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The developed procedure was very effective to capture the fiber fragments. Thus, this test was viewed as invalid. Single filament tensile test Single filament tensile specimen Load Cell: 2. Test method for measuring the tensile strength of single fiber and specimen geometry. In the capturing of fracture fragments, a small rectangular plastic film 6.

The glycerin effectively damped the shock wave in the fracture of fiber and it usually fractured only at one location. Each fiber segment remained to its half of the mounting frame, which is important for later SEM examination.

Because the fracture of ceramic materials generally originates from the critical flaws, assuming those flaws in the fiber are distributed randomly in location, then the strength of the fiber is determined by the strength at its weakest point weakest link rule.

Test on randomly selected fibers will show a considerable dispersion in failure strengths because of the presence of flaws. The strength of fibers can be shown generally to follow the classical twoparameter Weibull distribution. The two-parameter Weibull theory of statistical fracture was applied to characterize the fracture behavior of brittle SiC fibers [23]. Actually, m is the slope in a two parameter weibull plot, which can be obtained by least squares fitting to the linear relationship of equation 2.

This is especially true with fine diameter fibers which are often degraded by an air test environment and can be easily fractured during grip and strain sensor attachment. Another problem is the accuracy of the creep strain, because in many cases it is hard to define the gauge length including the cold grip and hot grip during the tensile test. To avoid these problems, in this study, a modified bend stress relaxation BSR method was utilized to evaluate the creep resistance of SiC fibers, and attempts were made to relate the BSR with tensile creep for finediameter fibers.

An schematic illustration of the BSR test jig was shown in Figure 5. For evaluating the environmental effect on the creep resistance of SiC-based fiber, a modification High Temperature Mechanical Properties and Microstructure… 9 was made on the conventional method [25] as shown in Figure 5. This improvement makes the tested specimen to be sufficiently exposed to the test environment.

Comparison between conventional and modified bend stress relaxation test method.


In this method, the fiber with a length of cm are wound around the rod at a constant surface strain and held at desired temperature for given times in controlled environment. For small diameter fibers in ambient conditions, the bending modes of different applied strain can be achieved by tying the fiber into small loops with different radius, R0. The fiber loop is then subjected to a specific time t , temperature T , and environmental treatment.

After treatment, the applied stain is then removed by release the fiber loop from the test jig or broken the fiber loop at one point at room temperature. The stress relaxation-induced effects are measured in terms of the residual radius of fiber loops, Ra. If the fiber remains completely elastic during treatment, the broken loop will be straight with no curvature, i. If the creep-induced stress relaxation occurs, the Ra will be finite and typically will decrease with increasing the treatment time and exposure temperature.

To quantify the stress relaxation occurred during thermal exposure, a parameter m, stress relaxation parameter was defined, which is the ratio of final to initial stress at any local position in the fiber as illustrated in Figure 6.

The first assumption of linear strain dependence is generally valid for polycrystalline materials which stress relax due to grain boundary sliding mechanisms that are either elastically or diffusionally accommodated.

If these assumptions apply, the BSR m ration is independent of position and initial applied 10 Jianjun Sha strain. Schematic representation of the test principle of the bend stress relaxation originally developed in Ref.


In comparison to tensile creep test which conducted under a dead load with accessories for strain measurement and a defined gauge length, the BSR offers many advantages including the ability to simultaneously study many fibers of small diameter and short length under same time, temperature, and controlled environmental conditions. Here, it is obvious that stress relaxation parameter, m, can be determined based on the extent of permanent deformation occurred during stress relaxation.

An m value which approaches 1 indicates that no permanent deformation occurred during the high temperature exposure, while a m value of 0 indicates that the stress completely relaxed.

Hence, fibers are considered more thermally stable against creep as m values increase from 0 to 1 [25]. And also, it eliminates the need for furnace with long uniform hot zones, for mechanical grips, for remote sensors and for multiple experimental runs that are often required to establish time, temperature and stress dependencies and also to determine statistical variations.

Second, for polycrystalline fibers, which generally creep with stress power dependencies near unity, if the BSR m-ratios are independent of applied strains, and thus equal to those stress relaxation ratios that would be measured in a pure tensile test. Furthermore, by BSR test, it will be beneficial to understand the basic mechanisms which controlled the creep behavior of SiC fibers with fine diameters [25].

High Temperature Mechanical Properties and Microstructure… 11 2. Microstructural Characterization This section described some techniques that will help to clarify why the mechanical properties were changed and how the microstructure influenced the mechanical properties.

Optical microscopy with a video was used to examine the macrostructure of materials. It is also useful in the determination of fiber loop diameter in the BSR test. In BSR test, a photograph was taken of the loops before and after thermal treatment. The initial applied curvature R0 or residual curvature Ra was measured by fitting a circle on the fiber in the photograph, and then the curvature could be obtained by a graphic technique.

XRD is very useful in the identification of the crystal phase and the estimation of the crystallite size. X-ray scattering of the atom planes in the crystals gives a diffraction pattern characteristic of the crystal structure.

The relative intensity of the diffraction pattern varied with the diffraction plane to aid in structure identification. Comparing the experimental diffraction pattern to a known pattern allows the crystal structure to be identified. The specimen was prepared by attaching the powder sample on the glass slide with double-side adhesive tape.

The powder was obtained by pulverizing the fiber tow of about 0. During pulverizing, in order to prevent the spray of fiber fragments from pulverizing, the alcohol was mixed with powder to make viscous slurry. After careful milling and drying, powder was put on the glass slide with double-side adhesive tape, and it was pushed to be attached tightly. During scanning, the XRD operated at 40 kV and 20 mA was used to identify the crystal phase in the fibers.

FE-SEM Field-Emission Scanning Electron Microscope, model; JEOL JEM , which provides narrow probing beams as well as high electron energy resulting in both improved spatial resolution and minimized sample charging and damage, is a powerful weapon in the characterization of dimension and microstructure such as examination of surface morphologies and fractograph.

The selected fiber was attached on the specimen holder with the double-sided carbon tape. The fiber diameter was determined from SEM image with high magnification x Special care was taken in the register of fiber fragment so that the diameters represented the fibers that we want to investigate. To examine the fracture surface, firstly a technique described in section 2.

In order to take a high quality picture, the clean fracture surface is needed and it can be gotten by washing the fragment in ultrasonic bath contained alcohol for about 30 s. Each fiber segment for the successful tests was gripped with a narrow tip tweezers and broken off at the bonding point.

The clean segments were mounted on double-sided carbon tape applied to the circular side surface of cylinder specimen holder 10 mm copper cylinder in diameter. Usually about 10 segments were mounted with each pair of matching fiber fracture surfaces, and keep the fracture surface with a protruding length about 2 mm above the specimen holder surface.

And also, the fragments should be perpendicular to the horizontal surface of holder. Then, the fracture surfaces of the aligned fragments could easily be located, identified and imaged by SEM. Basic Characteristics 3. This is the likely situation for most polymer-derived SiC fibers that are processed using a spinning method [26]. Such errors cause additional scatter in the Weibull strength distribution, which results in a low value of Weibull modulus.

In such cases, to properly determine the fiber diameter in the assessment of fiber strength is important. The fiber diameter variation from fiber to fiber across a tow and along the single fiber length was assessed by image analysis from SEM.

Fiber Diameter Variation within a Tow To investigate the fiber diameter variation across a tow, a yarn of each fiber type was scattered and mounted on the plane surface of copper specimen holder, and then carbon tape was used to fix this fiber bundle. The picture was taken on these fibers one by one and the number of selected fibers is as large as possible. Figure 7 showed the fiber diameter variation across a fiber tow.

Mean diameter and standard deviation were also calculated in Table.

From this result, the HNLS fiber type showed smallest diameter variation across the tow, which indicated this fiber type has more uniform diameter within a cross section of its tow. The HNL fibers displayed a relatively wide fiber diameter variations within a tow Noting the average diameter for each fiber type is possibly different from batch to batch.

Table 2. Fiber Diameter Variation along the Fiber Length To investigate typical fiber diameter variation along a fiber length, three individual fibers with a length of 30 cm were pulled out randomly from a tow and cut sequentially into 1 cm segments.

Before pulling out the individual fiber, soaking the fiber tow in acetone for 2 days and followed by washing in boiling water for 1 minute. This step is quite necessary in aiding the fiber separation, pulling, and reduction of the handling damage to the fiber. For viewing by SEM, the 1 cm length segments were fastened sequentially on the flat specimen holder by carbon tape. For reducing the charging effects, the segments should be connected well with the specimen holder.

The diameter was determined directly from image taken by SEM. It should be noted that the ends of each segment were carefully retained in register so that the diameters represented the variation of the diameter along the fiber length at 1 cm interval. The fiber diameter variations at 1 cm interval along 30 cm length fiber for the randomly selected fibers were shown in Figure 8.

The HNL fibers exhibited a cyclic diameter variation with a repeat distance of about 15 cm. The two of three TySA fibers exhibited a very similar variation tendency in diameter, and the least variation ration of about 0. From A vg. In the case of HNL fiber, it has been reported by manufacturer that excess free carbon and amorphous phase existed at the grain boundaries. Tensile Properties and Fracture Surface In Figure 10, the tensile strength distribution and related tensile properties of three fiber types were shown in a two-parameter Weibull plot.

Two-parameter Weibull plot for three fiber types indicating the related tensile properties. The m values listed in Figure 10 are slightly lower than that of in Refs. As we know, the strength of ceramic fibers is associated with the gauge length weakest link rule and fiber diameter.

Long gauge length and poor uniformity of fibers might be responsible for low Weibull modulus. Furthermore, the performance of fibers also varied from batch to batch. To examine the fracture surface of individual fiber segments, each fiber segment for the successful test was gripped with narrow tip tweezers and broken off at the frame edge.

Generally, for the brittle materials, such as ceramics and glass, their fracture originated from the critical flaw surrounded by the mirror zone, mist zone and hackle zone.

Figure 11 is an schematic illustration of crack initiation and propagation route of the SiC-based fiber showing fracture mirror zone surrounding the critical flaw.

During microstructure observation, special care was taken on the features of critical flaw size rc , flaw type and mirror size rm. In the case of HNLS fiber, the critical flaws were mainly identified as the inner flaws inner pore or inclusion. Hackle zone Mist zone Mirror zone rm Critical flaw rc Figure Schematic illustration of fracture surface of fiber, showing fracture originated from critical flaw surrounding by mirror and hackle regions. Typical features of fracture surfaces for the SiC-based fibers are shown in the following SEM micrographs.

In Figures 12 a - b , a mating pair of fracture surfaces shows a surface critical flaw and the surrounding mirror, mist and hackle regions. Both sides of the mating fracture surfaces exhibited a well defined void. For this particular fiber, the diameter d , the mirror radius rm and the critical flaw radius rc were measured to be 15, 2.

Figures 12 c - d are a typical pair of mating fracture surfaces showing a critical flaw of inner pore type. It was very clear that the inner pore was observed in each surface. More attention was paid to HNLS fiber.

As observed on the fracture surface of HNLS fiber, most of critical flaws are inner flaws and identified as the second inclusions. Figure 12 e - f shows a pair of mating fracture surface of HNLS fiber exhibited a pore Figure 12 e and inclusion Figure 12 f on the opposite fracture surface.

For this type of flaw, it is possible that a second inclusion was pulled out freely from a pore. High Temperature Mechanical Properties and Microstructure… 1um 17 1um aa bb 1um 1um c a b d 1um 1um ea bf Figure The three line scans qualitatively show the variation of the C, O and Si concentrations in the region of the critical flaw.

The relative concentrations of C, O and Si are almost unchanged until the inclusion is encountered by the scan, where the C concentration abruptly increases and the Si concentration abruptly decreases at the region of the inclusion. The O concentration within the region of inclusion appears to somewhat lower than its base-line level.

The pertinent data for this particular fiber are given in the caption of Figure Based on this information, the inclusion region critical flaw appears to be a carbon-rich region. Inclusion with similar characteristics was also observed in previous study of HNL fiber [27].

Linking these micrographs to the production process, the defects such as voids or inclusions from impurities or un-melted precursors may exist in polycarbonsilane-derived fibers. These defects may generate local internal stress concentration during the process and lead to the crack formation under tension.

And also, under the same processing parameters, it is likely that the stress concentration will vary with fiber diameter, since it is easier to relax the stress concentration in a fiber with a smaller diameter. Wanger [30] suggested that the spinneret hole has laminar flow properties which change with diameters during the fabrication. This may also result in the flow density variation with varying fiber diameters.

In the case of TySA fiber with the small diameter about 7. The fracture surfaces showed a trans-crystallite fracture behavior. This fracture behavior could be partially related to the residual stresses caused by the addition of alumina in this fiber. Existence of residual stresses in the grain boundary of TySA fiber is quite possible because of significant mismatch in the coefficient of thermal expansion between SiC and Alumina SiC: 3. In TySA fibers, the change in the extension stability of micro-crack in the residual stress filed might improve the grain boundary strength.

The increase in grain boundary strength could explain the trans-crystalline fracture surface of TySA fiber. Typical fracture surface observation of TySA fiber showing trans-crystallite fracture.

In Figure 15, the fiber tensile strengths exhibit significant scatter. Nevertheless, the general tendency that fibers with larger diameters have lower strengths is consistent. The fiber with large diameter will be easy to cause the stress concentration around the defects. The similar phenomenon for Nicalon fibers was also observed in Ref. The values of K0 and n, determined from linear least squares fitting of the fiber tensile strength data to Equation 10 , are included in the plot for each fiber type.

The solid curve was obtained by fitting the data point in each plot with Equation Although the data is significant scatter, the general tendency that fibers with larger diameters have lower strength is consistent. Correlation between Tensile Strength and Mirror Size Critical flaw radius um Prior to understanding the correlation between the tensile strength and the mirror size, the dependence of critical flaw size on its corresponding mirror size was examined as shown in Figure The critical flaw radius vs.

This gave us an information that critical flaw size rc was linearly related to the mirror size rm. A linear relationship between the critical flaw size rc and the mirror size rm is commonly observed for brittle fracture of ceramics and glass [,]. Possibly, this difference could be attributed to the definition of mirror size and the crystallization of PCS-derived fibers. On the other hand, it could also be associated with the accuracy in estimating the actual flaw sizes from the SEM micrographs.

The best fit straight line has an slope of These values are very close to the The data scatter appeared in the plots of tensile strength vs. In general, for the ceramic fibers, the critical flaw density varied with varying fiber diameters, and the flaw within the fiber with larger diameter is more easily to cause the stress concentration.

Furthermore, the error in the diameter measurement of fibers will also cause the data scatter in the calculated fracture strength. And also, the strength of SiC fibers is sensitive to the surface critical flaw. For this batch of HNL and HNLS fibers, the fraction of critical flaws occurred at the fiber surface is relatively low, but mostly were distributed internally with only a slight preference for being located nearer to the fiber surface than to the fiber center.

Tensile strength vs. Fitting straight line represented slops of approximate 0. Fracture Toughness and Critical Fracture Energy Because the relation between the fracture strength and the critical flaw size observed the Griffith theory according to the results in Figure 16 and Figure 17, thus, the fracture mechanics principle could be applied to calculate the fracture toughness and the critical fracture energy.

The fiber tensile strength vs. The data were fit to a linear relation by regression analysis. The mirror constant Am, defined as the slope of the fit straight line, was determined to be 3.

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Since Am is an average value, the K1c-value determined for these fibers also is an average value. The Griffith theory presents a criterion for propagation of preexisting flaws that generally determines the failure of brittle materials. After enough energy has been supplied to the crack, it will propagate at velocity which increases as its length increases. Since the driving force depends on crack length, crack velocity will increase until it approaches a terminal velocity.

As the crack approaches the terminal velocity, the sum of the potential energy resulting from its increasing length and the kinetic energy resulting from its motion becomes greater than the energy that can be used to increase the velocity of the crack.

Small cracks are nucleated around the tip of the main crack, forming mist, but there is insufficient energy to propagate these secondary cracks very far. Limited velocity increases allow propagation of such secondary cracks to form hackle. Finally, when enough energy is available, the crack can branch macroscopically. Thus, the identification of correlation between performance and heat treatment temperature is essential for exploring the optimum condition for high performance CMCs fabrication and application.

Correlation between Tensile Strength, Crystal Size and Heat Treatment Temperatures Figure 19 shows average room temperature tensile strengths and apparent crystallite sizes for three types of fibers after heat treatment in Ar for 1 h at elevated temperatures.

Figure Tensile strength and its relation to the crystallite size of SiC fibers heat treated at elevated temperatures in Ar for 1 h. For the bulk materials with clean grain boundaries, the grain growth proceeds through large grains incorporating the small one by grain boundary diffusion. Especially, for grains with small size, the grain boundary diffusion operates much more readily.

On the other hand, the residual trace oxygen may play a role in the Si and C grain boundary transport by accelerating diffusion [39], because the oxygen is not necessarily eliminated from the fiber as reported in the literature [42], even for the HNLS fiber which was fabricated at very high temperature.

Considering the starting temperature for grain coarsening in Figure 19, the grain size might be related primarily to the maximum temperature at which the fibers were fabricated. From the Figure 19, it can be seen that the crystallite size increased when heat treatment temperature is above the fabrication temperature as expected. On the other hand, the thermally activated diffusion play an important role on the grain coarsening of SiC materials at high temperatures.

If we make a further comparison in crystallite size between two Nippon Carbon fibers again, we can see that a large difference in crystallite size was observed for two fibers het treated at same temperature. As above mentioned, the HNL fiber has an small starting grain size, which was expected to have a high diffusivity at grain boundaries and result in a large grain size as heat treating at high temperatures.

However, an unexpected phenomenon was observed between two Nippon Carbon fibers. This can be attributed to the excess carbon in HNL fiber.

TEM observation revealed that heat treatment of the HNL fiber results in a gradual organization of the free carbon phase in terms of the size of the carbon layer and the number of stacked layers as increasing temperatures [39]. Takeda et al. Grain growth is suppressed with increase in excess carbon.

In other studies [], the carbon suppressing growth and coalescence of the SiC microcrystals was also observed. The grain growth has a significant effect on the strength of SiC-based fibers. In both near stoichiometric fibers, strength degradations occurred at the temperatures where crystallite size began to increase. The growth of SiC crystals reduces the bonding forces at the grain boundaries.

Since the manufactures are always seeking the optimal fabrication temperature at which the superior thermal stability and excellent mechanical strength can be obtained simultaneously, thus, the upper fabrication temperatures are typically fixed by those temperature conditions 26 Jianjun Sha above which performance degradation of the fibers occurred. The dependence of strength on temperature in Figure 19 is in agreement with those of previous studies [18,,46]. Ichikawa et al.

TEM observation shows that this heat treated HNLS fiber has a SiC grain size of approximately nm, which is about 10 times larger than that of the asreceived fiber. As a result, the TySA fiber showed an excellent thermal stability in Ar atmosphere. The growth of the carbon corresponds to a decrease of localized spin centers [50]. The growth of the carbon grain might result in an increase of residual stress, however, this evidence is insufficient because the magnitude of residual stresses is strongly dependent on the volume fraction of carbon phase in a bulk material.

Microstructure Figure 20 shows SEM morphologies of the fibers after heat treatment at high temperatures in Ar for 1 h. The fibers showed a porous microstructure and large grains deposition on the fiber surface Figure 20 a. Obvious differences were found in subsequent observations of fracture surfaces. The critical flaw size and mirror size were measured.


The critical flaw sizes rc are: 0. In the case of HNLS fiber, the critical flaw sizes rc are: 0. The fracture origin and mirror zone are invisible on the fracture surface and fracture surface showed a trans-crystallite fracture behavior. The trans-crystallite fracture behavior could be partially related to a high compression residual stresses in SiC caused by addition of alumina in this fiber.

Existence of compression residual stresses in the grain boundary of TySA fiber is quite possible because of significant mismatch in the coefficient of thermal expansion between SiC and Alumina SiC: 3. In TySA fibers, the change in the extension stability of micro-crack in the compression residual stress filed might improve the grain boundary strength. The increase in grain boundary strength could explain the trans-crystalline fracture behavior of TySA fiber. Linking the tensile strength data in Figure 19 with the microstructure examination in Figure 20 again, the decomposition of amorphous phase, grain coarsening and residual stress at high temperatures in HNL fiber could be responsible for strength and microstructure degradation.

Observation of surface morphologies Figure 20 a - c and fracture surface Figure 20 d - f provided some information for the strength degradation of SiC fibers. The formation of porous structure in HNL fiber could be attributed to the rapid evolution of gases at the earlier stage of high-temperature exposure according reaction According to above result, the reaction 17 - 19 are quite dependent on the quantity of amorphous phase and content of carbon in SiC fibers.

Thus, thermal decomposition of the amorphous phase is almost negligible High Temperature Mechanical Properties and Microstructure… 29 in this fiber. As for the large grains deposited on the surface of HNLS fiber, it can be explained by reactions 18 - 20 , because the carbon layer on the surface of HNLS fiber 80 nm is thicker than that of HNL fiber 20 nm [42]. Concerning the origin of gas species, other mechanism could be responsible for it.

Active oxidation is quite possible. The transition from passive oxidation to active oxidation occurs at oxygen partial pressure of Pa and At same oxygen partial pressure level, the increase of heat treatment temperature will accelerate the transition from passive-to-active oxidation [51].

The small amount of alumina addition could inhibit the grain growth and enhanced the oxidation resistance. This higher stability can also be linked to the silica protective layer formed on the surface of fiber [48,54]. Combining the fracture properties with microstructure characterization of SiC-based fibers, we can not deny the existence of other degradation mechanisms such as contaminants during heat treatment and metallic impurities introduced during process [].

The existence of metallic impurities within the fibers is possible, because all these fibers are polymer derived. The metallic impurities can easily enter the fibers during the various steps of polymer handling and can cause rapid or abnormal grain growth in local areas.

There are at least two indirect observations supporting above mentioned mechanism: i Observation of fracture surface in Figure 20 for the HNL and HNLS fiber showed that the strength-limiting flaws after heat treatment are larger than the average grain size, indicating rapid defect growth in selected areas of the fiber and thus suggesting the possible existence of metallic impurites; ii the UF fiber showed high strength retention than HNL fiber [47]. This suggests that the UF fiber during processing did not introduce the metallic impurities to the degree that employed for the HNL fiber.

Heat treatments of the fibers above the processing temperature resulted in improved creep resistance as shown in Figure Such microstructural changes are expected to inhibit diffusion-controlled creep processes. This result indicated that the improved creep resistance depended on not only the crystallization and grain growth, but also the composition at grain boundaries. This implies that stability of Grain boundaries GB plays an important role on the creep resistance of SiC fiber.

This assumption was also demonstrated by TySA fiber. The enhanced creep resistance of the TySA fiber was obtained prior to increase its crystallite size. As a result of Al addition to TySA fiber, the complex oxide would be formed at GB by heat treatment and they can stabilize the grain boundary to improve the creep resistance.

The stability of GB could be affected by GB composition.

Jianjun Sha Stress relaxation ratio, m 1. Dependence of fracture toughness and critical fracture energy on heat treatment temperatures. High Temperature Mechanical Properties and Microstructure… 31 treatment temperature, but it did not show strong dependence on the heat treatment temperature.

For the as-received fibers, the carbon layer covered on the surface of fibers, can blunt the critical flaw and reduce the stress concentration on the surface flaw. However, this carbon layer can be removed by reaction with residual oxygen from fiber itself and atmosphere. In this case, the propagation of preexisting surface flaws will become easy. At fairly high strain rate 0.

However, the mechanical and thermal stabilities of SiC fibers as reinforcements in CMCs are very sensitive to their purity, crystallinity and service environments [] including thermal and loading history. For high temperature applications, the CMCs are often subjected to oxidative environments with different oxygen partial pressures.

As we know, the performance degradation of SiC materials in oxidizing environments strongly depends on the oxidation mechanism. Jacobson [61] has discussed the oxidation degradation mechanism of SiC materials in varied environments, but it is still insufficient because of the complexity of service environments. The key question concerns the oxidation kinetics: passive and active oxidation.

This topic has given rise to much controversy for SiC materials, because the temperature boundaries for the oxidation kinetics are quite dependent not only on the materials themselves purity and crystallinity , but also on the specific service environment exposure temperature, oxygen partial pressure and mechanical state.

Furthermore, rarely is one mechanism operative in performance degradation of SiC materials. In practice, several mechanisms operate simultaneously.


Therefore, for understanding the mechanical and thermal stabilities and failure mechanism of SiC fibers over a wide range of temperatures and varied environments, the part of this chapter reviewed the microstructure features and high temperature properties of SiC fibers under annealing and creep in various oxygen partial pressures at elevated temperatures, and attempted to clarify the correlation between the environment with mechanical and thermal stabilities.

Furthermore, the surface morphologies of fibers under annealing and creep were compared by the observation of field-emission scanning electron microscope FE-SEM. The SiO2 film formed uniformly on the surface of fibers during annealing at elevated temperatures Figures And also, no further oxidation between SiO2 film and SiC fiber surface was observed. Some patterns were also observed within the silica film formed on the surface of HNL fibers Figures 23 b-1 under creep condition.

No significant cracks were found within silica layer even under crept condition. Maeda et al. The actual kinetics involved at least four different parabolic stages for this material.

They attributed these to the various microstructural changes in the scale: crystallization of amorphous silica, transformation of those crystalline phase, and viscosity changes in the oxide scale due to migration of the additives. This indicated the oxidation of TySA fiber is much complex because of addition of alumina in this fiber. Under creep condition, pits appeared on the surface of fibers Figure 24 b This result indicated that stress applied by BSR test could enhance the oxygen attack.

The passive oxidation of TySA fiber at extremely low oxygen partial pressure appears to be attributed to the addition of a minute of alumina. Alumina, as a oxidation product of aluminum, reacts with the SiO2 film to form an alumino-silicates. The alumino-silicates have a high oxygen permeability, presumably enhancing the passive oxidation of fibers []. It is clear these huge crystals grew outward from the fiber surface.

The thermochemical correlation and oxidation dynamics for active and passive oxidation of silicon carbide have been investigated experimentally and theoretically in the literatures [51,]. These data were combined together and plotted into a new plot as shown in Figure The passive-to-active oxidation transition is strongly affected by factors such as the type of silicon carbide, temperature and oxygen partial pressure.

The oxygen partial pressures are, 2. The 36 Jianjun Sha oxygen partial pressures and test temperatures for this work were shown in Figure 26 by three lines. It is clear that proposed test conditions distributed in different regions. Air 1. E- 02 1. E- 03 1. E- 05 1. E- 01 Hi-NIcalon iC hn eid er 1. Based on the surface and cross section morphologies in Figure 23, it is thought that fibers were mainly oxidized in passive oxidation regime characterized by the formation of silica film when they exposed in air at high temperatures, but the passive oxidation was enhanced under crept condition due to increased oxygen permeation.

Cracks in the silica layer observed in Figure 23 are mainly due to the difference in coefficient of thermal expansion CTE between the fiber core and the silica layer.

Because the CTE of SiC fiber is less than that of Silica, on cooling, a tensile residual stress will be applied to silica layer. The passive oxidation formed SiO2 layer can refrain the further oxidation of SiC.

In low oxygen partial pressure atmosphere, the oxygen partial pressure for the transition from passive to active oxidation is a key point in the microstructure change of SiC materials. The pO2 value was about 0. From the observation of surface morphologies, we can see the active oxidation initialized at different temperatures in different atmospheres [66]. The temperatures for active oxidation in crept fibers, however, were shift to low values, indicating the transition from passive High Temperature Mechanical Properties and Microstructure… 37 oxidation to active oxidation was enhanced under creep condition.

The enhanced active oxidation in creep condition might be caused by the rupture of silica scale on the surface of fiber due to the stress applied by BSR test. Subsequently, a stress concentration would occur around the flaws generated by active oxidation or gas evolution, and then oxygen attack on the SiC fibers will be accelerated, leading to the formation of bubbles and pits.

Histórico de Todas as Dicas

The formation of bubbles in the silica scale may provide some indications of pressure buildup [72], especially, when the passive and active oxidation proceeded concurrently. Due to the near stoichiometric composition and high crystallinity of HNLS and TySA fibers, their active oxidation was gradual in comparison with other SiC fibers [,67,72].

Noteworthy is that the active oxidation is infinite slow if the oxygen partial pressure is very low [73]. Tensile Properties Figure 27 shows the dependence of mean strength on the testing environments. It should be noted during the specimen preparation that fibers with low strength became very difficult to set without breaking them. The mean strength we gave will consequently not take the weakest fibers into account no strength could be obtained.

Due to this shortcoming, overestimation of tensile strength is likely. The strength for the fiber annealed in air decreased only slightly Figure This is due to the SiO2 layer, which acts as environmental barrier coating, stopping the further oxygen attack. The fracture origins on the fracture surface of fibers were shown in Figure The fracture mainly originated from surface flaws after annealing in different atmosphere.

This means the surface of fibers was damaged by oxidation. The relative low strength retention for fiber annealed in UHP-Ar suggested changes in flaw population and flaw size. The low oxygen partial pressure accelerates the transition from passive to active oxidation resulting in a coarsening and pitting surface. At high temperature under low oxygen partial atmosphere, SiC is relatively easy to be oxidized actively.

High Temperature Mechanical Properties and Microstructure… 39 4. Namely, a weak dependence of BSR creep resistance on oxygen partial pressure is observed. As for the creep behavior of SiC fibers, generally, it can be explained by their oxygen content, grain size and second phase in the grain boundaries. The creep resistance of SiC fibers slightly decreased with increasing the oxygen partial pressure, which is likely controlled by the oxidation and grain coarsening. Under creep condition, grain coarsening of SiC has been observed for longer time creep experiments, which may contribute to a decelerating creep rate [].

The grain size increased with decreasing oxygen partial pressure has been observed in other studies [59,67]. Furthermore, under the creep test, the grain coarsening could be accelerated by applied stress [59]. The somewhat high creep resistance in low oxygen partial pressures might be attributed partially to the concurrent grain coarsening during BSR test.

On the other hand, the low creep resistance in air could be partially explained by the resistance effect of silica layer. The outmost SiO2 sheath formed in air after BSR test will counteract part of the initial applied stress leading to a low stress relaxation parameter m, especially, for the fibers with fine diameters.

Thermal Exposure Under Loading Most previous studies have been concerned with the degradation in mechanical properties of SiC fibers after high temperature exposures [58,]. There exist few studies on the performance change of SiC fibers in low oxygen partial pressure atmosphere under loading [79]. Therefore, for understanding the degradation mechanism of CMCs under loading in low oxygen partial pressure environments, the loading tests were performed on SiC fiber yarns by applying different dead loads at elevated temperatures in Ar atmosphere.

After each loading test, the room temperature tensile properties and microstructure were characterized to clarify the performance degradation mechanism. Tensile Properties Figure 30 showed the room temperature tensile strength distributions of SiC fibers with different conditions.

Likely, the time difference between 1h and 3h is too short. The combined effect of applied load and exposure temperature is more obvious in properties degradation of SiC fibers. The near-stoichiometric and high-crystallite SiC fiber, HNLS, degraded gradually in tensile properties with increasing load and temperature, as shown in Figure 30 e. This phenomenon seems to be related to the oxidation resistance and thermal stability of SiC fibers under loading.

Continued on next page. Room temperature tensile strength distributions for HNL and HNLS fibers after loading test with different conditions: a effect of loading, b effect of exposure temperature, c effect of exposure time, d combined effect of exposure temperature and loading, e tensile properties of SiC fibers with different conditions.

No obvious differences in the surface morphologies were observed between the as-received and annealed fibers Figures 31 a and c. A quite different morphology is apparent for these fibers. A rough surface with extensive grain growth and micro pores was observed in HNL fibers as shown in Fig 32 a. According to the observation of the fracture surface Figure 32 b , the HNL fibers fractured at an irregular groove, which extended from surface to interior and was significantly different from that of annealed fibers in shape and size Figure 31 d.

There are only a few individual large grains grown on the surface of HNLS fibers Figure 32 c , and the fracture of the HNLS fiber also originated from the irregular surface flaw Figure 32 d , but the flaw size is smaller than that of HNL fiber.

High Temperature Mechanical Properties and Microstructure… a c 43 Large grain pore Extensive grain growth 5 um 5 um b d Critical flaw Critical flaw 1 um 1 um Reprinted by permission of Elsevier from [79]. Fracture of HNL and HNLS fibers under loading mainly originated from an irregular surface flaw, and features are associated with brittle failure Figure 32 b and Figure 32 d. The irregular groove could be a stress corrosion crack SCC caused by combination of the oxidation and the applied load.

Generally, the oxidation behavior of silicon carbide at a high temperature depends on ambient oxygen partial pressure. Passive oxidation occurs at high oxygen partial pressures and results from the formation of SiO2 that grows on the surface of exposed fibers. Passive oxidation could protect the materials from further oxidation attack.

For active oxidation, pits and cracks occurred at low concentrations of O2 where the SiO2 formation rate is too low to seal the surface of materials. The passive-to-active oxidation transition is strongly affected by factors such as the type of silicon carbide, total gas pressure and gas flow rate, as well as temperature and oxygen partial pressure [51,70].

Thus, based on the surface morphologies and fractographs in Figure 31, it was thought that annealed fibers were mainly oxidized in passive oxidation mode characterized by formation of silica.

The passive oxidation formed a thin silica film on the surface of annealed 44 Jianjun Sha fibers and this thin silica film could protect the fiber from further oxidation attack resulting in no critical structural change. In this case, the thermal decomposition of amorphous phase which proceeded by reaction 17 should also be suppressed [67,], because the gas species cannot be removed fast enough through the silica film by diffusion or migration [67,70].

Under loading conditions, the rupture of silica scale on the surface of fiber would be easier and a stress concentration occurs around the flaws. At that time, the oxidation was enhanced by loading, leading to the stress corrosion crack, as observed in Figure 32 b and Figure 32 d. The total pressure in this study was Pa; hence, the pO2 value in furnace chamber was about 0.

A high total pressure will increase the limits for the transition from passive-to- active oxidation [51]. The grain growth also has a contribution to the strength degradation [69]. Because of the near stoichiometric composition and high crystallinity of HNLS fibers, its active oxidation was gradual in comparing to that of HNL fiber.

In previous studies [26,36,,], many researchers have observed that the strength of ceramic fibers was associated with critical flaws.

As observed in Figure 32, a clear mirror zone around the critical flaw on the fracture surface corresponds to the smooth propagation of a crack. In the present work, because most of the critical flaws are irregular in shape, it is very difficult to define the flaw size. Thus, for clarifying the degradation mechanism, the fracture mirror size on the fracture surface was measured.


There is a tendency that mirror size increased with increasing load and temperature. The increased mirror size could be related to the increased critical flaw size. Critical flaw size could be associated with the oxidation and creep resistance of SiC fibers under loading.

Having a low oxidation and creep resistance could easily cause the new flaw nucleation and growth and slow crack High Temperature Mechanical Properties and Microstructure… 45 propagation. Meanwhile, the excess carbon and amorphous phase in HNL fiber might also have a contribution to the flaw population and flaw size.

For easily understanding the oxidation corrosion, following mechanism was proposed. For annealed fibers, a thin silica film formed on the surface of fibers and prevented the gas species passing through the fiber surface.

This would be possible because of the use of the alumina furnace wall which resulting in the real oxygen partial pressure might be higher than the equilibrium pressure. In this case, the oxidation and the thermal decomposition of amorphous intergranular phase [67,82] were suppressed. The SiO and CO gases were transported through the fiber surface. SiC crystals grew on the fiber surface due to the reaction between SiO and free carbon. Meanwhile, the rupture of silica scale on the surface of fiber would be easier during loading test [70,83].

Once the flaw was produced by oxidation and decomposition of amorphous phase, the stress concentration occurred around these flaws and preexisted crack tip.

At that time, the active oxidation was enhanced by loading because of the rupture or removal of the protective silica film on the fiber surface, leading to the stress corrosion cracks, as schematically illustrated in Figure The stress corrosion cracks acted as the critical flaws in the fracture of fibers. The active oxidation is very detrimental to properties of materials and it proceeded mainly by reactions proposed in literatures [51,,67].

Schematic of strength degradation of SiC fibers under loading. In order to evaluate the creep resistance of these advanced SiC fibers, and also to clarify the creep mechanism to support continuing optimization of the properties, the apparent activation 46 Jianjun Sha energy of creep was calculated by applying a cross-cut method to the results of the long-term BSR tests [84].

Sixta et al. Besson et al. Honda et al. In both above-mentioned studies, these authors considered GBS accommodated by diffusion as the critical creep mechanism. Lane et al. The stress exponent increased from 1.

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These authors concluded: 1 the creep mechanism at low temperature is GBS accommodated by grain boundary diffusion, and at high temperatures the controlling mechanism becomes GBS accommodated by lattice diffusion; 2 the parallel mechanism of dislocation glide contributes increasingly to the total strain as the volume of precipitates declines as a result of progressive coalescence with increasing temperature.

As we know, creep of CMCs involves stress transfer between the matrix and fibers caused by their different creep rates, that may lead to fiber failures or matrix cracking []. When the matrix is elastic and creep resistant, fiber creep induces stress transfer from the fibers onto the matrix that may cause matrix cracking. Thus, the understanding of creeprelated properties of fibers is essential in the identification of time-dependent failure mechanism of CMCs. In literatures [25,], the creep behavior of SiC-based fibers has been investigated.

Based on the results of these studies, creep is a thermally activated process and controlling creep mechanism is grain boundary sliding. Particularly, in the tensile creep behavior of SCS CVD SiC fibers, some authors observed that the fibers exhibited only primary creep, which was characterized by a continuously decreasing creep rate for progressively longer times, and that the creep rate was proportional to an exponential power of time [].

The thermally activated creep can be described using traditional Bailey-type relationship []. Fortunately, the bend stress relaxation BSR [25] test has been demonstrated to be an effective method for comparing the relative creep resistance of a wide variety of ceramic fibers. And also, some researchers have been attempted to relate BSR data with that of tensile creep []. Their results showed good correlation between BSR and tensile creep data.

Bend Stress Relaxation and its Relation to the Tensile Creep In BSR test, the stress relaxation parameter, m, is defined as the ratio of final to initial stress at any local position in the fiber. This assumption is suitable to the advanced SiC-based fibers with a near-stoichiometric composition and high-crystalline structure.

If these assumptions apply, the BSR ratio m, obtained by a method illustrated in Figure 1, is independent of position and initial applied strain. It was also indeed that the negligible dependence of m ration on the initial applied strain was observed [66]. This result further supported that the above assumptions were appropriate. On the other hand, the negligible initial strain dependence of m also evidenced the grain boundary sliding GBS as the principal creep mechanism during the stress relaxation of SiCbased fibers.

The rest of BSR tests in this study were performed at a constant initial strain for times ranging from h at elevated temperatures in air. Furthermore, the apparent activation energy of thermally controlled creep was calculated by a cross-cut method from time-temperature dependence of stress relaxation parameter. Since the individual fiber type with uniform microstructure displayed a strainindependent m ratio, the predictions of tensile creep of fibers from BSR data should be possible according to the previous results [25,].

The time exponent, p, determined from a plot of log NCS vs log t, is constant. The duration of the stress relaxation tests was 1, 10, 25, 50 and h. The stress relaxation follows an S-shaped curve.

These results suggest that thermally activated process plays an important role in the creep behavior of SiC fibers. From these results, we can see the Q value increases with increasing the test temperature.

The same change in apparent activation energy was also found in earlier studies [90,]. The large activation energy at high temperature regions could be related to the concurrent microstructure change during BSR test, such as grain growth, the crystallization and loss of oxygen due to the decomposition of amorphous phase SiCxOy at grain boundary.

As we know, the creep behavior of ceramics can be explained by oxygen content and grain size. The crystallite sizes of as received fibers are 4. Furthermore, the grain growth could also be enhanced by applied stress.

Thus, the effect of grain growth must be taken into account at high temperature test. All the grain boundary sliding mechanisms have a negative grain size exponent, which means that smaller grains will result in a faster creep rate. Further comparison in Q value for heat treated HNL fiber supports above result. The creep parameters listed in Table 3 are the best fit in Figure The parameter, A0, is dependent on the temperature, while time exponent, p, is somewhat high in high temperature region.

The fibers exhibit very similar creep behavior, which suggests that they all creep mainly via a similar mechanism and the difference in the individual parameter given in Table 3 is due to compositional or microstructural difference among these fibers. For each fiber type, the tendency of NCS behaved similarly at different temperatures. Furthermore, since the bend stress relaxation tests were performed in air, the oxidation of surface of SiC fibers would enhance the surface stress relaxation [76].

Especially, for the long time BSR test of fine-diameter SiC fibers, the Silica layer carrying a part of initial applied stress by stress sharing mechanism might be possible, which results in an overestimation of NCS. This was reflected by somewhat high time exponent in high temperature region.

The result in present work is in agreement well with previous work [96], although the fiber types are different.

Further comparison was made in BSR test between present and previous work [96]. The parameters presented in Table 3 are the best fit of these curves. Figure 36 showed the tensile creep strain, which was predicted by BSR data. Generally, this prediction showed a similar time and temperature dependence with that of the primary creep stage. As pointed out by other researchers [96], for silicon carbide fibers which have uniform microstructures, the BSR predictions usually are very near the magnitude of tensile creep strain.

In addition, the actual tensile creep strains from Refs. It was observed that tensile creep strains from BSR predictions showed same order of magnitude with the actual tensile creep strains. Noting in many tensile creep tests of ceramic fibers, the data scattering in creep strain is significant. On the other hand, because of the lack of tensile creep data on advanced SiC-based fibers, this comparison was made among the studies with different fiber batch. The properties of SiC fibers are different from batch to batch and they are still in developing.

Further demonstration from tensile creep test with same fiber batch would be necessary. Nevertheless, in present work, the BSR data predicted the same time and temperature dependence of tensile creep for advanced SiC-based fibers. Table 3. Since the silicon self-diffusion coefficient is one order of magnitude slower than the carbon self-diffusion coefficient [], so the silicon is the controlling species during the diffusion process.

The strain rate for a pure diffusion lattice mechanism is given by Nabarro-Herring creep model []. The time exponent p and its slight variation at high temperature indicate that additional mechanism with similar activation energy might be operating besides GBS accommodated by diffusion.

The dislocation activity increased with the coalescence of the precipitate phase, because the interaction of dislocation with the precipitates makes it difficult to gild.

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