Advanced Powder Technology 25 (2014) 1754–1760
Contents lists available at ScienceDirect
Advanced Powder Technology
journal homepage: www.elsevier.com/locate/apt
Original Research Paper
Formation mechanism of B4C–TiB2–TiC ceramic composite
produced by mechanical alloying of Ti–B4C powders
M. Raﬁei ⇑, M. Salehi, M. Shamanian
Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
a r t i c l e
i n f o
Article history:
Received 1 February 2014
Received in revised form 28 May 2014
Accepted 3 July 2014
Available online 14 July 2014
Keywords:
Mechanical alloying
Ceramic composites
Thermodynamic analysis
Reaction mechanism
a b s t r a c t
In this study, the B4C–TiB2–TiC composite powder was synthesized by mechanical alloying (MA) of Ti–
B4C powder mixture. For this purpose, four powder mixtures of Ti and B4C powders with different molar
ratios were milled. In order to study the mechanism of Ti–B4C reaction during milling, structural changes
and thermal analysis of powder particles were studied by X-ray diffractometry (XRD) and differential
thermal analysis (DTA). Morphology and microstructure of powder particles during milling were studied
by scanning electron microscopy (SEM). It was found that during MA, after decomposition of the outer
layers of B4C particles, ﬁrst, C reacted with Ti and after that, B was diffused in Ti structure and TiC and
TiB2 phases were formed in gradual reaction mode. Also, the results of DTA and thermodynamic analysis
conﬁrmed the suggested mechanism for Ti–B4C reaction.
Ó 2014 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder
Technology Japan. All rights reserved.
1. Introduction
Boron carbide (B4C) is the third hardest material after diamond
and boron nitride at room temperature and its hardness is
retained at high temperatures [1,2]. B4C possesses unique physical and thermal properties such as high elastic modulus, high
melting point, good chemical stability and high neutron absorption cross section. In addition, B4C ceramics exhibit excellent
mechanical properties such as high hardness and high wear
resistance [3–7].
This material has also received attention for application in
nuclear fusion reactors because the low atomic numbers of boron
and carbon result in low X-ray absorption, and exposure to neutron
irradiation does not cause radioactive decay [8–11]. B4C is currently used in nuclear energy and high temperature thermoelectric
conversion. However, their widespread application has been
restricted mainly due to their low strength and fracture toughness
as well as poor sinterability due to a low self diffusion coefﬁcient.
Several additives have been investigated in order to promote the
sinterability and mechanical properties of B4C [3].
B4C ceramics with dispersed TiB2 particles have been investigated in order to improve both strength and toughness [12–15].
The improvement of fracture toughness, which was achieved by
TiB2 addition, has been explained in terms of microcrack formation
caused by the thermal expansion mismatch between dispersed
⇑ Corresponding author. Tel./fax: +98 312 5219109.
E-mail address: raﬁ[email protected] (M. Raﬁei).
particles and the matrix [13,14]. TiB2 has been proposed as a promising second phase for the improvement of B4C properties [15–18].
Pairing B4C and TiB2 in a composite can be beneﬁcial for the design
of thermoelectric materials [19,20] and composites with increased
fracture toughness and bending strength relative to pure boron
carbide [3,21]. Also, TiB2 was suggested as a possible grain growth
inhibitor for B4C [13].
In previous studies [22–24], some B4C based composites, e.g.
B4C/TiB2, B4C/TiB2/MB2, B4C/MB2, B4C/Al, B4C/SiC, have been
developed. Yamada et al. [3] fabricated B4C–20 mol.% TiB2 ceramic composites by reacting hot-pressing powder mixtures of
B4C, TiO2 and carbon black at 2000 °C. They reported that the
abnormal grain growth of B4C was inhibited, and the fracture
toughness was increased for some B4C–TiB2 specimens. It has
been reported that this improvement in toughness was due to
the formation of microcracks and the deﬂection of propagating
cracks caused by the thermal expansion mismatch between TiB2
particles and B4C matrix.
Wang et al. [25] synthesized dense TiB2/TiC composites by MA
of Ti and B4C powder blends and subsequent sintering. They indicated that diffusion played a major role in the formation of the TiC
and TiB2 phases. Also, TiC phase was formed earlier than TiB2
phase, due to the higher diffusivity of C than B into the Ti matrix.
Li et al. [26] prepared TiB2/TiC nanocomposite via MA of Ti and
B4C powders. It was found that the bulk of TiC and TiB2 was formed
abruptly by self-propagating reaction after 5 h of milling. The ﬁnal
product was composed of nano-sized TiC and micro-sized TiB2
particles.
http://dx.doi.org/10.1016/j.apt.2014.07.003
0921-8831/Ó 2014 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder Technology Japan. All rights reserved.
M. Raﬁei et al. / Advanced Powder Technology 25 (2014) 1754–1760
1755
Wang et al. [27] successfully prepared the TiB2–TiC nanocomposite by MA of Ti and B4C powder mixtures. They showed that
the formation of TiC and TiB2 was gradual. Also it was found that
the decomposition of B4C into amorphous B and C occurred during
MA.
Preliminary estimates of TiC–TiB2 composites fabricated from Ti
and B4C powders suggested that a 75% cost savings could be realized using B4C as a reactant instead of elemental B and C [28].
Therefore, in this work, the MA of 3Ti + xB4C (x = 1, 2, 3, 4) powder
mixtures and the formation mechanism of B4C–TiC–TiB2 composite
powder were studied.
2. Experimental
Ti and B4C powders, with the purity of 99.5 and 99%, respectively, were used as starting materials. The Ti powder particles
were irregular in shape with a size distribution of 50–100 lm
and B4C particles had a uniform angular shape with a size distribution of 150–200 lm. Ti and B4C powders were mixed with different
molar ratios according to the reactions (1)–(4). MA was carried out
in a high energy planetary ball mill, nominally at room temperature and under Ar atmosphere. The ball and vial materials were
hardened chromium steel. Ball to powder weight ratio and rotation
speed of vial were 10:1 and 300 rpm, respectively. The total powder mass was 20 g. Samples were taken at selected time intervals
and characterized by XRD in a Philips X’ PERT MPD diffractometer
using ﬁltered Cu Ka radiation (a = 0.1542 nm). Morphology and
microstructure of powder particles were characterized by SEM in
a Philips XL30. The mean powder particle size was estimated from
SEM images of powder particles by image tool software. Isothermal
annealing was carried out to study the thermal behavior of milled
powders. A small amount of powder was sealed and isothermally
annealed at 1200 °C for 1 h in a conventional furnace and then
cooled in air. The structural transitions occurring during annealing
were determined by XRD. DTA was done in 30–1200 °C temperature range with 20 °C min1 heating rate in order to study the thermal behavior of the milled powders.
3. Results and discussion
3.1. Structural changes
In order to study the formation mechanism of B4C–TiB2–TiC
ceramic composite, MA of Ti–B4C systems with different molar
ratios was done according to the following reactions:
3Ti þ B4 C ¼ 2TiB2 þ TiC
ð1Þ
3Ti þ 2B4 C ¼ B4 C þ 2TiB2 þ TiC
ð2Þ
3Ti þ 3B4 C ¼ 2B4 C þ 2TiB2 þ TiC
ð3Þ
3Ti þ 4B4 C ¼ 3B4 C þ 2TiB2 þ TiC
ð4Þ
3.1.1. MA of 3Ti–B4C system
Fig. 1 shows XRD patterns of 3Ti–B4C powder mixture at different milling times (according to the reaction (1)). As can be seen,
after 5 min of MA, only Ti and B4C peaks were seen. With increasing milling time up to 5 h, traces of TiC peaks appeared on XRD pattern. Also, one peak related to TiB phase was seen. After 10 h of
milling, the intensity of TiC peaks was increased due to the gradual
diffusion of C into the Ti structure to form TiC phase. Also, TiB2
phase was formed and the intensity of its peaks was increased
because of the gradual formation of it during MA. Wang et al.
[25] synthesized dense TiB2/TiC composites by MA and the subse-
Fig. 1. XRD patterns of 3Ti–B4C system at different milling times.
quent pressureless sintering and reported that the formation of TiC
and TiB2 compounds during MA was gradual in contrast to the conventional synthesis via the SHS reaction. At longer milling times,
with more diffusion of B into the Ti structure due to the MA process, TiB phase was disappeared and transformed to TiB2. After
30 h of MA, only TiB2 and TiC peaks were seen, indicating the complete formation of TiB2–TiC composite.
XRD patterns of 3Ti–B4C system after 30 h of MA and subsequent heat treatment are presented in Fig. 2. As can be seen, with
heat treatment, the intensity of TiC and TiB2 peaks was increased
due to the ordering of structure and strain removal. Also, Fe2B
phase was formed after heat treatment. The cause of Fe2B formation is related to the presence of Fe in powder mixture due to
the wear of steel bowl and balls with hard ceramic powders such
as B4C, TiC and TiB2 and the reaction of Fe with B during heating.
3.1.2. MA of 3Ti–2B4C system
XRD patterns of 3Ti–2B4C system at different milling times are
presented in Fig. 3. According to the reaction (2), it is expected that
MA of this system can lead to the formation of B4C–TiB2–TiC ceramic composite powder. As can be seen, after 5 min of MA, just Ti
and B4C peaks are seen. With increasing the milling time up to
2 h, the height of Ti and B4C diffraction peaks was decreased and
simultaneously, the traces of TiC peaks were seen in XRD pattern
due to the gradual reaction of Ti with B4C during MA. With further
Fig. 2. XRD patterns of 3Ti–B4C system, (a) after 30 h of milling and (b) after heat
treatment at 1200 °C for 1 h.
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M. Raﬁei et al. / Advanced Powder Technology 25 (2014) 1754–1760
The stages of B4C–TiB2–TiC ceramic composite powder formation in this system have been shown in Fig. 4. As can be seen, during MA process and at the early stages of milling, a thin layer of Ti
powder surrounds the hard B4C particles. Decomposition of B4C
particles at surface leads to the formation of free C and B. These
free C and B are not visible in XRD pattern. Wang et al. [27]
reported that this free C and B had an amorphous structure. C diffuses into the Ti structure because of the high density of interfaces
and the crystalline defects from MA process and TiC phase forms.
Then, B diffuses into the Ti structure and TiB2 phase forms and a
hard layer of TiB2 and TiC surrounds the B4C particles. After that,
this hard layer around B4C particles is broken and the surface layer
of B4C particles is surrounded by Ti layer. These stages are repeated
again.
Fig. 5 shows XRD patterns of 3Ti–2B4C system after 30 h MA
and heat treatment at 1200 °C for 1 h. As can be seen, after heat
treatment the intensity of TiB2 peaks was increased and their
width was decreased due to the relaxation and the removal of
residual stresses induced by MA process. As previously mentioned,
the intensity of TiC peaks is decreased with increasing the milling
time at longer milling times. After heat treatment of 30 h milled
powder, TiC peaks were sharp as a result of TiC particle growth
and the ordering of its structure. Also, the traces of Fe2B phase
are seen in heat treatment XRD pattern.
Fig. 3. XRD patterns of 3Ti–2B4C system at different milling times.
milling (5 h), TiB2 phase was formed. As can be seen, the formation
of this phase during MA Process occurred after TiC formation, similar to the previous system. MA for 10 h led to the complete disappearance of Ti peaks, especially (1 0 1) peak, and the gradual
increase in the intensity of TiC and TiB2 peaks, indicating that the
formation of TiC and TiB2 phases was in the gradual reaction mode
during MA. At longer milling times, the intensity of TiB2 peaks was
gradually increased due to the further formation of this phase with
increasing time. However, the intensity of TiC peaks, especially
(1 1 1) and (2 0 0) peaks, was decreased with further milling. This
can be attributed to the following reasons: (1) creating ﬁner TiC
particles at longer milling times as a result of the collision of TiC
particles with steel balls, and the existing B4C and TiB2 harder particles. Li et al. [26] reported that during MA of Ti and B4C powders,
the ﬁnal product was composed of nano-sized TiC and micro-sized
TiB2 particles due to the higher hardness of TiB2 as compared with
TiC, and a large number of TiB2 particles acted as small milling
balls in further milling process; (2) decreasing the crystalline order
of TiC phase as a result of mechanical impacts on its structure.
Finally, after 30 h of MA, only B4C, TiB2 and TiC peaks were seen,
indicating the complete formation of ceramic composite powder.
According to the above discussion, it is concluded that at early
stages of milling, C is diffused into the Ti structure and TiC phase
forms. The formation of this phase occurs before TiB2 formation
during MA process, because of the faster diffusion of C in Ti structure when compared with B in Ti structure [26]. Then, with the diffusion of B in Ti structure at longer milling times, TiB2 phase is
formed. The reaction path of this system can be summarized as
follows:
(1) Decomposition of surface layer of B4C particles to B and C:
B4 C ¼ 4B þ C
ð5Þ
(2) Diffusion of C in Ti structure and the formation of TiC:
Ti þ C ¼ TiC
ð6Þ
(3) Diffusion of B in Ti structure and the formation of TiB2:
Ti þ 2B ¼ TiB2
ð7Þ
3.1.3. MA of 3Ti–3B4C system
XRD patterns of 3Ti–3B4C powder mixture are shown in Fig. 6.
As can be seen, after 2 h of MA in this powder mixture, traces of
TiC peaks are shown on XRD pattern. With increasing the milling
time up to 5 h, the intensity of these peaks was relatively increased
and with further milling up to 30 h, these peaks were disappeared.
As previously mentioned, this is due to the formation of very ﬁne
TiC particles and also, the disordering of its structure during MA
process. Due to the lower volume percentage of TiC in this system
(15 vol.%), the intensity of TiC diffraction peaks was much lower
than that previously observed for 3Ti–2B4C system (20 vol.%).
Traces of TiB2 diffraction peaks were seen after 5 h of MA. With
increasing the milling time up to 30 h, their intensity was gradually increased, indicating the gradual reaction between Ti and B.
In this system, because of the higher volume percentage of B4C particles (50 vol.%), as compared with the previous system (35
vol.%), B4C peaks were clearer after 30 h of MA.
Fig. 7 presents XRD patterns of 3Ti–3B4C powder mixture after
30 h of MA and subsequent heat treatment at 1200 °C for 1 h. As
can be seen, after heat treatment of this powder mixture, sharp diffraction peaks of TiB2 and small peaks of TiC phases were seen on
XRD pattern. This conﬁrms that the disappearance of TiC peaks
after 30 h of MA was due to the very ﬁne particle size and its disordered structure during milling. After heat treatment, because of
particle growth and the ordering of structure, TiC peaks were
appeared on XRD pattern. Also, due to the entry of Fe contamination in powder mixture from steel bowl and balls during MA, FeB
phase was formed after heat treatment.
3.1.4. MA of 3Ti–4B4C system
Fig. 8 presents XRD patterns of 3Ti–4B4C after 30 h MA and heat
treatment at 1200 °C for 1 h. As can be seen, after 30 h of MA of this
powder mixture, no reaction was observed and only Ti and B4C
peaks were seen. However, in the previous systems, the reaction
of composite formation was completely occurred after 30 h of
MA. In this system, because of the high contents of B4C in the initial
powder mixture that acted as diluents, according to the reaction
(4), there was no reaction between Ti and B4C during milling and
MA only active powders. After heat treatment of this powder mixture, TiB2 and FeB phase were observed on XRD pattern and TiC
peaks were not seen. Gordienko [29] studied the thermodynamic
M. Raﬁei et al. / Advanced Powder Technology 25 (2014) 1754–1760
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Fig. 4. Schematic of reaction mechanism between Ti and B4C.
Fig. 5. XRD patterns of 3Ti–2B4C system, (a) after 30 h of milling and (b) after heat
treatment at 1200 °C for 1 h.
Fig. 8. XRD patterns of 3Ti–4B4C system, (a) after 30 h of milling and (b) after heat
treatment at 1200 °C for 1 h.
analysis of Ti–B4C reaction with different molar ratios and reported
that when the percentage of B4C was more that 37 wt.% TiC was
not formed.
3.2. Thermodynamic aspects
Fig. 6. XRD patterns of 3Ti–3B4C system at different milling times.
Fig. 7. XRD patterns of 3Ti–3B4C system, (a) after 30 h of milling and (b) after heat
treatment at 1200 °C for 1 h.
DG0298 and DH0298 values for reaction (1) were estimated using
thermodynamic data [30]. They were about 733.6 and
744.8 kJ mol1 respectively, indicating that this reaction can thermodynamically occur at room temperature and is highly exothermic.
The adiabatic temperature (Tad) is often used as a criterion for the
anticipation of modality of reactions [31]. Tad is the highest temperature that the reaction system can reach, provided that all the
released heat is used to increase the temperature of system after
the initiation of reaction [31]. It has been empirically demonstrated
that the reaction will be self-sustaining only if Tad P 1800 K [32]. Tad
values were estimated for reactions (1)–(4). Fig. 9 shows Tad changes
versus B4C content in the reactants. As can be seen in reaction (1), Tad
was about 3445 K, which was decreased with increasing B4C content
in reactants. This trend was due to the excess B4C that acted as diluents and therefore, Tad was decreased with increasing B4C content
in the reactants. According to Fig. 9 (blue line)1, it is concluded that
reactions (1)–(3) occur in combustive mode and reaction (4)
(Tad = 1700 K) occurs in the gradual mode. On the other hand, as previously discussed, all four reactions occur in the gradual mode during milling. By considering the suggested mechanism in this work for
the reaction between Ti and B4C (Fig. 4), because the reaction of Ti
with C occurred sooner than the reaction of Ti with B, Tad was also
calculated for the following reactions:
3Ti þ B4 C ¼ TiC þ 2Ti þ 4B
ð8Þ
3Ti þ 2B4 C ¼ B4 C þ TiC þ 2Ti þ 4B
ð9Þ
3Ti þ 3B4 C ¼ 2B4 C þ TiC þ 2Ti þ 4B
ð10Þ
1
For interpretation of color in Fig. 9, the reader is referred to the web version of
this article.
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M. Raﬁei et al. / Advanced Powder Technology 25 (2014) 1754–1760
gradual mode. This prediction is in agreement with the obtained
XRD results by considering Tad, and conﬁrms the validity of the
suggested mechanism for Ti–B4C reaction.
3.3. Microstructural and morphological observations
Fig. 9. Calculated adiabatic temperature versus B4C content at different Ti–B4C
systems.
3Ti þ 4B4 C ¼ 3B4 C þ TiC þ 2Ti þ 4B
ð11Þ
The results are presented in Fig. 9 (green line). As can be seen,
by considering this mechanism, Tad for all four reactions is lower
than 1800 K (red line) and therefore, these reactions occur in the
Fig. 10 shows SEM images of powder particles of 3Ti–2B4C system at different milling times. After 5 min of MA (mixing of powders), ﬁne Ti particles (particle size mainly less than 100 lm) and
coarse B4C particles (particle size mainly in the range of 150–
200 lm) were seen. With increasing the milling time up to 5 h,
the mean powder particle size was decreased due to the predominance of the fracturing of powder particles over the cold welding
process and also the reaction that occurred. With further milling,
some coarse powder particles were seen. With more accurate
assessment, it was indicated that these coarse particles at longer
milling times were agglomerated of very ﬁne particles as the high
surface area tended to stick together. The mean powder particle
size after 30 h of milling was in the range of 500–1000 nm. At
the early stages of milling, the particles had an irregular shape with
large particle size distribution, while at longer milling times, the
powder particles had spherical morphology, indicating the effect
of MA on changing powder morphology.
Fig. 10. SEM images of powder particles of 3Ti–2B4C system at different milling times.
M. Raﬁei et al. / Advanced Powder Technology 25 (2014) 1754–1760
1759
Fig. 11. Cross section SEM images of powder particles of 3Ti–2B4C system at different milling times.
with the maximum at about 1067 °C was seen. This peak was
related to the Ti–B4C reaction that occurred in solid state during
heating. As previously mentioned, due to the low adiabatic temperature (950 K) and the occurrence of this reaction in solid state,
a broad peak was seen, indicating that this reaction occurred gradually during heating. The same DTA results were reported by Zhang
et al. [33] and Liang et al. [34] in heating 3Ti–B4C system. This is in
good agreement with the obtained XRD patterns for this system
(Fig. 3) and conﬁrms the suggested mechanism for Ti–B4C reaction
in Section 3.1.2.
4. Conclusions
MA of Ti–B4C powder mixture with different molar ratios was
carried out and the following results were obtained:
Fig. 12. DTA curve of 1 h milled powder of 3Ti–2B4C system.
Cross sectional SEM images of powder particles are presented in
Fig. 11. After 5 min of MA, inhomogeneous distribution of Ti and
B4C powders was seen. With increasing the milling time up to
2 h, the distribution of powders was more uniform and a lamellar
structure with high interfaces between Ti and B4C was formed that
increased the diffusion rate of C, which was released from B4C into
the Ti structure. Further milling (5 h) led to the formation of the
composite structure of Ti, B4C and TiC. As can be seen in Fig. 11
(5 h), they are with different gray levels. At longer milling times
(30 h), as seen in high magnitude image, many particles were
observed, each of which was the composite of B4C–TiC–TiB2 with
very ﬁne composite particles in the range of 70–1000 nm. Dark
areas in this image are B4C and bright areas are TiC and TiB2.
3.4. DTA analysis
DTA curve of 3Ti–2B4C system after 1 h MA (before occurring
reaction) is presented in Fig. 12. Only one exothermic broad Peak
(1) MA of 3Ti–B4C system led to the formation of TiB2 and TiC
composite powder after 30 h of milling. Also, MA of 3Ti–
2B4C and 3Ti–3B4C systems led to the formation of B4C–
TiB2–TiC composite powders. There was no reaction in MA
of 3Ti–4B4C after 30 h of milling and after subsequent heat
treatment, the reaction took place.
(2) It was found that at the early stages of milling, a thin layer of
Ti powder surrounded the hard B4C particles. Decomposition
of B4C particle at surface led to the formation of free C that
was diffused into the Ti structure and TiC phase was formed.
Then, B was diffused into the Ti structure and TiB2 phase was
formed and a hard layer of TiB2 and TiC surrounded the B4C
particles. After that, this hard layer around B4C particles was
broken and the surface layer of B4C particles was surrounded
by Ti layer and these stages were repeated again.
(3) Thermodynamic analysis revealed that all four reactions in
this work had low adiabatic temperature and took place in
the gradual mode. This was in good agreement with XRD
results and conﬁrmed the suggested mechanism for Ti–B4C
reaction.
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(4) DTA analysis of 3Ti–2B4C powder mixture after milling for
1 h conﬁrmed that the reaction was in the gradual mode
with a broad reaction peak.
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