Synthesis of lead zirconate titanate from an amorphous precursor by mechanical activation

Article abstract of DOI:10.1016/S0925-8388(00)00917-8

Many of the chemistry-based processing routes for functional ceramics inevitably involve calcining the chemical-derived precursors at an intermediate/high temperature, in order to form the designed ceramic phase. This is very undesirable, although widely used, as the calcination can result in an extensive degree of crystal growth and particle coarsening at the calcination temperature and therefore ruins almost all the advantages offered by the chemistry-based processing routes, such as an ultrafine particle size and high sintering-reactivity. Using a specifically designed PZT precursor prepared by co-precipitation, it is demonstrated that the precursor-to-ceramic conversion can alternatively be realized by mechanical activation. In this connection, a single phase, nanocrystalline perovskite PZT powder has been successfully derived from an amorphous hydroxide precursor by mechanical activation. The resulting PZT powder was well dispersed, and the particle size was in the range of 30-50 nm, as observed using the scanning electron microscopy and transmission electron microscopy. This is in contrast to the poor particle characteristics, represented by very coarse and irregular particle and agglomerate sizes, for the powder derived from calcination at 750 °C. The activation-triggered PZT powder was sintered to a density of 97.6% theoretical density at 1150 °C for 1 h. Sintered PZT ceramic exhibits a dielectric constant of 927 at room temperature and a peak dielectric constant of approximately 9100 at the Curie point of 380 °C when measured at the frequency of 1 kHz.

Full text of DOI:10.1016/S0925-8388(00)00917-8

Journal of Alloys and Compounds 308 (2000) 139–146
L
www.elsevier.com/locate/jallcom
Synthesis of lead zirconate titanate from an amorphous precursor by
mechanical activation
*
Xue Junmin, John Wang , Toh Weiseng
Department of Materials Science, Faculty of Science, National University of Singapore, Singapore 119260, Singapore
Received 24 March 2000; accepted 25 April 2000
Abstract
Many of the chemistry-based processing routes for functional ceramics inevitably involve calcining the chemical-derived precursors at
an intermediate/high temperature, in order to form the designed ceramic phase. This is very undesirable, although widely used, as the
calcination can result in an extensive degree of crystal growth and particle coarsening at the calcination temperature and therefore ruins
almost all the advantages offered by the chemistry-based processing routes, such as an ultrafine particle size and high sintering-reactivity.
Using a specifically designed PZT precursor prepared by co-precipitation, it is demonstrated that the precursor-to-ceramic conversion can
alternatively be realized by mechanical activation. In this connection, a single phase, nanocrystalline perovskite PZT powder has been
successfully derived from an amorphous hydroxide precursor by mechanical activation. The resulting PZT powder was well dispersed,
and the particle size was in the range of 30–50 nm, as observed using the scanning electron microscopy and transmission electron
microscopy. This is in contrast to the poor particle characteristics, represented by very coarse and irregular particle and agglomerate sizes,
for the powder derived from calcination at 7508C. The activation-triggered PZT powder was sintered to a density of 97.6% theoretical
density at 11508C for 1 h. Sintered PZT ceramic exhibits a dielectric constant of 927 at room temperature and a peak dielectric constant of
|9100 at the Curie point of 3808C when measured at the frequency of 1 kHz.
2000 Elsevier Science S.A. All rights reserved.
Keywords: Lead zirconate titanate; Nanopowders; Mechanical activation; Dielectric properties; XRD
1. Introduction
position of the resultant PZT often occurs as a result of the
undesirable loss in lead content through volatilization of
PbO at elevated temperatures [4]. Many chemistry-based
processing routes, such as the oxalate route [5], co-precipi-
tation [6], alkoxide hydrolysis [7] and hydrothermal re-
action [8], have thus been devised and employed to prepare
ultrafine and sintering-reactive PZT powders. However,
almost all of these chemistry-based processing routes
require the precursors be calcined at a temperature in the
range of 600–9008C, in order to develop the designed PZT
perovskite phase. This often results in a considerable
degree of particle coarsening and agglomeration and thus
ruins the sintering-reactive nature of chemical-derived
precursors. In fact, the presence of hard particle aggregates
adversely affects the sintering and microstructural develop-
ment of PZT at the sintering temperature when differential
sintering among aggregates of differing green density
occurs.
Lead zirconate titanate, Pb(Zr,Ti)O3 (PZT) of perov-
skite structure, which was found to be piezoelectric in the
1950s, is widely used in various sensing and actuating
devices [1]. They are often fabricated by first preparing a
fine and homogeneous PZT powder, followed by shape-
forming and sintering at an appropriate temperature to
achieve the required sintered density, microstructure and
electrical properties [2]. The traditional method of syn-
thesizing PZT powder is through solid-state reaction,
which involves mixing the constituent oxides (PbO, TiO2
and ZrO2) in stoichiometric ratio, followed by calcination
for phase formation at an elevated temperature (.7008C)
[3]. However, this often leads to an incomplete reaction,
and in particular, particle coarsening and agglomeration in
the resulting PZT powder as a result of the multiple
interfacial diffusions and reactions involved at the calcina-
tion temperature. Furthermore, nonstoichiometry in com-
Following the pioneer work of Benjamin [9], who
devised mechanical alloying for synthesizing novel
nanocrystalline metal and alloying powders, there has been
a surge of interest in mechanical activation over the past
*Corresponding author. Tel.: 165-874-2958; fax: 165-776-3604.
E-mail address: maswangj@nus.edu.sg (J. Wang).
0925-8388/00/$ – see front matter
PII: S0925-8388(00)00917-8
2000 Elsevier Science S.A. All rights reserved.

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X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
few years [10–14]. Very recently, mechanical activation
was successfully devised to synthesize several functional
ceramics of perovskite structure in the authors’ laboratory
[15]. Using the novel technique, the formation of desirable
perovskite phases from the constituent oxides was trig-
gered by mechanical energy, instead of one or more steps
of calcination at elevated temperatures. We have further
observed that the phase-forming reactions triggered by
mechanical activation of mixed oxides underwent two
basic steps: (i) a significant refinement in particle and
crystallite sizes and a degree of amorphization of mixed
oxides at the initial stage, and (ii) the nucleation and
subsequent growth of perovskite crystallites in the highly
activated matrix. Therefore, the phase-forming reactions
triggered by mechanical activation proceed via a complete-
ly different route from that in a conventional solid-state
reaction, where the process is controlled by one or more
interfacial reactions and diffusions at elevated temperature
[16].
In this work, we prepared a PZT precursor by co-
precipitating lead, zirconium and titanium hydroxides. The
precursor was deliberately treated at a relative low calcina-
tion temperature and therefore the nucleation and sub-
sequent growth of PZT crystallites were not provoked by
thermal activation. Therefore, a study can be made of the
effects of mechanical activation on the phase development,
aiming to investigate whether nanocrystallites of perov-
skite PZT can be triggered to form in an amorphous
precursor by mechanical activation. This is done by
comparing both the powder characteristics and sintering
behavior of the activation-derived PZT powder with those
of the PZT powder derived from conventional calcination
at a much higher temperature. The activation-triggered
PZT powder was further studied by sintering at various
temperatures, and the dielectric properties of sintered PZT
were characterized.
sequently titrated to an ammonia solution (2 M) to form
Zr–Ti-hydroxide hydrate precipitates, which were filtered
and washed repeatedly until tested free of chloride ions
using silver nitrate solution. An Zr–Ti oxynitrate aqueous
solution (10 wt%) was prepared by dissolving the Zr–Ti
hydrate precipitates in an appropriate amount of 3.0 M
HNO3. A pre-calculated amount of Pb(NO3)3 was first
dissolved in deionized water and was then combined into
the Zr–Ti oxynitrate solution, in order to prepare an
aqueous nitrate solution containing Pb21, Zr41 and Ti41
ions in the molar ratio of 1:0.52:0.48. The coprecipitation
of Pb, Zr and Ti hydroxides was carried out by slowly
adding the mixed nitrate solution to an ammonia solution
of pH 9, which was maintained by adding a few drops of
concentrated ammonia solution during the coprecipitation
process. The co-precipitates were filtered and washed
thoroughly, and were then dried at 1008C for 4 h in an
oven. Then the dried precursor powder was thermally
treated at 4008C for 5 h, with both heating and cooling
rates being fixed at 58C/min in order to obtain an
essentially amorphous PZT precursor.
A batch of 5 g of the precursor powder was loaded into
a wear-resistant vial of 40 mm in diameter and 63 mm in
length together with a stainless steel ball 20 mm in
diameter. A high energy shaker mill operated at |900 rpm
was used to perform the mechanical activation of the
powder for various time periods ranging from 5 to 20 h
[15]. To make a comparison with the PZT powder derived
from calcination at a higher temperature, the second batch
of 5 g of the precursor powder was further calcined at
7508C for 1 h, with both heating and cooling rates being
fixed at 58C/min. The activation-derived PZT powders
were compacted uniaxially in a hardened steel die of 10
mm in diameter at a pressure of 50 MPa, and then
isostatically pressed at 350 MPa. Sintering of the isostati-
cally pressed powder pellets was carried out in air at
temperatures in the range of 1000–12008C for a fixed
duration of 1 h, with both heating and cooling rates being
fixed at 58C/min. The PZT powder synthesized by calcina-
tion was also pressed into pellets at the same pressures,
and they were then sintered in air at temperatures in the
range of 1100–13008C for 1 h.
2. Experimental procedures
The PZT composition selected in this study,
Pb(Zr0.52Ti0.48)O3, is near the morphotropic phase bound-
ary. The starting materials used are commercially available
Pb(NO3)2 (.99% in purity, Merck, Germany), TiCl4
(.99% in purity, Hayashi Pure Chemical Industries,
Japan), and a ZrO(NO3)2 aqueous solution (20 wt% ZrO2,
MEL, Manchester, UK). A high purity ammonia solution
(ammonia concentration: 28.30–30.00 wt%, J.T. Baker,
USA) was used as the precipitant in preparing the PZT
precursor by co-precipitation. An appropriate amount of
chilled (48C) deionized water was added slowly into a cool
TiCl4 solution while being stirred. A pre-weighted amount
of ZrO(NO3)2 solution was then blended into the TiCl4 –
water mixture. The Zr–Ti-aqueous solution was sub-
The PZT powders derived from both mechanical activa-
tion and calcination were characterized for phases present
using an X-ray diffractometer (X’Pert, Philips). A BET
surface area analyzer (NOVA-2000, Quantachrome, USA),
transmission electron microscopy (TEM, JEOL 100CX)
and scanning electron microscope (SEM, XL30 FEG,
Philips) were employed to study their powder characteris-
tics. PZTs sintered at various temperatures were studied
for microstructure and sintered density, which was mea-
sured using Archimedes method in deionized water. Their
dielectric behaviors were characterized using a HP 4284A
LCR meter at 1 kHz at various temperatures from room
temperature up to 4508C.

X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
141
3. Results and discussion
demonstrates that the initial 5 h of mechanical activation
resulted in a significant amount of PZT phase in the
amorphous precursor. This is supported by a fall in the
specific surface area from 8.73 m2 /g for the precursor that
was not subjected to mechanical activation to 5.37 m2 /g
for the powder subjected to 5 h of mechanical activation.
The initial 5 h of mechanical activation led to the
crystallization of perovskite PZT phase in the highly
amorphous precursor powder and therefore the specific
surface area was substantially reduced. The intensity and
sharpness of perovskite PZT peaks are further enhanced
when the mechanical activation is extended to 10 h and
then to 15 h. An essentially single phase of perovskite PZT
is obtained when the mechanical activation is extended to
20 h. Apparently, the transitional PT (PbTiO3) and PZ
(PbZrO3) phases were not involved with increasing the
degree of mechanical activation, prior to the occurrence of
PZT phase in the amorphous precursor.
Fig. 1 shows the XRD trace of the precursor powder
derived from co-precipitation, together with those of the
powders subjected to 5, 10, 15 and 20 h of mechanical
activation, respectively. The powder that was not subjected
to any mechanical activation exhibits a strong and
broadened peak over the 2u angle of 248 to 328, indicating
that it was highly amorphous and nanocrystallites of
perovskite PZT phase were not developed by the thermal
treatment at 4008C for 5 h. The strong and broadened peak
over the 2u angle of 248 to 328 is due to the occurrence of
PbO of low crystallinity in the hydroxide precursor. In
contrast, when mechanically activated for 5 h, a few rather
sharpened diffraction peaks are established at the 2u angles
of 22.0, 31.4, 38.3, 44.9, 50.4, 55.5 and 64.48. They
correspond to the perovskite PZT (100), (110), (111),
(200), (210), (211) and (220) peaks, respectively. This
Fig. 2 is a plot of the DTA trace at a heating rate of
108C/min in air for the powder mechanically treated for 20
h, together with that of the precursor powder that was not
subjected to any mechanical activation. One apparent
endothermic peak and two exothermic peaks are clearly
observed in the trace of the precursor powder. The
endotherm at |4718C is believed to relate to the decompo-
sition of hydroxide precursor, which was not completed at
the thermal treatment temperature of 4008C. The small
exotherm at |5188C is due to the formation of intermediate
phases, such as PbTiO3 in the precursor with increasing
temperature. The relatively broadened exothermic peak at
|6198C corresponds to the formation and crystallization of
PZT perovskite phase with rising temperature. In contrast,
only a single endothermic peak occurs at |4638C in the
powder that was mechanically treated for 20 h, which can
Fig. 1. XRD patterns of the precursor powders mechanically activated for
Fig. 2. DTA trace of the precursor powder at a heating rate of 108C/min
various durations ranging from 0 to 20 h (j: PZT, d: PbO).
in air, together with that of the powder mechanically activated for 20 h.

142
X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
be accounted for by the presence of hydroxide residuals.
The absence of exotherms at higher temperatures, i.e. those
at 5188C and 6198C observed for the precursor that was
not subjected to any mechanical activation, suggests that
the formation and crystallization of PZT phase are largely
completed by the mechanical activation. This has been also
demonstrated by the XRD traces shown in Fig. 1.
Fig. 3(a,b) shows two bright-field transmission electron
microscope (TEM) micrographs and the selected area
diffraction patterns for the precursor powder and the
powder subjected to 20 h of mechanical activation. Ag-
glomerates of irregular shape and 100 to 200 nm in size
were observed in the precursor powder. The very diffusive
rings of the selected area diffraction pattern suggest that it
is highly amorphous in nature. In a remarkable contrast,
the powder derived from 20 h of mechanical activation
consists of nanosized particles of 30 to 50 nm in size
which are rounded in morphology. The selected area
diffraction pattern clearly demonstrates its nanocrystalline
nature. These TEM results clearly show that the 20 h of
mechanical activation led to the occurrence of well estab-
lished PZT crystallites, which occur as nanosized particles
of rounded morphology.
Fig. 4(a,b) are two scanning electron microscope (SEM)
micrographs showing the microstructure of the PZT pow-
der derived from 20 h of mechanical activation and that
derived from calcination at 7508C, respectively. As shown
in Fig. 5, XRD phase analysis confirmed that a single
phase of perovskite PZT was also acquired when the
amorphous precursor was calcined at 7508C. However,
there is a huge difference in powder characteristics be-
tween the two powders. The PZT powder derived from
mechanical activation for 20 h consists of fine particles (30
to 50 nm in size) of rounded morphology, although they
Fig. 3. TEM micrographs and the selected area diffraction patterns for the precursor powder (a), and the powder subjected to 20 h of mechanical activation
(b).

X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
143
occur as small particle agglomerates. In a remarkable
comparison, large particle agglomerates of 0.5 to 2 mm in
size and of irregular morphology dominate the calcination-
derived PZT powder. This is expected, as similar powder
characteristics have been observed in many of the
chemical-derived ceramic precursors when subjected to a
post-synthesis calcination at intermediate/high tempera-
tures [17,18]. In other words, the final calcination ruins
almost all the advantages offered by chemistry-based
processing, such as a fine particle size and sintering-
reactive nature. In contrast, we have demonstrated that the
mechanical activation can trigger the precursor-to-ceramic
conversion, and in particular, the ultrafine nature derived
from chemistry is preserved.
The sintered density (relative density: %theoretical
density) of PZTs synthesized by 20 h of mechanical
activation and calcination at 7508C, respectively, as a
function of sintering temperature is shown in Fig. 6. For
both types of PZT, the sintered density increases with
increasing sintering temperature initially and peaks at an
optimum temperature, before a decrease is observed with
further increase in temperature. For example, the sintered
density of the activation-derived PZT increases with rising
temperature over the range from 1000 to 12008C, where it
maximizes at 97.6% theoretical density. Further increasing
the sintering temperature above 11508C results in a slight
fall in sintered density, due to the exaggerated grain
growth as will be discussed later. The PZT derived from
calcination at 7508C is much less sintering reactive and
exhibits a much lower sintered density than that of the
activation-derived PZT powder at each sintering tempera-
ture, with the maximum density of 86.4% theoretical
density occurring at 12508C. Indeed, it demonstrated little
sintering shrinkage at temperatures below 11008C. In
Fig. 4. SEM micrographs showing the PZT powder derived from (a)
mechanical activation for 20 h, and (b) calcination at 7508C for 1 h.
Fig. 6. A graph of the relative density against sintering temperature for
the PZT derived from mechanical activation and calcination at 7508C,
respectively.
Fig. 5. XRD trace of the PZT powder derived from calcination at 7508C
(j: PZT).

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X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
contrast, a sintered density of 88%, the theoretical density,
was realized for the activation-derived PZT at a tempera-
ture as low as 10008C.
The higher sinterability of the PZT derived from 20 h of
mechanical activation when compared to that of PZT
derived from calcination at 7508C can be accounted for by
their huge difference in powder characteristics. As has
been well established, particle coarsening and agglomera-
tion always occur in a ceramic powder as soon as a
chemistry-derived precursor is subjected to calcination at
an intermediate/high temperature in order to realize the
required ceramic phase [17,18]. This is a consequence of
Fig. 7. SEM micrographs showing the fracture surfaces of PZT ceramics sintered at different temperatures (a) 10008C, (b) 10508C, (c) 11008C, (d) 11508C
and (e) 12008C, respectively.

X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
145
the very high sintering reactivity of ultrafine precursor
particles, characterized by a very high specific surface
area. In other words, the sintering reactive nature of a
chemical-derived precursor is often ruined by the final step
of calcination at an elevated temperature. The strength of
the resulting particle aggregates increases dramatically
with increasing calcination temperature. As shown in Fig.
4(b), large particle aggregates of irregular morphology
occurred in the calcined PZT powder, and they adversely
affected the sintering behavior and microstructural de-
velopment of PZT at the sintering temperature. This is a
result of the differential sintering between regions of
differing green density. In comparison, the activation-
derived PZT powder consists of rather well dispersed PZT
particles of 30 to 50 nm in size.
Fig. 7(a–e) are the fracture surfaces of the activation-
derived PZTs sintered at 1000, 1050, 1100, 1150 and
12008C, respectively. The material sintered at 10008C
consists of rounded grains with fairly uniform grain sizes
in the range of 1–2 mm. The microstructure is rather
porous, agreeing with the measured density of |88%
theoretical. When sintered at 10508C, a relatively dense,
mainly intergranular fractured surface is observed with a
few small pores being present at the grain junction
positions. This agrees well with the measured density of
|92% theoretical. Its average grain size is in the range of
2–4 mm. There is an apparent increase in the average grain
size (in the range of 5–8 mm) and sintered density (to
|96% theoretical) when the sintering temperature is raised
to 11008C. A percentage of large grains are fractured in a
transgranular manner. The material sintered at 11508C
demonstrates a dense fracture surface with few small pores
being present, as shown in Fig. 7(d). Its average grain size
is in the range of 10–15 mm with a number of grains being
fractured in a transgranular manner. The PZT sintered at
12008C shows a primarily transgranular fracture. This is
due to the exaggerated grain growth, which is responsible
for the decrease in sintered density as shown in Fig. 6. Fig.
8 plots the average grain size for the activation-derived
PZT as a function of sintering temperature over the range
of 1000–12008C, as worked out using the linear intercep-
tion technique on the polished surfaces. There is a steady
increase in grain size with increasing sintering temperature
from 1000 to 11508C, where a sharp rise in increase rate is
observed with further increase in sintering temperature up
to 12008C.
Fig. 8. The average grain size as a function of sintering temperature for
the PZT derived from the powder mechanically activated for 20 h.
temperature, such as in the sintered density and grain size
as shown in Figs. 6–8. Fig. 10 shows the dielectric
constant as a function of temperature at 1 kHz for the PZT
sintered at 11508C for 1 h derived from the powder
mechanically activated for 20 h. A peak dielectric constant
of |9100 at the Curie point of 3808C was measured at 1
kHz.
4. Conclusions
The occurrence of perovskite nanocrystallites of PZT
phase in an amorphous precursor prepared by co-precipi-
tation has been realized by mechanical activation instead
of the traditional calcination at intermediate/high tempera-
The room temperature dielectric constant at 1 kHz as a
function of sintering temperature for PZTs derived from
the powder synthesized by 20 h of mechanical activation is
shown in Fig. 9. As expected, the dielectric constant
increases with rising sintering temperature from 1000 to
11508C, where it peaks at 927. A further increase in the
sintering temperature leads to a slight fall in dielectric
constant. This dependence of dielectric constant on sinter-
ing temperature may easily be accounted for by the
microstructural change of sintered PZT with increasing
Fig. 9. Dielectric constant of PZT derived from mechanical activation as
a function of sintering temperature.

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X. Junmin et al. / Journal of Alloys and Compounds 308 (2000) 139–146
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tures. A predominant perovskite PZT phase can be
achieved by a mere 5 h of mechanical activation and an
essentially single PZT phase was obtained at 20 h of
mechanical activation. The activation-derived PZT powder
exhibits a particle size in the nanometer range of 30–50
nm, and they are rounded in morphology and well dis-
persed. This is in contrast to the coarse and irregular
particles and agglomerates observed for the powder de-
rived from calcination of 7508C. Accordingly, the PZT
derived from mechanical activation shows a much better
sinterability, and it can be sintered to a density of 97.6%
theoretical density at 11508C for 1 h. A dielectric constant
of 927 at 1 kHz and room temperature is measured for the
PZT sintered at 11508C for 1 h.
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