A Review Of Sic In Gas Turbine Engineering Essay

Published: Last Edited:

This essay has been submitted by a student. This is not an example of the work written by our professional essay writers.

Silicon carbide is a ceramic material with properties that make it a great choice for high heat and high stress applications.  One such application is in gas turbine engines for either energy production or propulsion via aircraft engines.  Current research into this topic has yielded results that, at a first glance, look promising for increasing the efficiency of these turbines.  However, during experimental tests it is clear that research has a long way to go in failure analysis and failure reduction before SiC turbine blades are ready for widespread commercial use.  In this paper, we will look at how the blades are made, the various failure mechanisms, and possible solutions to these failure mechanisms.


SiC turbine blades can be manufactured in a variety of different ways.  One of these ways is by sintering the SiC into the desired shape.  Sintering is a manufacturing technique that takes the sample in powdered or granular form with relatively high purity and consistency, places the powder into a mold, and then heats the mold at high temperature for a period of time.  This allows the particles to fuse into a single crystal in the shape desired.  According to Maccagno et al, there are indications that the addition of boron and carbon to the SiC mix can aid in the properties desired for the product.  For example, it was discovered that the addition of boron increased the diffusivity of the silicon and carbon atoms in the crystal.  Carbon additives form additional SiC and CO gas by reacting with leftover SiO2 molecules.  This has the added effect of reacting with any boron added to the mix to produce BO, which allows for the neutralization of boron in any future sintering processes.

One process for sintering by Yoon, H. et al used a β-SiC powder where the average particle size was approximately 30nm in diameter.  They tested the addition of Al2O3, Y2O3, and SiO2 as sintering additives.  Their additives had an average particle size of 1 μm.  The SiC and additives were then mixed thoroughly by a ball-mixing machine where the container and 10mm-diameter balls were made from zirconium.  They mixed the sample for 12 hours at 160 rpm.  The sample was then dried for 24 hours prior to sintering at varying temperatures for varying times.  The results of this can be seen in the following table:https://lh3.googleusercontent.com/bpvFoCP_pRnsNcZaBGWlFz_m8Q3ik3TwgjyTYWhDEnCII_IiuK5VLmrBx_X_JrUE-mEXmMWra75XQATlVdx_U3dzGyTsn1rJ5JLdThtmAlJxdLEFLoY

One design for a set of turbine blades has a core part where blades can be slotted into the correct position.  This is better than a single casting of the part because it allows for the selection of a desired crystal direction for the maximization of tensile strength and creep reduction.

The turbine blades have a dove tail to allow them to slide into place and also be easily replaced if damaged.  The turbine blades in this patent also make use of a two part system to reduce weight.  The blades are made from a core insert, which handles the stress needed to remain in the slot via the dovetail, and the outer shell, which helps to direct the air and has a modified crystal direction to accommodate the difference in shear stress direction and strength. https://lh5.googleusercontent.com/BTUa1g8eI5p3PZvTq4v9Fr117jy1L1aZN-z3y68pwOJLN6_V_LLkuEWY_0BqetX3ouJJiR2QC2DNKguBPfWmIJcTvJPbKLbGsgoDbgtYJKw3NWz-IH4

A melt infiltration process can also reinforce the part, giving it additional properties that enhance its ability to handle thermal stresses and increasing cracking strength.

Ceramic matrix composites such as in the core of these blades usually contain anywhere from 20-30 volume percent silicon in the SiC matrix.  This low SiC volume percent leads to a reduction in vibrations as it is typically less rigid than a part manufactured with higher volume percents.

These cores can also be modified by using microspheres of other materials such as carbon up to 50 volume percent.  These additives would typically need to be added to the mix prior to the melt infiltration.  The carbon additives change the modulus slightly while simultaneously reducing the density of the material.  This reduction in density makes it easier for the turbine to resist creep deformation due to high RPM while in use.

After both the inner and outer portions are manufactured, they are assembled into a single piece and cured in an autoclave.  Any remaining voids are filled with an additional silicon melt infiltration.

Failure mechanisms - Wilson

In the application, actually the turbine blade is made up of SiC composites. So the reason why the blade fails deals with the interaction of the composite material instead of just the SiC itself. There are several different failure modes in the real situation.

The first one is the fatigue of the material in the turbine blade. As is known to all that after repeated loading on one material, the material would fail due to the internal structure change. There are also several types of fatigue about the failure of the SiC composite. They are the low-cycle fatigue (LCF), high-cycle fatigue (HCF) and the thermo-mechanical fatigue (TMF). The LCF and the HCF are quite normal among alloy. LCF is referring to cycles less than 103 cycles and HCF is referring to cycles more than 104 cycles. The TMF, however, is the overlay of cyclical mechanical loading and the thermal loading. In the turbine environment where much combustion and mechanical loading happens, the blade would subject to such fatigue after a long time.

        The second failure mode to be worried about is the hot corrosion in the SiC composite. In the SiC composite, usually there are V2O5 and Na2SO4 content. They have quite low melting temperature. V2O5 is approximately 963 K and Na2SO4 is 305 K. However, in combustion environment of the turbine, the temperature would reach the level of over 1000 K. These salts would melt in such a high temperature. The reaction of sintered α-SiC with melted Na2SO4 would produces large amounts of SiO2 and thus leads to dramatic etching of the SiC substrate. These melts lead to rough friable scales in the substrate, which would cause pitting in the substrate. Moreover, it seems that carbon promotes the dissolution of SiO2 as CO and CO2 forms in the process. Actually, the corrosion effect in Na2CO3 is more severe than in Na2SO4.  In order to reduce the effect of hot corrosion with Na2SO4 , small amount of Na2CO3/NaNO3 are added to enhance the performance.

Cyclic Oxidation is another thing to worry about in the turbine blade. Actually, according to the study recently, the SiC ceramic itself shows quite good performance under the oxidation environment. As we can see in the graph about the oxidation performance with regard to the specific weight change after oxidized for long time.


However, oxidation and stress interactions are complex, especially in engine environment. Stresses would form microcracks and cavities and alter diffusivity. They all increase the effect of oxidation. At intermediate temperature (~500 to 1000 degree Celsius) region, the amount of SiO2, which would not etch the SiC substrate if not sintered, formed is insufficient to seal cracks in the SiC matrix. Moreover, in combustion, alkaline elements melting would dissolve the SiO2 film. Then the following oxidation would cause the SiC composites to degrade in their microstructure and mechanical properties. When the SiC composites are oxidized during thermal cycling, there are stresses because of the coefficient of thermal expansion (CTE) mismatch between the silica scale and the SiC substrate.

The CTE for different materialshttps://lh3.googleusercontent.com/njNL4k3ysxJ_Llm1wwHuqrJYt2lMwHLeQOde7uIihNlUG-enTwoMN4oAZLV8CKZOAvd3YOxqHgcRvUQ0hotg68vIpiPh6n1qTTb6NPbYDaFd5xld7ZY

There are many corrosion/erosion reactions happening in the engine. As the aircraft is in the air, the engine on both sides of the wings would take in much gas and some of them maybe corrosive to the turbine blade. Furthermore, there are some impurities in the manufacturing of the turbine blade that may also have influence on the corrosion. In a reference study on corrosion issues in gas turbines, it has covered the effect by the water vapor, CO2, other aggressive gases and some metal cation impurities that may decrease the performance of the SiC composite material. Relatively speaking, the effects of CO2 oxidation and aggressive gases such as HCl and SO2 actually are infinitesimal compared with water vapor. However, the water vapor in the air is kind of deadly contributing to the reduction of lifetime of the turbine blade. Water vapor makes the transportation of impurities to the ceramic surface easier, increases the intrinsic oxidation rate and leads to the volatilization of silica by forming Si(OH)4. Chemical reaction equation is SiO2(s) + 2H2O(g) = Si(OH)4(g). Not only the water vapor from exterior air should be one of the concerns, but also should the K or Na cation impurities be issues. These cation elements would open up the SiO2 network and accelerate the oxidation. SiO2 is more vulnerable to these impurities than other protective oxides. As has mentioned above, the turbine blade's performance would be greatly degraded by oxidation.


As the blade is usually under loading, the creep effect can't be neglected in this case. The creep is the slow deformation of material under loading. For the Sic ceramic, to achieve the life time of 10000 hours, it was estimated that static tensile  stress of 300MPa at 1038C, 250MPa at 1150C and 180 MPa at 1350C cannot be exceeded. The creep-strain vs. time curve is shown as follows. Actually, if failure does not happen at loading, SiC is not likely to creep-rupture. The slow crack growth initiated from the surface results in the failure of the SiC. The following graph is the creep behavior of SiC under different temperatures.https://lh5.googleusercontent.com/Bwara7ehc4bkB9zP99bnSNNxLJxwnn5RjHy023-87myPH2OqpP3Y6ylFeIPoHUuii-51zXJREkSPD_XZpsuWpSZgdiE6UIn0JRC8nA4Igka9EBpCMU0


In the engine environment, as the blade is rotating all the time and scratch the other surrounding materials. In that case, fretting/wear effect is something that must be taken into consideration. According to definition, fretting is the wear and sometimes corrosion at the rough surface. It leads to mass loss in the material. This damage is induced under load and in the presence of repeated relative surface motion.

In conclusion, all the factors contribute to the possible failure of SiC composite turbine blade material. However, effects such as fretting, creep and fatigue and only be reduced because most of the material are experiencing this change. Our main focus is on the hot corrosion, oxidation and erosion part that fail the engine. The lifetime of the turbine blade should thus be thoroughly analyzed and tested before put into use.

Mathematics of failures - Jack

SiC structures are manufactured using a sintering process. Unfortunately, this process limits both the size and the complexity of structures that can be made from the material. Despite this, NASA and other space agencies take good advantage of SiC structures due to both its chemical inertness and its excellent mechanical properties. As a result, heat exchangers and chemical reactors are built using brazed SiC components. Brazing is a process similar to soldering in which a filler material is heated beyond melting temperature and wets the parts to be bonded, after which it cools to form a solid joint. The BraSiC process utilizes filler composed of Silicon and a proprietary metal; brazing temperatures approach 1400°C. Observation of the failure of these joints is conducted along with an analysis of the mathematics involved. A four point bending experiment is conducted as shown in Figure 1. ASTM C1211-02 and NF EN 843-1 are consulted to ensure no issues arising from friction by using a custom fixture. The sample is a joint of two SiC bars brazed using BraSiC CEA1 and CEA1. The dimensions of samples are approximately 4mm x 3mm x 46mm, and joint thicknesses range from 3mm to 200mm. Modulus of bulk SiC is assumed to be 416GPa, and through nano-indentation it is determined that CEA1 and CEA2 moduli are 154GPa and 171GPa, respectively. Load is applied at a rate of .5mm/min. Failure is defined as failure of the joint itself or failure of the interface between the brazing and the SiC.

Figure 1: 4 point bending machine. Load applied from above onto disc on compression block.

Stress on tensile surface is calculated to be

where F is the maximum load, L is the outer span, l is the inner span, b is the sample width, and h is the sample thickness.

CEA1 shows good strength, with strengths in the range of 194±42MPa for thicknesses between 3mm and 19mm, and ranging within 115±42MPa for thicknesses approximately 93mm.

Braze size is small compared to beam thickness and length, as is the case in any commercial setting. Finite element analysis can be used to analyze the seam, but it requires the use of refined meshes. A macro (ignoring joint thickness) or micro (ignoring sample geometry) method can be used to avoid this complication. Using a combination of these two approaches, causes of fracture can be pinpointed. Assuming Lamé's coefficient of lS=84.4GPa and mS=179.3GPa for the bulk material and Lamé's coefficient of lL=228.6GPa and mL=57.2GPa for the brazed joint, the elastic problem in We can be found in the outer expansion

The first term has no discontinuities through G (line) and can be solved through finite element analysis without any modifications to the mesh. Near the origin (O), the term can be expanded to the outer expansion:

where T is the tension on the lower part of W0 and

The stress fields are

U0(O) is the vertical deflection at the origin, and the singularity theory indicates that Trt1(Q) is a set of the higher powers of r.

Dilation around O is used to zoom in on the origin, as that is the area where the crack is expected to occur. The following assumption (referred to as the inner expansion) is valid inside the joint:

This assumption coincides with the outer expansion when U0(O)=W0(y1,y2) and Trt1(Q)=eW1(y1,y2). Assuming d*/dxi=e-1d*/dyi at i=1 and i=2, it is safe to state:

W(hat) refers to the displacement field. To ensure its continuity and knowing the interface occurs at y1³.5 and y1£-.5, the following piecewise function of W1(y1,y2) is derived:

Wa refers to the left side of the SiC substrate, Wb refers to the joint, and Wc refers to the right side of the SiC substrate.

Figure 2: The second term of the inner expansion

As can be seen in figure 2, and ensuring continuity, it is calculated that:

Using previous derivations to ensure further clarity, it is found that:

No longer ignoring the boundary condition, it is clear that:


The above set of equations indicates that s12=s22=0 only when necking is identical in both SiC substrate and brazed solder joint, or when:

Obviously it is extremely unlikely that such an event would occur due to the vast difference in the mechanical properties of the brazed joint and the substrate. W(hat) requires an additional term to combat the boundary condition problem outlined in the equations above. This term is as follows for -.5£y1£.5:

Since the thickness of the joint is 1, the force from this is -TBx1, which does not allow the solution to approach zero at infinity. It behaves as the point force solution:

This force approaches zero in the first direction and approaches B in the second direction, which is a far more suitable solution. As F(y1,y2) is singular as r approaches both zero and infinity, W(hat) can be once more rewritten as:

Where W(frown) is equivalent to W(hat) and f(r)=0 when r<r1 and f(r)=1 when r>r2. f(r) eliminates the singularity of ln(r) and guarantees proper upper limit behavior. W(frown) behaves as 1/r with a mode of 1/rt-1(q) where:

At infinity, the second displacement field, W(frown), behaves as follows, where ~ is an operator for behaves as at infinity and D0 is the stress intensity factor:

The second displacement field and the stress intensity factor can be analyzed using finite element analysis. The outer expansion can now be rewritten as:

At the origin (G), U0 is an under-elongating factor, which leads to a jump in the horizontal direction of the joint, as seen in Figure 3. The point force returns where necking is inconsistent between joint and substrate.

Figure 3: Discontinuity and point force on first order outer expansion.

Matching and splitting the outer expansion gains the result:


with U(hat) being the displacement field of U. It follows the boundary conditions of 1/rt-1(q).

Crack propagation occurs most readily at the interface between joint and substrate. Most accurately predicted, cracks occur either on or within a small distance of the edge and follow the interface as seen in figure 4.

Figure 4: Starting vertical loading and that associated with W(hat).

Beginning mathematical modeling of cracking in the structure, crack length is assumed to be much less than solder layer thickness. The crack length with respect to layer thickness is x=l/e. W(hat) must now meet a new criteria that factors in the now propagating crack. These conditions are as follows

Where y1=-1/2 on both sides of the crack and 0<y2<x. Since the force is constant with these introductions, the boundary conditions are the same, and it can be determined that:

Originally, D0 is sufficient to calculate W(frown) due to the lack of a crack in the material, but upon introduction of a crack, the outer expansion must be modified such that D0=Dz. As a result, the behavior of a crack in the joint or the substrate of an SiC brazed material during initial propagation, but upon failure, the characteristics of the mechanical properties of the material will drastically change.

Assuming f to be the contour of the origin and n to be the normal towards the origin, the energy of crack nucleation can be modeled by:

If the integral is path independent, substituting in displacement fields yields:

The energy release rate can thus be derived:

Using expansions and splitting, the energy release rate can be simplified to:

Y(t1,t-1) is a constant, and much of the equation is dependent on crack length, so using finite element analysis, it is possible to further simplify g(x) to a linear relationship such that:

Failure is mathematically modeled to occur in two cases, one energy related and one stress related. Assuming GC to be toughness and sC to be tensile strength, these models are:

Substituting from previous derived equations,

Experimental analysis shows that crack length and solder thickness vary directly. As a result, the normalized crack length is a constant and so is g(x). As is seen in laminates, the energy based failure model dominates at low joint thickness and the stress based failure model dominates at higher joint thicknesses. The thickness of the joint that correlates to the change between stress and energy failure can be derived as follows and can be seen in Figure 5:

Figure 5: transition between energy and stress driven failure mechanisms.

Figure 6: collected 4 point bending failure data (CEA1-squares, CEA2-diamonds) with expected values plotted on line.

Figure 6 shows the data collected from the experiment using the two different grades of SiC solder tested (CEA1, CEA2). Extremely thin solder layers showed extremely high strengths, with some not failing at all during testing. T values approach sC=84MPa, which can be calculated at e0=92mm, which can be interpreted as the tensile strength of the solder-substrate joint and the transition joint length, respectively. The given fracture toughness of CEA1 and CEA2 is 1.5 MPa m.5, which is assigned as the upper possible value of the toughness of the solder-substrate joint for these calculations. An assumption can be made that the normalized crack length is equal to 4, which correlates to the minimum length required for crack propagation. Following e0, clearly the relationship between tensile stress and crack length is constant. For thicknesses values below e0, a more complex relationship exists, which can be modeled by:

Thus, it is found that the failure mechanism does not depend on the logarithmic point force that is generated by the necking differences in the substrate and the joint. A least square regime can be used to calculate z, but that requires a large number of tests. It can also be calculated knowing the stress T at failure for certain e.

The failure mode depends on the energy regime in low joint thicknesses and the stress regime in high joint thicknesses. The stresses left over from the welding/soldering process are ignored in the research conducted above. It is estimated that a simple shift in the failure curve would occur from these effects. Calculated values for z hold true because it is relative to the existing system that ignores internal stresses.citation

This research gives a good understanding of the behavior of cracks in brazed SiC material. It is clear that both the energy of the system and the applied stress are relevant when it comes to the failure mechanism. It is clear that especially in high stress applications such as in turbines, a smaller solder joint, in which the energy of the system dominates, is far more desirable. The mathematical derivations above can be used to determine the failure scheme for brazed SiC structures without extensive experimentation. This ensures that the research and processing methods for SiC parts is as low as is possible.

Citation: L.M. Nguyen, D. Leguillon, O. Gillia, E. Riviere, Bond failure of a SiC/SiC brazed assembly, Mechanics of Materials, Volume 50, July 2012, Pages 1-8, ISSN 0167-6636, 10.1016/j.mechmat.2012.03.001. (http://www.sciencedirect.com/science/article/pii/S0167663612000518)

CMCs with EBC, TPS - William

Potential ways to modify the property and avoid the failure of SiC:Modifications of SiC ceramic matrix composites equipped in gas turbine.

we consider initially several ways to make improvement of the SiC composites. Some of them are EBC and TPS, which is environmental barrier coating and thermal protection systems. Firstly there is introduction of two kinds of non-EBC methods.

1.Modification of silicon-carbon film properties under high energy ion beam irradiation

High ion beam irradiation is an efficient method introduced by X Redondas at el. The properties of amorphous hydrogenated silicon-carbon layers by can be changed  by using 4He+ ion irradiation [1].

The changes observed in the refractive index and thickness under the ion beam irradiation can be explained as a consequence of the densification of the films, because  the weakly bonded hydrogen remove and the structural changes,  mainly leaded to the formation of silicon-carbon bonds [2].


Fig1 shows the effects on the ion beam irradiation on the refractive index in different deposition conditions. It can be explained by the results of more stable bonded hydrogen in the samples with higher refractive index values. In this kind of material the refractive index is related to the composition and the porosity [3]. Meanwhile, the changes observed can also illustrate the thickness under the irradiation, which could be interpreted as the densification of the SiC. Because of the removal of the bonded hydrogen, the structure changes, leads to the formation of silicon-carbon bonds. Finally the consisting in the formation of silicon-carbon bonds after the Si-H and C-H bond breaking results in a densification of the irradiated area, and  then oxidation resistance is improved.[1]

2. Hot pressed 2D Tyranno-SA/SiC composites

Another highly technical method for improving SiC performance is Tyranno SA/SiC composites with or without carbon coating that are introduced by researchers in Kyoto University[4]

Hot pressing is an effective manufacturing technique.[5] But there are damages to fiber reinforcement such as the decomposition of silicate-oxygen-carbon phase and creep deformation which leads to the degradation.

Tyranno SA fibers do not degrade or modify composition at 1000℃. https://lh4.googleusercontent.com/d1TPJqS_6IjkyPPMef5cQ4c3zybFV8BIgXiZkfhHXZq9yNFzgX_hQ_nNLBfUE5ECtDCgNeMwjHi4AV2xg8tbBJXHz_vdROB99WHXWzCw86a-HQRLZ4I

Through first process by using non-coated fiber was beneficial for obtaining higher density to create prepared conditions for fibers coated in fiber. Then applying carbon coated fiber, composites would obtain fracture behavior that would not be destroyed.

Table.1 indicates that increasing temperature will increase the densification. Meanwhile, the mechanical properties are also improved.

Generally, ultimate strength and elastic modulus were improved and the fiber degradation can also be eliminated by the coating.

3.Several  ways mainly using EBC and TPS techniques.

3.1 C/C-SiC coated with SiC-B4C-SiC-cordierite oxidation protection system

It is known that corrosion/oxidation is two main failures of SiC composites.[6]

One of the promising protecting system is a SiC-B4C-SiC with multiple coating layers. The reason for adopting these materials is that the coefficient of thermal expansion is lower than the CTE of typical coatings such as mullite. Besides it also contains self-healing process. However there are also limitation because of the volatilization and moisture sensitivity against hot gas corrosion.[8] So in order to improve system, researchers from Stuttgart University[7] developed a new multilayer coating. It consists of a cordierite (2MgO Ã- 2Al2O3 Ã- 5SiO2) layer outside and a CVD-Bora SiC (SiC-B4C-SiC) layer inside [9].

From the oxidation behavior of the test it showed that the degradation was not observed in the SiC-cordierite system and this system reduced the mass loss by about 43%.[10] Meanwhile, the residual strength is also improved. The mass changes are indicated in fig 2.

The microscopic testing are operated to investigate the oxidation rate. Results showed that the cordierite layer behaves good adherence to the BoraSiC. From the SEM micrographs it also showed a crack density reduction of 58% at the SiC-cordierite interface.[7]


                                             Fig 2

In order to improve the properties of the EBC system, a thermal treatment at 1300 C for a short time is also useful for stabilization. Nevertheless,  the parameters of the crack density reduction is not very sufficient to achieve an optimized protection, so the further research work onto it should focus on those parameters.

3.2.Traditional and advanced EBC  applications:

Some efforts are undertaken by  researchers  carried out with NASA by the High Speed Research Program.

3.2.1.Baseline mullite/zirconia EBC:


One coating system consisting of mullite layers and additional top laryer of zirconia layer on SiC to solve the corrosion problems.[11]

The advantages of the system is that the coefficient of thermal expansion of the mullite in that closely similar with the expansion of SiC.[12] However, zirconia's coefficient of thermal expansion is double SiC's. This leads to cracking in the coating during cycling. Another major issue with the mullite coating is thermal sprary processing causes dissociate and deposit which also lead to cracks. fig.3 is the summary of the problems associated with the system.



4.2.2 improved EBC systems: CAS, yttrium silicate, and BSAS:

In addition, besides the expansion coefficient, it was also important to consider EBC systems being capable to act as a barrier with phase stability and low volatility under high temperature condition.[12] So with these goals,  another three EBC systems have been discussed. Calcium aluminosilicate:

The system[13]was consisting of 40% Al2O3, 36% silica, and 24% CaO by weight which is considered as non-stoichiometric calcium aluminosilicate (ns-CAS).

Thermal cycle testing of the system in steam shows that it is 10 times more resistant under steam conditions than  basic SiC after 250 cycles and 500 h at 1200℃. The expansion of ns-CAS was also measured reasonably matches that of SiC.[12]

There is a significant problem with this system, which is the heat treatment required after thermal spraying. The coating will form pores.[13] Additionally, the maximum temperature was probably limited to 1300℃. silicate:

Yttria-stabilized zirconia (YSZ) is also a promising coat candidate.[14]

The steam behavior of yttrium is 10 times more stable than SiC under conditions of 250 thermal cycles and 500 h at 1200℃. One critical disadvantages of YSZ is its large CTE, twice that of SiC or mullite. Besides, when exposuring for longer time under thermal cycling, the CTE may not match SiC that will also causes cracking and delamination, resulting in EBC failure.[12] Barium strontium aluminosilicate

The barium strontium aluminosilicate (BSAS) system[11]has received the most development and testing to date.

The BAS composition has celsian phase which is adopted in applications since it has similar expansion with SiC. Testing figures show that the BSAS system is also with the similar data as Calcium aluminosilicate and Yttrium silicate more stable in steam environment than normal SiC. Besides, the thermal conductivity is approximately 1.6W/mK at 1200 C, similar to that of zirconia based thermal barrier coatings.

Though this system does not exhibit problems of systems mentioned before. However, this generation EBC also has durability issues limit their use.[15] One of them due to the volatilization of the top layer in high velocity combustion environment. Another key issue is BSAS and silica will be able to operate chemical reaction by oxidation.

3.2.3.Rare earth silicate environmental barrier coatings for SiC/SiC composites

So based on the problems and limitation of previous coating candidates mentioned above. Rare earth silicates have been investigated to evaluate their potential as an advanced environmental barrier coating (EBC) having higher temperature capability.  

Researchers focus on finding some new coating substitution that has much higher temperature capacity and chemical and mechanical compatibility with BSAS coat. So some rare earth silicates are identified by NASA research center[16] due to the low CTE and strong phase stability.

There are typically two kinds of rare earth: disilicates have polymorphs like RE2Si2O7, RE: rare earth element of Y, Tm, Er[17] while Lu2Si2O7 has no polymorphs.[18]  monosilicates like RE2SiO5, Sc, Lu, Yb, Tm, Er, and Dy do not have polymorphs, while the other rare earth monosilicates have two polymorphs.[19]

Materials having polymorphs are often not desirable coating candidates, however, rare earth monosilicates with low CTE sometimes have phase stability. Fig.4 shows CTE of several coating materials.


So the main aim is to discuss the volatility of selected rare earth in high steam environments and their performance as EBCs.

From the volatility testing in water vapor it indicates that the rare earth Sc2Si2O7 has volatility similar to that of BSAS and Sc2SiO5 has lower volatility than BSAS. These can be shown in the Fig.5.


Meanwhile, the microscopic photographs of the EBC containing BSAS and rare earth silicates also show that they maintain excellent adherence and chemical compatibility.[16] Better oxidation protection were also illustrated.

However, potential disadvantages of rare earth are exposed simultaneously which is their tendency to through-thickness cracking. But these problems may not be  very concerning for these system since cracks typically stop within the intermediate coat.

In summary, rare earth monosilicates are less volatile than BSAS and rare earth disilicates in combustion environments. And thermal cycling tests have also indicated the potential of rare earth monosilicate EBC (Yb2SiO5, Sc2SiO5, Lu2SiO5) for gas turbine engines can provide higher temperature capability and durability than other EBCs.

Ben's Works Cited

Carper, Douglas et. al. "Hybrid Ceramic Matrix Composite Turbine Blades for Improved Processibility and Performance and Process for Producing Hybrid Turbine Blades." Patent 20,110,229,337.  22 September 2011.

Yoon, Han. "International Journal of Modern Physics B."International Journal of Modern Physics B. 24.15&16 (2010): 2928-2933. Print.

Maccagno, T.M. Canada. National Research Council Canada. Processing of Advanced Ceramics which have Potential for Use in Gas Turbine Aero Engines. 1989. Print.

Wilson's Citation in IEEE format

J. Sha, J.-s. Park, T. Hinoki 和 A. Kohyama, "Hot Corrosion, Oxidation and Their Effects on the Tensile Strength," Materials Transactions, pp. 1032-1035, 15 5 2005.

N. Jacobson, D. Fox 和 J. Smialek, "Corrosion Issues for Ceramics in Gas Turbines," 2004. [online]. Available: http://ntrs.nasa.gov/archive/nasa/casi.ntrs.nasa.gov/20010061369_2001093608.pdf.

A. Wereszczak 和 T. Kirkland, "CREEP PERFORMANCE OF CANDIDATE SIC AND Si3N4 MATERIALS," 13 1 1996. [online]. Available: http://www.osti.gov/bridge/servlets/purl/236225-S4B51Q/webviewable/236225.pdf..