In-situ formation of oxidation resistant refractory coatings on SiC-reinforced ZrB2 ultra
high temperature ceramics (UHTCs)
D.D. Jayaseelan*§a, E. Zapata-Solvas a, P. Brownb and W.E. Lee**a
a
Centre for Advanced Structural Ceramics (CASC), Dept. of Materials, Imperial College,
London, SW7 2AZ, UK
b
Dstl, Porton Down, Salisbury, Wiltshire, SP4 0JQ, UK
Abstract
In-situ oxidation resistant and refractory coatings have been generated on 20 vol. % SiCreinforced ZrB2 ultra high temperature ceramics (UHTCs) containing 10 wt. % rare earth
(RE) additives such as LaB6, La2O3 and Gd2O3 fabricated by spark plasma sintering.
Oxidation for 1h at 1600C in static air led to formation of a dense layer (up to 250 m thick)
of ZrO2 and RE-zirconates on the composite systems underneath which were intermediate
layers (50 ~ 100 m) containing heterogeneous crystalline oxides such as La2Zr2O7 and
amorphous silicate phases. The benefits of predominating solid oxidation products over
having substantial volumes of liquid present in aerospace leading edge applications are
discussed.
Key words: refractory oxide coating, UHTC, oxidation, zirconate, microstructure
* - Member, The American Ceramic Society
** - Fellow, The American Ceramic Society
§
- Author to whom correspondence should be addressed. e-mail : d.j.daniel@imperial.ac.uk
Crown copyright 2011. Published with the permission of the Defence Science and
Technology Laboratory on behalf of the Controller of HMSO.
1
I.
Introduction
Ultra high temperature ceramics (UHTCs) are candidate materials for a variety of
aerospace applications owing to their unique combination of properties, including high
melting temperature (> 3000C), high strength, and high elastic modulus.1-5 However,
oxidation resistance is a major issue in the development of UHTCs for aero-propulsion and
hypersonic flight applications. The oxidation resistance of monolithic MeB2 (Me = Zr and
Hf) ceramics depends on formation of MeO2 and B2O3. Below 1200C, liquid B2O3 fills the
pores between oxide grains, providing oxidation protection. However, above 1200C, MeB2
oxidises rapidly as a result of volatilisation of the protective B2O3 in the oxide scale 6-9 which
leaves behind a porous non-protective columnar-grained ZrO2 or HfO2 layer.
Much effort has been aimed at improving the oxidation resistance of ZrB2 and HfB2
based composites 6, 15-31 focussing on the addition of 20-30 vol. % SiC. Compared with liquid
B2O3, the borosilicate (BS) glass has higher melting temperature, higher viscosity, and lower
oxygen diffusivity, thus providing more effective oxidation protection
32,33
. SiC-containing
ZrB2 ceramics have relatively good oxidation resistance to 1500C. Above 1200C, the
addition of SiC provides improved oxidation resistance by formation of a prophylactic BS
glass coating
10-20
. Karlsdottir and Halloran
14
used high temperature optical microscopy to
reveal that convection currents were set up in the BS liquid leading to “volcanoes” and burst
bubble structures on the surface of the oxide scale. However, at very high temperatures (>
1800oC) the BS/SiO2 liquid is lost exposing the underlying boride to continuous active
oxidation. As a result there is much interest in developing a composite system with improved
oxidation resistance above this temperature. To further improve the oxidation resistance of
ZrB2/SiC, transition metal borides, such as CrB2, TiB2, TaB2, NbB2 and VB2, have been
added to the MeB2-SiC composite system 34. Improved oxidation resistance was related to the
2
presence of transition metal oxides in the borosilicate glass inducing its phase separation
(immiscibility) leading to increased liquidus temperatures and viscosities. These
characteristic features of immiscible glasses, are beneficial for decreasing oxygen diffusivity
and suppressing boria evaporation from the glass
34-35
. Others
32, 34-37
have studied the effect
of Si-containing additives, such as Si3N4 and silicides, on the oxidation behaviour of ZrB2
and HfB2 ceramics. Talmy et al.32 discussed the role of different silicides on the oxidation
behaviour of ZrB2. However, Opila et al.38 reported that addition of 20 vol. % TaSi2 to
ZrB2/20 vol. % SiC improved its oxidation resistance at 1627C in air. The improved
behaviour is attributed to the presence of Ta, not to the increased Si content arising from the
additional Si-containing component. In general, the oxidation behaviour of non-oxide
ceramics depends largely on the chemical composition and properties of the oxidation
products and on the combination of physical and chemical processes taking place on the
surface exposed to oxygen atmosphere
32
. Modification of the chemical composition of the
oxide layer, leading to decreased inward diffusion of oxygen, is an effective way of
controlling oxidation resistance of non-oxide ceramics. This modification can be
accomplished by changing the bulk ceramic composition. Zhang et al.
39
added WC to ZrB2
inducing liquid phase sintering (LPS) ultimately leading to formation of a two-layer dense,
protective, oxide scale; the outer layer comprising ZrO2, B2O3 and WO3 and the inner layer
ZrO2 and WO3.
There have been many previous attempts to improve the oxidation resistance of
ZrB2/HfB2 ceramics by e.g. increasing the viscosity of the liquid silica layer, increasing the
immiscibility of different liquid phases, or by development of dense ZrO2 layers via LPS thus
reducing oxygen penetration and diffusivity.40 However, in hypersonic space and military
applications, the temperature often exceeds 2000°C in combination with severe airflow. At
such high temperatures and in such harsh environments, the BS melt will be blown off
3
rapidly exposing the underlying layer. This may lead to further oxidation and UHTC material
recession. This work aims to develop solid refractory oxide protective layers by adding rare
earth (RE) borides or oxides to the baseline (ZrB2/20 vol.% SiC) UHTCs. Refractory oxide
layers will be formed by in-situ reaction 41 during the oxidation process.
II.
Experimental Details
ZrB2 powder (>99%, d50 ~ 45 μm,  = 6.85 g/cm3, Sigma Aldrich, Gillingham, UK) and
SiC powder (α-SiC, 99%, d50 ~ 0.7 μm,  = 3.217 g/cm3, Good Fellow Chemicals,
Huntingdon, UK) were used as starting materials to form a baseline ZrB2/20 vol.% SiC
(hereafter termed ZS20) UHTC. Both ZrB2 (space group, P6/mmm) and 6H-SiC (space
group, P63mc) have a hexagonal structure with lattice parameters of a=3.17 Å, c=3.53 Å and
a=3.08 Å, c=15.12 Å, respectively. To improve the oxidation resistance of ZS20, RE
additions in the form of LaB6 (>99%, d50 ~ 2-3 μm,  = 4.72 g/cm3, Sigma Aldrich,
Gillingham, UK), La2O3 (>99%, d50 ~ 2 μm,  = 6.51 g/cm3, Fluka chemicals supplied
through Sigma Aldrich, Steinheim, Germany) and Gd2O3 (>99%, d50 ~ 2 μm,  = 7.07 g/cm3,
Fluka chemicals supplied through Sigma Aldrich, Steinheim, Germany) were added to the
starting materials. The as-received ZrB2 powder was further dry-milled for 30 min in a
“swing-mill” shatter box using a steel container coated with Teflon to reduce the average
particle size. The average particle size (measured using a Malvern® laser-diffraction unit) of
the powder obtained after milling was 5.11± 0.5 µm. Longer milling times (> 30 min) did not
result in significant reduction in the particle size. Appropriate amounts of ZrB2, SiC and RE
additive were wet ball milled in a plastic container using ZrO2 balls in ethanol for nearly 12 h
and dried using a rotary evaporator. The dried powder compositions, namely ZS20/10 wt %
La2O3 (hereafter referred to as ZSLO), ZS20/10 wt % LaB6 (hereafter it referred to as ZSLB)
and ZS20/10 wt % Gd2O3 (hereafter it referred to as ZSGO) were densified in a spark plasma
4
sintering (SPS) furnace (FCT Systems, Germany). A 20 mm diameter graphite foil lined die
was used. The graphite die was covered with graphite felt to reduce heat loss and the
temperature monitored by an optical pyrometer which was sighted from top of the graphite
punch. Samples were sintered under 5 Pa vacuum between 1750C and 1850C for less than
10 min. A heating rate of 100oC/min was maintained to the sintering temperature and an
applied load of ~ 70 MPa was applied during sintering. Bulk density measurements were
carried out using the Archimedes method in water. Relative density was calculated by
dividing measured bulk density by theoretical density (TD) calculated by the rule of mixtures.
20 mm diameter and 5mm thick RE-doped samples were cut in half and placed on an
alumina boat so as to have a minimum point of contact between the sample and boat and
oxidised at 1600C for 1h in static air using a laboratory open hearth furnace. Oxidised
samples were cut, mounted in epoxy resin moulds and ground and polished using various
grades of media down to 1 m. For comparison, dense (~99% of T.D.) monolithic ZrB2 and
ZS20 were also fabricated using SPS at 1900C and 1850C, respectively with application of
identical heating rate, pressure and holding time and oxidation tests were performed on them
under the same conditions.
Phase analysis of sintered and oxidised samples was carried out by X-ray diffraction
(XRD) on a Philips PW7100 diffractometer using Cu-Kα radiation. International centre for
diffraction data (ICDD) cards used to identify phases were ZrB2 (01-075-1050), 6H-SiC (00029-1131), Gd2O3 (01-074-1987), La2O3 (01-083-1345), LaB6 (34-0427), La2Zr2O7 (170450), Gd2Zr2O7 (16-0799), t-ZrO2 (00-048-0224) and m-ZrO2 (00-037-1484).Thermal
analysis of sintered samples was carried out using a thermal analyser (Netzsch STA 449F1,
Germany) from room temperature to 1600oC at a rate of 10oC/min under flowing air (50
mL/min). At least 2 samples were tested in each composition using the thermal analyser
5
under identical conditions and no significant change was observed between them. Plan view
and cross-sections of the polished sintered and oxidised samples were observed in a scanning
electron microscope (SEM) fitted with a field-emission gun (FEG, model LEO15). Secondary
electron (SEI) and back-scattered electron images (BEI) were taken at an operating voltage of
20 keV, and a beam current 105 A and at a working distance between 10-15 mm. Chemical
analyses were carried out using an energy dispersive spectroscopy (EDS) unit (Oxford
Instruments, UK) attached to the FEG-SEM. For all samples, the working distance was ~ 10
mm. Additionally, the oxidised samples were imaged and the oxidised layers were
compositionally analysed in an FEI FIB-SIMS 200T focussed ion beam work station (FIB)
with secondary ion mass spectroscopy facility attached (SIMS). Gallium ions were used to
bombard the sample surface. The samples were attached to an aluminium sample holder with
silver tape to avoid charge build-up on the sample surface. Specimens for transmission
electron microscopy (TEM) observations were prepared from SPS materials using
conventional mechanical polishing and ion thinning. Ion thinning was performed using a
Gatan Model 691 precision ion polishing system (PIPS). TEM sections on the oxide layers
were prepared using a FIB work station operating with a gallium beam at 30 keV. Brightfield (BF) images and selected area electron diffraction (SAED) patterns were acquired using
a JEOL JEM-2000EX transmission electron microscope operating at 200 kV with an Oxford
Instruments EDS microanalysis system and SAED patterns were solved by the ratio method.
III.
Results
All samples, namely ZSLO sintered at 1850oC, ZSLB sintered at 1750oC and ZSGO
sintered at 1800oC, attained above 99 % of TD in less than 10 min. Phase analysis (Fig. 1)
confirmed that all starting materials remained after SPS. However, after oxidation for 1h at
1600C, phase transformation and in-situ reaction have occurred on the exposed surfaces of
6
these samples. XRD detected predominantly m-ZrO2 in monolithic ZrB2 and ZS20 samples
after oxidation at 1600C. Figure 2 shows XRD of oxidised surfaces of ZSLO, ZSLB and
ZSGO revealing that the sample surfaces comprised predominantly m-ZrO2. A trace of t-ZrO2
was also observed. In addition, in-situ formation of RE earth zirconate (RE2Zr2O7) pyrochlore
was observed during oxidation. Figure 3 shows ground and polished surface microstructures
of ZSLO [Fig. 3(a)], ZSLB [Fig. 3(b)] and ZSGO [Fig. 3(c)], after SPS. These reveal large
(4.0 – 7.0 µm) medium grey ZrB2 grains, finer (~1.0 µm) angular dark grey SiC grains and
light grey (> 2.0 µm) RE additive phases. In all three compositions, La2O3, LaB6 and Gd2O3
grains always appear agglomerated together with SiC forming an interconnected network and
are distributed homogeneously in the ZrB2 matrix. Figure 4 shows a bright-field (BF) image
of as-sintered ZSGO which is representative of all three compositions. Different grains are
labelled and the corresponding EDS are shown. In addition to the ZrB2, SiC and Gd2O3
starting materials, other phases present are ZrO2 (grain A) possibly arising from grinding
media and grain E in contact with SiC is a Gd-based silicate phase whose morphology
suggests it is glassy.
Figure 5 (a) shows SEI of a plan view of a ZrB2 sample after oxidation for 1h at
1600°C revealing the porous nature of the surface. EDS detects only Zr and O. In crosssection [Fig 5 (b)], a 40 m thick oxidized porous ZrO2 top layer is revealed on the
unaffected ZrB2. Figure 6 shows plan view and cross-section microstructures of baseline
sintered ZS20 after oxidation for 1h at 1600C. In the case of ZS20, the surface features and
nature of the oxidized layers are different. Fig. 6 (a) reveals fine (1-2 m) bright contrast
spherical particles in a dark grey matrix. These are secondary ZrO2 particles which most
likely precipitated from liquid silicate10, 26. The cross-section microstructure of ZS20 [Fig.6
(b)] shows two layers, an outer 5 m thick silicate layer containing ZrO2 particles and an
7
intermediate ~ 30 m thick porous ZrO2 layer between the outer silicate layer and the
underlying unaffected ZrB2-SiC region.
Figure 7 shows the surface and cross-section microstructures of ZSLO oxidized 1h at
1600C. The oxidized surface [Fig. 7 (a)] contains spherical pores (1 m) and grain boundary
cracks. The smooth and rounded nature of the microstructural features suggests liquid
formation during oxidation. EDS [Fig. 7 (b)] from the oxidized surface reveals the presence
of Zr, La, Si and O in the liquid. The cross-section [Fig. 7 (c)] shows two distinct oxidized
layers. The outer ~ 250 m thick layer is predominantly a continuous bright phase with
isolated dark regions present. EDS of the circled region 1 of bright phase in the outer layer
reveals La, Zr and O suggesting the presence of La2Zr2O7 as confirmed by XRD (Figure 2).
Figure 7 (d) shows BSI of the dark area of the circled region 2 in Fig. 7 (c) consisting of two
phases, possibly arising from phase separation of isolated droplets of liquid. EDS of circled
region 2 reveals the presence of La, Si and O. The rough intermediate layer above the
unaffected bulk material is ~ 50 m thick, EDS reveals Zr and O are the main elements in it.
Figure 8 shows a region of the electron image of the oxidised outer layer in ZSLO
obtained at a tilt angle of 45 in the FIB workstation. The electron image [Fig. 8 (a)] shows
bright and dark phases. SIMS analysis [Fig. 8 (b)] was carried out on some grains causing the
craters in Fig. 8 (a). SIMS clearly reveals the presence of La, Zr, and Zr-O in the bright grains
suggesting they are lanthanum zirconate grains. Figure 8 (c) shows a BF image of a TEM
section of a sample ion-milled from the interface of the outer layer and intermediate layer of
the oxidised region consisting of at least three different phases, namely La2Zr2O7, ZrO2 and
silicate glass. Figure 8 (d) was the SAED pattern of circled region d in Fig. 8 (c) tilted to the
[001] zone axis. The pattern can be indexed as pyrochlore cubic structure with a lattice
parameter = 10.768 Å corresponding to La2Zr2O7. Figure 8(e) shows the EDS taken on region
8
‘d’ in Figure 8(c). La, Zr and O are its main constituents suggesting the formation of
La2Zr2O7.
Figures 9 (a)-(b) show the microstructures of the oxidised surface and Figs. 9 (c)-(d)
detail from a cross-section of ZSLB. The oxidised surface is heterogeneous and covered with
a variety of phase morphologies. In some regions [Fig. 9 (b)], it appears to be compact with
grain boundary phases presumably derived from liquid. Large pits ~ 600 m in diameter are
formed on oxidation [Fig. 9 (a)]. Two layers can be seen at low magnification in Fig. 9 (c).
The outer layer is ~ 120 m thick and dense while the intermediate layer is ~ 50 m thick.
The continuous phase in the outer layer contains mainly Zr, La and O (EDS 1) while the
isolated dark regions (EDS 2) are mostly silica suggesting they derive from liquid since silica
would melt at this temperature. A flower-like pattern is occasionally observed in the
microstructure of the cross-section of the oxidised layers. Figure 9 (d) shows a BEI of such a
region in the top outer layer in cross section at higher magnification revealing several phases
around the light ZrO2 grains (10 m) including pockets of dark petal-like silica (EDS 2)
decorating the ZrO2, and between and around them a continuous bright phase (EDS 3)
containing nanoscale dark particles. EDS 3 reveals the bright phase is predominantly La and
Si with high O and low Zr presumably from neighbouring grains which along with its
morphology suggests it is a lanthanum silicate glass while the dark particles are silica glass.
La ions concentrate together at temperatures reaching 1600C to form lanthanum silicate
glass within silicate glass. This is another evident of phase separation of two immiscible
glasses in this study. EDS 4 is from the second layer and reveals the presence of Zr, Si, La
and O. The layer below the top layer is porous with a rough appearance.
Figure 10 shows the top outer exposed surface and cross-section microstructures of
ZSGO oxidized 1h at 1600C. SEI of the outer exposed surface reveals it is compact with few
9
shrinkage cracks forming on cooling [Fig. 10 (a)]. Figures 10 (b)-(c) are SEI images of the
cross-section microstructure showing three phases in the 100 µm thick outer oxidized layer
with black, dark, grey and bright contrast. EDS reveals that the light grey contrast grains
indicated by 1 are ZrO2, the bright phase indicated by 2 contains Gd, Zr, O and Si and the
black contrast grains indicated by 3 are predominantly silica. The intermediate layer is ~50
µm thick.
Figure 11 shows thermogravimetric analysis of ZrB2, ZS20 and RE-ZS20 carried out
under compressed air flow from room temperature to 1600C. There was 5.14 ± 0.2 %
increase in mass for ZrB2 and 2.91 ± 0.1 % for ZS20. LaB6 and La2O3 additions had 3.41 ±
0.2 % and 4.12 % ± 0.2 mass increase respectively, whereas Gd2O3 had only 1.8 ± 0.2 %
mass increase.
IV. Discussion
Ground and polished SEI surface microstructures of sintered samples of ZSLO, ZSLB
and ZSGO were similar. The small SiC particles (0.7 µm) are always in close proximity with
the small RE2O3 particles (2-3 µm) as would be expected when mixing two small size range
particles with one larger (ZrB2). Previous studies 41 suggested that during LPS of SiC, La2O3
reacts with the surface SiO2 on SiC to form a lanthanum silicate phase. It seems likely
therefore that at the high sintering temperature, ~1800C, used in the present study at least
some of the RE oxide reacts with the surface SiO2 on SiC to form a liquid grain boundary
silicate phase. Figure 4 suggests that in ZSGO, reaction has occurred with the Gd phase to
form a Gd-based silicate glass. Furthermore, the presence of < 1 vol. % of ZrO2 is likely to
arise from the ZrO2 grinding media used during milling. Furthermore, the surface oxide
impurities inherent to the starting ZrB2 powders may also be a significant source of ZrO2 and
should not be neglected. RE phases (La2O3, LaB6 and Gd2O3) were added to baseline UHTCs
10
with the intention that they react with ZrO2 to form RE-zirconates during oxidation. Although
the retention of the starting materials without any reaction occurring between them is
important, this does not rule out participation of any reacted phases such as RE silicate and
zircon in later oxidation reactions. The main reactions expected during oxidation are 11:
𝑍𝑟𝐵2(𝑠) +
5
2
𝑂2(𝑔) → 𝑍𝑟𝑂2 (𝑠) + 𝐵2 𝑂3 (𝑙)
𝐵2 𝑂3 (𝑙) → 𝐵2 𝑂3 (𝑔)
𝑆𝑖𝐶(𝑠) +
3
2
(i)
(ii)
𝑂2 → 𝑆𝑖𝑂2 (𝑙) + 3𝑂2 (𝑔)
(iii)
When adding LaB6 to ZS20, the following reactions are expected to occur during oxidation.
𝐿𝑎𝐵6 (𝑠) +
11
2
1
𝑂2 (𝑠) → 2 𝐿𝑎2 𝑂3 (𝑠) + 3𝐵2 𝑂3 (𝑙)
𝑍𝑟𝑂2 (𝑠) + 𝐿𝑎2 𝑂3(𝑠) → 𝐿𝑎2 𝑍𝑟2 𝑂7 (𝑠)
(iv)
(v)
When adding RE2O3 to ZS20, the following reaction occurs during oxidation.
𝑍𝑟𝑂2 (𝑠) + 𝑅𝐸2 𝑂3(𝑠) → 𝑅𝐸2 𝑍𝑟2 𝑂7 (𝑠)
(vi)
𝑆𝑖𝑂2 (𝑠) + 𝑅𝐸2 𝑂3(𝑠) → 𝑅𝐸 𝑠𝑖𝑙𝑖𝑐𝑎𝑡𝑒 𝑔𝑙𝑎𝑠𝑠
(vii)
(where RE = La and Gd in this study)
RE zirconates have the pyrochlore structure and melting temperatures  2300C (La2Zr2O7 2295 + 10 oC and Gd2Zr2O7 - 2450 + 10 oC).
Furthermore, formation of RE2Zr2O7 is expansive, i.e. if reaction (v) occurs during oxidation,
1 unit cell volume of ZrO2 + 1 unit cell volume of La2O3  1 unit cell volume of La2Zr2O7
11
i.e., 140.62 Å3 (for m-ZrO2) + 82.9 Å3  1262.37 Å3
indicating ~ a 5 fold increase in unit volume of the reaction product. Owing to this volume
expansion, RE-zirconate formation might help to close the pores being generated by
evaporation of volatile species such as B2O3.
During oxidation, ZrB2 oxidises to form ZrO2 and B2O3, where B2O3 later volatilises
leaving behind porous ZrO2 40 [Fig. 5 (b)]. Similarly, SiC oxidises to form SiO2 [Fig. 6 (b)].
These oxide phases, ZrO2 and SiO2, are available to react with RE2O3 to form RE2Si2O7 or
RE2Zr2O7. However, XRD shows only formation of RE2Zr2O7. It seems likely RE silicate
formed as a liquid at temperature and cooled to a glassy phase [Figs. 4(e), 7(d), 9(d) and
10(c)], and hence was not detected by XRD. The predominance of m-ZrO2 implies almost
complete transformation from t-ZrO2 has occurred, which again involves nearly ~ 3 vol. %
expansion.
Comparison of the cross-sections of monolithic ZrB2 (Fig.5 (b)), baseline UHTC
ZrB2/20 vol. % SiC, ZS20 (Fig.6 (b)) and ZS20 with RE additions, (Fig.7 (c)) for ZSLO,
(Fig.9 (c)) for ZSLB and (Fig.10 (b)) for ZSGO after oxidation for 1h at 1600C reveals that
ZS20 compositions with added RE have improved oxidation resistance for the following
reasons. In the case of monolithic ZrB2, a weight gain of over 5 % occurred (Fig. 11) with
thinner depth (40 m) of oxidation (Fig. 5 (b)) and no protective layer developed. During
oxidation of ZrB2 into ZrO2 and B2O3, liquid B2O3 is completely volatilized at these
temperatures leaving only porous ZrO2 grains.11
According to reaction (i), 1 molar mass of ZrB2 (112.84 g.mo1-1) gives rise to 1 molar
mass of ZrO2 (123.22 g.mo1-1) and 1 molar mass of B2O3 (69.62 g.mo1-1).
1 molar mass of ZrB2
=
112.84 g.mo1-1
12
1 molar mass of ZrO2
=
123.22 g.mo1-1
1 molar mass of B2O3
=
69.62 g.mo1-1
Now there arise two possible situations,
(a) If ZrB2 fully oxidises to ZrO2 and B2O3 and there is no escape of B2O3, there should
be a weight increase of 141%.
(b) On the contrary, if B2O3 escapes during oxidation, there should still be a weight
increase of 8 %.
Figure 11 reveals only 5.14 % increase in mass suggesting that incomplete ZrB2
oxidation has taken place after 1h at 1600oC. However, the porous nature of this layer
facilitates oxygen transport through the pore channels and further degradation of ZrB2 under
more severe oxidation conditions would be expected.
Baseline UHTC (ZS20) had a thin (< 5 µm) protective outer silica layer (Figure 6 (b)),
which hinders inward oxygen diffusion and the mass gain was ~ 2.91 % (Fig.11). In ZS20,
according to reactions (i)-(iii), there should be a mass gain of 140 % for the complete
conversion of the original specimen to condensed oxide phases and in which the volatile
species do not escape. However, in ZS20, besides reactions (i)-(iii), interaction between
reaction products like silica, boria and zirconia giving rise to zircon, BS and Zr-BS also need
to be considered. In RE added ZS20, the reactions during oxidation become more complex.
Previous reports
32-34
detected formation of liquid BS glass, which is less viscous than liquid
SiO2. Phase separation of this glass has been reported to help delay oxygen transport. In a
borosilicate or silicate glass containing transition metal cations, the tendency toward liquid
immiscibility is known to increase with increasing cation field strength of the transition
metal. This phase separation has been argued to result in increased viscosity, which has been
13
correlated to reduced oxygen diffusion rates. Nonetheless, this protective silica is liquid, and
hence could not be used at temperatures above 1800C in an aerospace environment with air
flow as the liquid would be blown off. In addition, the second ZrO2 layer is highly porous
enabling oxygen transport to the underlying composite.
RE-added ZS20, however, had thick (up to 250 µm) outer layers (Figures 7(c), 9(c and
10(b)) and weight increases up to 4.12 % (Fig.11) during oxidation for 1 h at 1600oC. A
schematic diagram comparing the cross-sections of monolithic ZrB2, ZS20 and ZS20 with RE
additives is shown in Fig.12. Among the RE additives, ZSGO had only 1.81% mass increase
and is protected by a thick (average thickness ~ 150 m) oxide layer (Fig.10 (b)) when
compared with the thicker [Fig. 7(c), up to ~ 250 m] but less protective (4.12% weight
increase, Fig.11) layers on La2O3 and LaB6 [(Fig.9 (3), ~ 125 m), (Fig.11, 3.41% weight
increase)]. The outer layers are dense and consist predominantly of crystalline refractory
oxides such as ZrO2 and RE2Zr2O7 [Fig.8] and localized silicate phase. Although, the mass
increase in RE added ZS20 is higher than in ZS20, the formation of dense solid layers is
likely to be more advantageous than solely formation of liquid silica. This is because at
temperatures above 1600C and high threshold velocity, liquid silica will be blown off by the
air flow of the atmosphere with which it is in contact paving the way for further attack of the
surface. Chemical composition and phase analyses carried out by EDS and XRD (Figure 2)
of the outer layers reveal that in general they contain mainly ZrO2, RE2O3 and RE2Zr2O7.
EDS analysis of the intermediate layers revealed RE, Zr, O, and Si. Hence it is likely that
Re2Zr2O7 and silicate phases formed below the top surface during oxidation. A possible
reaction sequence during oxidation is: ZrB2 and SiC oxidise to ZrO2, B2O3 and SiO2 phases
and CO2 gas is emitted. This leads to formation of BS liquid. Later ZrO2 and RE2O3 grains
dissolve in this borosilicate glass
24-26
increasing the liquid viscosity. At high temperature,
B2O3 volatilises leaving behind ZrO2 and RE2O3 grains in a silicate melt and a competition
14
exists between ZrO2 and RE2O3 to react with SiO2. This leads to the formation of both
RE2Zr2O7 and/or silicate phase(s) along with unreacted ZrO2 and BS glass. The distribution
of Zr in both the unreacted material and the oxide scale is continuous and homogeneous
indicating that the oxide scale was coherent and compact even though many voids were
present owing to shrinkage cooling. The overall effect of the RE addition is to significantly
alter the chemical composition and crystalline nature of phases forming on top oxidised layer.
IV.
Conclusions
In-situ oxidation resistant and refractory coatings have been generated on spark plasma
sintered 20vol.% SiC reinforced ZrB2 (ZS20) ultra high temperature ceramics containing
10wt.% of rare earth (RE) additives such as LaB6, La2O3 and Gd2O3 . Oxidation for 1 hour at
1600°C in static air led to formation of a dense surface layer (up to 250µm thick) of ZrO2 and
RE-zirconates. With melting points well above 1600°C both phases remained solid
throughout the oxidation process. Conversely, the oxidised surface of ZS20 without REadditives comprised a porous ZrO2 layer (up to 10µm thick) covered by amorphous silica
which was liquid at 1600°C. Owing to the low oxygen permeability of liquid silica its
presence suppresses excessive oxidation of ZS20 in static air at 1600°C. However, in a
hypersonic air stream this protective advantage would be quickly lost owing to liquid silica
removal by viscous flow. The oxygen permeabilities of RE-zirconates, although not as low as
liquid silica, are still substantially lower than ZrO2. This, combined with the high melting
point of RE-zirconates, suggests that RE additions may be a useful approach to improving the
oxidation resistance of UHTC’s at intermediate temperatures in hypersonic air. In terms of
weight change, the greatest improvement in the oxidation resistance of ZS20 at 1600°C
reported here was associated with the use of Gd2O3 additions.
15
Acknowledgement
The Authors’ acknowledge Prof. Mike Reece, Nanoforce, Queen Mary, University of
London, UK for providing access to the Spark Plasma Sintering facility. DDJ thanks the
Defence Science and Technology Laboratory (Dstl) for providing the financial support for
this work under contract number DSTLX-1000015413. EZS acknowledges the support of
‘Fundación Ramón Areces, Spain’ and the Centre for Advanced Structural Ceramics (CASC)
for his postdoctoral fellowship to stay at Imperial College London, UK.
16
References
1
W. G. Fahrenholtz, G. E. Hilmas, I. G. Talmy, and J. A. Zaykoski, "Refractory diborides
of zirconium and hafnium," J. Am. Ceram. Soc., 90 (5), 1347-1364 (2007).
2
S. Q. Guo, "Densification of ZrB2-based composites and their mechanical and physical
properties: A review," J. Euro. Ceram. Soc., 29 (6), 995-1011 (2009).
3
Zimmermann, J. W, Hilmas, G. E, Fahrenholtz, W. G, Monteverde, F, And Bellosi, A.,
Fabrication and properties of reactively hot pressed ZrB2-SiC ceramics. J. Euro. Ceram.
Soc 27 (7), 2729-2736 (2007).
4
R. Savino, M. D. Fumo, D. Paterna, and M. Serpico, "Aerothermodynamic study of
UHTC-based thermal protection systems," Aerospace Science and Technology, 9 (2), 151160 (2005).
5
M. M. Opeka, I. G. Talmy, E. J. Wuchina, J. A. Zaykoski, and S. J. Causey, "Mechanical,
thermal, and oxidation properties of refractory hafnium and zirconium compounds," J.
Euro. Ceram. Soc., 19 (13-14), 2405-2414 (1999).
6
T. A. Parthasarathy, R. A. Rapp, M. Opeka, and R. J. Kerans, "Effects of phase change
and oxygen permeability in oxide scales on oxidation kinetics of ZrB2 and HfB2," J. Am.
Ceram. Soc., 92 (5), 1079-1086 (2009).
7
T. A. Parthasarathy, R. A. Rapp, M. Opeka, and R. J. Kerans, "A model for the oxidation
of ZrB2, HfB2 and TiB2," Acta Materialia, 55, 5999-6010 (2007).
8
F. Monteverde and A. Bellosi, "Oxidation of ZrB2-based ceramics in dry air," J.
Electrochem., Soc. 150 (11), B552-B559 (2003).
9
W. C. Tripp and H. C. Graham, "Thermogravimetric study of oxidation of ZrB2 in
temperature range of 800oC to 1500oC," J. Electrochem. Soc., 118 (7), 1195-1199 (1971).
17
10
S. N. Karlsdottir, J. W. Halloran, and A. N. Grundy, "Zirconia transport by liquid
convection during oxidation of zirconium diboride-silicon carbide," J. Am. Ceram. Soc.,
91 (1), 272-277 (2008).
11
W. G. Fahrenholtz, "Thermodynamic analysis of ZrB2-SiC oxidation: Formation of a SiCdepleted region," J. Am. Ceram. Soc., 90 (1), 143-148 (2007).
12
X. H. Zhang, P. Hu, and J. C. Han, "Structure evolution of ZrB2-SiC during the oxidation
in air," J. Mater. Res. 23 (7), 1961-1972 (2008).
13
F. Monteverde and L. Scatteia, "Resistance to thermal shock and to oxidation of metal
diborides-SiC ceramics for aerospace application," J. Am. Ceram. Soc. 90 (4), 1130-1138
(2007).
14
S. N. Karlsdottir and J. W. Halloran, "Rapid oxidation characterization of ultra-high
temperature ceramics," J. Am. Ceram. Soc. 90, 3233-3238 (2007).
15
J. C. Han, P. Hu, X. H. Zhang, S. H. Meng, and W. B. Han, "Oxidation-resistant ZrB2-SiC
composites at 2200 degrees C," Composites Science and Technology 68 (3-4), 799-806
(2008).
16
S. N. Karlsdottir and J. W. Halloran, "Oxidation of ZrB2-SiC: Influence of SiC content on
solid and liquid oxide phase formation," J. Am. Ceram. Soc., 92 (2), 481-486 (2009).
17.
A.Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, "Oxidation of zirconium diboride-silicon
carbide at 1500oC at a low partial pressure of oxygen," J. Am. Ceram. Soc., 89 (10), 32403245 (2006).
18.
P. Hu, X. H. Zhang, and S. H. Meng, "Oxidation behavior of zirconium diboride-silicon
carbide at 1800oC," Scripta Materialia, 57 (9), 825-828 (2007).
19.
J. C. Han, P. Hu, X. H. Zhang, S. H. Meng, and W. B. Han, "Oxidation-resistant ZrB2-SiC
composites at 2200oC," Composites Science and Technology, 68 (3-4), 799-806 (2008).
18
20.
C. M. Carney, P. Mogilvesky, and T. A. Parthasarathy, "Oxidation behavior of zirconium
diboride silicon carbide produced by the spark plasma sintering method," J. Am. Ceram.
Soc., 92 (9), 2046-2052 (2009).
21.
M.J.H Balat, "Determination of the active-to-passive transition in the oxidation of silicon
carbide in standard and microwave excited air," J. Euro. Ceram. Soc., 16, 55-62 (1996).
22.
G. Magnania, A. Brillanteb, I. Bilottib, L. Beaulardic and, and E.Trentinic, "Effects of
oxidation on surface stresses and mechanical properties of liquid phase pressurelesssintered SiC–AlN–Y2O3 ceramics " Mat. Sci. and Engg.: A., 486, 381-388 (2008).
23.
A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, "Evolution of structure during the
oxidation of zirconium diboride-silicon carbide in air up to 1500oC," J. Euro. Ceram. Soc.,
27 (6), 2495-2501 (2007).
24.
S. N. Karlsdottir and J. W. Halloran, "Formation of oxide films on ZrB2-15 vol% SiC
composites during oxidation: evolution with time and temperature," J. Am. Ceram. Soc. 92
(6), 1328-1332 (2009).
25.
S. N. Karlsdottir and J. W. Halloran, "Formation of oxide scales on zirconium diboridesilicon carbide composites during oxidation: relation of subscale recession to liquid oxide
flow," J. Am. Ceram. Soc. 91 (11), 3652-3658 (2008).
26.
S. N. Karlsdottir, J. W. Halloran, and C. E. Henderson, "Convection patterns in liquid
oxide films on ZrB2-SiC composites oxidized at a high temperature," J. Am. Ceram. Soc.
90 (9), 2863-2867 (2007).
27.
A.Pavese, P. Fino, C. Badini, A. Ortona, and G. Marino, "HfB2/SiC as a protective coating
for 2D C-f/SiC composites: Effect of high temperature oxidation on mechanical
properties," Surface & Coatings Technology, 202 (10), 2059-2067 (2008).
19
28.
W. B. Han, P. Hu, X. H. Zhang, J. C. Han, and S. H. Meng, "High-temperature oxidation
at 1900oC of ZrB2-xSiC ultrahigh-temperature ceramic composites," J. Am. Ceram. Soc 91
(10), 3328-3334 (2008).
29.
J. W. Hinze, W. C. Tripp, and H. C. Graham, "High-temperature oxidation behaviour of a
HfB2 + 20 vol.% SiC composite," J. Electrochem. Soc., 122 (9), 1249-1254 (1975).
30.
S. S. Hwang, A. L. Vasiliev, and N. P. Padture, "Improved processing, and oxidationresistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids," Mat.
Sci. and Engg. A, 464 (1-2), 216-224 (2007).
31.
C. M. Carney, "Oxidation resistance of hafnium diboride-silicon carbide from 1400 to
2000oC," J. Mater. Sci., 44 (20), 5673-5681 (2009).
32.
I. G. Talmy, J. A. Zaykoski, and M. M. Opeka, "High-temperature chemistry and
oxidation of ZrB2 ceramics containing SiC, Si3N4, Ta5Si3, and TaSi2," J. Am. Ceram. Soc.
91 (7), 2250-2257 (2008).
33.
D. W. McKee, C. L. Spiro and E. J. Lamby, “The effect of boron additives on the
oxidation behaviour of carbons”, Carbon, 22 (6), 507-511 (1984).
34.
I.G. Talmy, J.A. Zaykovski, M.M. Opeka and S. Dallek, Oxidation of ZrB2 ceramics
modified with SiC and Group IV-VI transition metal borides. In: M. McNallan and E.
Opila, Editors, High Temperature Corrosion and Material Chemistry III, The
Electrochemical Society, Inc., Pennington, NJ (2001), p. 144.
35.
B.G. Varshal, "A structure model for immiscibility in silicate glass-forming melts," Glass
Phys. Chem., 19[2], 218-225 (1993).
36.
B.G. Varshal (ed), "Two-phase glasses: structure, properties, and applications," AN SSSR,
Nauka, Leningrad, p.11-33, 1991.
20
37.
J. Kuchino, K.Kurokawa, T. Shibayama, and H.Takahashi, "Effect of microstructure on
oxidation resistance of MoSi2 fabricated by spark plasma sintering," Vacuum, 73 (3-4),
623-628 (2004).
38.
D. Sciti, A. Balbo, and A. Bellosi, "Oxidation behaviour of a pressureless sintered HfB2MoSi2 composite," J. Euro. Ceram. Soc. 29 (9), 1809-1815 (2009).
39.
E. Opila, S. Levine, and J. Lorincz, "Oxidation of ZrB2- and HfB2-based ultra-high
temperature ceramics: Effect of Ta additions," J. Mater. Sci., 39 (19), 5969-5977 (2004).
40.
S. C. Zhang, G. E. Hilmas, and W. G. Fahrenholtz, "Improved oxidation resistance of
zirconium diboride by tungsten carbide additions," J. Am. Ceram. Soc. 91 (11), 3530-3535
(2008).
41.
E.Eakins, D.D. Jayaseelan and W.E. Lee, “Toward oxidation-resistant ZrB2-SiC ultra high
temperature ceramics” Metall. and Mater. Trans. A, 42A, 878-887 (2011).
42.
S. Tabata, Y. Hirata, S. Sameshima, N. Matsunga, and K. Ijichi, "Liquid phase sintering
and mechanical properties of SiC with rare-earth oxide," J. Ceram. Soc. Jpn., 114, 247252 (2006).
43.
D.D. Jayaseelan, S. Ueno, T. Ohji, and S. Kanzaki, "Sol-gel synthesis and coating of
nanocrystalline Lu2Si2O7 on Si3N4 substrate," Mater. Chem. and Phys., 84 (1), 192-195
(2004).
21
Figure Captions
Figure 1
XRD of spark plasma sintered ZS20 with various RE additives; no phase
transformation observed.
Figure 2
XRD of ZS20 containing La2O3 (ZSLO), LaB6 (ZSLB) and Gd2O3 (ZSGO)
oxidised for 1h at 1600C showing product oxides and in-situ formed zirconate
phases.
Figure 3
Microstructures of as fabricated ZS20 with RE additives (a) BEI of ZSLO, (b)
BEI of ZSLB and (c) BEI of ZSGO.
Figure 4
Bright-field TEM image of ZSGO showing a range of phases. EDS from the
labelled grains is shown.
Figure 5
ZrB2 oxidised for 1h at 1600C. (a) SEI of surface showing porous ZrO2
microstructure and (b) BEI of cross-section showing 2 regions; porous ZrO2 top
layer, I, and un-reacted ZrB2 layer. Representative EDS from regions 1 in (a) and
(b) is shown.
Figure 6
SEIs of ZS20 oxidised for 1h at 1600C. (a) Plan view of surface showing bright
contrast ZrO2 particles (II) in silicate matrix (I) and (b) Cross-section showing 3
regions; top protective SiO2 + ZrO2 layer, intermediate porous ZrO2 layer and
bottom un-reacted ZrB2/SiC region. EDS from regions I and II in (a) are shown.
Figure 7
SEIs of ZSLO oxidised for 1h at 1600C. (a) exposed surface showing many
cracks and pores. (b) corresponding EDS, (c) Cross-section showing 2 oxidised
layers; top protective La-Zr-O and intermediate ZrO2 layers on the un-affected
ZrB2/SiC and (d) BEI image of the dark circled in region 2 in (c) showing its biphasic nature. EDS from regions 1 and 2 in (c) are shown.
22
Figure 8
ZSLO oxidised 1h at 1600C (a) FIB image of oxidised layer showing different
phases, (b) SIMS of bright grains in (a), (c) BF-TEM image of a region in
oxidised layer, (d) SAED of grain d (La2Zr2O7) in (c) taken along the [001] zone
axis and (e) EDS taken on grain “d” in (c) showing the presence of Zr, La and O.
Figure 9
ZSLB oxidised for 1h at 1600C. (a) SEI of exposed surface showing large pit
and cracks, (b) SEI showing smooth surface; (c) SEI of cross-section showing
different layers, (d) BEI of a region in top protective layer showing a flower-like
pattern, contain three different phases. EDS were taken on regions marked (1-4).
Figure 10 ZSGO oxidised for 1h at 1600C. (a) SEI of exposed surface showing grey and
dark phases (b) BEI of cross-section showing different layers with irregular top
layer thickness, and (c) BEI of region 2 in top protective layer showing three
different phases. EDS were taken from regions marked (1-3).
Figure 11 Thermogravimetric analysis of (a) monolithic ZrB2, (b) ZS20, (c) ZSLB, (d)
ZSLO and (e) ZSGO from room temperature to 1600C.
Figure 12 Schematic comparsion of the cross-sections of (a) monolithic ZrB2, (b) ZS20, and
(c) ZS20-10% Re2O3 / ReB6 oxidised for 1h at 1600C.
23