Journal of the European Ceramic Society 22 (2002) 279–288

www.elsevier.com/locate/jeurceramsoc

Processing and properties of zirconium diboride-based composites

F. Monteverde, A. Bellosi*, S. Guicciardi
CNR-IRTEC, Research Institute for Ceramics Technology, Via Granarolo 64, 48018 Faenza, Italy

Received 11 January 2001; accepted 28 March 2001

Abstract
Two zirconium diboride-base composites were produced and characterised. The chosen starting compositions were: 55 wt.%
ZrB2+41 wt.%TiB2+4 wt.% Ni and 83 wt.% ZrB2+13 wt.% B4C+4 wt.% Ni. The microstructure and properties of these com-
posites were compared to those of a monolithic ZrB2+4 wt.% Ni material. In all cases, metallic Ni as the sintering aid promoted
the formation of the liquid phase which improved mass transfer mechanisms during sintering. From the powder mixture
ZrB2+TiB2, two solid solutions of Zr–Ti–B were obtained. In the case of the other mixture, B4C particles were dispersed in the
ZrB2 matrix. The composite materials have better mechanical properties than those of the monolithic ZrB2 ceramic; in particular
the fracture toughness and the flexural strength were almost doubled at room temperature. Long term oxidation tests indicated that
the ZrB2-based composites, particularly the composite containing B4C as the second phase, were more resistant to oxidation than
the monolithic ZrB2 due to the formation of surface oxide products which were protective against the complete degradation by
oxidation observed for the ZrB2 matrix material. # 2001 Elsevier Science Ltd. All rights reserved.
Keywords: B4C; Composites; Electrical properties; Mechanical properties; Thermal properties; TiB2; ZrB2

1. Introduction the sintering of ZrB2 powders is rather difficult. Rela-

tively high densities are achieved only by pressure-assis-
Zirconium diboride (ZrB2) ceramics have a high ted sintering procedures at temperatures higher than
melting point, high hardness, high electrical con- 1900 C, i.e. temperatures exceeding 70% of the absolute
ductivity, excellent corrosion resistance against molten melting temperature.1,5 As with TiB2 materials,6 the
iron and slags and superb thermal shock resistance. introduction of sintering aids such as Fe, Ni, Co, W, C,
They constitute a class of promising materials for high WC improves the final density and allows a lower densi-
temperature applications in several industrial sectors, fication temperature,1,7 increasing volume diffusion and
like foundry or refractory industries. Like TiB2, ZrB2 is retarding evaporation mechanisms. Attempts to densify
wetted by molten metals but is not attacked by them, ZrB2 without sintering aids or with the addition of boron
making it a useful material for molten metal crucibles, resulted in rather low final densities.8 The properties of
thermowell tubes for steel refining and parts of electrical the dense materials then become strictly dependent on
devices such as heaters and igniters.1 Applications are the starting powders and processing parameters as they
also found in the aerospace industry: hypersonic re-entry determine microstructural features such as grain size,
vehicles, leading edges, nose caps, rocket nozzle inserts volume and chemistry of the secondary phases, etc.
and air-augmented propulsion system components.1
4 As the reliability of ZrB2-based materials for electrical
To ensure that favourable properties are obtained, con- and mechanical applications is limited by their poor
trol of densification and microstructure is necessary properties of toughness, stress corrosion cracking and
because strength and corrosion resistance are adversely high temperature oxidation, many applications of these
affected by porosity in sintered bodies. Due to the high materials require the addition of a second reinforcing
melting point and high vapour pressure of the constituents, phase to form complex structures or alloying elements
to form solid solutions.9
17 Studies were made on cera-
* Corresponding author. Tel.: +39-546-699-759; fax: +39-546-
mic composites containing ZrB2 as a second reinforcing
46381. and electroconductive phase,9
12,18
24 thanks to the low
E-mail address: bellosi@irtec1.irtec.bo.cnr.it (A. Bellosi). electrical resistivity of ZrB2.25
27 Among the advantages
0955-2219/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved.
PII: S0955-2219(01)00284-9
280 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288

of electroconductive monolithic or composite ZrB2-based shown in Table 2. A Ni powder was selected as a sin-
ceramics, the possibility of machining them by electrical tering aid for the monolithic and composite materials.
discharge offers a powerful tool for the manufacture of The powder batches were homogenised by wet mixing
complex shaped components.26
29 Another critical aspect in ethanol using zirconia balls. The hot pressing cycles
of refractory borides is their low thermal stability in air at were carried out under vacuum using a BN-lined gra-
high temperature, although several studies have shown the phite die.
formation of a surface protective scale of borosilicate The microstructure was analysed by scanning electron
glasses.3 Again, attempts were made to improve the oxi- microscopy and EDX analysis on polished and fracture
dation resistance of ZrB2-based materials, through the surfaces. The crystalline phases were identified by X-ray
addition of appropriate additives.9,20,30
32 diffraction analyses.
This work aims to highlight the microstructure and The following mechanical properties were measured:
properties of two ZrB2-based ceramic composites elastic modulus (E) by the resonance frequency method
obtained from fine commercial ZrB2 powders and the on 2880.8 mm3 samples, 4-pt flexural strength () on
addition of TiB2 or B4C as secondary phases. Micro- 252.52 (lengthwidththickness) mm3 test bars,
structure, mechanical properties, electrical resistivity using an outer span of 40 mm and an inner span of 20
and oxidation resistance were evaluated and compared mm with a crosshead speed of 0.5 mm/min, from room
to those of monolithic ZrB2.33 temperature up to 1400 C; microhardness (HV1.0) by a
Vickers indenter with an applied load 9.81 N; fracture
toughness (KIC ) using two methods: (i) direct crack
2. Experiments measurement (DCM)34 with a load of 98.1 N, (ii) chev-
ron notched beam method (CNB)35 on 2522.5
ZrB2, TiB2 and B4C powders produced by H. C. (lengthwidththickness) mm3 bars with a machine
Starck (Germany) were selected as raw materials. crosshead speed of 0.05 mm/min.
Their main properties are summarised in Table 1: the The thermal expansion behaviour was tested in air up
chemical composition was supplied by the producer to 1400 C.
except the oxygen amounts which were determined by The electrical resistivity measurements were carried
LECO. The specific surface areas were measured by the out by a four probe DC method at room temperature,
BET method while grain size and morphology were inducing a current in bar specimens of 22.525 mm3.
evaluated by SEM analyses. The particle shape of all the The current and the voltage readings were detected at
powders was irregular and mainly acicular, with a large the same time in two different digital high-resolution
grain size distribution. multimeters. The resistivity values were determined
Compositions of the two ZrB2-based composites, of from the electrical resistance measurement taking
the monolithic material and the adopted parameters account of the test lead distance and cross-section area
necessary to obtain high densities by hot pressing are of the samples.

Table 1
Characteristics of the raw powders (from H. C. Starck)

s.s.a. Grain size Crystalline Significant impurities (wt.%)

(m2/g) (mm) phases
C O N Fe W Si Al Hf Other metals

ZrB2 1. 0.1–8 ZrB2 0.25 1.0 0.25 0.1 0.2

TiB2 4.0 0.1–4 TiB2, WC, TiO2a 0.14 2.5 0.19 0.12 2 0.1
B4C 13.3 1.5–4.5 B4C, Ca 1.7 0.7 0.05 0.15 0.5 0.1
a
Traces.

Table 2
Composition, hot pressing parameters and densities of the hot pressed samples

Sample Composition Hot pressing parameters Density

(wt.%)
T ( C) P (MPa) Time (min) (g/cm3) (%)

A ZrB2+4 wt.% Ni 1850 30 30 6.05 98.0

B 55 wt.% ZrB2+41 wt.%TiB2+4 wt.% Ni 1600 30 30 3.65 100
C 83 wt.% ZrB2+13 wt.% B4C+4 wt.% Ni 1870 30 10 5.16 99.6
 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288 281

Long term oxidation behaviour at 1000 C was studied analysis revealed areas containing metallic Ni with tra-
in a laboratory kiln with interruptions in the tests in order ces of oxygen as marked in Figs. 1b and 2a. Metallic
to measure weight change at fixed times. Microstructure nickel could not be detected on X-ray diffractograms, as
of the surface and polished cross-sections of the oxidised its main peak is superimposed by a peak of zirconium
samples were analysed by X-ray diffraction analyses, diboride.
scanning electron microscopy and EDX analyses. SIMS analyses, reported in a previous work,36
revealed the presence of phases such as Zr–B–O, Ni–B,
Zr–Ni–O in this system. These phases show some degree
3. Results and discussion of reactivity between the different elements during hot
pressing. Ni reacts with the O2 present as an impurity in
3.1. Microstructure of the hot pressed materials the starting ZrB2 powders; subsequently, the exchange
reactions with ZrB2 give Ni2B and ZrO2. As confirma-
Material A reached a final density of 6.05 g/cm3 tion, small ZrO2 particles were detected, mostly within
which is about 98% of the theoretical density. The ZrB2 grains (Fig. 1b). Moreover, (Zr-O) compounds,
shape and size of ZrB2 grains turned out to be quite probably Zr2O5,36 were also present in the spherical
homogeneous; Ni-rich grain boundary phases were inclusions distributed mainly at grain boundaries (Fig. 2a
located mainly at triple points (Fig. 1a and b). It has to and b). The internal part of these inclusions consisted of
be pointed out that in the back scattered electron ima- an aggregation of small zirconium oxide particles
ges (an example of which is shown in Fig. 1b), the con- embedded in a boron oxide-based glassy phase where
trast arises from the channelling of the accelerated zirconium, oxygen, and low amounts of impurities like
electrons as a consequence of the different orientation of Al and Si segregated, as revealed by EDX analyses. It
the ZrB2 crystals and not from the change of the mean can be hypothesised that the boron oxide-liquid phase
atomic number of the phases present, i.e. variations in formed at high temperature as a consequence of the
composition and stoichiometry. From X-ray diffraction
analyses, traces of Ni2B were identified and micro-

Fig. 1. (a) fracture surface of material A: ZrB2+4 wt.%Ni; (b) back Fig. 2. Back scattered electron images of a polished section of material
scattered electron image of a polished surface. The following details A. In picture a: (1) Ni-rich grain boundary phases; (2) ZrB2 grains; (3)
are marked: (1) Ni-rich grain boundary phase; (2) ZrB2 grains; (3) spherical glassy inclusions, containing clusters of precipitated ZrO2
spherical glassy inclusions; (4) ZrO2 particles. nanoparticles. The microstructural features are highlighted in picture (b).
282 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288

concentration of oxygen and other impurities and their with different volume fraction and stoichiometry (Fig. 3).
reaction with boron and that the glass formed during On the basis of the cell parameters, two solid solutions
cooling from the hot pressing temperature. Zirconium were identified, one rich in titanium, (Ti0.8Zr0.2)B2, and
oxide nanoparticles later precipitated in this glass. SIMS the other rich in zirconium, (Zr0.8Ti0.2)B2, in amounts of
analyses36 brought out two important points: (i) the 44 and 37 by vol.% respectively. Of the starting borides,
presence of several BxOy and ZrBxOy species, indicating an amount of about 13% of ZrB2 was left, while TiB2
boron’s tendency to form clusters and, (ii) the presence disappeared almost completely indicating that under the
of other contaminants (Na, Ca, K, Cr, Ti) in addition to adopted hot pressing conditions, the dissolution and
those indicated by the powder supplier. segregation of zirconium in TiB2 was favoured in com-
The above reported results indicate that during hot parison to the solubility of titanium in ZrB2. However, a
pressing another melt formed, besides the Ni-based complete solution did not occur as the processing tem-
melt: a liquid glassy phase formed by B2O3 and metal perature of 1600 C was well below 2000 C. Additional
cations, such as Zr and impurity elements. It is likely factors, which could have influenced the rate of forma-
that a liquid film of B2O3 formed at grain boundaries tion of solid solutions and their stoichiometry, were: the
due to the oxygen released from ZrB2 powder particles. difference in atomic radii between Ti and Zr (Zr being
B2O3 melts at 723 C and vaporises at temperatures about 10% larger than Ti38), the heating atmosphere,
above about 800 C, but due to the applied pressure the cooling rate after hot pressing (slow cooling may
during sintering and to the fact that the liquid and induce decompositions15), the relative amounts of the
vapour boron oxide-based phases were contained within starting phases, the impurities present in the starting
closed interstitial spaces among solid particles, B2O3 powders, and the addition of metallic sintering aids
could not volatilise and remained in a liquid state. which promote the formation of metal borides as sec-
Probably dewetting took place, as the liquid assumed ond phases.14,15,38 In the present case a low amount of
approximately the shape of a sphere. This means that Ni2B was actually found. Oxygen impurity leads to the
once formed, the liquid did not wet the zirconium formation of zirconia particles randomly distributed
diboride particles (depleted of the amount of the surface (Fig. 4a and b) at grain boundaries. The various phases
oxygen) and the interfacial area was reduced due to a and relative compositions are marked in the micro-
strong driving force for shape change towards the geo- graphs of Fig. 4a and in the EDX spectra of Fig. 4c and
metrical configuration with the lowest surface energy. d. During sintering, two phenomena affect the forma-
The microstructural features (grain size and shape tion and distribution of solid solutions:38 (i) a composi-
and grain boundary phases) confirmed also that the tional gradient in the solid solution grains, which should
presence of Ni promoted the formation of a liquid phase be limited to the grain surface for equilibrium reasons;
at high temperature, similar to the phenomena observed (ii) a graded concentration of solute at triple points and
during the sintering of TiB26 and in agreement with within ZrB2 grains, since segregation is enhanced by the
previous results on the effects of the addition of Fe to different diffusion coefficients. Both these features were
ZrB2.7 The Ni-based liquid phase has two main effects: observed in the micrographs of Fig. 4a and b. Increased
(i) it favours powder particle rearrangement, and (ii) solution concentration caused lattice dilatation. These
enhances mass transfer mechanisms such as activated microstructural features affect the material properties.
volume diffusion, surface diffusion and formation of a In particular, as lattice anisotropy and elastic strain
solid solution.37 The temperature of this liquid phase
formation, as indicated by the densification curve, is
about 1260 C. Evaporation phenomena of ZrB2 was
also retarded due to the presence of liquid phase and
enhanced grain growth was avoided. Being dependent
upon the wetting behaviour, which is influenced by the
surface oxidation of ZrB2 particles, some residual pores
which were not completely infiltrated by the liquid
phase remained at the triple junctions.
Composite B, produced by the mixture ZrB2+
TiB2+Ni, reached a density of 5.65 g/cm3 which is
higher than the theoretical one calculated on the basis of
the starting components. This was due to the formation
of (Ti, Zr)B2 solid solutions as the two starting diborides
are mutually and fully soluble at T>2000 C.16,38 The
X-ray diffraction spectra of the crystalline phases and
their relative amounts pointed out that the main part of Fig. 3. X-ray diffractograms of sample B (composite ZrB2/TiB2), evi-
the starting phases formed two solid solutions (Ti, Zr)B2 dencing the formation of two different Zr–Ti–B solid solutions.
 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288 283

Fig. 4. (a) Back scattered electon image of a polished surface of composite B. The following phases are indicated: (1) ZrB2; (2) TiB2; (3) (Zr,Ti)B2;
(4) (Ti, Zr)B2; (5) ZrO2. (b) Secondary electron image of the same area represented in (a). (c) Windowless EDX spectra from the two phases indi-
cated as 1 and 2 in picture (a). (d) Windowless EDX spectra of the two solid solutions indicated as 3 and 4 in picture (a).

energy increase in the solid solutions,38 they determine concurrent factors: (i) the relatively low dissociation
the grain boundary cohesion in polycrystalline materi- energy for B4C that favoured the formation of volatile
als. Grain size was not homogeneous, and the material carbon oxide species and in boron (boron oxide) spe-
contained some defects and microcracks (Fig. 4a and b). cies, preventing also the formation of zirconia particles,
It was not possible to confirm the amount of each phase (ii) the reaction between boron oxide and nickel (prob-
by image analysis because the correspondence grey ably nickel oxide) that resulted in liquid phases (above
level-phase was disturbed by the channelling effect of about 1300 C).40 These liquids either could have been
the electrons. squeezed out by the applied pressure due to their low
Composite C, produced with the addition of 13 wt.% viscosity or could have originated volatile species, being
of B4C to the ZrB2+4 wt.%Ni matrix material, had a the processing temperature (1870 C) very high.
final density of 5.16 g/cm3 which is about the 99.6% of
the theoretical density. The microstructure (Fig. 5a–c) 3.2. Mechanical properties
showed ZrB2 grains of 5–15 mm mixed with B4C grains
of smaller dimensions (1–10 mm). The only crystalline The mechanical properties of the three tested materi-
phases revealed by X-ray diffraction analyses were ZrB2 als are summarised in Table 3.
and B4C. Very low amounts of Ni-based phases were
only randomly found within B4C particle aggregates 3.2.1. Young’s modulus
(Fig. 5c). Therefore, the main part of nickel introduced This property was clearly dependent on the composi-
as sintering aids was lost during hot pressing due to two tion. The value measured in material A is higher than
284 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288

values found in literature.1,14,15 The addition of TiB2, and phase. In this respect, material B was tougher than
the consequent formation of (Ti,Zr)B2 solid solutions, material A and material C was tougher than material B.
decreased its value in agreement with literature data.14
16 While the toughness values measured by DCM and
CNB were in a very close agreement for the composite
3.2.2. Hardness materials, the CNB value measured on material A was
The sample with the highest value resulted the com- higher than the DCM value. As the crack pattern gen-
posite containing B4C, however all the measured values erated by indentation in this material (Fig. 6) is defi-
were in the range reported in literature.1,4,16 nitely far from any reasonable approximation, scarce
reliability should be attributed to the data collected with
3.2.3. Fracture toughness this method. A direct comparison of our toughness
The comparison of the fracture toughness values values with those reported in literature is blurred by the
revealed the toughening effect associated with the addi- different compositions and densities of the tested mate-
tion of a second reinforcing phase or to a solid solution rials as well as the numerous models, equations and
experimental parameters used to calculate the fracture
toughness by DCM. The authors are not aware of any
fracture toughness estimation by CNB of this kind of
material.

3.2.4. Flexural strength

At room temperature the composite materials had a
markedly higher strength than the monolithic ZrB2
material (A) by virtue of their higher fracture toughness.
In general the critical defects which were found as frac-
ture origins were mainly due to processing faults, parti-
cularly the scarce homogenisation of the starting
powders which resulted in aggregates or pores and voids
(an example is shown in Fig. 7). At high temperature,
material A showed an increase of flexural strength in the
temperature range 600–800 C, reaching values higher
than 600 MPa, which are excellent for this class of

Fig. 5. Composite C. (a) Secondary electron image of the fracture

surface; (b) back scattered electron image of a polished surface, evi-
dencing the distribution of B4C particles (dark grains); (c) high mag-
nification of polished surface: arrows indicate the presence of residual
Ni-rich phases. Fig. 6. Vicker’s indentation on the polished surface of material A.

Table 3
Thermo-mechanical properties and electrical resistivity of the ZrB2-based materials
p
Sample  (m.cm) E (GPa) l (10
6 / C) HV (GPa) KIC (MPa m)  (MPa)

800 1000 1300 DCM CNB RT 600 C 800 C 1000 C

A 7.42 496 7.49 7.70 7.51 14.40.8 2.80.3 3.380.42 37124 616 8 62437 2373
B 15.55 439 7.79 8.25 8.63 17.70.5 4.30.3 4.090.14 599167 — 58228 25014
C 16.08 448 7.20 7.72 8.25 19.21.1 4.5.3 4.530.24 64386 — 57399 14810
 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288 285

for composite B from 7.8 to about 8.610
6  C
1 and for

composite C from 7.2 to 8.310
6  C
1. Above 1300 C,
an evident decrease of the thermal expansion coefficient
confirmed the occurrence of the grain boundary phase
softening. The thermal expansion coefficients measured
for these materials were generally higher than values pre-
viously reported: 5.510
6  C
1 from room temperature
up to 1000 C.1

3.4. Electrical resistivity

The electrical resistivitiy of the borides of transition

metals depends on the purity and microstructural fea-
tures of the materials. For a monolithic ZrB2, a value of
Fig. 7. Example of a crytical defect (composite C), due to either 9.2 m.cm was reported.25 In our case the presence of
agglomerates in the raw powder or powder processing procedure. nickel-rich grain boundary phase lowered the resistivity
to 7.4 m.cm. In the composite B, the formation of
materials, followed by a steep decrease at 1000 C. The (Ti,Zr)B2 solid solutions increased the resistivity. The
trend for the composite materials was a slow strength fact that ZrB2 is a better conductor than TiB225 suggests
degradation up to 800 C followed by a more evident that electrical conduction in the composite of inter-
drop at 1000 C. At 1000 C the load-displacement mediate compositions may be primarily via the zirco-
curves were markedly non-linear for all the three mate- nium-rich-solid solution phase. Regarding composite C,
rials. The general high temperature behaviour of the the electrical resistivity increased, as B4C has an intrin-
ZrB2-based ceramics can be mostly related to the char- sic resistivity of about 510
1 .cm.
acteristics of the grain boundary phase of these materi-
als. When a metallic phase was present, as for material 3.5. Oxidation behaviour
A, its moderate softening at intermediate temperatures
was beneficial for strength. When the temperature was The temperature (1000 C) selected to compare the
raised further, the excessive softening was instead detri- oxidation behaviour of the two composites with that of
mental for the same property. In the case of the com- the reference material falls in the range previously
posite materials, the grain boundary phase was unable defined of rapid oxidation for ZrB2,39 during which a
to compensate for the natural strength degradation due protective layer of crystalline ZrO2 forms and the oxida-
to the raise in temperature. tion kinetics is parabolic.33,39 The oxidation of zirconium
diboride follows the reaction: ZrB2+5/2 O2!ZrO2+
3.3. Thermal expansion coefficient B2O3, where zirconia is always a solid amorphous or crys-
talline product and boric oxide can be solid, fluid or gas-
The linear thermal expansion coefficient of the three eous, depending on the temperature.
materials increased with temperature (Fig. 8). In the In our case, B2O3 should have been liquid at the test
range from room temperature to 800–1300 C, the temperature. Due to the presence of Ni-rich grain
increase for material A was from 7.5 to 7.710
6  C
1, boundary phases and/or of additional phases in the
composites, the oxidation involved more complex phe-
nomena in comparison with those of pure ZrB2.
The weight change in function of the oxidation time
for the three tested materials (Fig. 9) was substantial and
approximately linear up to about 40 h for the composite
containing TiB2. There was a weight loss for the compo-
site containing B4C. For sample A it was possible to
measure the weight change only up to 24 h: longer expo-
sure produced a complete oxidation of the test piece and
the consequent fragmentation of the material.
The weight change curves reflected the contributions
of two simultaneous processes. One leads to weight loss
by the vaporization of volatile phases like B2O3, CO and
CO2, and it is supposed to occur mainly in material C.
Fig. 8. Variation of the thermal expansion coefficient in function of The other one promotes weight gain through the for-
the temperature, for the three tested materials. mation of ZrO2, TiO2, crystalline and glassy B2O3 or
286 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288

Fig. 9. Weight change due to oxidation at 1000 C of the three tested

materials. Fig. 11. Cross-section of sample C oxidized for 10 h at 1000 C. Voids
due to the oxidation of B4C particles are evident.

borate-based glasses and is mainly active in material B.

Further additional phenomena are related to the for- 3.5.2. Material B
mation of nickel oxide from the oxidation of the grain The oxidation products were a mixture of titania, zir-
boundary phases; Ni-O species subsequently react with conia and borate-based glasses. Due to the linear kinet-
B2O3 causing the formation of a liquid phase according ics observed up to 40 h of permanence at 1000 C,
to the phase diagram NiO.B2O3.40 chemical reaction between oxygen and (Ti, Zr)B2 could
The oxidation behaviours of the three materials can have been the mechanism controlling the process. Dur-
be described as follows: ing this first step the surface oxide scale was porous and
allowed oxygen diffusion towards the reaction interface.
3.5.1. Material A Longer exposure to oxidation did not cause weight
The only crystalline oxidation product was zirconia. change: B2O3-based melt probably sealed the surface
The cross-section of the sample oxidised for 10 h porosity in the oxidation product layer and protected
(Fig. 10) showed that the oxidation involved both ZrB2 the bulk from further oxidation.
grains and grain boundary phase. The reaction at grain
boundaries was very fast; the formation of nickel oxide 3.5.3. Material C
induced its reaction with boron oxide to liquid and solid The overall weight loss was due to the formation of
NiO.B2O3.40 Consequently this liquid phase progres- volatile species: B2O3 and /or fluid glass coming from
sively penetrated through the grain boundary channels the oxidation of ZrB2, of B4C and of the Ni-rich phases
towards the inner bulk. ZrB2 grains, surrounded by an (although their amount is rather negligible). This mate-
oxygen-rich liquid phase, underwent a fast oxidation. rial revealed the highest resistance to oxidation among
These phenomena led to material degradation up to the the three tested materials. In fact the weight of the
failure. sample, after a slight decrease in the first period of about
30 h, was almost constant up to 100 h of exposure at
1000 C. In this regime surface reaction scale impeded
further oxidation; bulk material adjacent to the reaction
interface did not undergo strong degradation. The skele-
ton of ZrB2 grains was not affected by oxidation,
although B4C particles partially disappeared leaving
voids (Fig. 11). One of the factors that prevented oxygen
transport towards the bulk and, therefore, improved the
oxidation resistance of this composite, was the quasi-
absence of Ni-rich phases and their consequent fluid
oxidation products along grain boundaries.

4. Conclusions

Fig. 10. Cross-section of the sample A oxidized for 10 h at 1000 C.

Dense ZrB2-based composites, in the systems
The surface oxidation product and the reaction at grain boundaries ZrB2+TiB2 and ZrB2+B4C, were produced by hot
are revealed. pressing. In all the compositions the addition of metallic
 F. Monteverde et al. / Journal of the European Ceramic Society 22 (2002) 279–288 287

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