal.
8
0
0
0
WC-0.4VC-8Co
— (S2)
Bal.
8
0.4
0
0
WC-0.4Cr3C2-8Co
— (S3)
Bal.
8
0
0.4
0
WC-0.4TaC-8Co
— (S4)
Bal.
8
0
0
0.4
WC-0.8VC-8Co
A2 (S5)
Bal.
8
0.8
0
0
WC-0.8Cr3C2-8Co
A3 (S6)
Bal.
8
0
0.8
0
WC-0.8TaC-8Co
A4 (S7)
Bal.
8
0
0
0.8
WC-0.4VC-0.4Cr3C2-8Co
A5 (S8)
Bal.
8
0.4
0.4
0
WC-0.4VC-0.4TaC-8Co
A6 (S9)
Bal.
8
0.4
0
0.4
WC-0.4Cr3C2-0.4TaC -8Co
— (S10)
Bal.
8
0
0.4
0.4
2.2 Characterization Density of sintered samples was measured by Archimedes principle. TRS was measured on samples with dimensions of 15×5×3 mm using a three-point bending instrument (UTM5105-G, China). Span and loading velocity was 12 mm and 0.1 mm/min, respectively. Vickers hardness was measured using a hardness tester (MN-9631-130-C, China) with load of 30 kg and loading duration of 15 s. Fracture toughness was calculated based on the length of crack generated by indentation using the following equation [26]. KIC = 0.0028 · (HV·P/L)1/2
(1)
where KIC is fracture toughness (MPa·m1/2); HV is Vickers hardness (N/mm2), P is load (N); and L is total length of cracks (mm). At least five measurements were taken and average value was obtained. Phase in sintered samples was analyzed by X–ray diffraction (XRD, D8 Advance, Germany) with copper Kα radiation. A scanning electron microscope (SEM, Quant 250FEG, USA) was used for observing the microstructure on polished surfaces of sintered samples. Based on SEM images of polished surfaces, size distribution and average size of WC grains were calculated by measuring more than 400 grains using the image processing software (Image-Pro Plus 6.0). Element distribution in the microstructure was investigated using a SEM-energy dispersive spectrometer (SEM-EDS) analysis system. WC grains in sintered samples were extracted by placing samples in hydrochloric acid to remove Co phases, and then the WC grain morphology was observed by SEM.
3. Results and discussion 3.1 Microstructure and phase constitution analysis Fig. 2 shows SEM (BSE) images of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at 1250 °C. Bright areas represent WC phases and black areas are Co phases. It was found that Co phases were distributed between WC grains. There were a few Co pools and pores with irregular shapes and size of several hundred nanometers. Because Co segregation and insufficient wettability were typical characteristics of solid phase sintering, Co phases were not sufficiently diffused in spaces formed by WC skeleton [21,22], which led to formation of Co pools and pores. Lack of Co in the region near Co pool led to increase of WC grain contiguity [7]. In these regions, contact of WC grains promoted grain growth. This phenomenon was prominent in sample A1. As shown in Fig. 2b-f, although Co rich phases did not completely disappear with the addition of VC/Cr3C2/TaC, the microstructure became fine and grain shape transformed from faceted shape to blunt shape. In addition, from Fig. 2b-d, it was found that the addition of VC led to finer WC grains than addition of Cr3C2/TaC. Still Cr3C2 and TaC had important effect on grain refinement despite the existence of coarse grains.
Fig. 2. SEM (BSE) images showing microstructures of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5, (f) A6 and (g) EDS line analysis from point A to C in A1.
As the sintering temperature increased, high thermal energy was generated and dissolution-reprecipitation effect was promoted, which enhanced driving force for diffusion of Co phases [17]. Therefore, the increasing sintering temperature resulted in the densified microstructure, as shown in Fig. 3. It was apparent that WC grains in sample A1 sintered at 1350 °C almost transformed to faceted shape compared to that sintered at 1250 °C. Compared Fig. 2 with Fig. 3, limited but significant grain growth was clearly observed at 1350 °C regardless of chemical composition. However, obvious inhibition effects were still observed in samples with GGI as compared to sample A1, as shown in Fig. 3. Moreover, no abnormal grain growth was
observed in any of samples with the increase in temperature. Microstructures of samples sintered at 1300 °C were not shown because it was similar to that of samples sintered at 1350 °C.
Fig. 3. SEM (BSE) images showing microstructures of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1350 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5 and (f) A6.
Fig. 4 shows XRD patterns of composite powders and sintered samples. No obvious differences were observed in phases formed before and after sintering. There were only WC and Co phases in all samples. No graphite or η phase was observed in any of sintered samples. Diffraction peaks of VC, Cr3C2 and TaC were not detected due to the low content, which was also mentioned in many studies [5,23,27]. Moreover, as there are not different in phases formed at different sintering temperatures, XRD patterns of samples sintered at 1300 and 1350 °C are not given.
Fig. 4. XRD patterns of composite powders and sintered samples, (a) composite powders and (b) samples sintered at 1250 °C.
Fig. 5 shows size distribution and average size of WC grains in WC-8Co-(VC/Cr3C2/TaC) cemented carbides. It can be seen that size distribution became narrow and average grain size decreased with the addition of GGI. As shown in Fig. 5a, the grain size of sample A1 ranged from 100 to 900 nm and average grain size was 280 nm at sintering temperature of 1250 °C. By adding 0.8 wt.% VC/Cr3C2/TaC, the average grain size of samples A2, A3 and A4 reduced to 170, 230 and 250 nm, respectively. Compared to sample A1, the average grain size of sample A2 reduced by 40%, which was similar to results of some studies [5,12]. In these studies, 0.8 wt.% VC was added during the conventional liquid phase sintering. Therefore, inhibition effect of VC was not limited to liquid phase sintering stage. Even if at under-eutectic temperature, VC still had significant effect on suppressing grain growth. Moreover, as half of VC was replaced by Cr3C2/TaC, the average grain size of samples A5 and A6 increased compared to sample A2. Average grain size of sample A5 with 0.4 wt.% VC-0.4 wt.% Cr3C2 was 200 nm. Sample A6 with 0.4 wt.% VC-0.4 wt.% TaC had average grain size of 220 nm. This result showed that inhibition effect of VC was better than that of Cr3C2 and TaC.
Fig. 5. Cumulative size distribution and average size of WC grains in WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at (a) 1250 °C, (b) 1300 °C and (c) 1350 °C.
From Fig. 5b and c, it was found that proportion of coarse grains increased due to the grain growth with the increasing sintering temperature, resulting in the increasing average grain size. From 1250 to 1350 °C, the average grain size of samples increased by about 40-70%. At 1350 °C, the average grain size of sample A2 increased from 170 to 260 nm, but was still smaller than that of sample A1 sintered at 1250 °C. During the liquid phase sintering, the inhibition effect of GGI greatly depends on limiting dissolution-reprecipitation mechanism [12,18]. GGI can be dispersed
at WC surfaces to reduce surface energy of WC and solubility of WC in Co phases [16,28]. Meanwhile, GGI are dissolved in binder phase and further reduces solubility of WC [7,10]. During the spark plasma sintering, dissolution-reprecipitation effect are limited by low sintering temperature (1250 °C). However, VC still has significant effect on inhibiting grain growth, which is likely related to formation of V rich layers. Thin (V, W) Cx films are formed at WC/Co interfaces, which limits diffusion phenomenon involved in grain growth. 3.2 Effect of GGI on the WC grain morphology WC grain morphologies of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at 1250 °C are shown in Fig. 6. From Fig. 6a, it was found that WC grains of sample A1 were truncated triangular prisms, which is the equilibrium morphology of WC grains [29]. It is well known that WC has two (0 0 0 1) basal planes and three (1 01 0) prismatic planes. (0 0 0 1) and (1 01 0) planes were seen as triangular shapes and rectangular shapes, respectively. As shown in Fig. 6b, WC grains of sample A2 had triangular prismatic shapes with multi-steps based on basal planes. Multi-steps were composed of (0 0 0 1) and (1 01 0) planes parallel to basal and prismatic planes, respectively.
Fig. 6. SEM images showing WC grain morphologies of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) A1, (b) A2, (c) A3, (d) A4, (e) A5 and (f) A6.
The growth process of WC grains is summarized in Fig. 7. Fig. 7a shows growth process of WC grains in cemented carbides without GGI. Coalescence of WC particles occurs in the early sintering stage of nano-sized WC-Co powder [20]. After coalescence, WC grains exhibit round shapes and irregular shapes at low sintering temperature. As the temperature increases, diffusion mobility of atoms enhances as well as W and C diffuse through Co films at surfaces of round grains, resulting in faceted shapes of grains. Truncated triangular prisms are formed from faceting
of round grains driven by surface energy minimization [19].
Fig. 7. Schematic drawing showing growth mechanisms of WC grains, (a) WC-Co and (b) WC-Co-VC.
The existence of multi-steps indicates that grain growth is mainly controlled by 2D nucleation mechanism [30]. Generally, W and C atoms diffuse from the edges of (0 0 0 1) planes towards center and reach the center to form a nucleus. Then, the nucleus grows rapidly until it reaches edges. However, the addition of VC causes (V, W) Cx layers to form at (0 0 0 1) and (1 01 0) planes of WC grains [13-16]. (V, W) Cx layers enhance energy barrier of 2D nucleation in the edge region of (0 0 0 1) planes and hinder diffusion of W and C atoms. The new nucleus can appear at the center of (0 0 0 1) planes before growth completion of preceding nucleus, and then the existing nucleuses grow simultaneously on (0 0 0 1) planes [19]. Therefore, WC grains are stacked along [0 0 0 1] direction to form multi-steps based on (0 0 0 1) basal planes. Grain growth along [0 0 0 1] direction is effectively inhibited because (W, V) Cx layers are thicker on (0 0 0 1) planes than on (1 01 0) planes [3]. As Cr has high solubility in Co phases and WC dissolves in TaC to form (Ta, W) C solid solution, segregation layers with Cr and Ta are not formed at WC/Co interfaces. The multi-step
structure was not observed in samples A3 and A4, as shown in Fig. 6c and d. Compared to sample A1, grain morphologies of samples A3 and A4 tended to become round and blunt. From Fig. 6e and f, it was found that WC grains with multi-step structure in samples A5 and A6 reduced significantly when half of VC was replaced by Cr3C2/TaC. 3.3 Mechanical properties 3.3.1 Effect of sintering temperature on the mechanical properties Fig. 8 shows relative density and mechanical properties of WC-8Co-(VC/Cr3C2/TaC) cemented carbides sintered at different sintering temperatures. From Fig. 8a, it was found that almost all samples exhibited good relative density of above 98.5% at 1250 °C. Sample A1 had higher relative density of 99.5%. For samples A2, A3 and A4, the relative density was 99.4, 98.9, and 98.4%, respectively. With the addition of VC/Cr3C2/TaC, the relative density presented a slight variation. This variation was related to the different solubility of GGI in Co phases and the limited diffusion of Co phases at low sintering temperature [9,21,22]. As the sintering temperature increased, relative density of sample A1 was not improved but that of samples with GGI slightly increased. For samples with GGI, the GGI dispersing in binder phases reduced the solubility of WC in binder phases and the driving force of WC particles. As the temperature increased to 1300 and 1350 °C, the solubility of WC in binder phases increased. Meanwhile, the higher temperature provided a better driving force for the binder phases, causing the liquid phases to diffuse along the capillary formed by WC grains. Therefore, the dissolution-reprecipitation and rearrangement of carbides as well as the enhancing liquid phase fluidity caused the relative density to increase slightly.
100.00
(a)
(b)
2200
99.75
2150
99.50
2100
99.25
2050
99.00
2000
98.75
1950 A1 A4
98.50 98.25
1250
A2 A5
1300
A1 A4
1850 1250
1350
A1 A4
1300
1350
Temperature (°C)
12.0 11.5
A3 A6
1900
A3 A6
Temperature (°C)
(c)
A2 A5
A2 A5
A3 A6
(d)
2000
A1 A4
A2 A5
A3 A6
1900
11.0 1800 10.5 1700 10.0 1600 9.5 1500 9.0 1400 8.5
1250
1300
Temperature (°C)
1350
1250
1300
1350
Temperature (°C)
Fig. 8. (a) relative density, (b) hardness, (c) fracture toughness and (d) transverse rupture strength of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at different temperatures.
From Fig. 8b and c, it was clear that all samples sintered at 1250 °C achieved maximum hardness and fracture toughness. From 1250 to 1350 °C, the hardness presented a slow downward trend (reduced by 5-11%), while the fracture toughness exhibited few changes. Densification and grain size of cemented carbides have significant effect on their hardness [32]. Generally, the hardness increases as the relative density increases while decreases as the WC grain size increases (Hall-Petch relation). With the increasing temperature, the change in relative density was slight (Fig. 8a), while the WC grain size increased significantly (Fig. 5). As a result, the grain growth caused by the increasing temperature is the primary reason of the decrease in hardness. From Fig. 8d, it can be seen that the TRS of most samples reduced gradually with the increasing temperature, but that of sample A4 presented a deviation behavior. At 1250 °C, most
samples reached good densification which effectively decreased strain points and improved the TRS. From 1250 to 1350 °C, the fracture paths along WC/Co and WC/WC interfaces decreased due to the grain coarsening, resulting in the reduction of fracture energy. As a result, the TRS decreased gradually. However, an acceptable densification that was beneficial for the TRS was not achieved for sample A4 sintered at 1250 °C. The increasing temperature was effective on eliminating pores, causing the TRS of sample A4 to increase. 3.3.2 Effect of GGI content on the mechanical properties From analysis in Sec. 3.3.1, it is clear that good mechanical properties are reached at sintering temperature of 1250 °C. Therefore, the influence of GGI content on the mechanical properties of cemented carbides sintered at 1250 °C was further studied. Fig. 9 shows hardness and fracture toughness of samples sintered at 1250 °C. From Fig. 9a, it was found that the hardness increased gradually (1970→2080→2180 kg/mm2) with the increasing VC content (S1→S2→S5: 0→0.4→0.8 wt.%). By adding 0.4 wt.% Cr3C2/TaC, the hardness of samples S3 and S4 increased to 2020 and 2030 kg/mm2, respectively. However, as Cr3C2/TaC content increased to 0.8 wt.%, the hardness of samples S6 and S7 were not further improved. Moreover, as half of VC was replaced by Cr3C2/TaC, the hardness of samples S8 and S9 tended to decrease compared to sample S5. Because sample with VC had smaller grain size than that with Cr3C2/TaC, the sample with VC had higher hardness. It is well known that there is an antagonistic correlation between fracture toughness and hardness and improvement of hardness is generally accompanied by decrease of fracture toughness [8]. As shown in Fig. 9b, the fracture toughness decreased in various degrees with the addition of GGI. Sample S1 had higher fracture toughness (11.5 MPa·m1/2) compared to other
samples. However, sample S5 achieved minimum fracture toughness of 9.4 MPa·m1/2. Compared to sample S5, the fracture toughness of sample S8 and S9 increased because Cr3C2/TaC replaced half of VC. WC grain size has an important influence on fracture toughness of WC-Co when Co content keeps constant. When cracks propagate to WC grains with sub-micron and micron sizes, crack can pass directly through WC grains, which leads to the formation of transgranular fracture. As the intragranular strength is higher than intergranular strength, the transgranular fracture prevents effectively crack propagation, which is conducive to fracture toughness. However, with the decreasing grain size, cracks can propagate easily along grain boundaries and the transgranular fracture occurs rarely. In addition, the crack propagation morphologies can explain well for changes in fracture toughness, as shown in Fig. 10. From sample S1, it was clear that crack paths changed significantly during the crack propagation. The existence of micron-sized grains had a deflection effect on crack propagation [32]. Moreover, crack bridging increased fracture surface areas and enhanced energy consumption. As a result, crack deflection and crack bridging enhanced the resistance against crack propagation and improved fracture toughness. As shown in Fig. 10b-d, with the decrease of grain size, cracks extended mainly along grain boundary, which limited the contribution of crack deflection to fracture toughness. Therefore, in this study, sample without GGI had the higher fracture toughness due to its relatively larger grain size, while the fracture toughness decreased due to the decreasing grain size with the addition of GGI.
Fracture toughness (MPa·m½)
Vickers hardness (kg/mm2)
Fig. 9. (a) hardness and (b) fracture toughness of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co.
Fig. 10. SEM (BSE) images showing crack propagation morphologies of cemented carbides sintered at 1250 °C, (a) S1: WC-8Co, (b) S2: WC-0.4VC-8Co, (c) S5: WC-0.8VC-8Co and (d) S8: WC-0.4VC-0.4Cr3C2-8Co.
The influence of GGI on the TRS of samples is shown in Fig. 11. It can be seen that the TRS first increased and then decreased (1670→1810→1690 MPa) with the increasing VC content (S1 →S2→S5: 0→0.4→0.8 wt.%). As the WC grain size decreases, the volume fraction of total boundary areas between WC and Co increases. The increase of fracture path through WC/Co interfaces can bring about numerous fracture energies and enhance the TRS of cemented carbides
[33]. Therefore, if there are no pores, inclusions and abnormally large grains in the microstructure, the TRS is usually in inverse ratio to average grain size. However, compared to sample S1, the TRS of sample S5 did not increase significantly despite significant grain refinement. This phenomenon is likely related to VC additions. Generally, the addition of GGI can reduce WC/Co interface coherency, and low fraction of WC/Co interfacial coherency is detrimental to the TRS [4,34]. On the one hand, samples with VC has high stability of (V, W) Cx layers on WC surfaces and low fraction of WC/Co interface coherency [3], which decreases significantly WC/Co interfacial energy and results in the low separation energy. On the other hand, multi-step structure of WC grains causes stress concentration and increases fracture sensitivity [35], which makes it easy for cemented carbides to fracture along WC/Co interfaces. Therefore, the grain refinement, low fraction of WC/Co interface coherency and fracture sensitivity affected comprehensively the TRS of samples with VC. These factors worked together to cause the TRS to first increase and then decrease with the increasing VC content. The TRS of samples S8 and S9 increased by 17 and 5% compared to sample S5, respectively. As half of VC was replaced by Cr3C2/TaC, the V segregation and multi-step structure decreased, which was beneficial to improve the TRS. In addition, the TRS first decreased and then increased (1670→1540→1830 MPa) with the increasing Cr3C2 content (S1→S3→S6: 0→0.4→0.8 wt.%). Cr3C2 has weakest effect on reducing WC/Co interface coherency compared to other GGI [3]. As a result, the TRS of sample S6 increased by 13% compared to sample S1, which was a comprehensive outcome of grain refinement and large fraction of WC/Co interface coherency.
Transverse rupture strength (MPa) Fig. 11. Transverse rupture strength of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co.
Fig. 12 shows fracture surfaces of samples sintered at 1250 °C. For sample S1, transgranular fracture of coarsened WC grains and intergranular fracture of carbides were observed clearly, as shown in Fig. 12a. With the decrease of grain size due to the GGI additions, fracture surfaces presented that intergranular fracture became predominant while transgranular fracture decreased, as shown in Fig. 12b-f. In addition, some pores were observed on the fracture surface of sample S7, as shown in Fig. 12d. Pores are main fracture sources of cemented carbides which can cause stress concentration and produce cracks during deformation [31]. Therefore, the TRS of sample S7 was 1440 MPa and decreased by 17% compared to sample S1.
Fig. 12. SEM images showing fracture surfaces of WC-8Co-(VC/Cr2C3/TaC) cemented carbides sintered at 1250 °C, (a) S1: WC-8Co, (b) S5: WC-0.8VC-8Co, (c) S6: WC-0.8Cr3C2-8Co, (d) S7: WC-0.8TaC-8Co, (e) S8: WC-0.4VC-0.4Cr3C2-8Co and (f) S9: WC-0.4VC-0.4TaC-8Co.
Fig. 13 shows the relationship between mechanical properties of samples. It was found that distribution between hardness and fracture toughness exhibited an inverse relationship. Samples with mixed GGI were in the intermediary position (S8, S9 and S10). High hardness were achieved by using VC or mixed GGI with VC (S2, S5, S8 and S9). The addition of TaC neither provided high fracture toughness nor significantly improved hardness (S4 and S7).
Comprehensive mechanical properties were achieved by adding 0.4 wt.% VC-0.4 wt.% Cr3C2, with the hardness of 2110 kg/mm2, fracture toughness of 10.4 MPa·m1/2 and TRS of 1990 MPa. Properties of some WC-Co cemented carbides prepared by different GGI additives and sintering technologies are given in Table 2. It was clear that the cemented carbide prepared by spark plasma sintering in this work has a good comprehensive hardness and fracture toughness compared to other research results.
Fig. 13. Relationship between hardness, fracture toughness and transverse rupture strength of samples sintered at 1250 °C. S1: WC-8Co; S2: WC-0.4VC-8Co; S3: WC-0.4Cr3C2-8Co; S4: WC-0.4TaC-8Co; S5: WC-0.8VC-8Co; S6: WC-0.8Cr3C2-8Co; S7: WC-0.8TaC-8Co; S8: WC-0.4VC-0.4Cr3C2-8Co; S9: WC-0.4VC-0.4TaC-8Co; S10: WC-0.4Cr3C2-0.4TaC-8Co. Table 2 Properties and sintering technologies of some WC-Co cemented carbides. GGI additives (wt.%)
Co content (wt.%)
Grain size (nm)
Hardness (kg/mm2)
KIC (MPa·m1/2)
Sintering technology
0.9NbC [36]
12
320
1631
10.8
Vacuum sintering
0.5VC [6]
12
550
1490
10.9
Spark plasma sintering
1Cr3C2 [12]
12
214
1870
10.5
Hot isostatic pressing
0.26VC-0.45Cr3C2 [37]
6
210
2041
9.0
Liquid phase sintering
0.5VC-0.5Cr3C2 [21]
12
190
1925
10.2
Spark plasma sintering
0.6VC-0.3Cr3C2 [24]
10
–
2070
8.2
Hot isostatic pressing
4. Conclusions (1) Average WC grain size of cemented carbides decreased from 280 to 170, 230 and 250 nm with the addition of 0.8 wt.% VC/Cr3C2/TaC. The effect of VC on inhibiting grain growth was better than that of Cr3C2 and TaC. With the addition of VC, WC grain morphologies transformed from equilibrium to triangular prismatic shapes with multi-step structure because (V, W) Cx layers were formed at WC surfaces. However, grain morphologies became round and blunt with the addition of Cr3C2/TaC. (2) As sintering temperature increased from 1250 to 1350 °C, the hardness and TRS decreased gradually due to grain growth, while fracture toughness had few changes. At 1250 °C, with the addition of VC/Cr3C2/TaC, the smaller the grain size of sintered samples, the higher the hardness of sintered samples, whereas fracture toughness is in an inverse ratio to hardness. Grain size and WC/Co interfacial energy comprehensively affected the TRS of cemented carbides. (3) Excellent comprehensive mechanical properties can be achieved by using mixed GGI to balance and optimize the relationship between grain refinement and WC/Co interfacial energy.
Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 51775280 and No. 51675285) and the Jiangsu Provincial Six Talent Peaks Project (2016-HKHT-019).
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Highlights: • The effect of VC on inhibiting grain growth is better than that of Cr3C2 and TaC. • WC grains of alloys with VC are triangular prisms with multi-step structure. • VC/Cr3C2/TaC causes hardness to increase and fracture toughness to decrease. • Grain size and interfacial energy comprehensively affect transverse rupture strength.