Available online at www.sciencedirect.com
Ceramics International 41 (2015) 1986–1993
Inﬂuence of applied pressure during ﬁeld-assisted sintering
of Ti(C,N)–WC–FeAl based nanocomposite
, V.V.S.S. Srikanth
, S.V. Joshi
, J. Joardar
International Advanced Research Centre for Powder Metallurgy and New Materials (ARCI), PO Balapur, Hyderabad 500 005, India
School of Engineering Sciences and Technology, University of Hyderabad, Gachibowli, Hyderabad 500 046, India
Received 18 June 2014; received in revised form 18 August 2014; accepted 4 September 2014
Available online 16 September 2014
Ti(C,N)–WC–FeAl based nanocomposites are processed by ﬁeld-assisted sintering at 1500 1C. The phase and microstructural evolution during
the process under the inﬂuence of different applied pressures of 30, 50 and 100 MPa are studied using x-ray diffraction and scanning electron
microscopy. Lattice parameters of (Ti,W)(C,N) solid solution and binder phases after sintering are found to vary with applied pressure. The
nanocomposite grains are observed to possess a core-rim microstructure. Microstructural variations in terms of type, size and fraction of “core-
rim”structure as a function of applied pressure are investigated. The hardness and indentation fracture toughness values are in the range of 17.6–
18.4 GPa and 5.9–6.8 MPa√m, respectively. These values are comparable with those reported for Ti(C,N)-based composites with metal binder.
&2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: A. Sintering; B. Nanocomposites; C. Mechanical properties; E. Cutting tools
Ti(C,N) cermet based machining inserts are widely used for
high-speed semi-ﬁnishing and ﬁnishing operations on hardened
steel and cast iron [1–3]. Consumption of cermet inserts has
seen a gradual increase in recent years due to their superior
performance compared to conventional inserts made of cemented
carbides [4–8]. Improved performance of Ti(C,N) cermet based
inserts has been observed due to reﬁned structure when an
ultraﬁne grade hard phase is present in the microstructure [9–13].
It is envisaged that a further enhancement in the properties may
be attained with the incorporation of alternative binders as a
replacement for the usual metallic binders such as Ni and/or Co,
which are prone to oxidation at high temperatures generated at
the tool tip during machining [14,15].Inaddition,NiandCoare
also expensive as well as carcinogenic [16–18]. In this context,
aluminides with excellent oxidation and corrosion resistance at
high temperature have been studied as binders in TiC/Ti(C,N)
and WC-based composites [19–23].TiC–aluminide composites
have shown improved hardness, fracture toughness and oxidation
resistance and have comparable transverse rupture strength with
Apart from the composition of the cermet and binder consti-
tuents, sintering techniques and associated processing conditi-
ons are likely to play a major role in the development of the
microstructure. Sintering processes like ﬁeld-assisted sintering
technique (FAST) (also called spark plasma sintering (SPS)) and
hot isostatic pressing (HIP) have been used by several groups for
consolidation of ultraﬁne or nanostructured Ti(C,N)-based cermets
[5,26,27]. FAST/SPS, in particular, has added advantages of be-
ing a rapid process involving lower sintering temperature when
compared to conventional processes. Notable alteration in the
microstructure and properties of Ti(C,N)-based cermets processed
by SPS has been reported by various groups [28–30].These
studies indicate that microstructural control during fabrication
plays a key role in inﬂuencing the tribo-mechanical properties of
The microstructural evolution during FAST/SPS can be
controlled by altering sintering parameters like temperature,
pressure, holding period, rate of heating etc. For example,
Yong Zheng et al.  have shown changes in the thickness of
0272-8842/&2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Corresponding author. Tel.: þ91 40 2445 2411; fax: þ91 40 2444 2699.
E-mail address: firstname.lastname@example.org (J. Joardar).
the rim phase in TiC–TiN–WC–VC–Mo–Ni based cermets
consolidated by SPS at different temperatures and soaking
periods. Ping et al. , on the other hand, have reported
signiﬁcant inﬂuence of isothermal holding time and tempera-
ture during SPS on the degree of densiﬁcation as well as
hardness and transverse rupture strength of ultraﬁne TiC–TiN–
C–Ni cermets. Similarly, there are a few reports on TiC–
aluminide and WC–aluminide based composites processed by
FAST, which have shown improved properties [24,25,32].In
most of these studies on TiC- and WC- aluminide composites,
sintering by FAST/SPS was carried out only at certain set
of sintering parameters comprising of a particular tempera-
ture and pressure [25,32,33]. Alvarez et al.  have shown
changes in the density, shrinkage rate, secondary carbide
fraction, hardness and fracture toughness with variation in
the sintering temperature during FAST processing of TiCN–
composites. Not much information is avail-
able on possible inﬂuence of applied pressure during FAST/
SPS on the microstructure and mechanical properties of TiC/Ti
(C,N)/WC–aluminide based composites.
Here we report our ﬁndings on the role of applied pressure
on the microstructural development during FAST of Ti(C,N)–
WC–FeAl based nanocomposites. The concomitant changes in
the mechanical properties of the as-sintered nanocomposite
specimens are also discussed. To the best of our knowledge,
this is the ﬁrst such study carried out on Ti(C,N)–WC-based
nanocomposites with B2–FeAl binder.
The hard powder materials used in this study were Ti
) (average particle size of 100 nm; MK-Nano; 97%
purity) and WC (average particle size of 50 nm; MK-Nano;
99.5% purity). The FeAl binder powder (average particle size
of 100 mm) was synthesized from an elemental powder blend
of Fe (200 mesh; 99% purity; Alfa Aesar) and Al (325
mesh; 99.5% purity; Alfa Aesar) in appropriate proportion so
as to obtain the composition Fe–40at%Al. Blending was
performed for 24 h in a low-energy conventional horizontal
ball mill under toluene bath at a ball-to-powder weight ratio
(BPR) of 5:1. The powder blend was subsequently dried and
loaded in a horizontal tubular furnace and annealed under
Torr) at 950 1C for 2 h. Fig. 1 shows the x-ray
diffraction (XRD) proﬁles of the starting powders used for the
development of Ti(C,N)-based nanocomposites. The presence
of super-lattice (1 0 0) diffraction peak in the diffractogram
(Fig. 1) indicates B2-ordered structure of the as-synthesized
TiCN-based nanocomposite powders were prepared by blending
82 wt% of hard phase (TiCNþWC) with 18 wt% of as-
synthesized FeAl in a low-energy conventional horizontal ball
mill for 24 h under toluene bath at a BPR of 5:1. The powder
blend was dried, sieved and sintered using a FAST system (Dr.
Sinter SCM 1050; Sumitomo Coal and Mining) at 1500 1C.
Sintering was carried out at three different applied pressures for an
20 40 60 80
Fig. 1. XRD patterns of the starting powders.
Composition and sintering conditions of Ti(C,N)-based nanocomposites.
(TiCNþWC)–FeAl T-30 80–20 1500 30 10
(TiCNþWC)–FeAl T-50 80–20 1500 50 10
(TiCNþWC)–FeAl T-100 80–20 1500 100 10
40 60 80 100
Fig. 2. XRD patterns for Ti(C, N)-based nanocomposites by ﬁeld assisted
sintering as a function of pressure.
M.S. Archana et al. / Ceramics International 41 (2015) 1986–1993 1987
identical holding period of 10 min (Ref. Table 1 for sample
details). The sample temperature during sintering was measured
with a R-type thermocouple embedded in the graphite die and
located 3 mm away from the sample surface. Pressure during
sintering was applied from the beginning and maintained till the
end of the isothermal holding period. Sintering was carried out
under vacuum and the maximum temperature was achieved at a
heating rate of 100 1C/min.
The sintered samples were polished following standard
metallographic procedures for further studies. Phase fractions
in the binder and sintered samples are analysed by x-ray
diffraction (XRD: BRUKER D8 Advance model) with Cu Kα
radiation in step scan mode. Nelson–Riley method  was
used to calculate the precise lattice parameter of the starting
powders as well as of the sintered samples. Microstructural
characterization was carried out using ﬁeld emission scan-
ning electron microscopy (FESEM: Hitachi-S4300SE operat-
ing at 20 kV) in backscattered electron (BSE) mode. The
compositional analysis of the grains was conducted using an
energy dispersive x-ray spectroscopy (EDS) system attached to
the FESEM. Core size data was ascertained from BSE
micrographs by image analysis technique using ImageJ soft-
ware. The size analysis data was based on measurements on at
least 150 grains. In order to ascertain the inﬂuence of
microstructural variation on the mechanical properties of the
cermets, Vicker’s hardness and indentation fracture toughness
of the Ti(C,N) nanocomposites were determined. Vicker’s
hardness of the samples was estimated using a LECO Macro-
hardness tester with a load of 3 kg applied for 15 s. The
indentation fracture toughness (K
) of the composites was
evaluated from the indentation measurements following the
empirical relation :
KIC ¼0:03 P
where P,land adenote the load, crack length and diagonal
length of the indentation, respectively.
Intensity (arb. units)
35 36 41 42 43 43.5 44.0 44.5 45.0 45.5
Intensity (arb. units)
Fig. 3. Magniﬁed section of XRD the patterns indicating (a) peak shift in (1 1 1) and (2 0 0) peaks of Ti(C, N) and (b) absence of (1 1 0) peak of binder (B2–FeAl)
at high pressure of 100 MPa in Ti(C, N)-based nanocomposite.
Calculated lattice parameters of Ti(C,N) and FeAl binder phases in the precursor and as sintered nanocomposites.
Sample Lattice parameter (Å)
TiCN Binder-FeAl 110
111 200 220 311 222 400
Precursor 4.2787 4.2789 4.2762 4.2770 4.2773 –2.8903
T-30 4.2842 4.2873 4.2888 4.2877 4.2876 4.2879 2.8732
T-50 4.2883 4.2896 4.2914 4.2885 4.2897 4.2896 2.8747
T-100 4.2650 4.2706 4.2779 4.2784 4.2794 4.2795 –
M.S. Archana et al. / Ceramics International 41 (2015) 1986–19931988
3. Results and discussion
3.1. Phase analysis
Fig. 2 shows the XRD patterns of Ti(C,N)-based nanocom-
posites obtained by FAST at 1500 1C under different applied
pressures. The presence of FeAl binder phase is evident from
the XRD patterns of the samples sintered at lower pressures of
30 and 50 MPa. Magniﬁed sections of the XRD patterns for
Ti(C,N) and FeAl peaks are shown in Fig. 3. The ﬁgure reveals
signiﬁcant difference in the extent of peak shifts in case of
Ti(C,N) and FeAl. The Ti(C,N) peak in Fig. 3(a) shows a shift
towards lower 2θvalues for T-30 and T-50, indicating an
increase in the lattice parameter or in other words, lattice
expansion (Table 2). The higher applied pressure in case of
the T-100 sample, on the other hand, results in a shift of the
Ti(C,N) peak towards a higher 2θvalue indicating a decrease
in lattice parameter, thereby depicting lattice contraction.
Usually, the (Ti,W)(C,N) solid solution phase is expected to
form during sintering, which leads to complete elimination of
WC peaks in the XRD proﬁle [12,31]. Since W has a higher
atomic radius compared to Ti, formation of a substitutional
solid solution of W in Ti(C,N) is likely to introduce lattice
expansion or an increase in the lattice parameter, which is
evident from the observed shift in the Ti(C,N) peak position
towards lower 2θvalue (Fig. 3a).
Prior studies have reported complete dissolution of WC in
Ti(C,N) due to the shorter diffusion lengths when nanopowders
are used [37,38]. In the present study, it appears that at higher
applied pressure of 100 MPa, complete dissolution of Fe has
also taken place (in addition to W), as reﬂected by the absence
of FeAl peak in the XRD proﬁle in Fig. 3(b) otherwise noted at
lower applied pressures of 30 and 50 MPa (Fig. 3(b)). Since Fe
has a substantially lower atomic radius when compared to Ti, it
is likely to introduce signiﬁcant lattice contraction and drop in
lattice parameter, which overcompensates for any peak shift
caused by dissolution of W. Fig. 3bandTable 2 also show shift
in the (1 1 0) diffraction peak of FeAl binder phase towards
higher 2θvalue, indicating lattice contraction or a drop in lattice
parameter with applied pressure. Lattice contraction of FeAl
may be due to a signiﬁcant concentration of vacancies upon
cooling from high temperature .
3.2. Microstructural evolution
The microstructures of the TiCN–WC–FeAl based nanocom-
posites after FAST at 1500 1CareshowninFig. 4.The
micrographs clearly illustrate the role of applied pressure.
Undissolved or partially dissolved Ti(C,N) cores with size of
Fig. 4. FESEM-back scattered electron image of Ti(C, N)-based nanocompo-
sites (a) T-30, (b) T-50 and (c) T-100, as a function of pressure (White and
Black arrow indicates Ti(C, N)-rich and WC-rich core, respectively).
Metal content of different phases in Ti(C,N)-based nanocomposites processed
Sample Region Content (wt%)
Ti W Fe Al
T-30 Black core 79.6 17.9 2.5 –
White core 69.5 26.9 3.5 –
Grey rim 74.9 21.5 1.58 1.8
T-50 Grey core 75.3 21.8 2.8 –
White core 67.9 30.1 3.1 1.8
Light grey 75.1 25.8 3.4 2.4
T-100 Grey core 74.5 22.2 1.3 2
Light grey 72.4 20.6 2.9 3.9
M.S. Archana et al. / Ceramics International 41 (2015) 1986–1993 1989
observed in the T-30 samples corresponding to lowest applied
pressure during sintering (Fig. 4a). Most of the grains in T-30
sample have a typical “black core-grey rim”structure as usually
observed in conventional grade Ti(C,N) cermets [40,41].Itcould
be established from elemental analysis using EDS (plots not
shown) that the black core consists of Ti(C,N) with small amount
of dissolved W and Fe while the grey rim fraction could be
correlated to (Ti,W,Fe)(C,N) solid solution with W and Fe
content higher than the W level in the black core. Table 3 shows
the metal contents in Ti(C,N)-based nanocomposites as estimated
by FESEM–EDAX analysis. It is evident from Table 3 that W
level in grey rim is higher than that in the black core. In addition
to “black core-grey rim”,some“white core-grey rim”grains are
also present, albeit in small fraction. White core could be
correlated to (Ti,W,Fe)(C,N) solid solution with higher W content
when compared to the black core and grey rim (Table 3). Such
dissolution is possibly promoted by enhanced diffusion in the
presence of nanostructured grains with substantially higher frac-
tion of interface area when compared to coarse grade grains .
It may be noted that white core-grey rim structure has also been
observed in composites with excess secondary carbide additives
or ultraﬁne/nano Ti(C,N)—based cermets [38,40–42]. Zheng
et al.  andFengetal. have attributed such white core-
grey rim structure formation to fast dissolution of WC in
Ti(C,N) during solid state sintering followed by precipitation of
W-lean (Ti,W)(C,N) around the as formed W-rich (Ti,W)(C,N)
during the later stage of liquid phase sintering. As Ti(C,N)-based
cermet in this study is a nanocomposite a similar explanation can
be considered for the formation of white cores observed in
Fig. 4a. Presence of a thin light-grey region between the grains is
also evident in Fig. 4a. From EDS it is revealed that the light-grey
regions between the grains correspond to FeAl-rich binder phase.
Image analysis of the SEM micrographs revealed a core size of
150–400 nm for the black core of Ti(C,N) with dissolved W and
Fe while the size range of 270–500 nm was observed for the
white core comprising W-enriched (Ti,W,Fe)(C,N).
On increasing the applied pressure during FAST to 50 MPa,
the black core seems to disappear completely and transforms into
grey core with no trace of any rim (Fig. 4b). This indicates that
substantial dissolution of W and Fe in Ti(C,N) has taken place
with only marginal increase in the applied pressure to 50 MPa.
This is also supported by the decrease in lattice parameter values
of T-50 sample (Table 2). Small fraction of white core without
any rim phase is also observed in the T-50 sample. The particle
that appears black in the image was identiﬁed as Al-rich oxide
from FESEM–EDAX analysis. Similar oxides have been
reported in other Al containing Ti(C,N) cermets .
At higher applied pressure of 100 MPa (T-100) sample,
additional differences in the microstructure are quite apparent.
The microstructure in Fig. 4c shows the presence of grey core
surrounded by thin white rim with no distinct binder phase.
The thin white rim around the grey core is expected to have
excess precipitation of W-rich solid solution of (Ti,W,Fe,Al)
(C,N) on (Ti,W)(C,N) grains. Black areas identiﬁed as Al-rich
oxide is also observed.
The core size distribution also shows changes with applied
pressure. The size distribution of cores in the TiCN–WC–FeAl
nanocomposites, as determined from the BSE micrographs is
shown in Fig. 5. The core (black) size in T-30 is in the range of
150–400 nm with mean size of 250 nm. With an increase in
applied pressure to 50 MPa, there is not much variation in the
mean size of core (grey) 270 nm with maximum size ranging
up to 350 nm. At higher pressure of 100 MPa, even though no
further change in the mean core size is detectable, some of the
core shows signiﬁcant reﬁnement to 150–320 nm. Reﬁnement
Core size (
Core size (
0.0 0.1 0.2 0.3 0.4 0.5
0.0 0.1 0.2 0.3 0.4 0.5
0.0 0.1 0.2 0.3 0.4 0.5
Core size (
Fig. 5. Particle size distribution of the Ti(C, N)-rich core in Ti(C, N)-based
nanocomposites as a function of pressure.
M.S. Archana et al. / Ceramics International 41 (2015) 1986–19931990
of core size indicates increased dissolution of WC with
increase in pressure during liquid phase sintering that pre-
cipitates as white rim (Fig. 4c) and also dissolution of Fe and
Al in Ti(C,N) at high pressure.
3.3. Mechanical properties
values of different nanocomposites studied in
this work are tabulated in Table 4. To facilitate comparison,
hardness and fracture toughness values from previous reports
in the literature are also included. It is clearly evident from the
table that even though the sample processed at high pressure
(100 MPa) shows higher hardness compared to T-30 and T-50,
value is marginally inferior. The difference in fracture
toughness can be explained based on the typical intergranular
crack-path as observed in the nanocomposites after Vicker’s
indentation measurement (Fig. 6). Referring to the microstruc-
tures in Fig. 4, it is apparent that the presence of a thin region
of FeAl binder phase between the hard phase grains (which
also coincides with the intergranular crack-path), in case of
T-30 and also to some extent in T-50, has led to marginal
toughening of the composite compared to T-100. It has already
been mentioned that T-100 is devoid of such an intergranular
FeAl layer as all of its Fe and Al content diffuses into Ti(C,N)
at higher pressure. The absence of the binder phase in the
intergranular region leads to poor fracture toughness in T-100.
Similarly the presence of complex solid solution phase (Ti,W,
Fe,Al)(C,N) due to complete dissolution of binder phase expl-
ains the increase in hardness of T-100 sample.
Table 4 also suggests that the hardness and toughness val-
ues of the TiCN–WC–FeAl nanocomposites are comparable
with those of TiCN–Ni/Co based composites reported in the
literature [5,44,45,28]. Of course, higher hardness and fracture
toughness values have been reported for (Ti,W)C–FeAl
based composites [24,25]. However, it should be emphasized
binders in these composites have low melting
points (Table 4) and consequently, their stability at high
temperature is difﬁcult. In this respect, B2-ordered-FeAl with
higher melting point when compared to FeAl
superior oxidation resistance when compared to Co and Ni
is expected to perform better as a binder in Ti(C,N)-based
Different applied pressures during FAST were found to
inﬂuence phase and microstructural evolution in the Ti(C,N)-
based nanocomposites. Microstructural variations like type and
size of cores in the core-rim structure as well as changes in the
hardness and fracture toughness values as a function of applied
Comparison of hardness and fracture toughness values with previously reported values of Ti(C,N)-based nanocomposites processed by SPS/FAST.
Nanocomposite Size (mm) Sintering parameters Core size (mm) Hv (GPa) K
HP-(wt%)BP BP-M.P. (1C) Temp. (1C) Press. (MPa)
TiCN–WC–18FeAl 1400 0.1–0.05–100 1500 30 0.02 17.6 6.8 9
50 0.04 17.5 6.6
100 0.05 18.4 5.9
1160 0.02–0.02 1300(p) 80 –24 10.5 
854 –1300(p) 80 –25.5 8.5 
C–26Ni 1455 1.5–3.3–4
1370(p) 100 0.1 15.7 6.7 
1383(p) 100 0.06 18 7.5 
1350(p) 100 0.1 13.8 6.4 
TiCN–Ta–5Ni 1455 5–25–50
1500 30 –17.8 6.9 
(Ti,Ta)C–20Co 1495 45
1350(p) 30 1.6 17.4 3.2 
1350(p) 30 1.06 16.1 4.5
1150(p) 30 1.23 14.1 5.6
TiCN–15Co 1495 45–150
1300 80 4 17.1 5.5 
1400 80 5 15.2 7.2
C–12 Ni 1455 0.2–0.2–2.8–0.2 1250 30 0.4 16.7 
HP—Hard phase, BP—Binder phase, (p)—pyrometer reading.
-Size before milling.
Fig. 6. Indentation induced crack path in Ti(C, N)-based nanocomposites
sintered at 1500 1C under 100 MPa. Crack trajectory from right to left. White
arrow indicates bridging of crack by binder phase. Black arrow indicates inter-
M.S. Archana et al. / Ceramics International 41 (2015) 1986–1993 1991
pressure were noted. Highest hardness of 18.4 GPa was
observed in samples sintered under higher applied pressure of
100 MPa whereas highest fracture toughness of 6.8 MPa√m
was observed in samples sintered under lower applied pressure
of 30 MPa. The hardness and fracture toughness values of the
as sintered Ti(C,N) nanocomposites with FeAl binder were
comparable to Ti(C,N) based composites with Ni/Co metallic
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