Journal of the European Ceramic Society xxx (xxxx) xxx
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Journal of the European Ceramic Society
journal homepage: www.elsevier.com/locate/jeurceramsoc
Improved ageing-resistance and fracture toughness of zirconia-toughened
alumina bioceramics via composition and microstructure design
Yuhong Hong, Mingmin Bai *, Shaohua Wang, Qibing Chang, Xiaozhen Zhang, Xueguo Zhao,
Yongqing Wang *
School of Materials Science and Engineering, Jingdezhen Ceramic University, Jingdezhen 333403, China
A R T I C L E I N F O
A B S T R A C T
Keywords:
Zirconia-toughened alumina
Nano-crystalline PrAlO3
Praseodymium hexaaluminates plate-like crys­
tals
Low temperature degradation
Mechanical properties
Low-temperature degradation (LTD) is the main obstacle that hinders the application of zirconia-toughened
alumina (ZTA) ceramics as artificial hip joints. In this study, a new type of ZTA bioceramic with LTD resis­
tance, high fracture toughness, and superior wear resistance was prepared by the in-situ formation of plate-like
crystals and co-stabilised tetragonal phase zirconia (t-ZrO2). This new design is realised by synthesising nano­
crystalline PrAlO3 through a one-step solution combustion method and introducing it into a ZTA matrix to form
praseodymium hexaaluminate (Pr0.833Al11.833O19, PHA) plate-like crystals by solid-state sintering. PHA plays a
key role in toughening and Y–Pr co-doped t-ZrO2 slows down the low-temperature degradation of ZTA bio­
ceramics. After hydrothermal ageing, the combination strategies have a positive influence on the flexural
strength, fracture toughness, and Vickers hardness of the ceramics. This study provides a novel direction for
ensuring long-term safety of bioceramics.
1. Introduction
Zirconia-toughened alumina (ZTA) matrix composites have been
considered as candidates for replacing ultra-high molecular weight
polyethylene and cobalt-chromium alloys in the manufacturing of arti­
ficial hip joints owing to their good biocompatibility, excellent wear
resistance, and superior mechanical properties [1–3]. Yttria-stabilised
tetragonal zirconia (Y-TZP) is largely responsible for the improved me­
chanical performance of ZTA ceramics [4]. However, the lack of dura­
bility of Y-TZP in vivo limits the application of ZTA bioceramics [5,6].
That is because Y-TZP can spontaneously convert from the tetragonal
phase zirconia (t-ZrO2) to the monoclinic phase zirconia (m-ZrO2) in a
humid environment. This low-temperature degradation (LTD) behav­
iour not only reduces the toughening of ZTA but also causes volume
expansion, leading to cracking of the ZTA in vivo and a secondary sur­
gery [7,8]. Therefore, the inhibition of the LTD of ZTA ceramics with
guaranteed mechanical properties is of great importance for their bio­
logical applications.
In recent years, the synergistic toughening of ZTA by Y-TZP and
plate-like crystals has attracted considerable attention [9–11]. Syner­
gistic toughening not only substantially improves the toughness of ZTA
but also effectively reduces the Y-TZP content below the threshold for
continuous low-temperature ageing (approximately 17 vol%), thereby
protecting the artificial hip joint from corrosion in humid environments
in vivo [12].
Praseodymium hexaaluminate (Pr0.833Al11.833O19, PHA) has a
unique plate-like morphology owing to its specific anisotropy in growth
kinetics [13]. It has been reported that the in-situ growth of PHA
plate-like crystals from praseodymium nitrate successfully increased the
fracture toughness of ZTA [14]. Meanwhile, the inclusion of Y3+ ions
into Pr0.6Zr0.35Y0.05O2 stabilises Pr4+ and enhances the Pr4+/Pr3+
reversible redox behaviours in the sample [15]. Owing to its excellent
reversible redox cycle properties, Pr-doping plays a key role in elimi­
nating the interaction between the sample surface and water vapour
[16].
In general, the toughening of ZTA by the in-situ growth of plate-like
crystals occurs mainly through the solid-phase reaction of rare-earth
oxides with alumina. Although this process is simple, it has some un­
avoidable drawbacks. Owing to the low reactivity of rare earth oxides, a
small amount of intermediate product intermediates remain after sin­
tering, which adversely affects the mechanical properties of bioceramics
[17–19]. Moreover, if the plate-like crystals are formed at a high tem­
perature and the ZTA substrate is compacted before the plate-like
crystals are formed, there is insufficient growth space (small
* Corresponding authors.
E-mail addresses: bellebai2010@126.com (M. Bai), wyq8248@126.com (Y. Wang).
https://doi.org/10.1016/j.jeurceramsoc.2022.12.051
Received 9 August 2022; Received in revised form 9 December 2022; Accepted 20 December 2022
Available online 22 December 2022
0955-2219/© 2022 Elsevier Ltd. All rights reserved.
Please cite this article as: Yuhong Hong, Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2022.12.051
Y. Hong et al.
Journal of the European Ceramic Society xxx (xxxx) xxx
length-to-diameter ratio) for the plate-like crystals and thus, the ce­
ramics have a weak toughening effect [20].
Our previous research revealed that PrAlO3 is an essential interme­
diate product in the synthesis of plate-like PHA crystals from the solidphase reaction of Pr6O11 and Al2O3. The number and morphology of
the plate-like crystals, as well as the properties of the ZTA ceramic, are
directly influenced by the PrAlO3 grain size [14]. Therefore, we
hypothesise that if PrAlO3, an intermediate phase with a fine grain size
and high reactivity, is introduced into ZTA ceramics, it will reduce the
synthesis temperature of PHA plate-like crystals to allow sufficient
growth space and increase the length-to-diameter ratio of the plate-like
crystals, thus improving the synergistic toughening effect. More
importantly, this method can significantly reduce the Y-TZP content and
improve the hydrothermal ageing resistance of ZTA ceramics, thus
promising the long-term use of ZTA ceramics in humid environments.
In this study, nanocrystalline pure PrAlO3 was synthesised in one
step via a combustion method using a composite fuel. The effects of
different ratios of the composite fuel on PrAlO3 were investigated. The
mechanism of the in-situ growth of PrAlO3 in ZTA ceramics to grow PHA
plate-like crystals was elucidated. In addition, the mechanical properties
of ZTA–PHA before and after hydrothermal ageing were investigated in
detail. This data is important for evaluating the application prospects of
ZTA–PHA bioceramics.
→PrAlO3(s)
(1)
2.2.2. Preparation of Pr0.833Al11.833O19
The PrAlO3 prepared at x = 0.67 in Section 2.2.1 was mixed with
alumina by ball milling according to the molar ratio n(PrAlO3):n
(Al2O3)= 5:33. The obtained slurry was granulated using a spray dryer.
Bulk-shaped green compacts (4 mm × 6 mm × 20 mm) were prepared
by uniaxial pressing at 20 MPa followed by cold isostatic pressing at
200 MPa for 2 min. The obtained green compacts were sintered in an
electric furnace at 1100, 1200, 1300, 1400, 1500, and 1600 ◦ C in air for
2 h at a heating rate of 300 ◦ C/h. The specimens were subsequently
cooled down to 25 ◦ C at a rate of 400 ◦ C/h.
2.2.3. Preparation of zirconia-toughened alumina–Pr0.833Al11.833O19
The PrAlO3 prepared at x = 0.67 in Section 2.2.1 was mixed with
ZTA by ball milling. Two sets of formulae were designed with PrAlO3
contents of 0 and 2 wt% and the sintered samples were denoted as ZTA
and ZTA–PrAlO3, respectively. In a typical preparation procedure, 85 wt
% Al2O3 and 15 wt% ZrO2 were added to deionised water to prepare a
ZTA slurry. The slurry was granulated using a spray dryer. Bulk-shaped
green compacts (4 mm × 6 mm × 40 mm) were prepared by uniaxial
pressing at 20 MPa followed by cold isostatic pressing at 200 MPa for
2 min. The obtained green compacts were sintered in an electric furnace
at 1600 ◦ C in air for 2 h at a heating rate of 300 ◦ C/h. The specimens
were subsequently cooled down to room temperature at a rate of
400 ◦ C/h.
2. Materials and methods
2.1. Materials
The starting materials used in the experiments were alumina (Al2O3,
D50 =0.16 µm, purity 99.99%, Daimyo Chemical, Tokyo, Japan), zir­
conia (3Y-ZrO2, D50 =0.36 µm, purity >99.5%, Ceramica Materials Co.
Ltd., Jiangxi, China), praseodymium nitrate (Pr(NO3)3⋅6 H2O, Sigma­
–Aldrich, Shanghai, China), aluminium nitrate (Al(NO3)3⋅9 H2O, Sig­
ma–Aldrich, Shanghai, China), urea (CH4N2O, Sigma–Aldrich,
Shanghai, China), and glycine (C2H5NO2, Sigma–Aldrich, Shanghai,
China).
2.3. Characterisation
Phase identification of the samples was accomplished by X-ray
diffraction (XRD, Bruker D8 Advance, Germany) employing Cu-Kα ra­
diation (λ = 0.15418 nm) with a step size of 0.02◦ (2θ) and a scanning
rate of 5◦ /min ranging from 10◦ to 70◦ . Differential scanning calorim­
etry (DSC) and thermogravimetric measurement (TG) analyses were
performed using a thermal analyser (Netzsch STA 449 C, Germany) from
room temperature to 1400 ◦ C. Scanning electron microscopy (SEM,
JEOM-JMS-6700 F, Japan) and energy-dispersive spectroscopy (EDS,
JED-2200 T, Japan) were used to conduct the microstructural observa­
tions and compositional studies, respectively.
Before observation, thermal etching was done on the polished sam­
ples by heating them in air at 1300 ◦ C for 30 min. The morphology,
crystal structure, and boundary chemistry of ZTA–PrAlO3 were analysed
using high-resolution transmission electron microscopy (HRTEM, Talos
F200s, Thermo Fisher, USA) equipped with an energy-dispersive X-Ray
spectrometer (EDX) and selected area electron diffractometer (SAED).
To investigate the LTD behaviour, autoclave tests were conducted in
accordance with ISO 13356 at 134 ◦ C, 2 bar, and for 10 h in steam [21].
After hydrothermal ageing, the ZTA and ZTA–PrAlO3 samples were
designated ZTA–LTD and ZTA–PrAlO3–LTD, respectively. The volume
fractions of the monoclinic phases (Vm) were calculated using the for­
mula described by Toraya [22].
The Vickers hardness test (HV1, HV-1000, Haoxinda, China) was
conducted using a standard diamond indenter with a load of 9.8 N for
15 s. Seven samples were analysed to determine the mean values and
standard deviations. The volume density of the sintered composites was
measured using the Archimedes method and the relative density was
evaluated by considering the theoretical density of each composite [23].
Assuming the theoretical density of alumina, Y2O3 partially stabilised
t-ZrO2, and praseodymium hexaaluminate are 3.987, 6.05, and
4.18 g/cm3, respectively, the theoretical density of each composite was
calculated using the law of mixture [24].
Mechanical property tests were conducted using benchtop universal
2.2. Synthesis
2.2.1. Preparation of PrAlO3 by solution combustion
As a typical procedure, 0.18 mol of Pr(NO3)3⋅6 H2O, 0.18 mol of Al
(NO3)3⋅9 H2O, and 0.6 mol of glycine were dissolved in deionised water.
The prepared solution was heated in a water bath to 85 ◦ C while being
vigorously stirred. Following the evaporation of water from the solution,
a viscous gel was obtained. The gel was subsequently heated on a hot
plate maintained at 200 ◦ C. As the water evaporated further, the gel
swelled into foam with the evolution of the fumes. On continuous
heating, this foam bulged and burst into flames with rapid evolution of a
large amount of gas. The reaction was completed within a few minutes,
resulting in a fluffy pale-yellow foam-like powder. The powder was
sieved through a 200-mesh screen and crushed using an agate mortar.
The same procedure was used for other fuels. The combustion reaction
represented by Eq. (1) can be used to calculate the moles of fuel
required. The details of the samples with respect to the fuel type and
ratio are listed in Table 1.
Table 1
Composition of the reacting preparative mixture.
Samples
Pr (NO3)3
(mol)
Al (NO3)3
(mol)
Glycine
(mol)
Urea
(mol)
xa
a
b
c
d
0.18
0.18
0.18
0.18
0.
0.
0.
0.
0.6
0.4
0.3
0.15
0
0.3
0.45
0.675
1
0.67
0.5
0.25
a
18
18
18
18
10
xC2 H5 NO2(aq.) + 5(1 − x)CH4 N2 O(aq.)
3
+ (8 − 10/3x)N2(g) + (5 + 5/3x)CO2(g) + (10 − 5/3x)H2 O(g)
Pr(NO3 )3(aq.) + Al(NO3 )3(aq.) +
x is valued from Eq. (1).
2
Y. Hong et al.
Journal of the European Ceramic Society xxx (xxxx) xxx
testing equipment (WDW-20, China). Each specimen, measuring
4 mm × 3 mm × 35 mm, was ground and polished using emery papers
of grades 400, 800 and 1200 grit and diamond paste of 0.5 µm on a
polishing cloth. The edges of all the specimens were chamfered to
eliminate the stress concentration caused by machining defects. The
flexural strength was determined using three-point bending tests with
30 mm loading spans and a 0.5 mm/min displacement rate. The fracture
toughness of 3 mm × 4 mm × 35 mm test bars was determined using
the single-edge notched beam method with a 30 mm span and crosshead
speed of 0.05 mm/min [25]. A notch depth of approximately 2 mm was
cut at the centre of the surface (3 mm × 35 mm) of each bar using a
diamond wire with a thickness of approximately 125 µm. A minimum of
five specimens were tested under each experimental condition.
The friction and wear properties of the samples were analysed using
a ball-on-disc tribometre (UMT-2, CETR Corporation, United States).
Alumina balls 10 mm in diameter were used as the ball material. The
investigated ceramics ZTA–LTD and ZTA–PrAlO3–LTD were shaped into
discs. The diameter of the discs was approximately 19 mm and the
thickness was approximately 4.5 mm. The samples were polished to a
mirror state with a 2.5 µm diamond grinding paste to eliminate the in­
fluence of the surface roughness of the specimens on the friction and
wear properties.
The tribological tests at room temperature environments of all
ceramic samples were performed at a sliding velocity of 0.1 m/s for a
sliding time of 1 h under a sliding load of 50 N. The coefficient of friction
(COF), which is defined as the ratio of the friction force to the applied
compressive force, was measured online. After the test, threedimensional (3D) scanning of the wear track was observed using a
white light interferometer (Phase shift MicroXAM-3D, United States,
KLA-Tencor Corporation) and the specific wear rate was calculated.
The wear volume for all the samples was calculated by multiplying
the average (three measurements in every wear scar) cross-sectional
area of the scar with the respective circumferential length of the wear
track. The specific wear rate was calculated as the wear volume per
applied load and the sliding distance in mm3/(N⋅m). To identify the
wear mechanisms, the worn surfaces of the polymer and its counterparts
were observed by SEM.
3. Results and discussion
3.1. Effect of fuel on the synthesis of PrAlO3 by solution combustion
Owing to the different temperatures and combustion reactions,
different ratios of composite fuels have a significant impact on the
morphology and composition of the products. Fig. 1 shows the com­
bustion process and the powders obtained at different fuel ratios ac­
cording to Table 1. When glycine was used as the only fuel (Sample a in
Table 1), there was no open flame during the combustion process with
copious gas evolution and spongy powders were obtained (Fig. 1a).
When glycine and urea were used as fuel mixtures (Fig. 1b–d), the
combustion mode changed to volume combustion synthesis. A flame was
present during the entire combustion process and an incandescent re­
action front was observed, which is typical for combustion reactions. In
the case of x = 0.67 (Sample b in Table 1), the intense red flame lasted
for 12–15 s and mild eruption combustion occurred [26]. A yellowish
foam-like product was obtained. When x = 0.5 (Sample c in Table 1), an
eruption combustion reaction was noted [27] and the flame lasted for
approximately 18–20 s. At the conclusion of the combustion process,
slightly grey powders were produced, indicating the existence of carbon
residue from insufficient fuel oxidation [28]. With a further increase in
urea content (x = 0.25, Sample d in Table 1), violent eruption com­
bustion occurred and a large number of foam-like grey-yellow powders
were spewed out by the airflow.
The XRD patterns of the powders produced using the various com­
posite fuel combustion ratios are shown in Fig. 2. When glycine was the
only fuel (x = 1), the product had an amorphous structure, indicating
that the smouldering combustion reaction was insufficient for PrAlO3
formation. The XRD diffraction peaks of the sample with x = 0.67 are
consistent with those of PrAlO3 (JCPDS 97–009–0558). PrAlO3 is formed
directly from the combustion reaction without any additional heat
treatment. This result highlights the benefits of using a composite fuel
approach rather than the traditional single-fuel approach. Small
Fig. 1. Combustion processes and products of the combustion reaction using different ratios of composite fuels.
3
Y. Hong et al.
Journal of the European Ceramic Society xxx (xxxx) xxx
Fig. 2. X-ray diffraction (XRD) patterns of PrAlO3 prepared with different ratio of composite fuel.
Fig. 3. Standard enthalpy, standard Gibbs free energy, and adiabatic temperature of PrAlO3 prepared at diffident ratios of composite fuels.
impurity peaks were observed in the XRD patterns as the urea concen­
tration in the fuel increased. These observations suggest that a proper
ratio of the composite fuel during the combustion reaction is important
for the formation of pure PrAlO3.
To explain the purity of PrAlO3 formed at x = 0.67, the standard
enthalpyΔHθ298 , standard Gibbs free energyΔGθ298 , and adiabatic tem­
peratureTad (the highest temperature that may be reached under adia­
batic conditions without any energy loss) were calculated using Eqs.
(2–4) [29,30]：
)
)
∑(
∑(
ΔH θ298 =
ni ΔH θi,f ,298
−
ni ΔH θi,f ,298
(2)
react
Al(NO3 )3 ⋅9H2 O⟶60−
75◦ C
ΔGθ298 =
)
∑(
ni ΔGθi,f ,298
∫
ΔH θ298 =
Tad
298
∑(
ni Cpi
−
react
)
prod
)
∑(
ni ΔGθi,f ,298
prod
dT
(3)
(4)
where ni is the number of moles of each substance involved in the re­
action andCpi is the molar heat at constant pressure. The calculation re­
sults are shown in Fig. 3. The equations for the thermal decomposition of
Al(NO3)3⋅9 H2O and Pr(NO3)3⋅6 H2O are also given [31,32]:
prod
Al(NO3 )3 ⋅6H2 O⟶120−
135◦ C
◦
Al(NO3 )3 ⋅3Al(OH)2 ⋅2.5H2 O⟶200 C 4Al2 O3 ⋅3N2 O5 ⋅14H2 O
4
(5)
Y. Hong et al.
Pr(NO3 )3 ⋅6H2 O⟶35−
Journal of the European Ceramic Society xxx (xxxx) xxx
73◦ C
Pr(NO3 )3 ⋅4H2 O⟶73−
104◦ C
Pr(NO3 )3 ⋅2H2 O⟶104−
141◦ C
Pr(NO3 )3 ⋅H2 O⟶141−
291◦ C
Pr(NO3 )3 ⟶291−
426◦ C
PrONO3 ⟶426−
503◦ C
Pr6 O11
(6)
From a thermodynamic standpoint, the glycine single-fuel (x = 1)
combustion reaction is the most exothermic, with the lowest standard
Gibbs free energy and the highest Tad . In this situation, the synthesis of
PrAlO3 is expected to occur readily using the glycine single-fuel method.
Nevertheless, the actual temperature during the incomplete combustion
reaction was significantly lower than the adiabatic temperature
(Fig. 1a). Therefore, the heat emitted was insufficient to synthesise
PrAlO3, which is consistent with the XRD pattern (Fig. 2).
The primary causes for reaching the highest combustion temperature
are the overlap of the temperature ranges at which metal nitrates and
fuels decompose and the capacity of these metal nitrate-fuel systems to
produce the necessary concentration of combustible gases, such as NOx,
NH3, and CO [33]. The decomposition temperature of glycine is 240 ◦ C
[34]. According to Eqs. (5) and (6), the glycine single-fuel cannot
overlap the decomposition temperature intervals of metal nitrates.
Hence, the critical density of combustible gases cannot be attained. The
decomposition of urea (152 ◦ C) occurs at a lower temperature than
glycine [35]. In the urea–glycine fuel system, urea and Al(NO3)3 would
burn first, resulting in an increase in temperature that ignites the gly­
cine–Pr(NO3)3 combustion reaction and the creation of PrAlO3.
With a further increase in urea content, the urea–Al(NO3)3 com­
bustion reaction becomes more intense and a large volume of gas es­
capes, leading to a reduction in the density of the spontaneous
combustion gas, a decrease in the flame temperature, and a higher
amount of residual carbon in the products. Therefore, x = 0.67 was
selected as the best ratio of composite fuel.
As shown in Fig. 3, the ΔHθ298 and ΔGθ298 of the reactions with
different fuel ratios are both negative. From a thermodynamic stand­
point, all four reactions are considered spontaneous. The calculated
adiabatic temperature reached a maximum (2287.71 ◦ C) when glycine
was used as the fuel, whereas it decreased as the urea content increased.
3.2. Growth mechanism of Pr0.833Al11.833O19
To understand the thermochemical changes in PrAlO3 and Al2O3
during heating, thermal analysis was conducted on the reaction process
of the green compacts prepared in Section 2.2.2. At the same time,
another part of the green compact was heated to 1100–1500 ◦ C for 2 h
before cooling in the furnace. Fig. 4 shows the corresponding TG–DSC
curves, XRD patterns, and SEM images. The TG curves demonstrated a
total mass loss of less than 1% from room temperature to 1400 ◦ C
(Fig. 4a), indicating that the combustion reactions were nearly com­
plete. The thermal decomposition process comprises four major steps,
two endothermic peaks, and two exothermic peaks. The first endo­
thermic peak is at approximately 56.8 ◦ C on the DSC curve, which
corresponds to a mass loss of approximately 0.2% on the TG curve. This
endothermic peak was caused by the evaporation of absorbed water. The
endothermic peak observed at 235.4 ◦ C and a weight loss of 0.6% are
caused by the decomposition of the organic binder. In addition, the weak
and broad exothermic peaks at approximately 640.2 and 1196.2 ◦ C
could be due to the oxidation of praseodymium compounds and their
subsequent reactions to form to PHA, respectively. PHA synthesis is
controlled by diffusion and has no discernible temperature effects [36].
During the crystallisation process, the TG curve revealed the presence of
two minor mass gains that might be connected to a thermally induced
redox reaction with praseodymium [37].
The XRD patterns of the samples heated at different temperatures for
2 h are shown in Fig. 4b. After heating at 1100 ◦ C, the characteristic
peaks of PrAlO3 (JCPDS 97–009–0558) and Al2O3 (JCPDS
99–000–0826) were observed in the XRD pattern. When the temperature
increased to 1200 ◦ C, the main phases in the sample included PrAlO3
and Al2O3 and there also exist weak peaks of PHA (JCPDS
97–007–4317). This indicates that PHA can be formed at 1200 ◦ C when
PrAlO3 synthesised by the combustion method is added to the Al2O3
Fig. 4. (a) Thermogravimetric measurement (TG)–differential scanning calo­
rimetry (DSC) curves, (b) XRD patterns, and (c) a scanning electron micrograph
(SEM) of Pr0.833Al11.833O19 prepared with PrAlO3 and Al2O3 at 1600 ◦ C. The
inset shows the schematic representation of a hexagonal layered structure and
the morphology of Pr0.833Al11.833O19 pellets.
5
Y. Hong et al.
Journal of the European Ceramic Society xxx (xxxx) xxx
Fig. 5. (a) Cylinder model, (b) SEM-EBSD micrograph of the cross section of the selected area in Fig. 5a, (c) energy-dispersive spectrum (EDS) of the line scan of the
yellow arrow position, and (d) a schematic diagram of the solid-state diffusion reaction mechanism. The blue dashed lines are the approximate positions of the
interface. AGG represents the abnormal grain growth of Al2O3. PHA represents Pr0.833Al11.833O19.
matrix as a praseodymium source. The initial formation temperature of
PHA is 100 ◦ C lower than when praseodymium nitrate was used as the
praseodymium source with Al2O3 [14].
With a further increase in the heat-treatment temperature, the
characteristic peaks of PHA became sharper. At 1600 ◦ C, the diffraction
peaks of PrAlO3 and Al2O3 disappeared completely and the XRD pattern
matched the standard pattern of PHA. This suggests that a solid-state
reaction conducted for 2 h at 1600 ◦ C can generate pure PHA. Fig. 4c
clearly indicated that all the PHA plate-like crystals grow well with a
diameter of 18.27 ± 5.54 µm and aspect ratio of 6.40 ± 2.55. Owing to
the uneven packing of the plate crystals, many micropores were
discovered. Step-like cleavage was observed on the fracture surfaces.
To investigate the growth mechanism of PHA crystals during heat
treatment, duplex green compacts were created, as shown in Fig. 5a. To
create the interior component, a disc specimen was created using diepressed PrAlO3 powder, which was placed into a die with a diameter
greater than that of the disc specimen. Subsequently, the A12O3 powder,
in accordance with the stoichiometric ratio of PHA, was fed into the die
and reformed under pressure to yield a duplex cylindrical compact. The
compact was heated at a rate of 300 ◦ C/h to 1600 ◦ C in air, maintained
for 2 h at the peak temperature, and cooled naturally in the furnace.
After sintering, a microstructure investigation was performed on the
cross sections of the selected locations. As shown in Fig. 5b, there is a
clear interface between Al2O3 and PrAlO3. The dark area on the left
Fig. 6. XRD patterns of different ZTA samples before and after hydrothermal ageing. Inset with enlarged view in the range of 2θ = 28–32◦ .
6
Y. Hong et al.
Journal of the European Ceramic Society xxx (xxxx) xxx
signify Al2O3 with uniform grain sizes and there are abnormally grown
Al2O3 grains near the interface. The bright area on the right signifies
PrAlO3. Plate-like PHA crystals are observed in the PrAlO3 area.
Elemental line scanning was performed along the Al2O3 and PrAlO3
layers. As shown in Fig. 5c, only Al and O are present in the left region.
Both the interface and right areas contain Al, Pr, and O. However, the
right area has a stronger Pr signal than the interface. The above analysis
shows that when the Al2O3 and PrAlO3 phases are in contact at high
temperatures, Al3+ and Pr3+ diffuse to each other and the diffusion rate
of Al3+ is greater than that of Pr3+. A diagram of the growth mechanism
of the plate-like PHA crystals is shown in Fig. 5d. On the right side of the
interface, Al3+ and electrons migrate through the product to phase
boundary II, where oxygen is transferred through the gas phase. The
following reaction occurs at phase boundary II:
/
/
5 6PrAlO3 + 11Al3+ + 33e− + 33 4O2 = Pr0.833 Al11.833 O19
(7)
in Fig. 6. As shown in Fig. 6, the main crystalline phases of all the
samples were Al2O3 (JCPDS, 97–005–2024) and t-ZrO2 (JCPDS,
04–005–4504). PHA phases (JCPDS, 97–007–4317) were detected
before and after hydrothermal ageing of the sample with 2 wt% PrAlO3.
This indicates that the addition of 2 wt% PrAlO3 was able to form PHA
in-situ through a solid-phase reaction with ZTA at a sintering tempera­
ture of 1600 ◦ C, which is consistent with the results in Fig. 4b. Before
hydrothermal ageing, the characteristic monoclinic peak m(111) at
28.2◦ (2θ) was not observed for the ZTA and ZTA–PrAlO3 samples.
However, after hydrothermal ageing, a weak m-ZrO2 phase (JCPDS,
03–065–1023) appears near 2θ = 28.2◦ and 31.3◦ . According to calcu­
lations, the volume fractions of the m-ZrO2 phase in ZTA–LTD and
ZTA–PrAlO3–LTD was 13.71% and 6.04%, respectively [22].
The magnified plot in the range of 2θ = 28–32◦ shows that the main
diffraction peak of t-ZrO2 in the ZTA–PrAlO3 sample is shifted towards a
higher angle than that of the ZTA sample. According to the Bragg
equation, this indicates that the interplanar spacing of zirconia in the
ZTA–PrAlO3 sample was smaller than that in ZTA. This is due to the fact
that when ZrO2 is co-doped with Y–Pr, it leads to an elevated oxygen
vacancy concentration [39]. The higher oxygen vacancy concentration
allows for a certain degree of contraction of the Zr–O bond, which
counteracts the cell expansion caused by the larger ionic radius when
Pr3+ (1.126 Å) replaces Zr4+ (0.84 Å) [40]. Meanwhile, a portion of the
Pr3+ ions has a tendency to segregate at the ZrO2 grain boundaries
(Fig. 8c), which slows down the grain boundary mobility owing to
pinning [41]. In summary, the segregation of doped Pr3+ and the
reduction in the size of zirconia crystals improved the t-ZrO2 stability,
which restrained the t-m phase transformation during hydrothermal
ageing.
Fig. 7a and b display the typical TEM image of ZTA–PrAlO3 and
HRTEM images of PHA. The TEM image reveals the coexistence of
Al2O3, ZrO2, and PHA in the ZTA–PrAlO3 sample. The HRTEM image
On the left area of the interface, a small amount of Pr3+ ions migrated
to phase boundary I. Owing to the low solubility limit of Pr3+ in the
Al2O3 lattice, there was a strong trend for grain boundary segregation.
Segregation of Pr3+ results in an increase in the grain boundary diffu­
sivity of Al2O3 and abnormal growth of Al2O3 grains. This phenomenon
was also observed for Nd2O3-doped Al2O3 [38].
3.3. Effect of Pr0.833Al11.833O19 on the properties of zirconia-toughened
alumina
To study the effect of PHA plate-like crystals on the properties of ZTA
ceramic materials, the phase composition, microstructure, and me­
chanical properties of the ZTA and ZTA–PrAlO3 samples before and after
hydrothermal ageing were investigated. Fig. 6 shows the XRD patterns
of the ZTA and ZTA–PrAlO3 samples sintered at 1600 ◦ C. For compari­
son, the XRD patterns of ZTA–LTD and ZTA–PrAlO3–LTD are presented
Fig. 7. (a) TEM bright-field image of ZTA–PrAlO3 (dark zirconia grains, bright Al2O3 grains, and bright elongated Pr0.833Al11.833O19 grains are observed). (b)
Enlarged image of (a) with the SAED pattern. (c) STEM-HAADF and the mapping distribution images of the (d) Al, (e) Zr, (f) Pr, and (g) O elements of ZTA–PrAlO3,
respectively.
7
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Journal of the European Ceramic Society xxx (xxxx) xxx
Fig. 8. (a,b) STEM-HAADF micrographs illustrating the position of the EDS line scanning across the grain boundary between the PHA and ZrO2 grains. (c) Observed
signals of Pr and Zr along the scanning line.
Fig. 9. (a) Flexural strength, fracture toughness, and Vickers hardness of ZTA before and after hydrothermal ageing. (b) SEM image of ZTA and (c) SEM image
of ZTA–LTD.
shows no glassy phase at the interface between Al2O3 and PHA, implying
that the plate-like structure of the PHA phase is ascribed to the prefer­
ential orientational diffusion of Al cations (Fig. 5d), instead of the
presence of a liquid phase. The SAED pattern confirms that PHA grains
are crystalline and possess magnetoplumbite-type hexagonal crystal
structure, which is consistent with the XRD results (Fig. 6). Further
investigation of the microstructure and composition of ZTA–PrAlO3 was
conducted by STEM-HAADF. The corresponding micrographs and
associated EDX maps are shown in Fig. 7c–f. The micrograph of
ZTA–PrAlO3 displayed nearly a homogeneous microstructure
comprising well distributed plate-like PHA grains in a fine grained Al2O3
matrix and smaller ZrO2 grains located largely at triple junctions. EDS
analysis has revealed that aluminium is located in both rounded and
plate-like darker grains, zirconium is exclusively detected in the zirconia
grains (brighter grains), and plate-like crystals are enriched with pra­
seodymium (Pr) compared to the matrix.
Fig. 8 shows the position of the EDS line being scanned by STEMHAADF over a grain boundary between the PHA and ZrO2 grains. The
8
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Y. Hong et al.
Fig. 10. (a) Flexural strength, fracture toughness, and Vickers hardness of ZTA–PrAlO3 before and after hydrothermal ageing. (b) SEM image of ZTA–PrAlO3 and (c)
SEM image of ZTA–PrAlO3–LTD. (d) Grain size distributions of Al2O3 and ZrO2 in different ZTA samples after sintering.
decreasing concentration gradient of praseodymium across the interface
can be attributed to the diffusion-controlled reaction, which supplies the
driving force for further praseodymium diffusion via interfaces. As can
be seen, there was no obvious segregation of praseodymium at the grain
boundaries. Praseodymium was predominantly present in ZrO2 crys­
tallites to form a solid solution of Y–Pr–ZrO2.
The mechanical properties and microstructures of ZTA before and
after hydrothermal ageing are shown in Fig. 9. The flexural strength,
fracture toughness, and Vickers hardness of ZTA obtained in this work
were 672 MPa, 8.07 MPa•m1/2, and 18.95 GPa, respectively, which
decreased after hydrothermal ageing. During hydrothermal ageing,
Table 2
Theoretical volumes and relative densities of different zirconia-toughened
alumina (ZTA) samples.
Sample
Theoretical density (g/
cm3)
Volume density (g/
cm3)
Relative density
(%)
ZTA
ZTA–PrAlO3
4.20
4.22
4.14
4.19
98.62
99.45
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Fig. 11. SEM backscattered images after Vickers indentation on ZTA–PrAlO3 showing the action of various toughening mechanisms along the Al2O3/PHA/
ZrO2 interface.
some t-ZrO2 particles undergo volume expansion upon transformation.
The dilation and shear strains are part of the shape strain caused by the
martensitic transformation. These strains frequently cause grain
boundary microcracking (Fig. 9c), which severely degrades the me­
chanical characteristics of the material.
The mechanical properties and microstructures of the ZTA–PrAlO3
samples before and after hydrothermal ageing are shown in Fig. 10a–c.
Fig. 10d shows the grain size distributions of Al2O3 and ZrO2 in the ZTA
and ZTA–PrAlO3 samples. The relative densities of ZTA and ZTA–PrAlO3
were 98.62% and 99.45%, which are listed in Table 2. The flexural
strength of the ZTA–PrAlO3 sample was 660.15 MPa，which is close to
that of ZTA. The fracture toughness of ZTA–PrAlO3 increases compared
to that of ZTA, reaching 8.6 MPam1/2, as a result of the in-situ formation
of PHA, which adds new toughening mechanisms while reducing the
thermal-mismatch stress between alumina and zirconia [42]. It is well
known that the grain size and densification level of samples are key
factors in determining their hardness. The ZTA–PrAlO3 sample had a
smaller grain size (Fig. 10d) and a higher degree of densification
(Table 2) than the ZTA sample. Thus, the ZTA–PrAlO3 sample has a
lower defect size, a larger cross-sectional area to bear the applied load,
and the Vickers hardness increases from 18.95 (ZTA) to 19.34 GPa
(ZTA–PrAlO3). Interestingly, according to Fig. 10a, the mechanical
properties of the samples were significantly improved after hydrother­
mal ageing, with the flexural strength, fracture toughness, and Vickers
hardness increasing from 660.15 to 700.39 MPa, 8.6–9.49 Mpa•m1/2,
and 19.34–19.89 GPa, respectively.
The improvement in flexural strength after hydrothermal ageing was
mainly affected by the fracture surface energy [43]. The volume
expansion caused by the phase transformation of t-ZrO2 during hydro­
thermal ageing led to compressive stress at the crack tip, which
improved the fracture surface energy. Therefore, higher stress is
required for further crack propagation. Similarly, the accumulation of
compressive stress on the surface increases the Vickers hardness of the
sample after hydrothermal ageing.
The internal stress interacts with the weakened interface between the
matrix and plate-like crystal during hydrothermal ageing because of the
t-m transformation, which can change the crack paths, such that the net
crack driving force is reduced and the fracture toughness is improved.
This process is revealed by the crack propagation induced by Vickers
indentation in the ZTA–PrAlO3–LTD sample (Fig. 11a). The dark poly­
gons, bright polygons, and pale rectangles represent Al2O3, t-ZrO2, and
plate-like PHA grains, respectively. The crack initiated from the tip of
the indentation (Fig. 11b) and expanded forward. From the analysis of
the cracking expansion process, it can be concluded that the synergistic
toughening effect is not a simple summation of the plate-like PHA grains
and t-ZrO2 action but rather of many factors, including the stressinduced t-m phase transformation, crack branching (Fig. 11c and i),
crack bridging (Fig. 11d, g, and j), crack deflection (Fig. 11f), and platelike crystal pull-out (Fig. 11h) to toughen the material, which is superior
to the single toughening action. The crack finally terminated within the
ZrO2 grain approximately 17 µm from the starting point (Fig. 11e).
Therefore, the fracture toughness is significantly increased to
9.49 MPa•m1/2 after hydrothermal ageing (Fig. 10a).
The COF, specific wear rate, 3D morphology of the wear trajectory,
and surface morphology after wearing the ZTA–LTD and ZTA–PrA­
lO3–LTD samples were investigated to explore the effect of PHA platelike crystals on the sliding friction and wear properties. The fluctua­
tions in the COF values with loading time are shown in Fig. 12a. The COF
curve of the ZTA–LTD sample displayed a considerable amplitude and
oscillation, while the ZTA–PrAlO3–LTD sample remained stable after
only a brief running-in period. The average COF values of the ZTA–LTD
and ZTA–PrAlO3–LTD samples are 0.69 ± 0.17 and 0.83 ± 0.03,
respectively, as shown in Fig. 12a. The specific wear rate of the ZTA–LTD
sample was 1.54 × 10− 10 mm3/(N⋅m), which indicates a mild wear
regime [44]. Almost no wear was observed in the ZTA–PrAlO3–LTD
sample after 1 h of dry-sliding friction. The specific wear rates of both
10
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Journal of the European Ceramic Society xxx (xxxx) xxx
Fig. 12. (a) Coefficient of friction and (b) average coefficient of friction and specific wear rate between the different samples with a Al2O3 ball. (c) Optical profiles,
(d) three-dimensional morphologies and (e) depth profiles of the wear track of different samples after friction with a Al2O3 ball: (1) ZTA–LTD and (2)
ZTA–PrAlO3–LTD.
11
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Fig. 13. SEM morphology images of the wear track and local details of different samples, namely (a) ZTA–LTD and (b) ZTA–PrAlO3–LTD. (c) Schematic diagram of
the wear mechanism change process of a ZTA–LTD sample during sliding friction.
samples met the requirement, namely 1 × 10− 8 mm3/(N⋅m), of
ceramic-on-ceramic combinations used in artificial hip joints [45]. The
surface wear resistance is also represented by the width and depth of the
frictional wear track. The smaller the dimensions, the better the surface
wear performance is. The optical profile, 3D morphology, and depth
profile of the wear tracks of the two samples are shown in Fig. 12c–e.
After 1 h of dry-sliding friction, several furrows parallel to the sliding
direction were observed on the surface of the ZTA–LTD sample
(Fig. 12c1 and d1). However, no continuous wear traces were visible on
the surface of the ZTA–PrAlO3–LTD samples (Fig. 12c2 and d2). The
depth profile of ZTA–LTD shows more fluctuation than ZTA–PrA­
lO3–LTD (Fig. 12e), which is consistent with the images in Fig. 12c and
d. Therefore, it can be concluded that the surface of the ZTA–PrA­
lO3–LTD sample has a much better wear-resistant behaviour than
ZTA–LTD. Thus, it is expected to improve the service performance of
bioceramic implants.
The surface topographies of the different samples after wear are
shown in Fig. 13a–b. In the first, second, and third columns, the worn
shapes of the wear track are shown at low, medium, and high magnifi­
cation, respectively. A schematic illustration of the wear mechanism
change process of the ZTA–LTD sample during sliding friction is pre­
sented in Fig. 13c. At the beginning of the wear process, the compressive
stress that was formed on the surface of the ZTA–LTD sample owing to
the t-m phase transition in the hydrothermal process hindered crack
growth caused by tensile stress after frictional contact (Fig. 13c1). The
number of cracks increases with wear, thereby weakening the binding of
m-ZrO2 grains to the matrix and making it easier to pull them out as wear
debris (Fig. 13c2). The debris deformed plastically under heavy loads,
forming a friction layer with parallel grooves along the direction of the
sliding friction (Fig. 13a1). This increases the actual contact area of the
friction surface and lowers the friction coefficient (Fig. 12a). However,
the friction layer had several voids and cracks that could easily come off
during contact, causing three-body debris and accelerating wear
(Fig. 13a3). The wear surface of the ZTA–PrAlO3–LTD sample was
smoother than that of the ZTA–LTD sample, with only minor scratches
and shallow surface pits (Fig. 13b). Thus, the ZTA–PrAlO3–LTD sample
exhibited greater wear resistance under dry sliding conditions than the
ZTA–LTD sample. This may be due to the following reasons: i) ZTA–­
PrAlO3 is insensitive to hydrothermal ageing and fewer microcracks
appear on the free surface of the ZTA–PrAlO3–LTD sample than the
ZTA–LTD sample; and ii) the good combination of ZTA and PHA in­
terfaces significantly increases the energy required for crack
propagation.
4. Conclusion
ZTA–PHA composites with low-temperature degradation resistance
and high fracture toughness were successfully synthesised by introduced
PrAlO3 which was obtained by one-step solution combustion. The
following conclusions were drawn:
(1) The type and proportion of composite fuel in the solution com­
bustion method had a significant impact on the combustion state
and purity of the product. When glycine and urea were used as
the fuel mixture, the metal nitrates matched the fuel decompo­
sition temperature to produce spontaneous combustion gas at a
critical concentration, leading to open flame combustion, and
pure PrAlO3 was directly formed through the combustion
reaction.
(2) Al3+ diffused into PrAlO3 and formed PHA plate-like crystals at
1200 ◦ C. Pr3+ ions segregated at the Al2O3 grain boundary,
leading to abnormal growth of Al2O3 grains.
(3) PrAlO3 reacted in-situ with the ZTA matrix to produce PHA platelike crystals, which improved the fracture toughness. Y–Pr co12
Journal of the European Ceramic Society xxx (xxxx) xxx
Y. Hong et al.
doped t-ZrO2 slowed down the low-temperature degradation of
ZTA bioceramics. As hydrothermal ageing proceeded, the flexural
strength and fracture toughness increased from 660.15 MPa and
8.6 MPa•m1/2 to 700.39 MPa and 9.49 MPa•m1/2, respectively,
which was closely related to the synergistic toughening of platelike crystals and t-ZrO2.
(4) Compared with ZTA–LTD, ZTA–PrAlO3–LTD had a much higher
wear resistance. The specific wear rates of the ZTA–LTD and
ZTA–PrAlO3–LTD samples met the requirements of ceramic-onceramic combinations used in artificial hip joints.
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Declaration of Competing Interest
The authors declare that they have no known competing financial
interests or personal relationships that could have appeared to influence
the work reported in this paper.
Acknowledgements
This work was supported by the National Key Research and Devel­
opment Program of China (2018YFC1903406), National Natural Science
Foundation of China (52102023), Jiangxi Provincial Natural Science
Foundation (20202BABL214008), and Education Project of Jiangxi
Province in China (GJJ211309).
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