thermal and loading effects on...

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94•GT-397

THERMAL AND LOADING EFFECTS ON MECHANICAL PROPERTIES OF A HOT ISOSTATICALLY PRESSED Si 3 N4

Jagannathan Sankar, Jayant Neogi, Suneeta S. Neogi, Marvin T. Dixie, and Ranji Vaidyanathan Department of Mechanical Engineering

North Carolina A&T State University Ill 111111 11111 111 111111 Greensboro, North Carolina ,

ABSTRACT

The effect of thermal soaking on the mechanical properties of a candidate material for advanced heat engine applications namely, hot isostatically pressed (111Ped) silicon nitride (GTE-PY6) are reported here. Pure uniaxial tensile tests conducted at room and at elevated temperatures indicated that the tensile strength of this material dropped significantly after 1000°C. The residual tensile strength of PY6 material after thermal soaking at 1200° and 1300°C was also investigated. Test results showed that thermal soaking at 1200° and 1300°C increased the residual tensile strength. The thermal soaking time had a greater effect on the residual tensile strength at 1300°C. Tensile creep tests performed at 1200° and 1300°C showed that the steady state creep rate was influenced by both the temperature and the applied stress. The higher stress exponent in HIPed as compared to a sintered silicon nitride shows higher creep resistance in the case of HIPed materials.

INTRODUCTION

The use of ceramic materials offers a number of advantages including higher-temperature operation, decreased weight, greater thrust-to-weight ratio, lower life-cycle cost, and reduced dependency on strategic materials. Significant progress has been made in recent years in developing hot isostatically pressed (HEPed) silicon nitride (Si,N 4) structural ceramics for heat engine applications. These materials have high strength and are oxidation- and thermal-shock resistant. Further, for their intended applications, a number of requirements must be met, namely, minimal creep, high temperature durability, resistance to time-dependent failure, high strength at room and elevated temperatures, and good fracture toughness (Hecht el al., 1992).

Because silicon nitride is a material that does not sinter easily in the "pure" state, additives such as Y 203 , MgO, and A1203 , or a mixture of these are used to promote liquid-phase sintering. The retained intergranular vitreous phase inherent in this process effectively controls high-temperature behavior (Wiederhom et al., 1993). An important structural material, fl-Si,N„ is an example of this type of a ceramic. Hot isostatically pressed (HIPed) fl-Si,N, is dense and contains a few percent of a glass phase in the grain boundaries. The glass phase may be a-Y2S1207 (triclinic yttrium disilicate), N-apatite (a hexagonal yttrium-silicon oxynitride) or other such compounds (Wiederhorn et al., 1993). Porosities can occur in either HIPed or sintered Si,N, if the process is not optimized. This paper reports the mechanical behavior of a dense material.

The glassy phase affects the mechanical properties adversely at high temperatures. At

Presented at the International Gas Turbine and Aeroengine Congress and Expositor.• The Hague, Netherlands — June 13-1 13, 1994

Copyright © 1994 by ASME

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higher temperatures the glass softens and degrades the strength of the HIPed material. The modulus of rupture (MOR) shows a peak before declining as the temperature is increased (Tsai et al., 1982). It was observed by Tsai et al. (1982), that the thermodynamics of crystallization of a glass which is segregated on a microscopic scale in graiQboundaries, is different from that of bulk glass of the same composition. Recently, Wiederhorn et al. (1993) have reported that thermal soaking at elevated temperatures ( 1200 0-1400°C) resulted in the devitrification of the intergranular phases in a HIPed silicon nitride (NT-154). Liu et al. (1990) have observed that treatment by thermal annealing enhanced the resistance to creep deformation in NT-154.

This paper summarizes the effect of thermal soaking on tensile mechanical properties. The differences in mechanical behavior between HIPed and sintered Si 3N4 ceramics are discussed. Some of the microstructural details related to the failure mechanism of this material system are also reported here.

EXPERIMENTAL PROCEDURE

Materials

The material, GTE-PY6 (injection molded and HIPed 0- Si 3N4 with addition of 6 wt% Y203), a commercial grade of silicon nitride was procured from GTE Labs, Inc., Waltham, MA, in the form of cylindrical buttonhead tensile specimens. The specimens were approximately 165 mm (6.496 inches) long and had a nominal gage dimension of 25.0 mm (0.984 inches) in length by 6.36 mm (0.2504 inches) in diameter. All specimens were ground in the longitudinal direction and had a surface finish of 0.5 Am 20 tt inches). The mechanical properties obtained for HIPed silicon nitride were compared with GTE SNW-1000, a commercial grade sintered silicon nitride. These sintered specimens contained a nominal chemical composition of 12-13 wt% Y203 , 2-3 wt% A1203 and the rest silicon nitride. The PY6 and SNW-1000 specimen configurations are given elsewhere (Dixie, 1993; Sartkar et al., 1991).

Test Equipment and Mechanical Investigations

A servo-hydraulic mechanical test system' equipped with a self-aligning hydraulic grip developed by ORNL (Liu et al., 1985) that produces near-zero bending moments (C 0.05% of the applied tensile stress) was used for conducting the uniaxial tensile, creep and fatigue tests. All tests were performed in load control mode. The automated system consisted of a test processor interfaced to the control electronics, a computer and its associated peripheral devices, and software to operate the system.

A box shaped, low-profile, resistance-heating furnace with silicon carbide heating elements2 was used to heat the specimens. The ends of the specimen were connected to superalloy extension rods (Rene 41), which were water-cooled. The specimens were coated with

'MTS 880 System, MTS, Minneapolis

'Model 3450, ATS Corporation, Butler, PA

2


a boron nitride lubricoats to avoid bonding between the pullrod sleeves and the specimen. The longitudinal deformation of the specimen was measured using a non-contact laser

telemetric system'. A laser beam was directed through a window in the furnace towards the gage section of the specimen. The specimens were instrumented with a set of silicon nitride flags attached to the gage section of the specimen using a high temperature alumina based adhesives. Any change in the distance between the flags was sensed by the laser extensometer and was converted into an analog signal. The signal was then converted into digital elongation using an AID channel.

All elevated temperature tests were conducted in air, after 4 hours of heating to ensure thermal equilibrium within the furnace prior to application of the load. This procedure was followed to minimize the errors introduced in strain measurements due to air density fluctuations in the furnace (Sankar et al., 1990).

Microstructural and Fracture Analyses

Fractography studies were conducted using a scanning electron microscope (SEM) 6. The sections used for analyses in the gage section were, a) fracture surface, b) a polished section 2 mm below the fracture surface along the transverse (perpendicular to stress axis) direction and c) a polished section 5 mm below the fracture surface along the longitudinal (parallel to stress axis) direction. Molten potassium hydroxide was used to etch the polished specimens (— 20 seconds).

RESULTS AND DISCUSSION

Tensile tests

Tensile tests were performed at four different temperatures, namely, 20°, 1000°, 1200° and 1300°C. A stressing rate of = 50 MPahnin (equivalent to a cross head speed of 0.004 cm/sec) was used in all the tests. The tensile tests were essentially conducted to compare the PY6 material with previous batches of PY6 tested at Oak Ridge National Laboratory (Ferber and Jenkins et al., 1992) and University of Dayton Research Institute (Hecht et al., 1989). Table 1 lists the tensile strength and modulus for PY6 at different temperatures. Figure 1 A shows the effect of temperature on the tensile strength of GTE PY6 Si 3N4. For comparison purposes, similar data reported by Ferber and Jenkins et al., (1992) and Hecht et al., (1989) for this material are also included in this figure. Good agreement between the sets of data obtained is observed. Based on the present work it can be observed that while the fracture strength decreases

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with increasing test temperature, a similar reduction in modulus is absent (Table 1). Figure IA also contains the tensile strength versus temperature data for GTE SNW-1000 (Sankar et al., 1993). Figure 1B compares the effect of temperature on tensile strength, Young's modulus and strain. The above parameters were non-dimesionalized by taking the ratio with their values obtained to failure, at 1300°C.

Table I. Tensile strength and Young's modulus at different temperatures.

Temperature (°C)

Tensile Strength (MPa)

Young's Modulus (GPa)

20 621 306

1000 451 300

1200 370 297

1300 250 295

0 200 400 600 $OO 1000 1200 1400

0 250 500 750 1000 1250

Temperature ( C)

Temperature ( C)

Figure 1. (A) Comparison of tensile strength measurements for (a) NC A&T SU (PY6), (b) ORNL (PY6), (c) UDRI (PY6) and (d) NC A&T SU (SNW-1000) and (B) Effect of temperature on tensile strength, Young's modulus and strain.

It can be observed that at room temperature, the PY6 is stronger than SNW-1000. The PY6 tensile specimens failed from both surface and volume flaws which was similar to earlier studies by Sankar et al. (1993) on sintered S1 3N4. At higher temperatures the glassy phases might soften and degrade the strength of both the Hll'ed and sintered material. Studies showed that at all temperatures, the tensile specimens failed mostly from radial grind marks in the gage section produced by machining (Figure 2A). A number of impurities were also observed at the fracture initiation site in PY6 Si 3N4, predominantly containing iron with some chromium (Figure 2B).

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Iron impurities ( — 40 to 80 ppm) form iron silicides and produce microcracking as a result of the difference in the coefficients of thermal expansion between the silicon nitride and the iron silicide (2.6 versus 7 x 1041°C,) (Hecht etal., 1992). A polished and etched cross-section of an as-received tensile specimen and cross-sections (parallel and perpendicular to the tensile stress axis) of an as-tested tensile specimen (Figures 3A, B and C) revealed that a large distribution of cavities was present in the material. Figure 4 reveals a very high concentration of cavities at the fracture initiation site (surface) as compared to the bulk (Figure 3B).

Figure 2. Tensile tested specimen at 1200°C: (A) Typical surface-flaw-initiated failure and, (B) impurity particles at fracture origin.

Figure 3. Micrographs of polished and etched cross-sections of: (A) an as-received specimen, (B) and (C) - a tensile tested specimen at 1200°C perpendicular to tensile axis and parallel to tensile axis respectively.


Figure 4. Micrograph of a polished and etched cross-section of a tensile tested specimen at 1200°C (perpendicular to stress axis) showing higher concentration of cavities at the surface,

Thermal Soaking/Tensile Tests

A total of ten specimens were thermally soaked for different periods of time and at two different temperatures 1200° and 1300°C. Due to the limited availaibility of samples only two samples were tested for each condition. The term thermal soaking/tensile test signifies holding a specimen at a particular temperature for a prolonged period of time and then performing the tensile test on the specimen at the same soaking temperature without cooling the specimen to room temperature. Table 2 lists the residual tensile strength results after thermal soaking at 1200° and 1300°C for soaking times ranging from 0 to 300 hours. The '0 hours' soaking corresponds to pure tensile results (Table 1) at elevated temperatures. Figure 5 represents the same data. It can be observed that when thermal soaking was performed for 28 hours, there was an increase in the residual tensile strength. Increasing the thermal soaking time beyond 28 hours did not increase the tensile strength any further. Recently, Wiederhorn et al. (1993) have reported that thermal soaking at elevated temperatures ( 1200°-1400°C) resulted in the devitrific,ation of the intergranular phases in a HIPed silicon nitride (NT-154). From the results obtained, it could be speculated that the increase in residual tensile strength at 1200°C was due to the devitrification of the intergranular phases. A similar effect at 1300°C was probably not observed because the test temperature was beyond the softening point of the intergranular phases.

Figures 6 (A) and (B) show high resolution micrographs of the as-received and tested (thermal soaking/tensile test at 1200°C for 28 hours) PY6 material. From figure 6(A), it can be seen that there is some amorphous phases present at triple point junctions and grain boundaries. During thermal soaking, devitrification occurs rapidly, resulting in nearly complete crystallization of the glass contained within the multi-grain junctions through out the sample (Figure 6(B)). This may be one of the reasons for the increase in the residual tensile strength during thermal soaking. The strength improvement is only moderate since the PY6 material contained only a few percent of glassy phase.

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Table 2. Effect of thermal soaking time and temperature on residual tensile strength.

Thermal soaking time

(Hrs.)

Average residual tensile strength (MPa)

Soaking temperature (1200°C)

Soaking temperature

(1300°C)

0 370 250

28 436 240

66 404 _

300 421 290

0 27.7 66 300 Thermal Soaking Time (hrs.)

Figure 5. The effect of thermal soaking time and temperature on residual tensile strength.


Figures 6. High resolution micrographs of: (A) as-received and, (B) a tested (thermal soaking/tensile test at 1200°C, soaking time of 28 hours) PY6 material.

Micrographs of the etched surfaces of the material, both as-received and tested (thermal soaking/tensile test at 1200°C, soaking time of 300 hours) showed that the silicon nitride grains varied in both size and morphology. Grain shapes tended to be either equiaxed or acicular. The acicular grains were usually 1-5 Am long with widths of 1 Am or less in the as-received condition while in the tensile tested specimens the acicular grains were 2-8 Am long with widths of 2.5 Am or less. The equiaxed grains had dimensions either in the subrnicron range or in the 1-3 Am range for both conditions. Figures 7 (A) and (B) show micrographs of transverse sections for both conditions.

Figure 7. Representative view of microstructure showing variable size and morphology of: (A) an as-received and, (B) a tested (thermal soaking/tensile test at 1200°C, soaking time of 300 hours) PY6 material.

8


Creep Tests

Creep tests were conducted at 1200° and 1300°C, and at various loads. A total of ten (10) specimens were tested. Table 3 summarizes the results obtained after creep test. Figure 7 shows a typical creep curve obtained at 1200°C and at an applied load of 60% af (crf is the failure strength at the particular temperature). The creep strain rate was measured by approximating steady state behavior with a line.

Table 3. GTE-PY6 creep test summary

Applied stress (MPa)

Steady state strain rate (1/sec)

Rupture time (Hrs.)

Test conducted at 1200°C

161 4.69E-10 66 (NF)

174 2.43E-07 1.0

186 1.86E-09 270 (NF)

202 1.12E-09 1050 (NF)

217 2.14E-07 0.7

Test conducted at 1300°C

176 9.98E-07 0.17

176 1.09E-06 0.7

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Figure 8. TyPical creep curve for GTE-PY6


As shown in Figure 9 typical volume-initiated failures with a well defined mirror region were observed in creep tested specimens. A number of inclusions were observed at the fracture initiation sites. These inclusions were found to be iron-containing compounds (Dixie, 1993).

Figure 9. Typical volume-flaw-initiated failure in a creep tested specimen.

The tests conducted at 1200°C (Figure 8) clearly exhibited a dominant steady-state region but did not exhibit a tertiary region. However, these tests did not exhibit the expected trend of increasing strain rate with increasing stress and instead, were somewhat erratic in nature. This erratic nature has been observed in previous studies of this material (Ferber and Jenkins, 1992) and is probably due to the variation in the crystallinity of the inter-granular phase and Fe-bearing impurities.

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Figure 10 compares the applied stress versus steady state creep data for HIPed PY6 and sintered SNW-1000 silicon nitrides at 1200°C. A stress exponent 'n' of 9.36 was obtained for PY6, while SNW-1000 exhibited a stress exponent of 4.32. This shows that at applied stresses, PY6 experiences lower steady-state strain rates or improved creep resistance as compared to SNW-1000. Ferber and Jenkins (1992) have reported stress exponent values of 16 at 1150°C, 5.6 at 1260°C and 4.7 at 1370°C. The value of 'le of 9.36 obtained in the present study for tests at 1200°C appropriately falls between the range of these values. Ferber and Jenkins (1991) have suggested that the transition of high values of 're (16 at 1150°C) to medium values (5.6 at 1260°C) probably results from a change in the failure mechanisms from 1150° to 1260°C. The applied stress versus failure time data for the specimens indicated (Figure 10) that the failure mechanism at 1200°C is similar to that at 1150°C. From Figure 11 it can be evident that slow crack growth controls the failure at 1200°C. Hecht et al (1992) and Ferber and Jenkins (1992) have suggested that slow crack growth of pm-existing flaws or defects as well as the extension of these flaws control the failure in the temperature range of 1000°4200 °C. The failure at 1260° and 1370°C is controlled by creep failure.

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FAILURE TIME (Hrs.)

Figure 11. Applied stress versus Failure time data for GTE-PY6.

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Figure 12. Rupture time versus Strain rate for GTE-PY6 Si 3N4 (1150°, 1260° and 1370°C tests by ORNL; 1200°C tests by NC A&T SU).

A Monkman Grant type relation can be derived for GTE PY6-Si 3N4, as shown in Figure 12. The slope of the rupture time versus strain rate relation for the specimens tested at 1200°C was 0.97, (7-- 1.0), which is consistent with the values for other structural ceramics (Ferber and Jenkins, 1992).

CONCLUSIONS

This study provides an insight into the mechanical behavior of a HIE'ed Si 3N4 ceramic at elevated temperatures. The tensile strength of GTE-PY6 was found to drop considerably at high temperature probably due to the softening of the intergranular phase. A significant difference in the mechanical behavior was observed between IllPed and the sintered Si 3N4 ceramics. This could be due to the difference in the amount of densification aids and the processing methods.

It was observed that thermal soaking at 1200°C increased the residual tensile strength. Increasing the thermal soaking time beyond 28 hours did not increase the residual tensile strength any further. It is believed that the soaking time did not have a significant effect on the residual tensile strength due to the rapid devitrification of the intergranular phases. Similar soaking treatments at 1300°C did not produce the same effect on the residual tensile strength

A creep stress exponent 'n' of 9.36 was obtained for GTE-PY6 versus 4.32 for SNW-

12


1000 at the same applied stresses at 1200°C. This showed that GTE-PY6 has improved creep resistance compared to SNW-1000 at 1200°C. The Monkman Grant relation holds good for GTE-PY6 and the slope of the rupture time versus strain rate relation was found to be consistent with the values for other structural ceramics.

The tensile specimens failed at surface flaws, while the creep specimens failed due to flaws in the volume. Most of the surface and volume failures were at the site of inclusions. These inclusions were found to be primarily iron containing compounds. The improvement in residual tensile strength due to thermal soaking could play a vital role in the performance of Si 3N4 ceramics for long-term heat engine applications.

Acknowledgements

This research was sponsored by the U. S. Department of Energy, Office of Transportation Technologies, as part of the Ceramic Technology for Advanced Heat Engines Project of the Advanced Materials Development Program, under contract DE-AC05-84R21400 with the Martin Marietta Energy System, Inc.

The authors wish to thank Dr. Ray Johnson, and Dr. Ken Liu of Oak Ridge National Laboratory for their continuing support and many helpful technical suggestions throughout this program.

References

Dixie, M. T., 1993, "Mechanical Properties Investigation of GTE-PY6 Silicon Nitride Material at Elevated Test Temperatures," M.S. Thesis, North Carolina Agricultural and Technical State University, Greensboro, NC.

Ferber, M. K., and Jenkins, M. G., 1992, "Evaluation of the Elevated Strength and Creep-Fatigue Behavior of HIPed Silicon Nitride," Journal of American Ceramic Society, Vol. 75, No. 9, pp. 2453-2462.

Hecht, N. L., Goodrich, S. M., Chuck, L., McCollum D. E., and Tennery, V. J., 1992, "Mechanical Properties Characterization of One SiC and Two Si 3N4 Commercially Available Ceramics," American Ceramic Society Bulletin, Vol. 71, No. 4, pp. 653-659.

Liu, K. C., and Brinkman, C. R., 1985, "Tensile Cyclic Fatigue of Structural Ceramics," Proceedings of the Twenty Third Automotive Technology Development Contractors' Coordination Meeting, Society of Automotive Engineers, Warrendale, PA, P-165, pp. 279-284.

Sankar, J., Kristmaraj, S., Vaidyanathan, R., and Kelkar, A. D., 1993, "Elevated temperature behavior of sintered silicon nitride under pure tension, creep and fatigue," Life prediction methodologies and data for ceramic materials, ASTM STP 1201, C. R. Brinkman and S. F. Duffy, eds., American Society for Testing Materials, Philadelphia, PA (in press).

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Sankar, J., Kelkar, A. D., Vaidyanathan, R., and Gao, J., 1991, "Creep Testing of S/s1W-1000 Sintered Silicon Nitride," Proceedings of the Annual Automotive Technology Development Contractors' Coordination Meeting, Society of Automotive Engineers, Warrendale, PA, P-256, pp. 293-305.

Tsai, R. L., and Raj, R., 1982, "Creep Fracture in Ceramics Containing Small Amounts of a Liquid Phase," Acta Metall., Vol. 30, pp. 1043-1058.

Vaidyanathan, R., Sankar, J., and Avva, V. S., 1987, "Testing and Evaluation of Si 3N4 in Uniaxial Tension at Room Temperature," Proceedings of the Bventy Fifth Automotive Technology Development Contractors' Coordination Meeting, Society of Automotive Engineers, Warrendale, PA, P-209, pp. 175-186.

Wieclerhorn, S. M., Hockey, B. J., Cranmer, D. C., and Yeckley, R., 1993, "Transient Creep Behavior of Hot Isostatically Pressed Silicon Nitride," Journal of Materials Science, Vol. 28, pp. 445-453.

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