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HIGH TEMPERATURE RELIABILITY OF
ADVANCED CERAMICS IN GAS TURBINES
Abstract-
This paper describes high-temperature reliability, particularly
creep and creep rupture behavior of three engineering
ceramicssilicon nitride, silicon carbide, and alumina-based
silicon-carbide-particulate ceramicswhich are considered the most
potential candidates for the use of blades of high-efficiency
ceramic gas turbine. The structural reliability of silicon nitride
is very often limited due to the softening of glassy phases formed
at grain boundaries. On the other hand, silicon carbide, which
generally does not contain glassy phase at the grain boundaries,
shows excellent creep resistance even at very high temperatures.
Finally, it is shown that creep resistance of alumina can be
markedly improved by dispersing nano-sized silicon carbide
particles into the grain boundary.
Introduction:
Because of their excellent resistance to tensile creep, advanced
ceramics have become a leading candidate for use of
high-temperature structural applications such as turbine blades and
nozzles. For example, current commercial grades of silicon nitride
now have a creep resistance that far exceeds that of commercial
high-temperature metallic alloys. This behavior is illustrated in
Fig following figure,
where the 1000-h lifetimes of two grades of silicon nitride are
compared with that of a single crystal alloy currently used in gas
turbine. For a given applied stress, the silicon nitrides are
capable of operating at the temperatures 300C higher than the
alloys. A ceramic gas turbine where such advanced ceramics are used
in high-temperature components such as turbine blades and nozzles
makes it possible to increase the turbine inlet temperature (TIT)
up to 1300 ~1400C, resulting in high thermal efficiency. For
example, in the Japanese national energy conservation project
entitled "300 kW Industrial Ceramic Gas Turbine (CGT) Research and
Development Project," a ceramic gas turbine with high thermal
efficiency has been developed to promote the high- performance
co-generation systems. This R&D project was initiated in 1988
with 10-yr scheme, and consisted of two major activities: ceramic
component fabrication technology, and component assembling
technology. The former has dealt with the development of new
ceramic materials, mainly silicon nitride, with high structural
reliability, and also the fabrication of components with
complicated shape such as turbine blades. The target properties in
this material development are the strength higher than 400 MPa with
the Weibull modulus higher than 20 at 1500C and the fracture
toughness larger than 8 MPa.m1/2 at room temperature. In the
latter, the developed ceramic components have been assembled, and
the performance of the ceramic gas turbine has been evaluated. The
target thermal efficiency is more than 42 percent with TIT of
1350C, axial output of 300 kW, and very low NOx emission.
Regenerative ceramic gas turbine
By adopting the developed silicon nitride to the components
including turbine blades, nozzles, combustor liners, and nose
cones, TIT can be increased without cooling; this leads to high
thermal efficiency of about 40 percent, as shown in Fig. 3,
FIG: Thermal efficiency curve
where the results are compared with those of a metal gas
turbine.
When ceramic materials are used for a high-temperature
structural component such as a gas turbine blade, the durability of
more than 100,000 h is required in many cases. Particularly at high
temperatures, creep and creep rupture behavior is most important
for the durability of the materials. However, it is practically
impossible to perform a thorough, long-term creep test of ceramic
materials at high temperatures for 100,000 h. Therefore, attention
is directed to the accurate estimation of long-term lives from the
results of short-term tests. Most of the creep tests conducted so
far on ceramics were at a level of several hundred hours, which are
too short to estimate the aforementioned long-term durability. The
collection of long-term creep rupture data is definitely required
to assess the extrapolation methods and to raise reliability in the
estimated durability. The problem in long-term creep testing
includes durabilities of a heat element, a thermo-couple, etc. For
example, damage sometimes occurs in a molybdenum disilicide heat
element in use for several thousand hours. The lifespan of a
thermo-couple with reliable accuracy is a few thousands hours. At
the National Industrial Research Institute of Nagoya, long-term
creep testing of ceramics at a level of 10,000 h has been conducted
using a special creep testing facility, where such parts are
renewable even during high-temperature operation [1]. This paper
deals with creep and creep rupture data accumulated so far for the
three representative engineering ceramics: silicon nitride, silicon
carbide, andalumina-basedsilicon-carbide-particulate ceramics.
Silicon Nitride
Silicon nitride has been recognized as one of the most promising
ceramic materials for high-temperature structural components for
nearly two decades, and high-temperature strength has been
substantially improved, as shown in Fig. 4. At high temperatures,
the strength is degraded and the structural reliability is very
often limited due to the softening of glassy phases, which are
formed at grain boundaries as a result of processing with sintering
additives. There are two regions in a delayed-fracture mechanism
map of silicon nitride at the temperatures above 1200C: slow crack
growth failure and creep damage rupture is shown. The former is a
fracture that occurs when a crack grows subcritically from a
pre-existing flaw and reaches the critical size. This is
predominant in the high-stress, short-term life region. The latter
is due to the formation of a macrocrack with the critical size by
cavity nucleation and coalescence. This prevails in the low-stress,
long-term life region. Generally, long-term durability for the
practical service is estimated from the short-term data. The
difference between these two fracture mechanisms is understood in
terms of creep rate properties, creep life properties,
microstructural changes, etc. The transition from the slow crack
growth fracture to the creep damage rupture one occurs when the
applied stress decreases below about 200 MPa.
fig 4 shows improved strength of ceramics at high
temperature
This section deals with three grades of silicon nitride; one is
hot-pressed silicon nitride doped with 5 wt percent yttria and 3 wt
percent alumina (hereafter referred to as SN A). A glassy phase is
present continuously around the silicon nitride grains, and the
thickness of the glassy phase between two grains typically ranged
from 1 to 3 nm, as shown in Fig. 5(a). The second is HIP-ed silicon
nitride doped with 5 wt percent ytterbia, etc. (hereafter referred
to as SN B). The third is a HIP-ed material that has been newly
developed in the aforementioned CGT project (hereafter referred to
as SN C). In this silicon nitride, very little glassy phase exists
at the interfaces, Fig. 5(b).
The creep curves of silicon nitride at high temperatures
generally consist of three regimes: transient, steady-state, and
accelerated creep regimes, similar to the case of metals. In
general, steady-state creep regimes are apparently observed in the
curves at 200 MPa or lower stresses; however, at the
stresses higher than 250 MPa, the failure is caused by slow
crack growth and occur in the transient creep regime. It is known
that in the sample where stable steady-state creep regimes are
evident, facet-sized cavities are sporadically distributed [2]
HYPERLINK
"http://scitation.aip.org/journals/doc/JERTD2-ft/vol_123/iss_1/64_1.html"
\l "R3" \n _blank[3] HYPERLINK
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\l "R4" \n _blank[4] HYPERLINK
"http://scitation.aip.org/journals/doc/JERTD2-ft/vol_123/iss_1/64_1.html"
\l "R6" \n _blank[6]. The SN C (CGT new material) where the grain
boundary is significantly strengthened, showed very good durability
even at 1400C, and the creep rupture hardly occurred at stresses
below 200 MPa. Figure 6 shows the tensile creep curves at 1400C
under 200 and 250 MPa; the creep strain was very limited, and
solely transient creep was observed until it failed. The TEM
(transmission electron microscopy) study revealed that almost no
cavitation was generated during the creep.
The creep life for structural ceramics is conventionally
expressed as
where tf is the time-to-failure, is the applied stress, and CL
and N are constants. The exponent N determines the stress
dependency of the life, and it is often referred to as a fatigue
exponent. Figure 7 shows the stress dependencies of the lives for
the SN A, SN B, and SN C. The plots indicate some changes in the
slopes around 200 to 250 MPa, and the estimated fatigue exponent
was 10 or higher in the high stress range and 2 to 3 in the low
stress range. The occurrence of the creep damage rupture
substantially limits the high-temperature structural reliability of
silicon nitride. It should be also noted that a longer creep life
in the low stress range tends to be erroneously extrapolated from
the short-term results obtained in the high stress range. In the SN
C, creep rupture hardly occurred at the stresses below 200 MPa.
It is well known that the creep life of silicon nitride follows
well the relation proposed by Monkman and Grant [9], and it is
expressed as follows
where d/dt is the steady-state strain rate, m is the strain rate
exponent, and CMG is a constant. This is a very useful methodology
to estimate the creep life. The relation between the creep life and
the strain for the SN A, SN B, and SN C is shown in Fig. 8. For the
SN A, the data fall well on a line with 1 for the m value,
irrespective of the applied stress or temperature. On the other
hand, for the SN B and SN C, both of which have stronger grain
boundaries, there are significant temperature dependencies of CMG
and m, and the m value tends to become larger than 1. The validity
of the Monkman-Grant relation for the SN A can be endorsed by the
dependence of the time for facet-sized cavity formation, tp, on the
strain rate. The TEM studies for the SN A specimens, most of which
were fractured in the steady-state creep regimes, revealed solely
the sporadic distribution of facet-sized cavities. Thus, tp can be
considered to be nearly equal to the total time-to-failure. Under
the constrained conditions, a product of tp and d/dt depends
primarily on the geometric parameters ,and then can be regarded as
a constant, leading to m = 1. However, the creep ruptures of the SN
B and C are governed by the subcritical crack growth emanating from
a pre-existing flaw, rather than the cavity formation. This
tendency becomes strong when the creep life is short at high
applied stresses, resulting in m>1.
Silicon Carbide
Silicon carbide generally does not contain glassy phases at
grain boundaries, even when doped sintering additives such as
alumina .Due to this rigid interface, the strength is not degraded
at very high temperatures; see Fig. 4. Because of the good
high-temperature mechanical properties as well as good corrosion
resistance, silicon carbide is one of the most important candidate
materials usable at high temperatures around 1400C. In this
section, creep and creep rupture behavior of silicon carbide doped
with 5 wt percent alumina 1400C is described. The TEM observation
revealed that there is no glassy phase at the interfaces between
two silicon carbide grains; even if any glassy phase is present,
its thickness is in the order of atomic dimensions, as shown in
Fig. 9.
Then, the measured creep rate of this material at 1400C, 200 MPa
is as small as 61012/s. No cavity is formed during creep, though
creep deformation should be controlled by grain boundary diffusion
[12], and creep failure is caused by slow crack growth from a
pre-existing flaw. The crack grows
subcritically along grain boundaries with diffusional process.
In this case, it is possible to apply a diffusive crack growth
model to the results, since this model is based on the assumptions
that a crack propagates along grain boundaries by surface and grain
boundary diffusion, the grains on both sides of the boundary behave
elastically, and the crack grows along grain boundaries in a steady
state at a fixed velocity. The crack velocity, V in this model is
given by the following relation:
V= CD[0.59(k/kG) + { 0.35 (k/kG) 2 --1}1/2]1/2
where CD is a constant, K is the applied stress intensity, and
KG is the critical K value predicted by Griffith's theory for
propagation of an atomistically sharp crack in an interface. KG is
related to true surface energy, and it is normally one order of
magnitude lower than the fracture toughness measured by using a
fracture mechanics test specimen. In this -K relation, there is a
threshold stress intensity, Kth, defined as Kth = 1.69KG. Below
Kth, the applied stress is not sufficient to drive the crack.
Applied stress, , can be related to K by K=Ya1/2 where Y is a
coefficient depending on the crack geometry and a is the crack
size. Combining this relation and Eq. (3) yields time-to-failure,
tf, as a function of (Fig. 10). The line is the tfcurve estimated
from Eq. (3). The curve agrees well with the plotted data. The
curve also can predict a threshold applied stress below which the
pre-existing crack does not grow and then the delayed fracture does
not occur. The predicted value is 165 MPa. This stress can be
considered as a safety applied stress, and the value is very
important in making a component design with this material. It is
interesting to compare the present results to those shown in Fig.
7. In the case of silicon nitride, the short-term results obtained
in the high stress range tend to estimate a longer creep life in
the low stress range. On the other hand, creep fractures of silicon
carbide are controlled by the diffusive crack growth from a
pre-existing flaw at all the stress levels, due to the absence of a
glassy phase in it. Then, in the stress-life diagram, the creep
life in the low stress range is much longer than that estimated
from the data in the high stress range.
Figure 10.
Alumina/Silicon Carbide Nanocomposite
In the field of creep of metals, it is well known that
high-melting oxide particles on the grain boundaries significantly
inhibit diffusional creep. For example, this is seen in
oxide-dispersion-strengthened (ODS) superalloys. This section
describes that the dispersion of silicon carbide particles with
nanometer size into an alumina matrix (alumina/silicon carbide Nan
composites) gives rise to significant improvements in mechanical
properties, particularly in creep resistance, which is very similar
to the case of metals [14] HYPERLINK
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\l "R15" \n _blank[15] HYPERLINK
"http://scitation.aip.org/journals/doc/JERTD2-ft/vol_123/iss_1/64_1.html"
\l "R16" \n _blank[16]. It has been revealed that, for example, the
minimum creep rate of alumina/17 vol percent silicon carbide
nanocomposite was about three orders of magnitude lower than that
of monolithic alumina [15]. The creep life of the nanocomposite was
10 times longer and the creep strain at fracture was eight times
smaller than those of the monolith at 1200C and 50 MPa. Another
feature is that the nanocomposite tended to show almost solely
transient creep until failure, while slightly accelerated creep as
well as steady-state creep was observed in the monolith. The TEM
study revealed that the intergranular silicon carbide nanoparticles
inhibited grain boundary sliding of the alumina/silicon carbide
nanocomposite, as shown in Fig. 11
As grain boundary sliding proceeded, the pinning effect of the
particle increased, resulting in the occurrence of transient creep.
Along with grain boundary sliding, the particles plunged into
adjacent alumina grains. As the byproducts of this process, the
small intergranular cavities were generated around the particles.
These cavities induced intergranular crack formation, leading to a
final creep failure. It has been also revealed that the interface
between the intergranular silicon carbide particles
and the alumina matrix was much stronger than the
alumina/alumina interface [17]. The rigid bonding at the
alumina/silicon carbide interfaces caused the inhibition of vacancy
nucleation and annihilation at the interface, which remarkably
improved creep resistance of the nanocomposite.
Since the intergranular silicon carbide nanoparticles have been
found to inhibit grain boundary sliding in the creep of the
alumina/silicon carbide nanocomposite, a threshold stress model can
be anticipated to be operating in this system, as is similar to the
ODS superalloys [18]. The minimum creep rates of the nanocomposite
and monolith at 1200C are shown in Fig. 12, as a function of the
applied stress. The plots of the monolith followed well one
straight line. For the nanocomposite, however, the minimum creep
rates at 35 and 40 MPa were remarkably lower than the line
extrapolated from the data in the stress range above 50 MPa. These
results strongly suggest the presence of a threshold in the creep
of the nanocomposite. The threshold stress can be very crudely
estimated to be in the range of 20 to 35 MPa. The creep lives, tf,
of the nanocomposite and monolith are shown in Fig. 13, as a
function of the applied stress. Again, while the plots for the
monolith followed well one straight line, those for the
nanocomposite were deviated from a straight line at the stresses
below 50 MPa; the creep lives of the nanocomposite at 35 and 40 MPa
were substantially longer than those estimated from the higher
stress range, although their measurements were interrupted after
10,000 h.
Figure 12.
Figure 13.
Ashby [19] proposed a model that the particles pin the motion of
the dislocations to estimate theoretically the threshold stress.
This mechanism is based on an assumption that vacancies can be
absorbed and emitted only at the edges of grain boundary
dislocations. Thus, the continuous provision of vacancies requires
movement of the dislocation. When a dislocation moves in the
particle matrix interface, it will interact with the particles
whose modulus, lattice parameter, or chemical composition differ
from those of the matrix. This interaction causes a pinning force
on the dislocation, resulting in a threshold stress, tr, given by
tr = CAGbB/ D where CA is a constant, G is the shear modulus, bB is
the Burgers' vector of grain boundary dislocation (bB = ~1.61010
m), and D is the particle spacing on the grain boundaries. CA is
0.8 for the case of classic Orowan's calculation of the pinning of
lattice dislocation by hard particles. However, at high
temperatures where climb is active, CA can be considered to range
from 0.3 to 0.8, depending on intrinsic local mobility at
particle-matrix interface (the kinetics of atom rearrangement).
Then, tr, estimated by Eq. (4), ranges from about 18 to 47 MPa.
This shows a good agreement with the experimentally estimated
threshold stress, 20~35 MPa. Equation (4) shows one advantage of
nano-particle dispersion; at the same volume fraction, the
threshold stress, tr, tbecomes larger when the particle spacing
(and then particle size) becomes smaller.
Summary
This paper describes high-temperature reliability, particularly
creep and creep rupture behavior of three engineering ceramics:
silicon nitride, silicon carbide, and alumina-based
silicon-carbide-particulate ceramics, which are considered the most
potential candidates for structural application below 1500C. A
ceramic gas turbine where such advanced ceramics are used in
high-temperature components such as turbine blades and nozzles
makes it possible to increase the turbine inlet temperature up to
1300~1400C, resulting in substantially high thermal efficiency.
Finally, it should
be noted that, though nonoxide systems such as silicon nitride
and silicon carbide can be viable for structural application below
1500C, oxide-based ceramics are needed to be used for the service
above 1500C, which will be required in the near future,
Nomenclature
a =crack size
bB =Burgers' vector of grain boundary dislocation
CA =constant
Cd =constant
CL =constant
CMG =constant
D =particle spacing on grain boundaries
d/dt =steady-state strain rate
G =shear modulus
K =applied stress intensity
KG =critical K value predicted by Griffith's theory for
propagation of atomistically sharp crack in interface
Kth =threshold stress intensity
m =strain rate exponent
N =constant
tf =time-to-failure
tp =time for facet-sized cavity formation
V =crack velocity
Y =coefficient depending on crack geometry
