17_Colombo Ceramic Foams MDI

Mechanical Properties of Silicon Oxycarbide Ceramic Foams

Paolo Colombo*,†

Dipartimento di Chimica Applicata e Scienza dei Materiali, Universita di Bologna, Viale Risorgimento 2,40136 Bologna, Italy

John R. Hellmann* and David L. Shelleman*Materials Research Institute, Department of Material Science and Engineering, The Pennsylvania State University,

University Park, Pennsylvania 16802

The mechanical properties of ceramic foams obtained througha novel process that uses the direct foaming and pyrolysis ofpreceramic polymer/polyurethane solutions were investigated.The elastic modulus, flexural strength, and compressivestrengths were obtained for foams in the as-pyrolyzed condi-tion; values up to 7.1 GPa, 13 MPa, and 11 MPa, respectively,were obtained. The strength of the foam was virtually un-changed at temperatures up to 1200°C in air; however,long-term exposure at 1200°C led to a moderate degradation instrength, which was attributed to the evolution of intrastrutporosity during the oxidation of residual free carbon, as well asdevitrification of the foams struts.

I. Introduction

CELLULAR ceramics possess properties that are of interest forseveral engineering applications. These properties include low

bulk density, low net thermal conductivity, low dielectric constant,high specific strength and stiffness, excellent resistance to thermo-chemical corrosion, and excellent thermal shock resistance.1–4

The properties of ceramic foams are influenced by their relativedensity and morphology (open cell versus closed cell), as well asthe properties of the struts that comprise the cell walls.1,2 Numer-ous methods for manufacturing cellular ceramics with open- andclosed-cell morphologies have been developed over the last twentyyears. Ceramic slurry deposition on sacrificial polymer foampreforms followed by sintering,5 direct foaming from polymericprecursors or ceramic sols,6,7 and chemical vapor deposition(CVD) of ceramics onto carbon foam preforms8 are among thosemethods that have been demonstrated for the manufacture ofceramic foams with requisite microstructures and macrostructuresfor good thermomechanical behavior.

Recent results in our laboratories on the direct foaming ofpreceramic polymer/polyurethane solutions, followed by pyrolysisat elevated temperature in an inert gas ambient, have proven usefulfor fabricating silicon-oxycarbide-based foams to near-net shape,and with excellent room-temperature mechanical properties.9 Theapproach offers significant advantages over other ceramic foamfabrication techniques in that it is a simple, one-step process thatallows the fabrication of foams to complex net shapes by castinginto molds. Selection of the appropriate preceramic polymer (andfillers) offers the ability to fabricate a variety of ceramics in this

fashion, including a wide range of nitrides, carbides, and oxides.Tailoring the pyrolysis conditions and the chemistry of thepolyurethane backbone allows manipulation of the resulting stoi-chiometry, carbon content, interstrut and intrastrut porosity, andsubsequent mechanical, elastic, and chemical characteristics of thefoam.9–11

However, to date, little has been reported on the thermochemi-cal and thermomechanical durability of silicon-oxycarbide-basedfoams that have been produced using the direct foaming process.The objective of this study was to examine the interdependence offoam structure, strength, elastic modulus, fracture toughness, andoxidation resistance at room temperature and after high-temperature oxidative exposure at 800° and 1200°C.

II. Experimental Procedure

(1) MaterialsIn this study, the typical polyurethane component in the

preceramic polymer polyurethane (PU) foam was obtained from amixture of two polyether polyols (Tercarol� 3 (hydroxyl numberof 56 mg KOH/g, viscosity of 500 mPa�s at 25°C, molecularweight (MW) of 3000) and Tercarol� 1 (hydroxyl number of 168mg KOH/g, viscosity of 280 mPa�s at 25°C, MW � 1000), eachsupplied by EniChem, San Donato Milanese, Italy), amine cata-lysts (Niax A1 and Niax A-33), surfactants (polydimethyl siloxanecopolymers SC250 and SH205, Goldschmidt Italia, Pandino (CR),Italy), a blowing agent (dichloromethane, CH2Cl2), and polymericMDI isocyanate (Tedimon� 31 (polyphenylmethane polyisocya-nate), viscosity of 180–250 mPa�s at 25°C, with 31% NCO;EniChem). This mixture, when used by itself, produces a semirigidPU foam.

The preceramic polymer used in this study was a thermosettingmethyl-hydroxyl-siloxane (SR350, General Electric, SiliconeProducts Division, Waterford, NY), which can cross-link at roomtemperature through the amine-catalyzed condensation of SiOOHgroups. Its pyrolysis in an inert atmosphere yields a SixOyCz

ceramic, with a weight loss of �16%.12 Silicon oxycarbide (SiOC)ceramics prepared from preceramic polymer routes are thermo-chemically stable, at least up to 1200°C in air, and show compo-sitional or microstructural modifications only after extended peri-ods of time at elevated temperature.13 Although CH2Cl2 waschosen as the solvent, it also serves as a physical blowing agent forthe PU system and is compatible with both the PU precursors andthe preceramic polymer. After blowing, the PU is phase-separatedas small islands (typically 5–10 �m in size) embedded in a siliconeresin matrix. During pyrolysis, the polyurethane decomposes(leaving a limited free-carbon residue11), and the preceramicpolymer is converted to a ceramic material.

(2) Preparation and Characterization of FoamsA solution of SR350 polymer in dichloromethane was added to

the polyol mixture (polyol � amine catalysts � surfactant) and

R. Raj—contributing editor

Manuscript No. 187946. Received February 9, 2001; approved May 21, 2001.Invited paper for the 2nd International Workshop on “Ultrahigh Temperature

Polymer Derived Ceramics” (Boulder, CO, July 23–29, 2000).*Member, American Ceramic Society.†Author to whom correspondence should be addressed.

Ultrahigh-Temperature Ceramics

J. Am. Ceram. Soc., 84 [10] 2245–51 (2001)

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excess dichloromethane was evaporated. Then, a stoichiometricamount of isocyanate was added and the expansion reaction of themixture was started by mechanical stirring. The final preceramic-polymer:polyurethane weight ratio in the foam was 1:1. Because ofthe presence of amine catalysts, a more-or-less complete crosslink-ing of the silicone resin occurs via a condensation of silanols andthe release of water. The silicone resin was preconditioned byheating for various amounts of time at 100°–150°C, to control theamount of SiOOH groups that are present, which affects the risingprofile of the foams (chemical expansion).

Different final foam densities (�0.27–0.56 g/cm3) were ob-tained by varying the amount of dichloromethane and catalysts, themold temperature (which was 17°–35°C), and the polyetherpolyols ratio in the PU precursor mixture (which affects theviscosity and the rising characteristics). Foaming occurred withina very short time after stirring, and the foams were conditioned for24 h at 30°C before pyrolysis for 1 h at 1200°C in a flux ofnitrogen gas (99.99% pure, with a heating rate of 2°C/min).

The foam morphology was characterized using scanning elec-tron microscopy (SEM) (Cambridge Instruments, Cambridge,U.K.), and the geometrical features were derived using imageanalysis software (Image-Pro Plus, Media Cybernetics, SilverSpring, MD). The bulk (apparent) density of the foams wascomputed from the weight:volume ratio; the true density wasmeasured on powdered samples using a helium pycnometer. Thecrystalline phase assemblage was characterized by X-ray diffrac-tometry (XRD) (Model 1730/1820, Philips Research Laboratories,Eindhoven, The Netherlands), using CuK� radiation (40 kV, 40mA). The coefficient of thermal expansion (CTE) was measuredusing a dilatometer (Model 402E, Netzsch, Selb, Germany) inflowing nitrogen gas or in air, at a heating rate of 10°C/min.Simultaneous differential thermal analysis and thermogravimetry(DTA/TGA) were performed using a thermobalance (ModelSTA409, Netzsch) in different atmospheres (air or nitrogen, at aheating rate of 10°C/min).

The mechanical and elastic behavior was investigated in air onsamples in the as-pyrolyzed condition at room temperature (sam-ples labeled “RT”), 800°C (samples labeled “800°C”), and 1200°C(samples labeled “1200°C”), as well as after a long-term oxidationtreatment (12 h at 800°C (samples labeled “HT 800°C”) or 12 h at1200°C (samples labeled “HT 1200°C”) in static air), to assess theeffect of exposure to aggressive environments on the foams. Themechanical property test specimens were sectioned from panels(dimensions of 20 cm � 14 cm � 2 cm) that were obtained bycasting the solution into a closed aluminum mold. The specimenswere cut before pyrolysis to avoid shape distortions that sometimesoccur during the pyrolysis of a large panel. The flexural strengthwas measured by four-point bending (inner span of 10 mm, outerspan of 20 mm) on specimens with dimensions of 45 mm � 6mm � 6 mm, using a universal testing machine (Model 4202,Instron, Danvers, MA) with a cross-head speed of 0.5 mm/min.The crushing strength was measured by compression testing, usingalumina loading rams and a cross-head speed of 0.5 mm/min, onsamples with a nominal size of 6 mm � 6 mm � 12 mm. Strengthdata were analyzed using two-parameter Weibull statistics, with amaximum likelihood approach to estimating the characteristicstrength and Weibull modulus.14

The fracture toughness was measured using the single-edge-notched-beam method in three-point bending (span of 30 mm),with a notch depth that was �50% of the specimen thickness. Theelastic modulus was measured at room temperature in air, usingultrasonic sound velocity and acoustic resonance methods.15

III. Results and Discussion

(1) Physical and Morphological CharacterizationTo produce a sufficient number of specimens for testing,

different panels of the same nominal density were fabricated.Weight loss during pyrolysis in nitrogen was �55%, whichindicated that residual carbon that was derived from the decom-position of the PU backbone is retained in the foams (the ceramic

Table I. Summary of Mechanical and Elastic Properties of Different Foams†

Density (g/cm3) �F (MPa)��F

(MPa�m3/mF) mF

�C(MPa)

��C(MPa�m3/mC) mC

Elastic modulus,E (GPa)

Fracture toughness,KIC (MPa�m1/2)

Room temperature0.274 0.009 1.81 0.24 (5) 1.89 7.6 1.31 0.51 (9) 1.46 2.7 1.60 0.30 (5) n/m0.400 0.011 3.69 0.28 (9) 3.79 15.0 3.32 0.74 (10) 3.60 4.4 3.54 0.50 (9) 0.16 0.01 (8)0.447 0.01 4.63 0.89 (5) 4.94 4.4 3.76 1.52 (9) 4.23 2.8 4.30 0.50 (5) n/m0.450 0.027 3.66 0.57 (10) 3.88 6.3 3.85 0.94 (10) 4.20 3.6 4.94 0.70 (10) 0.16 0.01 (5)0.500 0.015 7.00 1.27 (10) 7.50 5.0 7.30 1.64 (10) 7.92 4.4 5.57 0.13 (10) 0.25 0.01 (5)0.526 0.01 5.95 0.69 (6) 6.21 7.2 8.45 2.07 (10) 9.25 3.7 5.24 0.27 (6) n/m0.550 0.012 11.17 1.05 (10) 11.59 10.1 8.98 3.04 (10) 10.60 4.6 7.40 0.37 (10) 0.24 0.07 (5)0.558 0.013 6.43 0.74 (8) 6.69 9.1 9.90 1.78 (9) 9.97 3.0 6.05 0.48 (5) n/m

800°C0.400 0.011 4.45 1.11 (10) 4.84 4.3 3.19 0.99 (5) 3.50 2.9 n/m n/m0.450 0.027 4.41 0.41 (6) 4.19 4.8 3.43 0.96 (5) 3.74 3.1 n/m n/m0.500 0.015 8.55 1.24 (9) 9.04 6.6 4.79 0.62 (5) 5.02 5.7 n/m n/m0.550 0.012 12.08 2.67 (8) 13.05 4.3 7.76 0.39 (5) 7.87 19.6 n/m n/m

1200°C0.400 0.011 4.72 0.62 (5) 4.91 8.8 6.01 1.02 (5) 6.34 6.1 n/m n/m0.450 0.027 4.94 0.27 (5) 5.01 15.6 6.60 0.49 (5) 6.75 13.6 n/m n/m0.500 0.015 9.00 0.84 (5) 9.29 9.2 9.30 0.79 (5) 9.56 10.3 n/m n/m0.550 0.012 10.75 1.39 (7) 11.28 7.0 10.19 0.87 (5) 10.48 10.4 n/m n/m

HT 800°C0.400 0.011 3.73 0.97 (5) 4.05 3.5 2.98 0.67 (5) 3.21 3.9 3.29 0.20 (5) n/m0.450 0.027 4.08 0.48 (5) 4.25 7.4 3.33 0.69 (5) 3.65 3.7 3.53 0.35 (5) n/m0.500 0.015 7.62 0.82 (5) 7.96 10.5 5.06 0.67 (5) 5.30 6.7 5.02 0.28 (5) n/m0.550 0.012 13.13 2.70 (5) 14.13 6.1 8.57 0.70 (5) 8.82 8.8 7.08 0.29 (5) n/m

HT 1200°C0.400 0.011 2.29 0.31 (5) 2.40 6.2 2.33 0.75 (5) 2.57 2.8 2.28 0.30 (5) n/m0.450 0.027 2.36 0.21 (5) 2.42 13.8 2.88 0.68 (5) 3.11 3.7 2.66 0.27 (5) n/m0.500 0.015 4.85 0.75 (5) 5.11 5.6 4.31 0.75 (5) 4.50 6.5 3.56 0.31 (5) n/m0.550 0.012 7.47 0.85 (5) 7.76 7.7 5.61 0.54 (5) 5.80 8.1 4.95 0.38 (5) n/m

†The value after the plus-or-minus sign () indicates the standard deviation; value within the parentheses is the reported number of specimens tested. Flexural and compressivecharacteristic strength (�F and �C, respectively) were determined by the maximum likelihood method, using WeibPar software.14 “n/m” � not measured.

2246 Journal of the American Ceramic Society—Colombo et al. Vol. 84, No. 10

yield for pure SR350 polymer is 84% at 1200°C), in agreementwith prior work.11 Carbon analyses, which were performed usinga carbon determinator (Model CS-244, LECO, St. Joseph, MI)with an oxidation accelerator, confirmed that the SiOC foamcontained more carbon (29.6 0.9 wt%) than the SiOC ceramicobtained from the direct pyrolysis of the SR350 silicone resin(14.1 0.7 wt%.)

The variability in density of the pyrolyzed samples was small(based on 50 samples per unit of density, a standard deviation of�2.5% was observed), which indicated that the sample fabricationprocess is quite reproducible. The measured densities are given inTable I.

The typical cell morphology, as a function of density, isillustrated in Fig. 1. The average cell size varies linearly with thefoam density, from 100 40 �m for the 0.55 g/cm3 samples to600 145 �m for the 0.27 g/cm3 samples. The cells are regularlyshaped, with no evidence of preferential orientation, and seem tohave interconnected pores; i.e., the foams can be considered to bepredominantly open-celled. Some closed cells are present (�10%–20%), and the number of closed cells increases as the foam densityincreases. The cell walls and the struts are mostly dense, with onlyoccasional closed porosity present. No macroscopic cracks ordefects are observed in the ceramic foams after pyrolysis. The CTEof the ceramic foams (at 20°–1200°C) is 3.50 � 106 0.19 �106 K1. All the SiOC foams that have been pyrolyzed at1200°C are X-ray amorphous.

(2) Mechanical PropertiesThe flexural and compressive strengths (�F and �C, respec-

tively) for the samples, as a function of temperature, as well asafter long-term oxidative exposure, are given in Table I. Theroom-temperature compressive and flexural strengths, as a func-tion of bulk density, are presented in Weibull fashion in Fig. 2.Although at least thirty samples are needed to determine accurateWeibull parameters, we were limited in the number of samplesavailable for this study. As such, the Weibull data listed in thisstudy were obtained solely to compare sample behavior and is notintended to represent true material parameters.

In agreement with prior studies,1,2,16 the strengths increase asthe bulk density increases. Figure 3 presents the flexural andcompression relative-strength data, as a function of the relativedensity, and is compared with the behavior that is expected forcompletely open-cell or closed-cell foams, according to modelsproposed by Ashby,1 Gibson and Ashby,2 and Brezny andGreen.16 We used the pycnometrically determined value of 2.12g/cm3 for the density, 97.9 GPa for the Young’s modulus (E),and 153 MPa for the modulus of rupture (MOR) as referencevalues for the SiOC strut and cell-wall material.12

Ashby’s theory1,2 indicates that the relative strength of acellular material is related to its relative density via the expression

�

�s� C��

�s�m

Fig. 1. Morphologic characteristics of the foams produced in this study (densities of (a) 0.40, (b) 0.45, (c) 0.50, and (d) 0.55 g/cm3).

October 2001 Mechanical Properties of Silicon Oxycarbide Ceramic Foams 2247

where C is a dimensionless constant and the exponent m is 3⁄2 or 2,depending on if the cell morphology is open cell or closed cell,respectively. Similarly, the relative elastic modulus is related torelative density as

E

Es� C��

�s�n

where n � 2 for open-cell foams or n � 3 for closed-cell foams.A linear-regression fit of the room-temperature strengths yields

exponents of 2.3 and 3.6 for the bending strength and compressionstrength (with correlation coefficients of R � 0.92 and 0.97),respectively. Similarly, a plot of relative elastic modulus versusrelative density for the as-pyrolyzed specimens at room tempera-ture yields a relative density exponent of 1.83, with R � 0.97. In

Fig. 2. Strength distributions as a function of foam density for samples measured in (a) flexure and (b) compression. F is the failure probability.

Fig. 3. Comparison of room-temperature flexural and compressive relative-strength data for behavior expected for open-cell versus closed-cell foams (afterAshby,1 Gibson and Ashby,2 and Brezny and Green16).


Ashby’s model,1,2 the morphology of the open-cell foams consistsof interconnected ligaments (struts), and the cells do not have anycell walls (thus, the structure is totally open); such a morphologyis generally obtained in ceramic foams that are produced using thereplica method (so-called “reticulated foams”). In contrast, in thesame model, the morphology of closed-cell foams consists of cellsin which the cell walls are completely retained and do not containany openings (cell windows). In our case, the morphology of thefoams obtained from the direct foaming method, as shown in Fig.1, is much more similar to a closed foam, in that the cell walls arelargely retained. However, the foams can be considered to have aninterconnected porosity (and, thus, can still be considered to beopen cell), because some openings (cell windows) are present inmost of the cell walls. Thus, it is reasonable that the strength datafor our samples are more descriptive to what the model predicts forclosed-cell foams.

The lack of a rigorous fit to Ashby’s theory1,2 may be due toseveral factors, such as the presence of macrostructural inhomo-geneities (a distribution of cell sizes), the presence of mixed closedand open cells, the volume fraction of solids in the cell faces(hybrid cell walls—i.e., those containing some openings—cer-tainly contribute to a different stress distribution than if the foamhas a completely open-cell or closed-cell structure), the presenceof closed voids or defects in the struts, or variations in the (strut)strength or strut microstructure, relative to the density.16 Forexample, the unexpectedly low flexural strength that has beenrecorded for some specimens with high density (0.558 g/cm3) canbe attributed to the presence of an abnormally large amount ofvoids or microdefects in the struts, which are derived from thedecomposition of the PU component in the foam. This observation,in turn, is due to a change in the type and ratio of polyols (and,

thus, compatibility with the silicone resin) in the PU mixture, tovary the rheological and rising characteristics of the startingsolution.

The high-temperature flexural and compressive strengths of thematerials in the as-pyrolyzed condition and after oxidative expo-sure for 12 h at 800° and 1200°C are shown in Fig. 4. (To improvethe readability, the data relative to the foam with a density of 0.274g/cm3 have been omitted in the figure. They are reported in TableI.) The trend of increasing strength with increasing bulk density ismaintained at high temperature, in both the as-pyrolyzed andoxidized specimens. No significant modification in strength isobserved, relative to the room-temperature, as-pyrolyzed strengthsfor samples tested at 800° and 1200°C. The higher strength that isexhibited in compression by the foams when tested at 1200°Cmight be attributed to viscoelastic deformation phenomena thatpossibly occur in the material. Oxidative exposure at 800°C yieldsno strength modification; however, the 1200°C oxidized speci-mens do experience a reduction of �30% in both �F and �C, aswell as in E (see Table I).

The mechanical properties of the foams produced with thismethod are superior to those reported in the literature for porousceramics of similar density that have been produced using con-ventional technology (e.g., reticulated foams17). This result may bedue to the lack of macroscopic defects in the struts and cell walls,such as the hollow features that are typical of conventionallymanufactured reticulated foams.4 In fact, materials that have beenobtained by gelcasting also display better mechanical properties,because of the dense structure of the ceramic struts.18

TGA studies (Fig. 5) suggest that the strength degradationduring long-term oxidative exposure at 1200°C may be attributedto the loss of carbon from the struts and cell walls of the foams.

Fig. 4. (a) Flexural strength and (b) compression strength of SiOC foams, measured at various temperatures and after oxidation treatments, as a functionof bulk density.


The nitrogen TGA results in Fig. 5(a) reveal that the polymer-to-ceramic conversion was completed with the processing schedulethat was adopted for the pyrolyzed specimens; in fact, the weightloss in nitrogen was �1% for all temperatures and times investi-gated. TGA performed in air showed a weight loss of �1.5 wt%during the reheating of pyrolyzed specimens in the temperatureinterval of 600°–800°C, which is linked to the burnoff of some ofthe residual free carbon, in agreement with earlier work.19 How-ever, the weight loss at 800°C occurs quickly and then ends (Fig.5(b)) with no further change in weight up to times of 12 h. Incontrast, a significantly larger weight loss (2.6%) is observed forsamples that have been exposed to 1200°C in air, with a noticeablereduction in weight loss rate at long times, possibly because of thecompetition between CO2 evolution and the formation of apassivating silica scale on the foam struts (CO2 evolution has beenobserved when Fourier transform infrared (FTIR) analysis wasperformed on the gas that was released during an oxidation test19).Carbon analysis has indicated carbon contents of 29.2 0.5 and24.7 0.6 wt%, after the 800° and 1200°C treatments, respec-tively. Compared with the starting carbon content of 29.6 0.9wt% in the as-pyrolyzed specimens, it is apparent that the residualfree carbon is consumed more aggressively at 1200°C than at800°C, at least for oxidation times of up to 12 h. SEM fracto-graphic analysis of the foams has revealed an increase in the strutporosity in oxidized specimens (Fig. 6(a)–(c)), because of theburnout of PU-derived carbon-rich islands. Moreover, the forma-tion of cracks in the struts is observed for samples that have beenoxidized at 1200°C; this phenomenon is possibly due to either theevolution and release of CO2 during oxidation or fracture of thesilica scale during rapid cooling from the oxidation temperatures.

Powder XRD revealed that the samples oxidized at 800°Cremained X-ray amorphous; however, the foams began to devitrifyduring oxidation at 1200°C. Minor quantities of cristobaliteformed in SiOC foams that were oxidized for long times at1200°C. Notably, the same treatment that was performed on a bulkSiOC ceramic that was obtained from the direct pyrolysis ofSR350 silicone resin did not produce the same change incrystalline-phase assemblage; i.e., the specimen remained amor-phous.19 The difference in devitrification behavior between bulkSiOC and the foams may possibly be attributed to the oxidation ofresidual carbon that remained from the PU precursor, which yieldsheterogeneous nucleation sites for the crystallization of silica.

A possible solution to the observed limited stability of the SiOCceramic foams when exposed to an oxidative environment at hightemperature would be the use of a smaller amount of PUprecursors when fabricating the material. This practice, in fact, hasbeen recently demonstrated,19 wherein foams that were preparedwith only �30% of the amount of PU used in these experimentsshowed no decrease in strength after exposure in air at 1200°C for12 h.

IV. Summary

In summary, we have demonstrated a novel process for thenear-net-shape fabrication of silicon oxycarbide (SiOC) cellularceramics via the direct foaming and pyrolysis of preceramicpolymer/polyurethane solutions. The approach offers substantialflexibility in producing foams with controlled cell size andmorphology, bulk density, and residual carbon content. Thestrength of the foams produced in this manner is superior to that ofmost foams produced commercially and is virtually unaffected byshort-term high-temperature exposure in air. Modifications instrength are observed only after long-term exposure at 1200°C in

Fig. 5. Weight loss of a SiOC foam pyrolyzed sample, as a function ofheating temperature or heating time (at 800° and 1200°C). Analysis wasperformed in nitrogen and air.

Fig. 6. SEM micrographs of (a) the strut microstructure in the as-pyrolyzed foam, (b) the strut after oxidation at 800°C, and (c) the strut afteroxidation at 1200°C. Foam density was 0.5 g/cm3.


air. The modification in strength is attributed to the evolution ofintrastrut porosity that is due to oxidation of residual free carbon,accompanied by devitrification of the SiOC and/or the passivatingsilica scale that forms during oxidation.

Acknowledgments

The authors wish to thank Dr. M. Modesti (University of Padova, Italy) forformulating and tailoring the composition of the polyurethane mixture used in thisstudy, and Mr. N. Phelps for the help in the characterization of the specimens. Theyare also indebted to Prof. D. Green (Pennsylvania State University) for the numeroushelpful discussions.

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17_Colombo Ceramic Foams MDI