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(Ti,W)(CN)-based cermets
Journal of Alloys and Compounds 657 (2016) 671e677
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Carbide/binder interfaces in Ti(CN)e(Ti,W)C/(Ti,W)(CN)-based cermets Choongkwon Park, Sangwoo Nam, Shinhoo Kang* Department of Materials Science and Engineering, Seoul National University, Gwanak-gu, Seoul, 151-742, South Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 24 August 2015 Received in revised form 12 October 2015 Accepted 13 October 2015 Available online xxx
Partially solutionized carbide cermets (PSCs) were prepared for cutting tool applications by replacing a portion of Ti(CN) with (Ti0.88W0.12)C or (Ti0.88W0.12)(C0.7N0.3) in a conventional Ti(CN)-based cermet. The PSC containing (Ti0.88W0.12)C exhibited high fracture toughness with no loss of hardness. The elimination of Ti(C0.7N0.3) cores in the microstructure reduced the strain at the core/rim interfaces, leading to the increase in fracture toughness. The PSC containing (Ti0.88W0.12)C showed coherent relationships at the core/rim and rim/binder interfaces, contributing to the improved mechanical properties. In contrast, nitrogen in (Ti0.88W0.12)(C0.7N0.3) resulted in a different microstructure and properties. The major factors determining the mechanical properties of the cermets are discussed in terms of the carbide/binder interfaces and the thermal stability of the added carbides. © 2015 Elsevier B.V. All rights reserved.
Keywords: Ti(CN) cermet Mechanical properties Microstructure Interface HRTEM
1. Introduction Ti(CN)-based cermets have received much attention as promising alternatives to conventional WCeCo cemented carbide. They exhibit excellent wear-resistance, good hardness, a low friction coefficient, good deformation resistance, good chemical stability at high temperatures, and natural abundance [1e4]. However, the relatively low fracture toughness of Ti(CN)-based cermets limits their wide use as commercial cutting tools. Secondary carbides, such as WC, Mo2C, NbC, and TaC, have been used to improve mechanical properties [3e12]. Adding secondary carbides to conventional Ti(CN)-based cermets (CTCs) produces a core/rim structure through the dissolutioneprecipitation of raw carbides during liquid-phase sintering [7e10]. Thus, CTCs contain a variety of interfaces in their microstructures, including core/inner rim, inner rim/outer rim, and outer rim/binder interfaces, large fractions of which can impair the mechanical properties of the materials [3,9,13]. Core/rim interfaces generally have a high lattice strain due to the large difference in their lattice parameters [9]. Moreover, the mechanical properties of metal-ceramic composites are highly dependent on the strength of ceramic/metal interfaces
* Corresponding author. E-mail address: [email protected] (S. Kang). http://dx.doi.org/10.1016/j.jallcom.2015.10.121 0925-8388/© 2015 Elsevier B.V. All rights reserved.
[14e16]. TiC-based complete solid-solution carbide cermets (CSCs) that consist of a single carbide phase and a binder phase to eliminate the core/rim interfaces have been reported [12,17e19]. They exhibited improved fracture toughness, but suffered from low hardness. In our previous reports, partially-solutionized carbide cermets (PSCs) were firstly introduced, which were fabricated by blending CTC with (Ti1-xWx)C complete solid-solution phase, and showed both improved hardness and toughness [20]. According to the results, PSCs exhibited peak hardness and toughness values with the addition of 50 wt.% (Ti0.93W0.07)C solid solution although the PSC showed the largest particle size and the smallest core fraction among PSCs. The mechanism for the improvements has been remained unclear, and it is necessary for these cermet systems to understand it for the field of the cermet research. In this study, the effect of (Ti0.88W0.12)(C1yNy) addition on the microstructure and mechanical properties of Ti(CN)-based cermets is investigated in detail in terms of the core/rim/binder interfaces and the thermal stability of the added carbides. To the best of our knowledge, we firstly suggest the effect of the interfacial characteristics in an atomic scale on mechanical properties in Ti(CN)based cermet systems. For this investigation, compositions similar to those of commercial cermets, i.e., Ti(C0.7N0.3)e WCeMo2CeNbCeNi/Co, were chosen and two PSCs were prepared by adding (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solutions to
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contrast microstructures and mechanical properties. 2. Experimental procedure (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solution powders were synthesized via high-energy ball milling and subsequent carbothermal reduction, as described in our previous reports [17,18]. TiO2 (particle size: 20 mm; purity: 99%; SigmaeAldrich, St. Louis, MO, USA), WO3 (particle size: 20 mm; purity: 99%; SigmaeAldrich), and graphite (particle size: 1.65 mm; purity: 99%, SeungLim Carbon Metal, Ansan, Korea) were used as raw materials. The amount of each raw material was determined by the following reaction equations.
ð1 xÞTiO2 þ xWO3 þ ð3 þ xÞC ¼ ðTi1x Wx ÞC þ ð2 þ xÞCOðgasÞ (1) N2
ðgasÞ
ð1 xÞTiO2 þ xWO3 þ ð3 þ x yÞC!ðTi1x Wx Þ C1y Ny
þ ð2 þ xÞCOðgasÞ (2) The raw materials were milled with a planetary high-energy ball mill (Pulverisette 7, Fritsch, Idar-Oberstein, Germany) at a speed of 250 rpm for 20 h. A WCeCo-coated stainless steel jar and WCeCo balls were used for milling with a ball-to-powder weight ratio (BPR) of 40:1. The milled powder mixtures were carbothermally reduced at 1300 C for 2 h in a graphite vacuum furnace in a vacuum (103 torr) for (Ti0.88W0.12)C and under 10 torr of nitrogen flow for (Ti0.88W0.12)(C0.7N0.3). The heating rate was 10 C/min. The raw materials used to prepare the PSCs are listed in Table 1. They were conventionally mixed to reach the target compositions listed in Table 2. Ti(C0.7N0.3) and (Ti0.88W0.12)C/ (Ti0.88W0.12)(C0.7N0.3) solid solution powders in fixed quantities were the major components, and WC, Mo2C, and NbC were added as secondary carbides along with Ni and Co as binder metals. (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solution powders were used to produce PSC1 and PSC2, respectively. Overall compositions were the same except for the total nitrogen content (Table 2). The mixtures were then milled in ethanol using WCeCo balls in plastic bottles. The BPR was 10:1, and rotation was at 200 rpm for 20 h. Paraffin (4 wt %) was added for ease of compaction. The mixed slurries were dried and then sieved (125 mesh) to remove agglomerates. The powder mixtures were pressed into discs under a pressure of 150 MPa. The green discs were conventional vacuum sintered in a graphite vacuum furnace (~103 torr) at 1510 C for 1 h. For microstructural observations, as-sintered samples were mechanically ground and then polished using diamond slurry (1e6 mm). The porosity levels of samples were measured according
to the ASTM (American Society for Testing and Materials) specifications to provide the reliability for the discussions below. Powder X-ray diffraction (XRD) patterns were obtained with a diffractometer (D8-Advance, Bruker, Billerica, MA, USA) with monochromatized Cu-Ka radiation (l ¼ 1.54056 Å). The lattice parameters of the specimens were calculated by Rietveld refinement of the XRD results. Scanning electron microscopy (SEM) observations were performed by field emission SEM (FE-SEM; MERLIN Compact, ZEISS, Oberkochen, Germany). Conventional transmission electron microscopy (TEM), highresolution TEM (HRTEM), and scanning TEM high-angle annular dark-field (STEM-HAADF) observations were carried out using a transmission electron microscope (F20 FEI Tecnai, FEI, Hillsboro, OR, USA) operated at 200 KV. Energy dispersive X-ray spectroscopy (EDS) analyses were also performed using an EDX RTEM X-ray detector combined with a STEM-HAADF detector (F20 FEI Tecnai). The carbide sizes and the binder fractions of the sintered specimens were estimated using isolution DT 10.3 software equipped in optical microscopy (Image & Microscope Technology, Daejon, Korea) based on the back-scattered electron (BSE) SEM images. More than 300 carbide particles were measured to obtain statistically meaningful results. Hardness was measured using a Vicker's hardness indentation tester with a 30 kg load and a 15 s loading time. Fracture toughness was calculated by measuring the crack length generated by the Vickers indentation using the expression derived by Shetty et al. [21]. 3. Results and discussion 3.1. Powders and mechanical properties of PSC cermets Fig. 1 shows FESEM-SE images and XRD patterns for assynthesized (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) powders. Both the (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) powders have an irregular shape, and the former is larger than the latter. The XRD results in Fig. 1(c) show that both powders have a NaCl-structure and consist of a single-phase. The XRD peaks of (Ti0.88W0.12)C were at smaller angles than those of (Ti0.88W0.12)(C0.7N0.3) and Ti(C0.7N0.3). The average particle sizes estimated with specific surface areas obtained by BET measurements were ~400 nm for (Ti0.88W0.12)C and ~40 nm for (Ti0.88W0.12)(C0.7N0.3), which are in good agreement with the FESEM images. The mechanical properties, amount of binder phase, and porosity levels of PSC1 and PSC2 cermets are summarized in Table 3, which is recently reported in our report [22]. PSC1 (Hv, 14.0 ± 0.1 GPa; K1C, 12.2 ± 0.2 MPa m1/2) is much tougher, with no deterioration in hardness, than PSC2 (Hv, 13.7 ± 0.2 GPa; K1C, 9.8 ± 0.2 MPa m1/2) and commercial. The porosity levels of PSC1 and PSC2 were excellent as A1B1. Comparing mechanical properties of PSC1 and PSC2, the higher fracture toughness of PSC1 is found not a function of binder fraction. It is because PSC2 exhibited much
Table 1 Powder specifications. Constituents
BET SSA(m2/g)
Size (mm)
Purity (wt %)
Manufacturer
Ti(C0.7N0.3) WC Mo2C NbC Ni Co (Ti0.88W0.12)C (Ti0.88W0.12)(C0.7N0.3)
1.85 e e e e e 2.72 26.51
1.50 2.73 1.85 1.45 2.40 2.30 (0.354b) (0.038b)
99.7 99.7 98.8 99.8 98.8 99.9 e e
Treibacher Industrie AG TaeguTec PPM Ltd.a PPM Ltd. PPM Ltd. Cobatech Corp. Synthesized Synthesized
a b
Pacific Particulate Materials Ltd. Particle size calculated by BET-SSA value.
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Table 2 Compositions of two cermets (wt.%). Specimens
(Ti0.88W0.12)C
(Ti0.88W0.12) (C0.7N0.3)
Ti(C0.7N0.3)
WC
Mo2C
NbC
Ni/Co
PSC1 PSC2
32 e
e 32
27.2 27.2
5.6 5.6
9.6 9.6
5.6 5.6
20 (Ni/Co ¼ 1) 20 (Ni/Co ¼ 1)
Fig. 2. XRD patterns for as-sintered cermets for (a) PSC1 and (b) PSC2. Lattice parameters of the rim and binder phases are shown [22].
higher binder fraction than PSC1 (Table 3), indicating that other factors control the toughness. Also, the mechanical properties of PSCs strongly depend on the composition of the added solidsolution carbides, especially the N content. To investigate the effect of the carbide addition on the mechanical properties of the cermets, we examined the microstructure of the PSC1 and PSC2 in detail. 3.2. Core/rim microstructures of PSC cermets
Fig. 1. XRD patterns and FESEM-BSE images for as-synthesized solid solution powders of (a) (Ti0.88W0.12)C and (b) (Ti0.88W0.12)(C0.7N0.3). (c) XRD patterns for (Ti0.88W0.12)C (red line) and (Ti0.88W0.12)(C0.7N0.3) powders (blue line) (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
Fig. 2 is the XRD profiles of the PSC1 and PSC2 along with reference data for Ti(C0.7N0.3), Ni, and Co from the JCPDs cards [22]. PSC1 consists of a single carbide phase with no other phases (Fig. 2(a)). In contrast, the peaks for the carbonitride phases in PSC2 are separated into two groups. The low-intensity peaks match those for undissolved Ti(C0.7N0.3), while the high-intensity peaks, presumably corresponding to the (Ti,W,Mo,Nb)(CN) solid-solution carbonitride rim phase, formed through dissolutioneprecipitation of the raw materials, which are located at lower angles (Fig. 2(b)). In both cermets, the peaks for the solid-solution carbonitride and binder phases are shifted to lower angles compared with the reference data for Ti(C0.7N0.3) and Ni/Co. The lattice parameters of the solid-solution carbonitride and binder phases calculated by Rietveld refinement are shown next to each XRD profile in Fig. 2. The greater peak shifts for the carbonitride phase in PSC1 than PSC2 is due to different dissolution rates of Ti(C0.7N0.3), (Ti0.88W0.12)C, and (Ti0.88W0.12)(C0.7N0.3) as explained in elsewhere [9,10,22,23]. However, the peak shifts for the binder phases are larger for PSC2
Table 3 Black core fractions, black core sizes, carbonitride sizes, binder fractions, mechanical properties, and porosities for PSC1, PSC2 and commercial CTCs [22]. Specimens
Black core fraction (%) Black core size (mm) Carbonitride size (mm) Binder fraction (%) Hardness (GPa) Fracture toughness (MPa$m1/2) Porosity
PSC1 2.5 PSC2 18.2 Commercial CTCs 16.0
0.63 0.67 0.62
1.19 0.82 0.85
16.8 21.9 20.5
14.0 ± 0.1 13.7 ± 0.2 13.8e14.1
12.2 ± 0.2 9.8 ± 0.2 8.8e9.2
A1B1 A1B1 A1B1
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Table 4 Binder compositions analysed by STEM-HAADF/EDS for PSC1 and PSC2. Specimen
PSC1
Element
Content (wt %)
Uncert. (%)
Content (wt %)
PSC2 Uncert. (%)
Co Ni Ti W Mo Nb C N
44.40 41.98 1.21 2.25 9.12 0.08 0.41 0.51
0.22 0.21 0.03 0.07 0.16 0.05 0.04 0.04
33.01 29.51 1.38 16.49 16.52 0.47 1.03 1.55
0.27 0.26 0.05 0.30 0.32 0.05 0.20 0.10
than for PSC1 due to the low chemical affinities of W and Mo with N in the PSC2, which we have reported elsewhere [12,22,24,25]. Table 4 shows the STEM-HAADF/EDS results, confirming that the binder phase of PSC2 has higher W, Mo, and Nb contents than that of PSC1. These increase the volume fraction of the binder phase, but the large fraction of binder phase did not increase the toughness of PSC2 (Table 3). The FESEM-BSE images for PSC1 and PSC2 are shown in Fig. 3(a) and (b), respectively. The cores in PSC1 appear as grey contrast with inner and outer rims; they were confirmed by STEM-HAADF/EDS as
Fig. 3. FESEM-BSE images for as-sintered (a) PSC1 and (b) PSC2. White, black and red arrows indicate inter-granular, trans-granular and inter-core/rim failure modes, respectively (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
raw (Ti,W)C. Some black Ti(CN) cores are also visible because of the similar phase stabilities of Ti(C0.7N0.3) and (Ti0.88W0.12)(C0.7N0.3), which is in good agreement with the XRD results in Fig. 2 [22]. In addition, many Ti(C0.7N0.3) cores in PSC2 were completely or partially denuded from the rim phase with some white core/grey rims as indicated in Fig. 3(b) [10,26]. Fig. 3 also shows the crack paths induced by indentation method in PSC1 and PSC2. Three kinds of crack propagation mode were observed, i.e. inter-granular, trans-granular and inter-core/rim failure modes, which were marked by white, black and red arrows, respectively. The cracks in PSC1 mainly propagated in transgranular failure mode and the cracks traversed even small particles located at the direction of crack propagation (black arrows in Fig. 3 (a)). In contrast, inter-granular and inter-core/rim failure modes are dominant in PSC2 (red and white arrows in Fig. 3(b)). It is true that larger carbide size (>30%) and low fraction of hard phase in PSC1 would facilitate trans-granular failure when crack propagates. However, the crack paths in PSC1 also indicate that the rim/binder interfaces in PSC1 could be much stronger than those of PSC2. Therefore, the interfacial characters between core/rim and rim/binder were characterized to understand the factors influencing the higher toughness of PSC1. 3.3. Rim/binder interfaces of PSC cermets Fig. 4 shows bright-field TEM (BF-TEM) images of PSC1, which contains various interface regions of interest observed by HRTEM. More than 20 interfaces of >10 particles were observed from each specimen. Some representative HRTEM images are shown in Fig. 4(AeD) of the regions that are enclosed by white squares in the BF-TEM images in Fig. 4(a) and (b). Each inset in Fig. 4(AeD) shows the corresponding FFT patterns. The lattice planes on the HRTEM images were indexed by comparing the interplanar distances (d) measured from the HRTEM images with the values calculated from the XRD results. The interplanar distances, d{111}C, d{200}C, and d {111}B were 2.4917, 2.1578, and 2.0653 Å, respectively, where subscripts C and B indicate carbonitride and binder, respectively. The carbonitride particles in PSC1 consisted of faceted surfaces and rounded corners. Fig. 4(AeC) shows the facets of the carbonitride particles terminating parallel to the {111}C planes, and the interfaces between the carbonitride and binder appear coherent or semi-coherent with ~30% probability of overall interfaces selected randomly. For NaCl-structured carbides, such as TiC, {111} surfaces are less stable than {100} surfaces as vacuum surfaces because of their low packing density [27] and dipolar nature [28]. However, it is often reported that NaCl-structured TiC and Ti-based carbonitride particles in a metal matrix terminate commonly with {111} surfaces [16,29,30]. Ti-terminated {111} surfaces are known to have a high local density of states around the Fermi level (EF) [31,32], leading to the high electrical conductivity and metallic properties. Mitra et al. reported that the {111} surfaces of TiC precipitates are stabilized by electronic interactions with Al atoms in an Al metal matrix, resulting in strong metallic bonding [16]. A similar explanation can be applied to our PSC cermets. Fig. 4(A) shows HRTEM images of the interface between the carbonitride facet and binder boundary, where the interface is characterized as (11-1)C//(200)B. In the image, the carbonitride particle and binder are oriented with its [011] zone axis (see FFT pattern). The interface is asymmetric, but the two phases at the interface are found to have the interfacial relationship of (11-1)C// (100)B. Thus, the d-spacings for (200)C and (-11-1)B planes shown in Fig. 5(a) determines the level of mismatch. Based on this, the lattice mismatch between the (200)C and (-11-1)B planes is calculated as 4.4%, where semi-coherent interfaces tend to form, involving dislocations. However, the two phases are found perfectly matched
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Fig. 4. TEM and HRTEM images for PSC1 and FFT pattern: (a) and (b) show low magnification BF-TEM images for the typical rounded carbonitride grains. AeD show the HRTEM images for the interface areas enclosed by the white squares in (a) and (b). Insets are the corresponding FFT patterns.
Fig. 5. Schematic illustration of the orientation relationship. (a) Coherent relationship of (11-1)C//(200)B. (b) Stacking faults by Shockley partial dislocation of the ɑo/6〈112〉 Burgers vector.
and coherent in long-range, since the d-spacings for the (200)C planes and the (-11-1)B planes around the interface are measured the same (2.0668 Å). This indicates that at the interface, the lattices of the carbonitride are distorted to adjust the lattice parameters. In addition, stacking faults, which exist three atomic layers away from the rim/binder interface, are observed in the carbonitride. The faults are highlighted by yellow dotted lines (see the characteristic FFT pattern). The carbonitride can accommodate the local distortion at the interface up to the third atomic layer (part of the layer is outlined by a red dotted box in Fig. 4(A)), but beyond this layer, the stacking fault generation is energetically favourable. For ease of visualization, the interface and the stacking faults are shown in Fig. 5(a) and (b), respectively. Each stacking fault is generated by a Shockley partial dislocation with a ɑo/6〈112〉 Burgers vector, and
the total Burgers vector is equal to b ¼ ɑo/2〈112〉. The presence of three stacking faults completely restores the original stacking sequence of the carbonitride lattice and releases the lattice strain. Therefore, the (11-1)C//(100)B interface becomes coherent. Fig. 4(B) shows representative carbonitride/binder interfaces at the round corner of a carbonitride particle. Steps at the interface were observed that consisted of two {111} faces, and are indicated by black dotted lines in the figure. The interfacial relationships at the steps were also similar to those at the faceted interfaces; thus, the interfaces were coherent/semi-coherent. However, the (11-1)C interface shown in Fig. 4(C) is parallel to the high-index planes of the binder phase, and the binder lattice was distorted around the interface to maintain the lattice matching. The orientation angle of the (-11-1)B planes to the
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interface [(11-1)C] was q ¼ 64 . In this orientation, the misfit between the (200)C and (-11-1)B planes was calculated as 15%, which is the highest misfit we have observed. Around the interface, the (-11-1)B planes were slightly bent to adjust the lattice mismatch. In addition, edge dislocations in the binder lattice were observed at the interface, where a typical dislocation is indicated by a white arrow. The arrays of the binder and carbonitride lattice planes are locally changed around the defects. Therefore, the lattice parameters are adjusted at the interface by the distortions, resulting in a semi-coherent interface. Similar observations have been reported for several ceramic/metal composites such as TiC/Al [16] and WC/ Co [33]. Fig. 4(D) shows the HRTEM image of the representative (Ti,W)C grey core/rim interface. The lattice orientations between the two phases were similar (d(200)grey core ¼ 2.145 Å and d(200)rim ¼ 2.156 Å), and the interface was highly coherent (d ¼ 0.5%). These results are in good agreement with the XRD results in Fig. 2(a), indicating that the grey core/rim carbonitride particles are free from lattice strain. Fig. 6 shows BF-TEM images for typical faceted carbonitride particles in the PSC2 binder. HRTEM images of the typical smooth interface regions (enclosed by white squares in Fig. 6(a) and (b)) are presented in Fig. 6(A) and (B), respectively, where the insets are the corresponding FFT patterns. The internal angles of the carbonitride particles in PSC2 were mostly measured as 70.53 and 109.47, which are the angles observed when 3-D shape of particles is an octahedron terminated by {111} facets [30], and the edges of the carbonitride particles in PSC2 were highly sharp in contrast to those in PSC1 (Fig. 6 (a) and (b)). HRTEM images in Fig. 6(A) and (B) prove
that the facets of the carbonitride particles in PSC2 are terminated by {111} planes. HRTEM images also show that the carbonitride particles were oriented with the electron beam parallel to [011] zone axis, but the binder phases were tilted far away from the [011] zone axis. It indicates that the carbonitride and binder phases in PSC2 were just randomly oriented at their interfaces and PSC2 did not show any preferred orientation relationship between them similar to that shown by PSC1. Therefore, the difference in the carbonitride/binder interfacial relationships between PSC1 and PSC2 can generate the great difference in their mechanical properties, especially toughness. It is well known that the strength and toughness of the ceramic metal composites are strongly dependent on the interfacial characteristics between ceramic and metal matrix [33e35]. The coherency state at the core/rim/binder interfaces in PSC1 can enhance their bond strength, leading to its high toughness. The formation of coherent interfacial relationships in PSC1 depends on the lattice parameters of the core, rim and binder phases, as well as the electronic interactions between the carbonitride and binder phases. 4. Conclusions We prepared PSC cermets by adding (Ti0.88W0.12)C or (Ti0.88W0.12)(C0.7N0.3) solid solution to the composition of a conventional Ti(CN)-based cermet, and then compared their mechanical properties. The elimination or minimization of highly strained Ti(CN) core/rim interfaces through adding (Ti0.88W0.12)C
Fig. 6. TEM and HRTEM images for PSC2 and corresponding FFT patterns: (a) and (b) are low magnification images showing the typical faceted carbonitride grains in the binder phase. A and B show the HRTEM images for the regions enclosed by the white squares in (a) and (b). Insets show the corresponding FFT patterns.
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solid solution significantly toughened the PSCs. The solid-solution carbonitride in PSCs commonly exhibited {111} plane faceting at the interface with the binder phase. Coherent interfaces with a (111)C//(200)B relationship or semi-coherent interfaces were often observed at the interfaces in PSC1, whereas incoherent interfaces were observed in PSC2. The high fraction of the binder phase in PSC2 did not result in an obvious improvement in the toughness. Thus, we conclude that coherent or semi-coherent interfaces in PSC1 are one of the important factors which contribute to the high fracture toughness of PSC1 more than the binder fraction in the system. Furthermore, we reconfirmed that the thermal stability of the Ti(CN), (Ti0.88W0.12)C, and (Ti0.88W0.12)(C0.7N0.3) phases and the chemical affinities of W, Mo, and nitrogen in the system are the main causes of this change in toughness. Acknowledgement This work was supported by Industrial Strategic Technology Development Program, Grant No. 10043203, funded by the Ministry of Trade, Industry & Energy (MI, Korea). We are thankful to Soo J. Shin, Ki H. Kim and Hyun H. Lee at Nanotech Inc. for helpful discussion. We also acknowledge the Research Institute of Advanced Materials, Seoul National University for the use of facilities. References [1] V. Sarin, Comprehensive Hard Materials, Elsevier Science Ltd, Waltham, MA, 2014. [2] S. Zhang, Titanium carbonitride-based cermets: processes and properties, Mater. Sci. Eng. A 163 (1993) 141e148. [3] P. Ettmayer, H. Kolaska, W. Lengauer, K. Dreyer, Ti(C,N) cermetsdmetallurgy and properties, Int. J. Refract. Metals Hard Mater. 13 (1995) 343e351. [4] Y. Peng, H. Miao, Z. Peng, Development of TiCN-based cermets: Mechanical properties and wear mechanism, Int. J. Refract. Metals Hard Mater. 39 (2013) 78e89. [5] M. Humenik, N.M. Parikh, Cermets: i, fundamental concepts related to microstructure and physical properties of cermet systems, J. Am. Ceram. Soc. 39 (1956) 60e63. [6] S. Kang, Some issues in Ti(CN)-WC-TaC cermets, Mater. Sci. Eng. A 209 (1996) 306e312. [7] F. Qi, S. Kang, A study on microstructural changes in Ti (CN)eNbCeNi cermets, Mater. Sci. Eng. A 251 (1998) 276e285. n, Effect of carbon content on the microstructure and [8] J. Zackrisson, H.-O. Andre mechanical properties of (Ti, W,Ta,Mo)(C,N)e(Co,Ni) cermets, Int. J. Refract. Metals Hard Mater. 17 (1999) 265e273. [9] S.Y. Ahn, S. Kang, Formation of core/rim structures in Ti(C,N)-WC-Ni cermets via a dissolution and precipitation process, J. Am. Ceram. Soc. 83 (2000) 1489e1494. [10] S.Y. Ahn, S.W. Kim, S. Kang, Microstructure of Ti(CN)eWCeNbCeNi cermets, J. Am. Ceram. Soc. 84 (2001) 843e849. [11] J. Wang, Y. Liu, Y. Feng, J. Ye, M. Tu, Effect of NbC on the microstructure and sinterability of Ti(C0.7N0.3)-based cermets, Int. J. Refract. Metals Hard Mater. 27 (2009) 549e551.
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[12] S.M. Rafiaei, J.-H. Kim, S. Kang, Effect of nitrogen and secondary carbide on the microstructure and properties of (Ti0.93W0.07)CeNi cermets, Int. J. Refract. Metals Hard Mater. 44 (2014) 123e128. [13] J. Jung, S. Kang, Effect of nano-size powders on the microstructure of Ti(C,N)xWC-Ni cermets, J. Am. Ceram. Soc. 90 (2007) 2178e2183. [14] M. Ashby, F. Blunt, M. Bannister, Flow characteristics of highly constrained metal wires, Acta Metall. 37 (1989) 1847e1857. [15] F.-S. Shieu, R. Raj, S. Sass, Control of the mechanical properties of metalceramic interfaces through interfacial reactions, Acta Metall. Mater. 38 (1990) 2215e2224. [16] R. Mitra, W. Chiou, M. Fine, J. Weertman, Interfaces in as-extruded XD Al/TiC and Al/TiB2 metal matrix composites, J. Mater. Res. 8 (1993) 2380e2392. [17] S. Park, S. Kang, Toughened ultra-fine (Ti,W)(CN)eNi cermets, Scr. Mater. 52 (2005) 129e133. [18] S. Park, Y. Kang, H. Kwon, S. Kang, Synthesis of (Ti,M1M2)(CN)eNi nanocrystalline powders, Int. J. Refract. Metals Hard Mater. 24 (2006) 115e121. [19] J. Kim, M. Seo, S. Kang, Microstructure and mechanical properties of Ti-based solid-solution cermets, Mater. Sci. Eng. A 528 (2011) 2517e2521. [20] J. Kim, S. Ahn, S. Kang, Effect of the complete solid-solution phase on the microstructure of Ti(CN)-based cermet, Int. J. Refract. Metals Hard Mater. 27 (2009) 224e228. [21] D. Shetty, I. Wright, P. Mincer, A. Clauer, Indentation fracture of WC-Co cermets, J. Mater. Sci. 20 (1985) 1873e1882. [22] C. Park, S. Nam, S. Kang, Enhanced toughness of titanium carbonitride-based cermets by addition of (Ti,W)C carbides, Mater. Sci. Eng. A (2015), http:// dx.doi.org/10.1016/j.mesa.2015.10.025. [23] S. Ahn, S. Kang, Effect of various carbides on the dissolution behavior of Ti(C0.7N0.3) in a Ti(C0.7N0.3)e30Ni system, Int. J. Refract. Metals Hard Mater. 19 (2001) 539e545. [24] A. Doi, T. Nomura, M. Tobioka, Thermodynamic Evaluation of Equilibrium Nitrogen Pressure and WC Separation in Ti-W-C-N System Carbonitride, High. Temp. High Press 18 (1986) 443e452. [25] Y. Kang, S. Kang, WC-reinforced (Ti,W)(CN), J. Eur. Ceram. Soc. 30 (2010) 793e798. [26] S. Kim, J.-M. Zuo, S. Kang, Effect of WC or NbC addition on lattice parameter of surrounding structure in Ti(C0.7N0.3)eNi cermets investigated by TEM/CBED, J. Eur. Ceram. Soc. 30 (2010) 2131e2138. [27] A. Arya, E.A. Carter, Structure, bonding, and adhesion at the TiC(100)/Fe(110) interface from first principles, J. Chem. Phys. 118 (2003) 8982e8996. [28] P. Tasker, The stability of ionic crystal surfaces, J. Phys. C Solid State Phys. 12 (1979) 4977. [29] R. Mitra, W. Chiou, J. Weertman, M. Fine, R. Aikin, Relaxation mechanisms at the interfaces in XD™ Al/TiC p metal matrix composites, Scr. Metall. Mater. 25 (1991) 2689e2694. [30] Q. Yang, W. Xiong, G. Zhang, B. Huang, Grain growth in Ti(C,N)-based cermets during liquid-phase sintering, J. Am. Ceram. Soc. 98 (2014) 1005e1012. [31] A. Bradshaw, J. Van der Veen, F. Himpsel, D. Eastman, Electronic properties of the clean and hydrogen-covered TiC(111) Ti-terminated polar surface, Solid State Commun. 37 (1981) 37e40. [32] A. Fujimori, F. Minami, N. Tsuda, Electronic properties of TiC(100) and polar TiC(111) surfaces, Surf. Sci. 121 (1982) 199e217. [33] X. Song, Y. Gao, X. Liu, C. Wei, H. Wang, W. Xu, Effect of interfacial characteristics on toughness of nanocrystalline cemented carbides, Acta Mater. 61 (2013) 2154e2162. [34] R. Mitra, Y. Mahajan, Interfaces in discontinuously reinforced metal matrix composites: an overview, Bull. Mater. Sci. 18 (1995) 405e434. [35] X. Wang, X. Song, X. Liu, X. Liu, H. Wang, C. Zhou, Orientation relationship in WC-Co composite nanoparticles synthesized by in situ reactions, Nanotechnology 26 (2015) 145705.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Carbide/binder interfaces in Ti(CN)e(Ti,W)C/(Ti,W)(CN)-based cermets Choongkwon Park, Sangwoo Nam, Shinhoo Kang* Department of Materials Science and Engineering, Seoul National University, Gwanak-gu, Seoul, 151-742, South Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 24 August 2015 Received in revised form 12 October 2015 Accepted 13 October 2015 Available online xxx
Partially solutionized carbide cermets (PSCs) were prepared for cutting tool applications by replacing a portion of Ti(CN) with (Ti0.88W0.12)C or (Ti0.88W0.12)(C0.7N0.3) in a conventional Ti(CN)-based cermet. The PSC containing (Ti0.88W0.12)C exhibited high fracture toughness with no loss of hardness. The elimination of Ti(C0.7N0.3) cores in the microstructure reduced the strain at the core/rim interfaces, leading to the increase in fracture toughness. The PSC containing (Ti0.88W0.12)C showed coherent relationships at the core/rim and rim/binder interfaces, contributing to the improved mechanical properties. In contrast, nitrogen in (Ti0.88W0.12)(C0.7N0.3) resulted in a different microstructure and properties. The major factors determining the mechanical properties of the cermets are discussed in terms of the carbide/binder interfaces and the thermal stability of the added carbides. © 2015 Elsevier B.V. All rights reserved.
Keywords: Ti(CN) cermet Mechanical properties Microstructure Interface HRTEM
1. Introduction Ti(CN)-based cermets have received much attention as promising alternatives to conventional WCeCo cemented carbide. They exhibit excellent wear-resistance, good hardness, a low friction coefficient, good deformation resistance, good chemical stability at high temperatures, and natural abundance [1e4]. However, the relatively low fracture toughness of Ti(CN)-based cermets limits their wide use as commercial cutting tools. Secondary carbides, such as WC, Mo2C, NbC, and TaC, have been used to improve mechanical properties [3e12]. Adding secondary carbides to conventional Ti(CN)-based cermets (CTCs) produces a core/rim structure through the dissolutioneprecipitation of raw carbides during liquid-phase sintering [7e10]. Thus, CTCs contain a variety of interfaces in their microstructures, including core/inner rim, inner rim/outer rim, and outer rim/binder interfaces, large fractions of which can impair the mechanical properties of the materials [3,9,13]. Core/rim interfaces generally have a high lattice strain due to the large difference in their lattice parameters [9]. Moreover, the mechanical properties of metal-ceramic composites are highly dependent on the strength of ceramic/metal interfaces
* Corresponding author. E-mail address: [email protected] (S. Kang). http://dx.doi.org/10.1016/j.jallcom.2015.10.121 0925-8388/© 2015 Elsevier B.V. All rights reserved.
[14e16]. TiC-based complete solid-solution carbide cermets (CSCs) that consist of a single carbide phase and a binder phase to eliminate the core/rim interfaces have been reported [12,17e19]. They exhibited improved fracture toughness, but suffered from low hardness. In our previous reports, partially-solutionized carbide cermets (PSCs) were firstly introduced, which were fabricated by blending CTC with (Ti1-xWx)C complete solid-solution phase, and showed both improved hardness and toughness [20]. According to the results, PSCs exhibited peak hardness and toughness values with the addition of 50 wt.% (Ti0.93W0.07)C solid solution although the PSC showed the largest particle size and the smallest core fraction among PSCs. The mechanism for the improvements has been remained unclear, and it is necessary for these cermet systems to understand it for the field of the cermet research. In this study, the effect of (Ti0.88W0.12)(C1yNy) addition on the microstructure and mechanical properties of Ti(CN)-based cermets is investigated in detail in terms of the core/rim/binder interfaces and the thermal stability of the added carbides. To the best of our knowledge, we firstly suggest the effect of the interfacial characteristics in an atomic scale on mechanical properties in Ti(CN)based cermet systems. For this investigation, compositions similar to those of commercial cermets, i.e., Ti(C0.7N0.3)e WCeMo2CeNbCeNi/Co, were chosen and two PSCs were prepared by adding (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solutions to
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contrast microstructures and mechanical properties. 2. Experimental procedure (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solution powders were synthesized via high-energy ball milling and subsequent carbothermal reduction, as described in our previous reports [17,18]. TiO2 (particle size: 20 mm; purity: 99%; SigmaeAldrich, St. Louis, MO, USA), WO3 (particle size: 20 mm; purity: 99%; SigmaeAldrich), and graphite (particle size: 1.65 mm; purity: 99%, SeungLim Carbon Metal, Ansan, Korea) were used as raw materials. The amount of each raw material was determined by the following reaction equations.
ð1 xÞTiO2 þ xWO3 þ ð3 þ xÞC ¼ ðTi1x Wx ÞC þ ð2 þ xÞCOðgasÞ (1) N2
ðgasÞ
ð1 xÞTiO2 þ xWO3 þ ð3 þ x yÞC!ðTi1x Wx Þ C1y Ny
þ ð2 þ xÞCOðgasÞ (2) The raw materials were milled with a planetary high-energy ball mill (Pulverisette 7, Fritsch, Idar-Oberstein, Germany) at a speed of 250 rpm for 20 h. A WCeCo-coated stainless steel jar and WCeCo balls were used for milling with a ball-to-powder weight ratio (BPR) of 40:1. The milled powder mixtures were carbothermally reduced at 1300 C for 2 h in a graphite vacuum furnace in a vacuum (103 torr) for (Ti0.88W0.12)C and under 10 torr of nitrogen flow for (Ti0.88W0.12)(C0.7N0.3). The heating rate was 10 C/min. The raw materials used to prepare the PSCs are listed in Table 1. They were conventionally mixed to reach the target compositions listed in Table 2. Ti(C0.7N0.3) and (Ti0.88W0.12)C/ (Ti0.88W0.12)(C0.7N0.3) solid solution powders in fixed quantities were the major components, and WC, Mo2C, and NbC were added as secondary carbides along with Ni and Co as binder metals. (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) solid solution powders were used to produce PSC1 and PSC2, respectively. Overall compositions were the same except for the total nitrogen content (Table 2). The mixtures were then milled in ethanol using WCeCo balls in plastic bottles. The BPR was 10:1, and rotation was at 200 rpm for 20 h. Paraffin (4 wt %) was added for ease of compaction. The mixed slurries were dried and then sieved (125 mesh) to remove agglomerates. The powder mixtures were pressed into discs under a pressure of 150 MPa. The green discs were conventional vacuum sintered in a graphite vacuum furnace (~103 torr) at 1510 C for 1 h. For microstructural observations, as-sintered samples were mechanically ground and then polished using diamond slurry (1e6 mm). The porosity levels of samples were measured according
to the ASTM (American Society for Testing and Materials) specifications to provide the reliability for the discussions below. Powder X-ray diffraction (XRD) patterns were obtained with a diffractometer (D8-Advance, Bruker, Billerica, MA, USA) with monochromatized Cu-Ka radiation (l ¼ 1.54056 Å). The lattice parameters of the specimens were calculated by Rietveld refinement of the XRD results. Scanning electron microscopy (SEM) observations were performed by field emission SEM (FE-SEM; MERLIN Compact, ZEISS, Oberkochen, Germany). Conventional transmission electron microscopy (TEM), highresolution TEM (HRTEM), and scanning TEM high-angle annular dark-field (STEM-HAADF) observations were carried out using a transmission electron microscope (F20 FEI Tecnai, FEI, Hillsboro, OR, USA) operated at 200 KV. Energy dispersive X-ray spectroscopy (EDS) analyses were also performed using an EDX RTEM X-ray detector combined with a STEM-HAADF detector (F20 FEI Tecnai). The carbide sizes and the binder fractions of the sintered specimens were estimated using isolution DT 10.3 software equipped in optical microscopy (Image & Microscope Technology, Daejon, Korea) based on the back-scattered electron (BSE) SEM images. More than 300 carbide particles were measured to obtain statistically meaningful results. Hardness was measured using a Vicker's hardness indentation tester with a 30 kg load and a 15 s loading time. Fracture toughness was calculated by measuring the crack length generated by the Vickers indentation using the expression derived by Shetty et al. [21]. 3. Results and discussion 3.1. Powders and mechanical properties of PSC cermets Fig. 1 shows FESEM-SE images and XRD patterns for assynthesized (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) powders. Both the (Ti0.88W0.12)C and (Ti0.88W0.12)(C0.7N0.3) powders have an irregular shape, and the former is larger than the latter. The XRD results in Fig. 1(c) show that both powders have a NaCl-structure and consist of a single-phase. The XRD peaks of (Ti0.88W0.12)C were at smaller angles than those of (Ti0.88W0.12)(C0.7N0.3) and Ti(C0.7N0.3). The average particle sizes estimated with specific surface areas obtained by BET measurements were ~400 nm for (Ti0.88W0.12)C and ~40 nm for (Ti0.88W0.12)(C0.7N0.3), which are in good agreement with the FESEM images. The mechanical properties, amount of binder phase, and porosity levels of PSC1 and PSC2 cermets are summarized in Table 3, which is recently reported in our report [22]. PSC1 (Hv, 14.0 ± 0.1 GPa; K1C, 12.2 ± 0.2 MPa m1/2) is much tougher, with no deterioration in hardness, than PSC2 (Hv, 13.7 ± 0.2 GPa; K1C, 9.8 ± 0.2 MPa m1/2) and commercial. The porosity levels of PSC1 and PSC2 were excellent as A1B1. Comparing mechanical properties of PSC1 and PSC2, the higher fracture toughness of PSC1 is found not a function of binder fraction. It is because PSC2 exhibited much
Table 1 Powder specifications. Constituents
BET SSA(m2/g)
Size (mm)
Purity (wt %)
Manufacturer
Ti(C0.7N0.3) WC Mo2C NbC Ni Co (Ti0.88W0.12)C (Ti0.88W0.12)(C0.7N0.3)
1.85 e e e e e 2.72 26.51
1.50 2.73 1.85 1.45 2.40 2.30 (0.354b) (0.038b)
99.7 99.7 98.8 99.8 98.8 99.9 e e
Treibacher Industrie AG TaeguTec PPM Ltd.a PPM Ltd. PPM Ltd. Cobatech Corp. Synthesized Synthesized
a b
Pacific Particulate Materials Ltd. Particle size calculated by BET-SSA value.
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Table 2 Compositions of two cermets (wt.%). Specimens
(Ti0.88W0.12)C
(Ti0.88W0.12) (C0.7N0.3)
Ti(C0.7N0.3)
WC
Mo2C
NbC
Ni/Co
PSC1 PSC2
32 e
e 32
27.2 27.2
5.6 5.6
9.6 9.6
5.6 5.6
20 (Ni/Co ¼ 1) 20 (Ni/Co ¼ 1)
Fig. 2. XRD patterns for as-sintered cermets for (a) PSC1 and (b) PSC2. Lattice parameters of the rim and binder phases are shown [22].
higher binder fraction than PSC1 (Table 3), indicating that other factors control the toughness. Also, the mechanical properties of PSCs strongly depend on the composition of the added solidsolution carbides, especially the N content. To investigate the effect of the carbide addition on the mechanical properties of the cermets, we examined the microstructure of the PSC1 and PSC2 in detail. 3.2. Core/rim microstructures of PSC cermets
Fig. 1. XRD patterns and FESEM-BSE images for as-synthesized solid solution powders of (a) (Ti0.88W0.12)C and (b) (Ti0.88W0.12)(C0.7N0.3). (c) XRD patterns for (Ti0.88W0.12)C (red line) and (Ti0.88W0.12)(C0.7N0.3) powders (blue line) (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
Fig. 2 is the XRD profiles of the PSC1 and PSC2 along with reference data for Ti(C0.7N0.3), Ni, and Co from the JCPDs cards [22]. PSC1 consists of a single carbide phase with no other phases (Fig. 2(a)). In contrast, the peaks for the carbonitride phases in PSC2 are separated into two groups. The low-intensity peaks match those for undissolved Ti(C0.7N0.3), while the high-intensity peaks, presumably corresponding to the (Ti,W,Mo,Nb)(CN) solid-solution carbonitride rim phase, formed through dissolutioneprecipitation of the raw materials, which are located at lower angles (Fig. 2(b)). In both cermets, the peaks for the solid-solution carbonitride and binder phases are shifted to lower angles compared with the reference data for Ti(C0.7N0.3) and Ni/Co. The lattice parameters of the solid-solution carbonitride and binder phases calculated by Rietveld refinement are shown next to each XRD profile in Fig. 2. The greater peak shifts for the carbonitride phase in PSC1 than PSC2 is due to different dissolution rates of Ti(C0.7N0.3), (Ti0.88W0.12)C, and (Ti0.88W0.12)(C0.7N0.3) as explained in elsewhere [9,10,22,23]. However, the peak shifts for the binder phases are larger for PSC2
Table 3 Black core fractions, black core sizes, carbonitride sizes, binder fractions, mechanical properties, and porosities for PSC1, PSC2 and commercial CTCs [22]. Specimens
Black core fraction (%) Black core size (mm) Carbonitride size (mm) Binder fraction (%) Hardness (GPa) Fracture toughness (MPa$m1/2) Porosity
PSC1 2.5 PSC2 18.2 Commercial CTCs 16.0
0.63 0.67 0.62
1.19 0.82 0.85
16.8 21.9 20.5
14.0 ± 0.1 13.7 ± 0.2 13.8e14.1
12.2 ± 0.2 9.8 ± 0.2 8.8e9.2
A1B1 A1B1 A1B1
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Table 4 Binder compositions analysed by STEM-HAADF/EDS for PSC1 and PSC2. Specimen
PSC1
Element
Content (wt %)
Uncert. (%)
Content (wt %)
PSC2 Uncert. (%)
Co Ni Ti W Mo Nb C N
44.40 41.98 1.21 2.25 9.12 0.08 0.41 0.51
0.22 0.21 0.03 0.07 0.16 0.05 0.04 0.04
33.01 29.51 1.38 16.49 16.52 0.47 1.03 1.55
0.27 0.26 0.05 0.30 0.32 0.05 0.20 0.10
than for PSC1 due to the low chemical affinities of W and Mo with N in the PSC2, which we have reported elsewhere [12,22,24,25]. Table 4 shows the STEM-HAADF/EDS results, confirming that the binder phase of PSC2 has higher W, Mo, and Nb contents than that of PSC1. These increase the volume fraction of the binder phase, but the large fraction of binder phase did not increase the toughness of PSC2 (Table 3). The FESEM-BSE images for PSC1 and PSC2 are shown in Fig. 3(a) and (b), respectively. The cores in PSC1 appear as grey contrast with inner and outer rims; they were confirmed by STEM-HAADF/EDS as
Fig. 3. FESEM-BSE images for as-sintered (a) PSC1 and (b) PSC2. White, black and red arrows indicate inter-granular, trans-granular and inter-core/rim failure modes, respectively (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
raw (Ti,W)C. Some black Ti(CN) cores are also visible because of the similar phase stabilities of Ti(C0.7N0.3) and (Ti0.88W0.12)(C0.7N0.3), which is in good agreement with the XRD results in Fig. 2 [22]. In addition, many Ti(C0.7N0.3) cores in PSC2 were completely or partially denuded from the rim phase with some white core/grey rims as indicated in Fig. 3(b) [10,26]. Fig. 3 also shows the crack paths induced by indentation method in PSC1 and PSC2. Three kinds of crack propagation mode were observed, i.e. inter-granular, trans-granular and inter-core/rim failure modes, which were marked by white, black and red arrows, respectively. The cracks in PSC1 mainly propagated in transgranular failure mode and the cracks traversed even small particles located at the direction of crack propagation (black arrows in Fig. 3 (a)). In contrast, inter-granular and inter-core/rim failure modes are dominant in PSC2 (red and white arrows in Fig. 3(b)). It is true that larger carbide size (>30%) and low fraction of hard phase in PSC1 would facilitate trans-granular failure when crack propagates. However, the crack paths in PSC1 also indicate that the rim/binder interfaces in PSC1 could be much stronger than those of PSC2. Therefore, the interfacial characters between core/rim and rim/binder were characterized to understand the factors influencing the higher toughness of PSC1. 3.3. Rim/binder interfaces of PSC cermets Fig. 4 shows bright-field TEM (BF-TEM) images of PSC1, which contains various interface regions of interest observed by HRTEM. More than 20 interfaces of >10 particles were observed from each specimen. Some representative HRTEM images are shown in Fig. 4(AeD) of the regions that are enclosed by white squares in the BF-TEM images in Fig. 4(a) and (b). Each inset in Fig. 4(AeD) shows the corresponding FFT patterns. The lattice planes on the HRTEM images were indexed by comparing the interplanar distances (d) measured from the HRTEM images with the values calculated from the XRD results. The interplanar distances, d{111}C, d{200}C, and d {111}B were 2.4917, 2.1578, and 2.0653 Å, respectively, where subscripts C and B indicate carbonitride and binder, respectively. The carbonitride particles in PSC1 consisted of faceted surfaces and rounded corners. Fig. 4(AeC) shows the facets of the carbonitride particles terminating parallel to the {111}C planes, and the interfaces between the carbonitride and binder appear coherent or semi-coherent with ~30% probability of overall interfaces selected randomly. For NaCl-structured carbides, such as TiC, {111} surfaces are less stable than {100} surfaces as vacuum surfaces because of their low packing density [27] and dipolar nature [28]. However, it is often reported that NaCl-structured TiC and Ti-based carbonitride particles in a metal matrix terminate commonly with {111} surfaces [16,29,30]. Ti-terminated {111} surfaces are known to have a high local density of states around the Fermi level (EF) [31,32], leading to the high electrical conductivity and metallic properties. Mitra et al. reported that the {111} surfaces of TiC precipitates are stabilized by electronic interactions with Al atoms in an Al metal matrix, resulting in strong metallic bonding [16]. A similar explanation can be applied to our PSC cermets. Fig. 4(A) shows HRTEM images of the interface between the carbonitride facet and binder boundary, where the interface is characterized as (11-1)C//(200)B. In the image, the carbonitride particle and binder are oriented with its [011] zone axis (see FFT pattern). The interface is asymmetric, but the two phases at the interface are found to have the interfacial relationship of (11-1)C// (100)B. Thus, the d-spacings for (200)C and (-11-1)B planes shown in Fig. 5(a) determines the level of mismatch. Based on this, the lattice mismatch between the (200)C and (-11-1)B planes is calculated as 4.4%, where semi-coherent interfaces tend to form, involving dislocations. However, the two phases are found perfectly matched
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Fig. 4. TEM and HRTEM images for PSC1 and FFT pattern: (a) and (b) show low magnification BF-TEM images for the typical rounded carbonitride grains. AeD show the HRTEM images for the interface areas enclosed by the white squares in (a) and (b). Insets are the corresponding FFT patterns.
Fig. 5. Schematic illustration of the orientation relationship. (a) Coherent relationship of (11-1)C//(200)B. (b) Stacking faults by Shockley partial dislocation of the ɑo/6〈112〉 Burgers vector.
and coherent in long-range, since the d-spacings for the (200)C planes and the (-11-1)B planes around the interface are measured the same (2.0668 Å). This indicates that at the interface, the lattices of the carbonitride are distorted to adjust the lattice parameters. In addition, stacking faults, which exist three atomic layers away from the rim/binder interface, are observed in the carbonitride. The faults are highlighted by yellow dotted lines (see the characteristic FFT pattern). The carbonitride can accommodate the local distortion at the interface up to the third atomic layer (part of the layer is outlined by a red dotted box in Fig. 4(A)), but beyond this layer, the stacking fault generation is energetically favourable. For ease of visualization, the interface and the stacking faults are shown in Fig. 5(a) and (b), respectively. Each stacking fault is generated by a Shockley partial dislocation with a ɑo/6〈112〉 Burgers vector, and
the total Burgers vector is equal to b ¼ ɑo/2〈112〉. The presence of three stacking faults completely restores the original stacking sequence of the carbonitride lattice and releases the lattice strain. Therefore, the (11-1)C//(100)B interface becomes coherent. Fig. 4(B) shows representative carbonitride/binder interfaces at the round corner of a carbonitride particle. Steps at the interface were observed that consisted of two {111} faces, and are indicated by black dotted lines in the figure. The interfacial relationships at the steps were also similar to those at the faceted interfaces; thus, the interfaces were coherent/semi-coherent. However, the (11-1)C interface shown in Fig. 4(C) is parallel to the high-index planes of the binder phase, and the binder lattice was distorted around the interface to maintain the lattice matching. The orientation angle of the (-11-1)B planes to the
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interface [(11-1)C] was q ¼ 64 . In this orientation, the misfit between the (200)C and (-11-1)B planes was calculated as 15%, which is the highest misfit we have observed. Around the interface, the (-11-1)B planes were slightly bent to adjust the lattice mismatch. In addition, edge dislocations in the binder lattice were observed at the interface, where a typical dislocation is indicated by a white arrow. The arrays of the binder and carbonitride lattice planes are locally changed around the defects. Therefore, the lattice parameters are adjusted at the interface by the distortions, resulting in a semi-coherent interface. Similar observations have been reported for several ceramic/metal composites such as TiC/Al [16] and WC/ Co [33]. Fig. 4(D) shows the HRTEM image of the representative (Ti,W)C grey core/rim interface. The lattice orientations between the two phases were similar (d(200)grey core ¼ 2.145 Å and d(200)rim ¼ 2.156 Å), and the interface was highly coherent (d ¼ 0.5%). These results are in good agreement with the XRD results in Fig. 2(a), indicating that the grey core/rim carbonitride particles are free from lattice strain. Fig. 6 shows BF-TEM images for typical faceted carbonitride particles in the PSC2 binder. HRTEM images of the typical smooth interface regions (enclosed by white squares in Fig. 6(a) and (b)) are presented in Fig. 6(A) and (B), respectively, where the insets are the corresponding FFT patterns. The internal angles of the carbonitride particles in PSC2 were mostly measured as 70.53 and 109.47, which are the angles observed when 3-D shape of particles is an octahedron terminated by {111} facets [30], and the edges of the carbonitride particles in PSC2 were highly sharp in contrast to those in PSC1 (Fig. 6 (a) and (b)). HRTEM images in Fig. 6(A) and (B) prove
that the facets of the carbonitride particles in PSC2 are terminated by {111} planes. HRTEM images also show that the carbonitride particles were oriented with the electron beam parallel to [011] zone axis, but the binder phases were tilted far away from the [011] zone axis. It indicates that the carbonitride and binder phases in PSC2 were just randomly oriented at their interfaces and PSC2 did not show any preferred orientation relationship between them similar to that shown by PSC1. Therefore, the difference in the carbonitride/binder interfacial relationships between PSC1 and PSC2 can generate the great difference in their mechanical properties, especially toughness. It is well known that the strength and toughness of the ceramic metal composites are strongly dependent on the interfacial characteristics between ceramic and metal matrix [33e35]. The coherency state at the core/rim/binder interfaces in PSC1 can enhance their bond strength, leading to its high toughness. The formation of coherent interfacial relationships in PSC1 depends on the lattice parameters of the core, rim and binder phases, as well as the electronic interactions between the carbonitride and binder phases. 4. Conclusions We prepared PSC cermets by adding (Ti0.88W0.12)C or (Ti0.88W0.12)(C0.7N0.3) solid solution to the composition of a conventional Ti(CN)-based cermet, and then compared their mechanical properties. The elimination or minimization of highly strained Ti(CN) core/rim interfaces through adding (Ti0.88W0.12)C
Fig. 6. TEM and HRTEM images for PSC2 and corresponding FFT patterns: (a) and (b) are low magnification images showing the typical faceted carbonitride grains in the binder phase. A and B show the HRTEM images for the regions enclosed by the white squares in (a) and (b). Insets show the corresponding FFT patterns.
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solid solution significantly toughened the PSCs. The solid-solution carbonitride in PSCs commonly exhibited {111} plane faceting at the interface with the binder phase. Coherent interfaces with a (111)C//(200)B relationship or semi-coherent interfaces were often observed at the interfaces in PSC1, whereas incoherent interfaces were observed in PSC2. The high fraction of the binder phase in PSC2 did not result in an obvious improvement in the toughness. Thus, we conclude that coherent or semi-coherent interfaces in PSC1 are one of the important factors which contribute to the high fracture toughness of PSC1 more than the binder fraction in the system. Furthermore, we reconfirmed that the thermal stability of the Ti(CN), (Ti0.88W0.12)C, and (Ti0.88W0.12)(C0.7N0.3) phases and the chemical affinities of W, Mo, and nitrogen in the system are the main causes of this change in toughness. Acknowledgement This work was supported by Industrial Strategic Technology Development Program, Grant No. 10043203, funded by the Ministry of Trade, Industry & Energy (MI, Korea). We are thankful to Soo J. Shin, Ki H. Kim and Hyun H. Lee at Nanotech Inc. for helpful discussion. We also acknowledge the Research Institute of Advanced Materials, Seoul National University for the use of facilities. References [1] V. Sarin, Comprehensive Hard Materials, Elsevier Science Ltd, Waltham, MA, 2014. [2] S. Zhang, Titanium carbonitride-based cermets: processes and properties, Mater. Sci. Eng. A 163 (1993) 141e148. [3] P. Ettmayer, H. Kolaska, W. Lengauer, K. Dreyer, Ti(C,N) cermetsdmetallurgy and properties, Int. J. Refract. Metals Hard Mater. 13 (1995) 343e351. [4] Y. Peng, H. Miao, Z. Peng, Development of TiCN-based cermets: Mechanical properties and wear mechanism, Int. J. Refract. Metals Hard Mater. 39 (2013) 78e89. [5] M. Humenik, N.M. Parikh, Cermets: i, fundamental concepts related to microstructure and physical properties of cermet systems, J. Am. Ceram. Soc. 39 (1956) 60e63. [6] S. Kang, Some issues in Ti(CN)-WC-TaC cermets, Mater. Sci. Eng. A 209 (1996) 306e312. [7] F. Qi, S. Kang, A study on microstructural changes in Ti (CN)eNbCeNi cermets, Mater. Sci. Eng. A 251 (1998) 276e285. n, Effect of carbon content on the microstructure and [8] J. Zackrisson, H.-O. Andre mechanical properties of (Ti, W,Ta,Mo)(C,N)e(Co,Ni) cermets, Int. J. Refract. Metals Hard Mater. 17 (1999) 265e273. [9] S.Y. Ahn, S. Kang, Formation of core/rim structures in Ti(C,N)-WC-Ni cermets via a dissolution and precipitation process, J. Am. Ceram. Soc. 83 (2000) 1489e1494. [10] S.Y. Ahn, S.W. Kim, S. Kang, Microstructure of Ti(CN)eWCeNbCeNi cermets, J. Am. Ceram. Soc. 84 (2001) 843e849. [11] J. Wang, Y. Liu, Y. Feng, J. Ye, M. Tu, Effect of NbC on the microstructure and sinterability of Ti(C0.7N0.3)-based cermets, Int. J. Refract. Metals Hard Mater. 27 (2009) 549e551.
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