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Advances in Ceramic Matrix Composites XI


Ceramic Transactions Volume 175

Proceedings of the 107th Annual Meeting of The American Ceramic Society, Baltimore, Maryland, USA (2005)

Editors Narottam P. Bansal

J.R Singh Waltraud M. Kriven

Published by The American Ceramic Society

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IV Advances in Ceramic Matrix Composites XI

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Contents

Preface ix

Ceramic Fibers

Poly[(Alkylamino)Borazine]-Derived Boron Nitride Fibers for Composite Applications 3 Samuel Bernard, Sylvain Duperrier, Fernand Chassagneux, David Cornu, and Philippe Miele

Processing

Multilayered Materials by ICV1 in Non-Oxide Self-Healing Ceramic Matrix Composites for High Temperature Applications 13 L. Vandenbulcke

Processing of Oxide/Oxide Composite Components for Efficient Energy Conversion Applications 27

Cengiz Kaya

From Polysaccharides to SiSiC Composites by 3D Printing* 37 Nahum Travitzky, Katrin Zimmermann, Reinhold Melcher, and Peter Greil

Characterization

Ultrasonic NDE of Reaction Bonded Ceramics 49 P.G. Karandikar and M.K .Aghajanian

On the Use of Digital Image Correlation to Analyze the Mechanical Properties of Brittle Matrix Composites 6

3

François Hild, Jean-Noël Périé, Jacques Lamon,and Matthieu Puyo-Pain

Multiscale X-Ray CMT of C/C Composite Preforms: A Tool for Properties Assessment* 77 O. Coindreau, G.L. Vignoles, and J.-M. Goyheneche

Microstructural Investigations of Reinforcing Materials in Zinc Phosphate Composites 85 Charles A. Weiss, Jr., Humberto Benitez, Melvin C. Sykes, and Philip G. Malone

Oxide Composites

High-Temperature Thermal Conductivity of Alumina-Reinforced Zirconia Composites 95 Narottam P. Bansal and Dongming Zhu

Dielectric Behavior in Ni^CoQ^Mn^Fe^O^g+PZT Composites 107 S.V. Suryanarayana, S. Narendra Babu, and T. Bhimasankaram

Advances in Ceramic Matrix Composites XI v

Mechanical Properties

Interlaminar Tension/Shear Properties and Stress Rupture in Shear of Various Continuous Fiber-Reinforced Ceramic Matrix Composites 119

Sung R. Choi and Narottam P. Bansal

Fatigue Behavior of Nextel™720/Alumina (N720/A) Continuous Fiber Ceramic Composite - Effects of Temperature and Steam Environment 135

C. A. Eber, M.B. Ruggles-Wrenn, and S. Mall

Microstructure and Mechanical Properties of Polymer-Derived Al203-SiC Micro-Nano Composites 151

Du§a n Galusek, Jaroslav Sedlâôek, and Ralf Riedel

High Temperature Creep of Yttria Tetragonal Zirconia Nanocrystals: The Role of Yttrium Segregation at the Grain Boundaries 161

C. Lorenzo, D. Gomez, A. Dominguez, and J. Routbort

Creep-Rupture Behavior of Nextel™720/Alumina (N720/A) Continuous Fiber Ceramic Composite - Effects of Temperature and Steam Environment 169

L.B. Harlan, M.B. Ruggles-Wrenn, and S. Mall

Damage Morphology of C/C-SiC Composites Under Impact Tests 181 V.K. Srivastava

Geopolymers and Geopolymer Matrix Composites

On Mix Compositions of Fly Ash Based Inorganic Polymeric Materials 191 Peijiang Sun and Hwai-Chung Wu

Nanostructural Design of Multifunctional Geopolymeric Materials 203 Peter Duxson, Grant C. Lukey, and J.S.J. van Deventer

Thermal Conversion and Microstructural Evaluation of Geopolymers or "Alkali Bonded Ceramics" (ABCs) 215

M. Gordon, J. Bell, and W.M. Kriven

Disposition of Water in Metakaolinite Based Geopolymers 225 D.S. Perera, E.R. Vance, K.S. Finnie, M.G. Blackford, J. V. Hanna, D.J. Cassidy, and C.L. Nicholson

High-Temperature Deformation of a Geopolymer 237 F. Gutierrez-Mora, A. Dominguez-Rodriguez, K.C. Goretta, D. Singh, J.L. Routbort, G.C. Lukey, and J.S.J. van Deventer

Modeling Si/Al Ordering in Metakaolin-Based Geopolymers 245 John L. Provis, Peter Duxson, Grant C. Lukey, and Jannie S.J. van Deventer

vi Advances in Ceramic Matrix Composites XI

Matrix and Interphase Design of Geopolymer Composites 253 Balakrishnan Nair, Qiang Zhao,Tina Rahimian, Reid F. Cooper, and Perumalsamy N. Balaguru

Index 265

* This paper was presented at the 29th International Conference on Advanced Ceramics and Composites, Cocoa Beach, FL, January 23-28,2005

Advances in Ceramic Matrix Composites XI vii

Preface

Advanced structural materials are needed for high-temperature applications in industries such as aerospace, power generation, and transportation. Ceramic-matrix composites are strong, tough, environmentally stable, light in weight, and have the ability to withstand high operating temperatures. This makes them viable candi-date materials for high temperature structural applications.

An international symposium on recent advances in ceramic-matrix composites and geopolymers was held dur-ing the 107th Annual Meeting, Exposition, and Technology Fair of The American Ceramic Society at Baltimore Marriott Waterfront, Baltimore, MD April 10-13, 2005. The objective of this symposium was to pro-vide an international forum for scientists, engineers, and technologists to discuss and exchange ideas on state-of-the-art ceramic-matrix composites and geopolymers. A total of 52 papers, including invited talks, oral pre-sentations, and posters were presented indicating continued interest in the scientifically and technologically in these important fields.

Researchers from 15 countries (Australia, China, France, Germany, India, Italy, Japan, New Zealand, South Korea, Slovakia, Spain, Taiwan, United Kingdom., Ukraine, and the United States) participated which reflects the international nature of this symposium. The speakers represented universities, industry, and government research laboratories.

These proceedings contain contributions on various aspects of ceramic-matrix composites and geopolymers that were discussed at the symposium. Twenty two papers describing the latest developments in the areas of ceramic fibers, processing and fabrication, characterization, oxide and non-oxide composites, mechanical behavior, fiber-matrix interface, geopolymers and geopolymer composites, etc. are included in this volume.

The editors wish to extend their gratitude and appreciation to the authors for their cooperation and contribu-tions, to the session chairs for their time and efforts in keeping the sessions on schedule, and to the reviewers for their useful comments and suggestions; without the contributions of all involved, this volume would not have been possible. Financial support from the Engineering Ceramics Division and the American Ceramic Society is gratefully acknowledged. Thanks are due to the staff of the meetings and publications departments of The American Ceramic Society for their invaluable assistance and for efficiently coordinating the review of the manuscripts.

It is our earnest hope that this volume will serve as a valuable reference for the engineers, scientists, and other technical people interested in different aspects of ceramic-matrix composites and geopolymers.

Narottam P. Bansal J. P. Singh Waltraud M. Kriven

Advances in Ceramic Matrix Composites XI IX

Ceramic Fibers

To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the property of The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, without the express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.

POLY[(ALKYLAMINO)BORAZINE]-DERIVED BORON NITRIDE FIBERS FOR COMPOSITE APPLICATIONS

Samuel Bernard*, Sylvain Duperrier, Fernand Chassagneux, David Cornu, and Philippe Miele Laboratoire des Multimatériaux et Interfaces (UMR CNRS 5615) University Claude Bernard - Lyon 1 43Bddu 11 novembre 1918 Villeurbanne, France, 69622

ABSTRACT A series of experimental boron nitride (BN) fibers were prepared from a B-

alkylaminoborazine-based polymer (=poly[(methylamino)borazine]) according to the Polymer-Derived Ceramic (PDC) route. The polymer was melt-spun in N2 into mono- and multifilaments, prior the curing of the resulting green fibers in NH3 at 400°C and the subsequent pyrolysis of the as-cured fibers in NH3 (1000°C), then N2 up to 1800°C to generate BN fibers. It was established that the melt-spinning operation could deliver BN fibers with different microtexture/microstructure, and therefore with different mechanical behavior depending on the spinning parameters. The relationship between mechanical properties and fiber microtexture/microstructure could be studied in the present paper. Melt-spinning operation into multifilaments provided either low-modulus fibers with a featureless cross-sectional microtexture as glassy-like materials or fibers with a microtextural skin-core heterogeneity decreasing the fiber strength. The disordered microstructure of low-modulus BN fibers consisted of disoriented nanosized grains mixed in an amorphous matrix. In contrast, the melt-spinning operation into monofilament produced high-modulus fibers with a coarse-grained microtexture. In such samples, the extended grains were ordered along the fiber-axis increasing the crack propagation along the cleavage basal planes, and therefore decreasing the failure strain of fibers.

INTRODUCTION Hexagonal boron nitride (h-BN) represents an advanced ceramic material with a layer

anisotropic structure and attractive properties such as high stiffness and toughness along the basal layers. It also exhibits a good oxidative resistance up to T - 1000°C, a good thermal stability up to T ~ 2500°C in an inert atmosphere and a low coefficient of thermal expansion (CTE) along the basal layers.1 Based on these properties, h-BN should be promising for preparing continuous fiber-reinforced ceramic-matrix composites (CFCCs) intended for high temperature applications. For CFCCs fabrication, high-modulus and strength oxidation resistant fibers with small diameter are required. Additionally, reinforcing fibers must be capable of retaining the structure, stiffness and strength under processing (matrix deposition) and service conditions. Keeping these in view, the preparation of a new generation of BN fibers-reinforced BN composites appears to be an excellent opportunity to replace the traditional carbon/carbon (C/C) composites which are very sensitive to oxidative and hydrolytic environment above 400°C. In addition, BN/BN composites could be used in radiation-transparent structures (low dielectric constant of h-BN) as well as in aerospace applications requiring, among others, ultra-light weight in accordance with its low density (2.27) compared with that of SiC or oxide-based materials, for example.2 Above all, the layered structure of h-BN as

Advances in Ceramic Matrix Composites XI 3

matrix could protect the fibers against the notch effect arising from matrix microcracking, deflect the microcracks parallel to the fiber surface and allow fiber sliding. In that case, the deposition of BN interphases to weaken the fiber-matrix bonding will be not necessary. The first objective for preparing such composite materials concerns the development of BN fibers with controllable mechanical properties. Preceramic polymers are ideally suited for the preparation of ceramic fibers using spinning and pyrolysis procedures.3*4 In our lab, we have prepared BN fibers from B-alkylaminoborazine-derived polymers (=poly[(alkylamino)borazines]).5"8 In particular, the poly[(methylamino)borazine] is seen as a melt-spinnable polymer for providing BN fibers with high mechanical properties, fine diameters and a low density. For example, in a previous paper, we showed that this polymer could deliver BN fibers with a tensile strength of 2 GPa, a modulus of 440 GPa, a density of 1.85 and a well-ordered and oriented microstructure.8

Permanent advances in polyMAB-derived BN fibers are important in the aim of improving the performances of the CFCCs. With this aim in mind, our work has been mostly dedicated to the study of the precursor synthesis, and to the numerous reactions occurring during the polymer-to-ceramic conversion. These investigations allowed to introduce significant improvements in the preparation of BN fibers. By varying the experimental conditions of melt-spinning, the present paper also shows that different types of fibers can be developed. In this context, the present study aims at the effect of the as-obtained microtextures/microstructures on the mechanical properties of BN fibers.

EXPERIMENTAL SECTION General comments

All synthesis reactions were carried out in Ar, whereas N2 and NH3 with electronic purity were used during the fiber preparation. Tensile tests and diameter measurements were achieved from 50 filaments with a gauge length of 10 mm. Diameters were measured by laser interferometry and mechanical properties were determined using a standard tensile tester (Adamel Lhomargy DY 22). Modulus and strains were averaged from the 50 tests and the strength distribution was described by Weibull statistics.9 Strength were averaged for a failure probability P=0.632. XRD was performed using a Philips apparatus (CuKa radiation; X = 1.5406 À at 40 kV and 30 m A). Fibers were crushed, prior characterization. SEM (Hitachi S800) was used to observe the cross-sectional microtexture of fibers. An Au/Pd film was deposed on fibers, prior observation. TEM was investigated using a Topcon EMB-002B microscope. Samples were embedded in a resin and cut into thin foils with an ultramicrotome. Foils were then set on microgrids to observe the longitudinal microtexture. Fiber preparation

A same lot of polyMAB was used in the present paper. Its synthesis and characterization was previously described. Green fibers were prepared in a glove-box in N2 by the melt-extrusion of the polymer followed by the stretching of the resulting mono- or multifilament by a spool. The as-spun fiber wound on the spool was transferred into a silica furnace to achieve the curing and pyrolysis processes in NH3 (25°C-1000°C, 0.8°C.min~l, dwell time of lh). After such heating, the fiber was transferred into a graphite furnace to undergo heat-treatment (lOX.min1) in N2 up to 1800°C (dwell time of lh). As-pyrolyzed BN fibers were white colored and their typical elemental composition (N2.8B3) showed that pure boron nitride was produced. The level in oxygen was extremely low (<2wt%).

4 Advances in Ceramic Matrix Composites XI

RESULTS AND DISCUSSION Melt-spinning of polyMAB-based fibers

Figure 1 details the melt-spinning as well as the process variables which are controlled to provide the different microtextures in the final BN fibers. Typically, the polyMAB is molten into an extruder at 180°C, then compacted, before pushing it through a filter for removing unexpected impurity. After passing through the filter, the melt is driven through a spinneret with one or seven capillaries, 200 |im in diameter for each. As the melt exits, it cools and emerges as an endless mono-Cone capillary) or multi- (seven capillaries) filament which is finally continuously wound on a spool. It should be noted that the spinning temperature is closely related to the polymerization degree of the polymer.

Piston, piston rate (m.min1)

Filter

Spinneret

Capillary, flow rate (m.min"1)

Spool, winding speed (m.min1)

Fig. 1. Melt-spinning process and process variables for the preparation of polyMAB fibers.

At typical extrusion conditions, the induced stretching inherently involves a significant reduction in the fiber diameter with distance from the spinneret. The rapid drawdown has been confirmed through measurements of the change in filament diameter during the stretching as shown in Figure 2.

10 15 20 25 30

Distance from (lie spinnerette [mm]

-1

■

----

—m— Flow rate = 2.05 m.

, —•—Flow rate = 2.77 m.

V —A— Flow rale = 3.70 m.

\*V.

nin '/Winding speed = 330 m.min '



Fig. 2. Variation of monofilament diameter with axial distance from the spinneret.


Fig. 2 shows that the filament solidifies above a distance of 15 millimeters from the spinneret. It is also seen (Table I) that, at a given flow rate, higher winding speeds result in smaller filaments. Again, Table I shows that, if the winding speed is held constant, the decreased flow rate results in an increased rate of drawdown, and therefore in a decreased monofilament diameter.

Table I. Effects of winding speed at a given flow rate on monofilament diameter Flow rate (m.min~ )

2.30 2.30 2.30 1.85 1.40

Winding speed (m.min" ) 132 157 170 132 132

Drawdown ratio* 7.5 8.4 9.3 9.1 10.0

Filament diameter (pm) 26.5 23.8 22.0 22.5 20.0

^Drawdown ratio is the diameter of the capillary divided by the final filament diameter.

The empirical results described above for monofilaments are also valid for multifilaments. Seven capillaries were used to form green multifilaments. However, it should be mentioned that the internal pressure, which can not exceed 45 daN, is closely related to the piston rate, the melt-viscosity of the polymer and the number of capillary: the highest the number of capillary, the lowest is the internal pressure. As the number of capillary is increased from one to seven, the internal pressure into the extruder is decreased, so that a higher latitude in the choice of the flow rate of the polymer can be found for spinning multifilaments. As an illustration, Table II gives the flow rate and winding speed values as well as the diameters of green fibers. Fibers lg and 2g are multifilaments, whereas fibers 3g are monofilaments.

Table II. Melt-spinning conditions used for the preparation of green fibers Preceramic green fibers

Flow rate (m/min) Winding speed (m/min)

Drawdown ratio Diameter (urn)

lg 0.66 97 12.1 16.5

2g 0.66 69

10.3 19.4

38 1.8 300 13

15.4

Poly mer-to-ceramic conversion The series of green fibers were cured and subsequently pyrolyzed to produce BN fibers

according to an identical procedure (see experimental section). During heat-treatment, fibers were maintained along the fiber-axis on the spool to prevent their crimping due to the important shrinking effects (mass loss and density increase) which occur during the polymer-to-ceramic conversion. The importance of the tension was previously reported to decrease significantly the fiber diameter, and therefore to improve the tensile strength by producing straighter fibers.6 Green fibers were cured in NH3 up to 400°C to render them infusible, then heated in NH3 from 400 to 1000°C to remove the majority of carbon-based groups bearing by the polymer. An additional heat-treatment was carried out in No from 25 to 1800°C to complete the ceramic conversion and obtain small-diameter purely BN fibers. As-obtained fibers display two distinct mechanical behaviors as illustrated in Figure 3.


Fig. 3. Difference in the flexibility of BN fibers.

Fibers lc are easily handled due to a high flexibility and are ideally suited for CFCCs, since they should be easily weavable to produce net-shape fiber preforms that would be infiltrated by a molecular precursor to generate the matrix. In contrast, poor flexible fibers 2c and 3c must be handled with great care, since they easily break under flexure or torsion.

Characterization of as-prepared BN fibers The differences in fiber handle reflect different mechanical properties (Table III).

Table III. Average diameters and mechanical properties of as-obtained BN fibers. BN Fibers Diameter (urn) Strength a (GPa) Modulus E (GPa) Strain e (%)

lc 2c 3c

7.8 9.8 7.4

1.5 1.0 1.4

105 235 400

1.30 0.43 0.31

Fibers lc display low Young's modulus and high failure strain, whereas fibers 2-3c are characterized by low failure strains. As the same polymer was used and the same curing and pyrolysis conditions were applied, these results could be unambiguously correlated to the flow rates and winding speeds fixed during the melt-spinning operation. It is clear that the change in the spinning conditions of polyMAB provide three types of BN fibers as reported in Table I. Fig. 4 shows SEM cross-sections of BN fibers from lc (Fig. 4a) with a glassy-like microtexture to 3c (Fig. 4c) with a coarse-grained microtexture through a skin-core heterogeneity in 2c (Fig. 4b).

Fig. 4. SEM micrographs of the as-obtained BN fibers lc (a), 2c (b), and 3c (c).

The surface morphology of the latter has been also investigated (Fig. 4b). The glassy-like microtexture of lc, similar to that of amorphous multi-phase SiBCN fibers, reflects a low grain size


with no orientation effects10, whereas the grainy microtexture of 3c suggests the polycristalline nature of the materials as seen in AI2O3 (Nextel® type) ceramic fibers.2 Therefore, as a similar fiber strength is obtained in lc and 3c, it is established that the Young's modulus of BN fibers is directly related to the granular nature of the microtexture: the poorest microtexture definition, the lowest is the Young's modulus. Besides, the BN fibers 2c are composed of grains extending in a radial direction over the skin region, whereas coarse grains are randomly-oriented in the core region. It is even shown that the propagation of cracks easily occurs on the surface of fibers, since mostly, cracks are seen to run along the edge of radially-oriented grains and propagated through the whole length of the fibers (Fig. 4b). Therefore, it is not surprising to find lower strength in such fibers.

SEM observations are highly reflected in XRD plots of samples presented in Fig. 5. Phase identification is achieved by locating the diffraction peaks of h-BN at 26.76° (002), 41.60° (100), 43.87° (101), 50.15° (102), 55.16° (004), 75.93° (110), 82.18° (112) and 85.52° (105).n

Fig. 5. XRD patterns of ground-up BN fibers from lc to 2c.

XRD analysis of as-pyrolyzed BN fibers shows some of the characteristic peaks of h-BN such as the (002), (004), (10) and (11) reflections. However, spectrum of lc shows broad peaks corresponding to short nanosized BN crystals (Lc = 31 Â and La = 86 Â) as well as (001) peaks shifted to lower diffraction angles of h-BN reflecting a large interlayer spacing. In particular, the doo2 value (3.45 Â in lc), which can be compared with 3.33 Â for h-BN crystals, reflects a disordered structure of the material. BN crystallinity and ordering are more pronounced in samples 2-3c. Comparing the corresponding grain size and interlayer spacing doo2 values of the respective powders, it is seen that crystal growth (U = 117 and La = 257 Â in 3c) and crystallinity (doo2 = 3.35 Â in 3c) are improved in fibers 2-3c. Moreover, the presence in the spectra of 2-3c of some major h-BN peaks, i.e., (101), (112) and (105), which did not appear in samples lc, reflects such characteristics.

TEM investigations are in good agreement with previous data, since the study confirms that the grain size and the microstructural ordering improve with increasing the Young's modulus from lc to 3c (Fig. 6). In addition, the HRTEM image of the low-modulus fibers lc (Fig. 7a) shows a mixture of poor crystallized regions and nanocrystalline phases in which the nanosized grains are buckled. Their poorly ordered stacking sequence is similarly to those seen in r-BN.12 We can even suggest that cross-linkages between neighboring small crystallites constituted an amorphous BN phase, and this was probably thought to be a factor promoting high strength.


Fig. 6. TEM micrographs of the as-obtained BN fibers lc (a), and 3c (b).

In contrast, the HRTEM image of the high-modulus fibers 3c shows local hexagonal regions embedded in a /-BN matrix (Fig. 7b). Moreover, HRTEM studies show the improvement of the (002) layer orientation along the fiber-axis from lc to 3c. As B and N atoms are covalently bonded in these layers with an in-plane average B-N bond energy of -500.75 kJ/mol13, it is also established that the modulus of fibers is correlated to the orientation level of the basal layers. In return, such structures are seen to involve faster crack propagations along the (002) layers leading to lower failure strains.

Fig. 7. HRTEM micrographs of the as-obtained BN fibers lc (a), and 3c (b).

Therefore, it could be concluded that high modulus was relatively straightforward to achieve, but usually at the expense of low strain-to-failure and brittleness. Nevertheless, it should be mentioned that the level of crystallinity is generally linked to the chemical and thermal stability: the highest level of crystallinity, the best is the stability. In this context, the use of lc-reinforced CFCCs in corrosive environment should be limited by the poor chemical/thermal stability (low crystallinity level) of fibers. Further works are in progress to study the thermal/chemical stability of BN fibers.

CONCLUSIONS The present study showed that a large variety of BN fibers could be prepared by varying the

melt-spinning conditions of a borazine-based polymer. Different types of microtexture/microstructure affecting the mechanical properties were developed in BN fibers. It was shown that BN multifilaments consisted in either featureless cross-sectional microtextures with a poor grain orientation along the fiber-axis as glassy-like materials or microtextures in which a cross-sectional skin (radial orientation)-core (random orientation) heterogeneity was identified. In contrast,


the melt-spinning operation into monofilament produced fibers with a coarse-grained microtexture in which grains were nearly oriented along the fiber-axis. It was shown that modulus was closely related to the nature of the microtexture/microstructure: (1) the BN multifilaments with a low modulus consisted in a mixture of disordered nanocrystals and amorphous regions, whereas (2) the spinning into monofilaments provided high-modulus fibers with large grains nearly oriented along the fiber-axis in an improved microstrutural ordering. The development of such flat layer microtexture increased the crack propagation along the cleavage basal layers causing a general decrease in the strain. The strengths were similar in both cases. Besides, the development of a skin -core textural heterogeneity in some of BN multifilaments decreased the fiber strength and flexibility.

REFERENCES *R. Haubner, M. Wilhelm, R. Weissenbacher and B. Lux, "Boron Nitrides - Properties,

Synthesis and Applications"; pp. 1-45 in High Performance Non-oxide Ceramic II, Vol. 102, Structure Bonding. Edited by M. Jansen, Springer-Verlag Berlin Heidelberg, 2002.

2A. R. Bunsell and M.-H. Berger, "Fine Diameter Ceramic Fibers", J. Eur. Ceram. Soc, 20, 2249-2260 (2000).

3M. Peuckert, T. Vaahs, and M. Brück, "Ceramics from Organometallic Polymers," Adv. Mater., 2, 398-404 (1990).

4J. Bill, and F. Aldinger, "Precursor-Derived Covalent Ceramics," Adv. Mater., 7, 775-87 (1995).

5B. Toury, S. Bernard, D. Cornu, F. Chassagneux, J.-M. Létoffé, and P. Miele, "High-Performances Boron Nitride Fibers Obtained from Asymmetric Alkylaminoborazine," J. Mater. Chem., 13, 274-279 (2003).

6S. Bernard, K. Ayadi, J.-M. Létoffé, F. Chassagneux, M.-P. Berthet, D. Cornu, and P. Miele, "Evolution of Structural Features and Mechanical Properties During the Conversion of Poly[(methylamino)borazine] Fibers into Boron Nitride Fibers," J. Sol. State. Chem., Ill, 1803-10 (2004).

7S. Bernard, D. Cornu, P. Miele, H. Vincent, and J. Bouix, "Pyrolysis of Poly[2,4,6-tri(methylamino)borazine] and its Conversion into BN Fibres," J. Organomet. Chem., 657, 91-97 (2002).

8S. Bernard, F. Chassagneux, M. P. Berthet, H. Vincent, and J. Bouix, "Structural and Mechanical Properties of a High-Performance BN Fibre", J. Eur. Ceram. Soc, 22, 2047-2059 (2002).

9K. Goda, and H. Fukunaga, "The Evaluation of the Strength Distribution of Silicon Carbide and Alumina Fibers by a Multi-Modal Weibull Distribution," J. Mater. ScL, 21,4475-80 (1986).

10S. Bernard, M. Weinmann, P. Gerstel, P. Miele, and F. Aldinger, "Boron-modified Polysilazane as a Novel Single-Source Precursor for SiBCN Ceramic Fibers: Synthesis, Melt-spinning, Curing and Ceramic Conversion", J. Mater. Chem., 15, 289-299 (2005).

T,R. S. Pease, "An X-Ray Study of Boron Nitride", Ada. Cryst., 5, 356-361 (1952). 12J. Thomas Jr, N. E. Weston, and T. E. O'Connor, "Turbostratic Boron Nitride, Thermal

Transformation to Ordered-Layer-Lattice Boron Nitride", J. Am. Chem. Soc, 84, 4619-4622 (1963). I3M. Côté, P. D. Haynes and C. Molteni, "Boron Nitride Polymers: Building Blocks for

Organic Electronic Devices", Phys. Rev. B, 63, 125207-1-4 (2001).


Processing

To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the property of The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, without the express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.

MULTILAYERED MATERIALS BY ICVI IN NON-OXIDE SELF-HEALING CERAMIC MATRIX COMPOSITES FOR HIGH TEMPERATURE APPLICATIONS

L. Vandenbulcke LCSR - CNRS, 1C Avenue de la Recherche Scientifique, 45071 Orléans cedex 2, France

ABSTRACT To manufacture outstanding multilayered non-oxide CMCs, an accurate CVI processing is

necessary for controlling the uniformity, the microstructure of each material and the interface bonding between both rigid carbide multilayers including SiC, BXC, Si-B-C materials and relatively flexible interphases. The incorporation of several interphases in the matrix is also important when a thermal expansion mismatch between different components must be accommodated. All the above materials can be incorporated in a multilayered matrix to minimize the stresses in all components and the microcracking of the matrix, also to obtain efficient crack deviation as far as possible from the fibers when microcracks occur. The best results can be expected if the fiber/interphase bonding is increased, for example with Si-B-N, if the thickness of the interphases is small and their structure well-controlled, as with BN. Under sufficient loading, microcracks allow penetration of oxygen in the composite. Minimization of the internal oxidation however can be obtained by an appropriate design of the matrix so as: - to deviate the cracks far from the fibers at interfaces between/inside rigid layers or interphases - to fill these cracks with glasses in a large temperature range. The self-sealant properties of the boron and silicon compounds are very efficient because of the formation of borosilicate glass at low and intermediate temperatures and silica-rich oxide at high temperature. After Cf/SiC and SiCf/SiC composites being manufactured at the end of the seventies, the processing of the self-sealant SiBC-based materials and the recent tailoring of the whole composite have permit a great improvement in the preservation of the CMCs properties under high temperature fatigue, thermal and mechanical cycling for long duration, in oxidative environments. These relatively "smart CMCs" allow therefore a breakthrough to applicability in the energy, space and aeronautic domains.

INTRODUCTION Non-oxide ceramic matrix composites (CMCs), mainly those which are composed of a

carbide matrix reinforced with carbon fibers, referred as Cf/SiC, or silicon carbide fibers, SiCf/SiC composites, have been studied for many years since the pioneering works of European groups 1 2 . Their non-brittle properties result from the appropriate design of the fiber/matrix interface. Classically a compliant interphase with a fairly low shear stress is formed during processing or intentionally deposited between the fibers and the matrix, this third component of the composite preventing the failure of the reinforcement fibers by deflecting any microcrack of the matrix induced by loading at a high enough level. This damaging mode induces the non-brittle character of these ceramics, but it favors the in-depth penetration of oxygen. The lack of preservation of the thermomechanical properties of these composites in oxidative environments, under high temperature fatigue conditions and thermal cycling has been a serious drawback for more than two decades for their use in many important applications in the fields of energy,


aeronautics and space. The main goal of lot of research was to improve the oxidation resistance of the various components of the CMCs and to reduce the oxidation penetration in the parts.

Here we are examining the various possibilities that we have studied for more than fifteen years to preserve the non-oxide CMCs properties by using silicon and boron containing compounds and more complex ternary or quaternary systems which include these two elements in monolayers or as multilayered materials. Some of these systems lead to limitations of particular deposition or infiltration processes that are pointed out. The chemical vapor deposition (CVD) and the isothermal chemical vapor infiltration (ICVI) of SiC, BXC, the ternary Si-B-C system and the corresponding multilayered materials are presented. All these materials are shown to permit a limitation of the oxidation penetration in the parts by a self-healing mechanism. BN has been studied as a more oxidation resistant interphase. The results obtained on the Si-B-N ternary materials are also presented and their resistance to oxidation is compared to those of BN. For each component of the composites, examples showing the influence on the final properties of the composite materials are reported. The great improvement in the preservation of the multilayered self-healing CMCs properties is proved from various studies carried out in oxidative environments, under high temperature fatigue, thermal and mechanical cycling. The optimization of the processing of the self-sealant SiBC-based materials is discussed.

EXPERIMENTAL TECHNIQUES The experimental devices for depositing thin films by CVD and ICVI have been described

in details previously 3' 4. Both are hot wall reactors with a graphite susceptor heated by a high frequency generator. However they differ by their size and configuration. In the CVD reactor 3, the gas mixtures were introduced from the bottom to associate the natural and forced convection. It allowed depositing SiC, BXC and Si-B-C films, alone, or as multilayered films on 3D C/C, 2D and 3D C/SiC and 2D SiC/SiC composites. The size of the samples was small (20 x 10 x 2-3 mm). They were used for oxidation tests. The ICVI reactor has been especially designed to avoid homogeneous nucleation when depositing BN and Si-B-N. The different reactants or mixtures of reactants were introduced from the top, separately through small holes to produce a pressure drop which limited retrodiffusion in a heating zone where the temperature was maintained at a value that permitted to avoid the NH4CI and adducts formations. The size of the isothermal deposition zone was much greater than in the CVD reactor (120 mm in diameter and 200 mm in length). It is obvious that this reactor could be used for the ICVI of the carbide films.

The chemical compositions of the films were determined by electronic microprobe analysis and nuclear methods as reported previously 5. The microstructure of the films was studied by scanning and transmission electron microscopy (SEM and TEM).

The procedure of the cyclic oxidation tests has been described earlier 6. The tests were carried out at temperature in the 723 - 1773 K range in flowing air, with six cycles of one hour followed by cycles of six hours, the total duration depending on the weight loss. The samples were put in or removed from the furnace in about 30 s, corresponding to an important thermal shock that produces many microcracks, even delamination of 2D composites in some cases.

Some composites were also tested under creep by using the four point bending technique at different stresses ranging from 150 to 250 MPa, in the 873 - 1123 K temperature range, for a maximum period of about 300 hours in air. Other mechanical results were obtained more recently in various conditions by several authors on different self-healing matrix composites 7"12.


RESULTS AND DISCUSSION CVD and ICVI of SiC, BXC and Si-B-C ternary system

Many studies have reported the CVD and the ICVI of SiC from methyltrichlorosilane (MTS) and hydrogen mixtures. Our deposition conditions were selected from the thermodynamic approach and the experimental work of Christin et al. '. These conditions lead to a badly crystallized mixture of cubic and hexagonal pure SiC with fine grains.

Concerning the boron carbide deposition, previous theoretical and experimental studies have shown that, under conditions of high supersaturation in reactants and relatively moderate temperature (1273-1373 K), boron carbide deposits of various compositions can be obtained ,3. The deposits are badly crystallized or amorphous and nearly all compositions from B50C2 to B4C can be deposited from BCI3-CH4-H2 mixtures. It has been shown that less stable hydrocarbons allow depositing C(B) materials with boron concentrations in the 5-20 at % range ,4. The co-deposition of the Si-B-C ternary system needed a thorough specific study. A thermodynamic approach was first undertaken to predict the composition of the Si-B-C-H-Cl system at equilibrium and to determine the most interesting domain of the experimental deposition conditions to be studied 3, especially the inlet composition in MTS, BCI3 and H2. The conditions controlled principally by mass transfer were used to verify the thermodynamic approach. Fig. 1 shows that non-uniform coatings were usually obtained in these conditions, as shown by the variations of their thickness and composition as a function of the distance h along a small sample.

Fig. 1: Thickness and chemical compos-ition (at %) uniformities in silicon I 1, boron 1=3and carbon I11111IL as a function of the inlet ratio [MTS]/[BC13]; (T « 1400 K, P = 40 kPa, [H2]/[MTS] = 20) ; (from Goujard and Vandenbulcke , 6) .

Then experimental conditions that allowed a kinetic control were searched for depositing Si-B-C coatings with both good thickness and composition uniformities ,5"17. This was obtained easily for boron-rich materials. However in such ternary system, the mass transfer-kinetic control of the process may vary as a function of the inlet composition. Therefore different sets of process conditions should be used to deposit uniform coatings of all compositions. Finally uniform layers with a large composition range could be deposited, for example with various B-contents ,7.


CVD and ICVI of BN and Si-B-N ternary system, microstructure The procedures for depositing BN from BCI3-NH3-H2 mixture and Si-B-N layers by

adding MTS to these reactants have been described in detail previously 4. This study was characterized by the simultaneous processing of BN on bulk substrates (CVD) and fibrous preforms (ICVI) that allowed gaining more insight into the deposition mechanism which influences the growth conditions and the deposit structure , 8 , 1 9. The influence on the size of the coherent domains, on their orientation relatively to the direction of growth and the interlayer distance in BN was evidenced ,8, ,9. A fairly good hexagonal BN microstructure with an inter-reticular distance of 0.35 nm, close to the theoretical value of 0.33 nm, was especially obtained by ICVI. A complete rearrangement of the microstructure by annealing has also been shown by Lacrambe 20 and confirmed, after the same growth conditions than ours, by Le Gallet et al. 2 . This control of the process and the possibility offered by annealing allow one to select the microstructure and then the interfacial and bulk properties of this material, density, coefficient of thermal expansion (CTE), Young's modulus... and its ability to accommodate the CTE mismatch of other components of the composites, especially when carbon fibers are used for the reinforcement. The CVD and ICVI of the ternary Si-B-N system were also thoroughly studied 22 as for Si-B-C, starting with a thermodynamic approach. Figure 2 presents the results of the thermodynamic calculations in terms of deposition domains of the different solid phases as a function of the inlet concentrations in MTS, BCI3 and NH3 (a residual concentration of oxygen was also considered in these calculations). The different inlet compositions studied are mentioned on this diagram.

[C"H3~STCT3]

Fig. 2: Deposition diagram of the various solid phases calculated at equilibrium from MTS-BCI3-NH3-(H2-O2) mixtures and inlet compos-itions which were used experim-entally; (T = 973 K, P = 1.33 kPa, [MTS+BC13+NH3]= 1 mole, H2 = 5 moles, O2 = 0.01 mole).

oc i 3 '

Different microstructures, different amount of carbon and oxygen incorporated in the deposits and different stabilities in wet air were obtained as a function of the deposition conditions 22. It can be reported here that a hexagonal microstructure is maintained when the silicon content is low and the deposits are stable in air when the S1-S2 gas phases are used. For inlet gas composition of type T, stable deposits can be obtained with a silicon concentration of 15-25 at % and low oxygen content, but they are amorphous 22.

C+BN+Si3N4 +Si2N20+SiC

C+BN+Si3N4+Si2N20

^ ^ J \ B N + S I 0 2 + S i 2 N 2 0

r-V\ >• V\. =4

C+BN+Si02

C+BN+Si6,+Si,N,0 C+BN+Si2N20


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