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Materials Science and Engineering A 528 (2011) 2517–2521

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

icrostructure and mechanical properties of Ti-based solid-solution cermets

iwoong Kim, Moonsu Seo, Shinhoo Kang ∗

epartment of Materials Science and Engineering, Seoul National University, 56-1 Shilim-dong Kwanak-ku, Seoul 151-742, Republic of Korea

r t i c l e i n f o

rticle history:eceived 30 April 2010eceived in revised form2 November 2010ccepted 24 November 2010

a b s t r a c t

The microstructure and mechanical properties of various (Ti1−xWx)C–20 wt.%Ni cermets were investi-gated, where x varies from 0.07 to 0.3. Homogeneous solid-solution (Ti1−xWx)C powders were preparedby carbothermal reduction via planetary milling of Ti, TiO2, WO3 and carbon powder mixtures. The cer-mets made of the powders showed a simple core–rim structure consisting of solid-solution carbides. Thehardness of the solid-solution cermets is somewhat lower than that of conventional cermets, but they

vailable online 2 December 2010

eywords:i-based solid-solutionarbothermal reductionicrostructure

show greater toughness. The transverse rupture strength increases with increasing W content.© 2010 Elsevier B.V. All rights reserved.

echanical propertiesransverse rupture strength

. Introduction

TiC- and Ti(CN)-based cermets are commonly used in cuttingool applications [1–3]. These cermets have a typical core–rimtructure in which the core consists of undissolved TiC or Ti(CN)articles and the rim is a solid-solution phase produced duringhe sintering stage. In practical cutting performance, TiC- andi(CN)-based cermets show lower fracture toughness than doC–Co cermets. The low toughness is largely attributed to the

oor wettability of Ni binder on TiC and Ti(CN), and to the com-lex microstructures of the cermets along with a large number of

nterfaces.In general, Ti(CN)–MeC–Ni cermet systems have core–inner rim

nd inner rim–outer rim (carbide) interfaces in addition to carbideouter rim)–Ni binder interfaces. The interface between Ti(CN) corend inner rim phase can be highly strained, depending on the dif-erence in composition between the core and constituent inner rimhase. These interfaces are potential sites for crack initiation andropagation. The presence of many interfaces in the microstructureesults in reduced fracture toughness for cermets, thereby limitingheir industrial applications [2–7].

Many studies have sought to enhance the fracture tough-ess of cermet tools. Because the solid-solution rim phases arenown to improve the fracture toughness of various cermets, manyesearchers have tried to increase the volume fraction of the rim

∗ Corresponding author. Tel.: +82 2 880 7167; fax: +82 2 884 0074.E-mail address: shinkang@snu.ac.kr (S. Kang).

921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2010.11.076

phase by adding various secondary carbides [4–11]. Other stud-ies have sought to increase the toughness of cermets by producingsingle-phase and homogeneous solid-solution cermets, using TiCand WC powders [5,7–9,11–13]. However, these approaches havemet with only partial success, meaning that the issue of enhancedtoughness has become an ongoing challenge in the development ofcermet materials.

In the present study, we produced various (Ti1−xWx)C single-phase and homogeneous solid-solution powders from oxides andcarbon as reported before [14]. We characterized the resulting cer-mets containing varying amounts of W in solid solutions. We alsoinvestigated the microstructure evolution in the cermet systems,as well as their fracture behavior. The (Ti,W)C-based solid-solutioncermets show new and attractive microstructures and mechanicalproperties.

2. Experimental

The starting materials, anatase TiO2 (>99%, ∼40 �m, Aldrich), Ti(>99%, ∼75 �m, Cerac), WO3 (>99%, ∼20 �m, Aldrich), and C (>99%,∼1.7 �m, Lonza) powders, were mixed to produce solid-solutioncarbide powders (Ti1−xWx)C. The mole fraction of W was variedfrom 0.07 to 0.30. Titanium elemental powder and oxide powderswere mixed and milled using a planetary mill (Fritsch Pulverisette

7, Germany) at a speed of 250 rpm for 20 h; the ball-to-powderratio was 30:1. We used Ti elemental powder and TiO2 mixtureto produce TiO during high-energy ball milling, because TiO has agreater tendency to be carburized than does TiO2 or Ti2O3 [15,16].Tungsten carbide balls (WC–Co) were used as the milling medium,

2518 J. Kim et al. / Materials Science and Engineering A 528 (2011) 2517–2521

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to Ti and W. However, the free carbon content remains constantat ∼0.2 wt.%. Table 1 lists the average compositions of the solid-solution powders. The carbon stoichiometry, calculated based onthe C/N/O analysis, is in good agreement with data obtained froma previous study [20]. The EDS analysis reveals an increase in the

Table 1Target and measured compositions of solid-solution carbide powders.

Target Measured

Fig. 1. X-ray diffraction profiles of (Ti1−xWx)C solid-solution powders.

nd milling was performed in a tungsten carbide-coated steel bowl.olid-solution powders were prepared by carbothermal reductiont 1350 ◦C for 2 h in a vacuum graphite furnace. After reduction andarburization, the powders were mixed with 20 wt.% of Ni (>99%,4.2 �m, Novamet) in ethanol by ball milling for 24 h. The powdersere dried at 90 ◦C in an oven and compacted into disks by applyinguniaxial pressure of 125 MPa. They were then sintered at 1510 ◦C

or 1 h in a graphite vacuum furnace.The powders were characterized by X-ray diffraction (D-

AX2500-PC, Rigaku, Japan) with monochromatized Cu K� radia-ion (� = 1.5418 A). The powder compositions were analyzed withcarbon analyzer (EC-200AC, LECO, US), an oxygen/nitrogen ana-

yzer (TC-436, LECO, US), and inductively coupled plasma–atomicmission spectroscopy (ICP–AES; Optima-4300 DV, Perkin Elmer,SA). The specific surface area of the powders was obtained by2 adsorption fitted with BET isotherms (Micromeritics, ASAP010, Norcross, USA). The powder particle size was calculatedrom the surface area and density measurements using the formula= {6/(A�)}, where D is the average diameter of the particles, A is

he surface area, and � is the density.The microstructures of sintered samples were observed by

canning electron microscopy (SEM; JSM-6360, JEOL, Japan). Theompositions of the phases found in various cermets were obtainedsing SEM energy-dispersive X-ray spectroscopy (EDS; JSM-6360,

NCA, Japan). The hardness and toughness were measured by theickers indentation method using 30 kg loads over 15 s. The frac-

ure toughness was evaluated by the direct crack measurementethod using the expressions derived by Evans and Charles [17],nstis et al. [18], and Shetty et al. [19]. The transverse rupturetrength (TRS) was measured with a universal testing machineShinpoong GT-Trosee, Korea) at a rate of 0.01 mm/s.

. Results and discussion

.1. Characteristics of (Ti1−xWx)C powders

Fig. 1 shows the X-ray diffraction results for the powders synthe-ized by carbothermal reduction at 1350 ◦C for 2 h after high-energyall milling. For all compositions, the results demonstrate that theomplete solid-solution (Ti1−xWx)C phases formed without anyther phases. In the figure, the peak shifts to lower angles withncreasing W content in (Ti1−xWx)C powder. This shows an increasen the lattice parameter of (Ti1−xWx)C, indicating that hcp WC is noteadily solutionized within the fcc TiC lattice. However, this trend isn contrast with the previous study that reported a gradual decrease

n lattice parameters with increasing of W content in (Ti1−xWx)C20]. The parameters in the previous study were obtained fromolid-solution carbide powders which have low carbon contentsx < 0.8) compared to those of this study as will be discussed below.

Fig. 2. Variations in C, N, and O contents as a function of W content in various(Ti1−xWx)C powders.

Thus, the difference might be caused by the carbon contents in thecarbides, which would vary as a function of initial carbon content,milling technique, reduction temperature and further processingsuch as sintering. The particle sizes of (Ti1−xWx)C powders, as cal-culated based on the measured surface area, are 1.57, 1.69, and1.51 �m for x = 0.07, 0.12, and 0.3 in (Ti1−xWx)C, respectively. Theparticle size of the milled oxide powders was in nanorange. How-ever, after carbothermal reduction process at high temperature, thesolid-solution carbide powders agglomerated severely. Therefore,the calculated particle sizes of (Ti1−xWx)C powders from surfacearea analyses are agglomerates size.

Ti was added to TiO2 and milled together to induce the formationof the TiO phase. The milled powder changes to Ti(C,O) and eventu-ally to (Ti,W)C during reduction, via the reaction of Ti(C,O) with WCformed previously. Based on previous studies, the following reac-tion sequence is expected to occur with increasing temperatureduring the carbothermal reduction process [10,16]:

Ti + TiO2 ⇒ 2TiO(milling)

WO3 + 4C ⇒ WC + 3CO(g)(∼600 ◦C)

TiO + C ⇒ Ti(C0.5O0.5) + 0.5CO(g)(∼1100 ◦C)

(1 − x)Ti(C0.5O0.5) + xWC + (1 −x)C ⇒ (Ti1−xWx)C

+ 0.5(1 − x)CO(g)(1100–1300 ◦C)

The final compositions of the solid-solution carbide powderswere measured using SEM–EDS and C/N/O analyses. Fig. 2 showsthe C/N/O analysis results, revealing that oxygen and nitrogen con-tents decrease with increasing W content in the solid-solutionphase. This result may be related to the affinity of these elements

TiC (Ti0.982W0.018)C0.860

(Ti0.93W0.07)C (Ti0.906W0.094)C0.822

(Ti0.88W0.12)C (Ti0.870W0.130)C0.913

(Ti0.70W0.30)C (Ti0.695W0.305)C0.899

J. Kim et al. / Materials Science and Engineering A 528 (2011) 2517–2521 2519

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ig. 3. SEM–BSE (back-scattered electrons) micrographs of conventional TiC–xWCrrows indicate inter- and trans-granular failure modes, respectively.

content of the solid-solution powders, which is attributed to Wontamination from the milling balls.

.2. Microstructure of solid-solution cermets

The SEM images in Figs. 3 and 4 show the microstruc-ures and fracture behaviors of various cermets. To compare

Ti1−xWx)C–20Ni cermets with conventional cermets, we producediC–xWC–20 wt.%Ni cermets of the same elemental compositionssing a conventional powder metallurgy (CPM) technique [21].hat is, individual carbides such as TiC (>98%, <44 �m, Aldrich) andC (>99%, ∼1.87 �m, H.C. Starck) were mixed together with Ni

ig. 4. SEM–BSE (back-scattered electrons) micrographs of (Ti1−xWx)C–20Ni solid-solutiondicate inter- and trans-granular failure modes, respectively. These compositions are eig. 3(a)–(d).

t.%Ni cermets, for x values of (a) 0, (b) 15, (c) 25, and (d) 48 wt.%. White and black

(>99%, ∼4.2 �m, Novamet) in a horizontal ball mill and sintered atthe same conditions as those for the solid-solution cermets.

Fig. 3 shows a typical core–rim structure of the CPMTiC–xWC–20 wt.%Ni cermets along with crack paths developed asa result of Vickers indentation tests. The volume fraction of the rimphase increases with increasing W content. Fig. 3(b) and (c) clearlyshows (from the contrast in BSE mode) that the rim consists of

inner and outer rims with contrasting W contents. Irrespective ofW content, the CPM cermets contain undissolved TiC particles ashard cores. In contrast, there is no TiC core in the microstructuresof the solid-solution cermets shown in Fig. 4(a)–(d), even thoughthey unexpectedly contain a core–rim structure. The cores consist

n cermets, for x values of (a) 0, (b) 0.07, (c) 0.12, and (d) 0.30. White and black arrowsquivalent to the compositions of the conventional TiC–xWC–20 wt.%Ni cermets in

2520 J. Kim et al. / Materials Science and Engi

Table 2Average compositions of the core and rim phases (at.%) in cermets made of(Ti0.88W0.12)C powder.

Ti W

Core 87.097 (±0.345) 12.098 (±0.338)Rim 85.737 (±0.486) 14.268 (±0.478)

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obtained using the method of Evans et al.

ig. 5. W contamination in (Ti1−xWx)C–20Ni powders after various processing steps.

f undissolved (Ti1−xWx)C solid-solution phase, whereas the rim isnother solid-solution phase of slightly different composition. Thisbservation is evidenced by the SEM–EDS data listed in Table 2.

To understand the origin of the core–rim structure in the solid-olution cermets, we performed ICP–AES analyses to determinehe variations in powder compositions. Fig. 5 shows the changen the W-to-Ti ratio (wt.%) in the solid-solution powders. Initially,he carbothermally reduced powders have similar compositionso the target compositions. However, the W-to-Ti ratio shows aemarkable increase after horizontal ball milling to mix the car-ide powders with Ni binder. This result shows that additional Was obtained at the carbide surface during horizontal ball milling,

ausing a W-rich rim phase to form during sintering. Consequently,he rims that show bright contrast in Fig. 4 formed on partiallyissolved (original) solid-solution carbides, producing a core–rimtructure. The same explanation could apply to the microstructuresf CPM TiC-20 wt.%Ni cermets shown in Fig. 3(a).

In general, planetary milling provides a stronger impact thanoes horizontal ball milling. However, Fig. 5 shows less W gain dur-

ng planetary milling than during horizontal ball milling. This resultay reflect differences in milling conditions and the nature of pow-

ers to be milled (brittle oxides or hard carbides); i.e., planetaryilling was performed on dry and brittle oxide mixtures, whereaswet condition was employed to mix hard carbides with Ni in hor-

zontal ball milling. More WC seems to be removed from milling

Fig. 6. Mechanical properties of conventional powder metallurgy (CPM) and solid-

neering A 528 (2011) 2517–2521

balls in the latter case, and a greater increase in W is produced bywet ball milling than by dry planetary milling.

In terms of particle morphology, all the CPM cermets appear tohave irregular or rectangular particles with a wide size distribu-tion (Fig. 3). In contrast, the solid-solution cermets in Fig. 4 exhibitcarbides of relatively uniform morphology and a narrow size dis-tribution. In the solid-solution cermets, the particle shape changesfrom faceted to spherical with increasing W content. As shown inFig. 4(d), the carbide particles in the (Ti0.70W0.30)C–20Ni cermet arespherical and have a narrow size distribution. This finding is relatedto the thermodynamic stability of the carbide phase (Ti0.70W0.30)C,because the most stable phase would show the lowest degree ofcoarsening during high-temperature sintering (1510 ◦C).

3.3. Mechanical properties of solid-solution cermets

Fig. 6 shows changes in the mechanical properties of the solid-solution cermets as a function of W content in (Ti1−xWx)C. TheVickers hardness values of the cermets are somewhat lower thanor close to those of the CPM cermets, reflecting the absence of theTiC hard phase. Nevertheless, the solid-solution cermets provideexcellent hardness values, due to carbide size and their intrinsichardness. As shown in Figs. 3 and 4, the solid-solution cermetscontain smaller particles than do CPM cermets, thereby compen-sating for the loss of hardness caused by the occurrence of (Ti,W)Csolid-solution carbides as cores.

Both cermets show a gradual decrease in hardness with increas-ing W content. It has been reported that the rim phases in CPMcermets cause a loss of hardness, especially with increasing vol-ume fraction [22,23]. Fig. 3 shows that the rim fraction increases inthe CPM cermets with increasing WC content, resulting in reducedhardness. In the solid-solution cermets, hardness values are alsocontrolled by W amount in the (Ti1−xWx)C phase (Fig. 6).

Fig. 6(b) shows the changes in the toughness values of var-ious cermets. The toughness values of the cermets made fromsolid-solution carbides are clearly superior to those of the CPMcermets. The cermets that contain 12 or 30 at.% W in the carbideshave appreciably higher toughness than do corresponding CPMcermets. The (Ti0.70W0.30)C–20Ni cermet possesses an excellenttoughness of approximately 15.4 MPa m1/2, as measured using theexpression derived by Evans and Charles [17]. To obtain reliabletoughness values, we calculated the toughness of the cermets usingthe methods proposed by Anstis et al. [18] and Shetty et al. [19]. Thevalues obtained using these methods show similar trends to those

Again, this improvement in toughness must be related not onlyto the intrinsic properties of the solid-solution particles, but to theconstituent carbide sizes and uniformity in morphology. In a sep-arate study, the intrinsic toughness of the (Ti1−xWx)C phase was

solution cermets as a function of W content: (a) hardness and (b) toughness.

J. Kim et al. / Materials Science and Engi

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ig. 7. Transverse rupture strength (TRS) of (Ti1−xWx)C–20Ni solid-solution cermetsith respect to W content. TRS of WC–10 wt.%Co is shown for reference.

bserved to increase with increasing W content. The presence ofin the fcc TiC seems to enhance the potential of the carbides

o slip during deformation. Furthermore, the reduction in interfa-ial area that occurs in all cermet systems is an important factor inxplaining their mechanical behavior.

Generally, various interfaces occur in typical cermet systems,ncluding those between the inner rim and outer rim, and betweenhe rim and binder phase. Cracks are readily initiated at the highlytrained interfaces commonly observed at the core–rim interface inost cermet systems. Therefore, a reduction in the interfacial area

s an important factor in increasing the cermet toughness. Althoughore–rim interfaces exist in the solid-solution cermets shown inig. 4(a)–(d), the strain level is not as high as that in conventionalermets because the cores and rims have similar compositions.

Figs. 3 and 4 show the crack paths that developed at the endf indentation tests in various cermets. Both trans-granular (blackrrows in the figures) and inter-granular (white arrows) failureodes prevail throughout the microstructure. In the case of CPM

ermets, trans-granular failure is more frequent at low W contents,ut the incidence of inter-granular failure increases with increas-

ng W content. However, the crack paths in the (Ti1−xWx)C–20Niystem are mostly of inter-granular mode, propagating along theim–binder interface (Fig. 4). These findings indicate that the bondtrength of the core–rim interface in CPM cermets is weaker thanhat of the rim–binder interface. In contrast, in solid-solutionermets, the rim–binder interface is the weakest area and most sus-eptible to cracking, due to the absence or reduction of interfacialtrain at the core–rim interface.

The above result differs significantly from the cracking behav-or of ordinary structural ceramics. That is, oxide ceramics failnter-granularly along the glassy grain-boundary phase. The cermetystems, which have a metallic component as binder phase, com-only show the same inter-granular failure mode. However, the

arbide–binder interface is more metallic in the nature of its bond-ng, while the interfaces in ceramic systems are roughly covalent.hus, the interfacial strength in the cermet should be high relativeo that in ceramics. Accordingly, the solid-solution cermets providemproved toughness by removing or reducing the interfacial strains

resent at the core–rim structure.

Fig. 7 shows that TRS increases with increasing W content inolid-solution cermets. This trend is in line with the high tough-ess of solid-solution cermet systems. Of note, the solid-solutionermets yielded higher TRS values than did WC–Co in the case that

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neering A 528 (2011) 2517–2521 2521

x in (Ti1−xWx)C–Ni lies in the range of 0.12–0.3. Compared withCPM cermets and WC–Co material, the mechanical properties ofsolid-solution cermets show tremendous promise.

4. Conclusions

The microstructure and mechanical properties of Ti-based solid-solution cermets were investigated and compared with those ofcermets prepared by a conventional powder metallurgy (CPM)technique. The conclusions of this study are as follows.

(1) The (Ti1−xWx)C–Ni cermets of this study consist of solid-solution core and rim phases.

(2) Compared with CPM cermets, the morphology of carbidesin (Ti1−xWx)C–Ni cermets is relatively uniform and becomesspherical with increasing W content. The (Ti0.7W0.3)C carbideshows the smallest particle size, indicating that it is the moststable solid-solution carbide among those investigated.

(3) Cracks in the (Ti1−xWx)C–20Ni system tend to propagate inter-granularly, and indentation toughness and TRS increase withincreasing W content.

Overall, the present results demonstrate that Ti-based solid-solution cermets have attractive microstructures and mechanicalproperties. Compared with CPM cermets and even WC–Co materi-als, the solid-solution cermets show promise as replacements forcurrent structural materials.

Acknowledgments

This work was supported by basic research funding from KOSEF(No. R01-2008-000-20905-0). We acknowledge the experimentalhelp from Dr. S.-Y Ahn and S. Oh at Korloy Inc., Korea and the useof facilities at the Research Institute of Advanced Materials, SeoulNational University

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