MICROTEXTURAL AND MICRO STRUCTURAL EVOLUTION INPOLY[(ALKYLAMINO)BORAZINE]-DERIVED FIBERS DURING THEIR CONVERSIONINTO BORON NITRIDE FIBERS

Samuel Bernard, Fernand Chassagneux, David Cornu, and Philippe MieleLaboratoire des Multimateriaux et Interfaces (UMR CNRS 5615)University Claude Bernard - Lyon 143 Bd du 11 novembre 1918Villeurbanne, France, 69622

ABSTRACTBoron nitride (BN) fibers were efficiently prepared from a B-aminoborazine-based polymer

according to the Polymer-Derived Ceramic (PDC) route via melt-spinning and heat-treatments up to1800°C in a controlled atmosphere. The microtextural and microstructural changes in the materialduring the polymer-to-ceramic conversion were investigated by means of electron microscopy andXRD observations. The microtexture of the fibers was featureless as glassy-like materials whenfibers were exposed at temperatures below 1400°C. Mechanical properties of such amorphous fiberswere poor at these temperatures. Upon further heating to 1500°C, the microstructure changed fromdisordered nanocrystals embedded into an amorphous phase to a turbostratic phase. This amorphous-to-crystalline transition was accompanied with a large increase in the mechanical properties. At1800°C, the microtexture of the fibers was coarse-grained and was correlated to the identification ofa "meso-hexagonal" BN phase with basal layers almost aligned along the fiber-axis in the material.Poly crystalline BN fibers exhibited high mechanical properties (a = 1.4 GPa, E = 360 GPa) aftercuring of the polymer fibers at 400°C and subsequent pyrolysis of cured fibers at 1800°C.

INTRODUCTIONInorganic/organometallic preceramic polymers play a major role in the preparation of shape-

controlled non-oxide ceramics due to their adjustable processability.1"2 One significant advantage ofthe Polymer-Derived Ceramic (PDC) route is that flexible and small-diameter ceramic fibers withdesired properties can be prepared via spinning of tractable polymers, curing of the as-spun fibersand pyrolysis of the resulting cured fibers according to the processing scheme presented in Fig. 1.These small-diameter and flexible fibers are generally ideally suited for Continuous Fibers-reinforced Ceramic matrix Composites (CFCCs) since they are easily weavable to produce net-shapefiber preforms that are subsequently infiltrated by a matrix.3

Synthesis+PurificationT Spinningk.

Curing (Crosslinking)L

Pyrolysik.

Fig. 1. Processing scheme for the preparation of ceramic fibers from preceramic polymers

The preparation of amorphous silicon carbide fibers from polycarbosilane clearly illustrated thisi j 4method. This process is also related to that used for the preparation of carbon fibers from

polyacrilonitrile.5 These examples highlight the utility of the polymer precursor approach to fiber

To the extent authorized under the laws of the United States of America, all copyright interests in this publication are the propertyof The American Ceramic Society. Any duplication, reproduction, or republication of this publication or any part thereof, withoutthe express written consent of The American Ceramic Society or fee paid to the Copyright Clearance Center, is prohibited.

43

research. Unfortunately, such fibers are either mechanically and chemically unstable above 1000°Cin both oxidative and non-oxidative environment (SiC fibers) or poorly resistant to oxidation in airabove 400°C (carbon fibers). For CFCCs fabrication, high-modulus and strength oxidation resistantfibers with small diameter are ideal. Additionally, reinforcing fibers should be capable of retainingthe structure, stiffness and strength under processing and service conditions. Keeping these in view,we have investigated the preparation of polycristalline boron nitride (BN) fibers from B-aminoborazine-derived polymers (= poly[(alkylamino)borazine]).6"8 Indeed, BN fibers should be of amajor interest in aerospace application as a reinforcing phase in a new generation of composites(BN/BN composites). As examples, in accordance to the good resistance of h-BN against oxidationand a graphite-like structure9, these materials could favorably replace the traditional carbon/carboncomposites which are oxidized around 400°C. Furthermore, in contrast to a large majority of non-oxide ceramic fibers, BN fibers offer the possibility to be used in radiation-transparent structures dueto the low dielectric constant of /z-BN. As the requirements to obtain high mechanical properties arerather stringent, an understanding of the polymer-to-ceramic conversion is of critical importance. Inthe present paper, the microtexturaVmicrostructural changes occurring during the preparation of BNfibers were investigated by electronic microscopy and were correlated to XRD and tensile tests data.

EXPERIMENTAL SECTIONGeneral comments

All reactions leading to the preparation of the starting polymer were carried out in a purifiedargon atmosphere using standard Schlenk techniques as previously reported.8 Nitrogen and ammoniawith 99.999% purity were used during the fiber preparation.

Tensile tests and diameter were obtained from single filaments with a gauge length of 10 mm.50 single filaments were analyzed for each test. The diameter § of each filament was measured bylaser interferometry and an average diameter was deducted. Single filament tensile properties weredetermined using a standard tensile tester (Adamel Lhomargy DY 22). Young's modulus and failurestrain were averaged from the 50 tests taking into account the system compliance. The strengthdistribution was described by Weibull statistics10 and the average room temperature tensile strengtha was estimated for a failure probability P=0.632.

X-ray diffraction (XRD) measurements were performed using a Philips diffractometer (CuKaradiation; X = 1.5406 A at 40 kV and 30 mA). Fibers were crushed prior characterization.

Scanning electron microscopy (SEM) (Hitachi S800) was used to investigate the cross-sectional microtexture of fibers. An Au/Pd film was deposed on fibers prior observation.

Transmission electron microscopy (TEM) was investigated using a Topcon 002Bmicroscope. Samples were embedded in a resin and cut into thin foils with an ultramicrotome. Foilswere then set on microgrids to observe the microtexture in the longitudinal sectional thin specimens.

Fiber preparationThe synthesis and characterization of the poly[(2,4,6-trimethylamino)borazine] (polyMAB)

were previously described.8Endless as-spun fibers, 15 um in diameter, were prepared by the extrusion of the polyMAB

around 180°C followed by the drawing of the emerging monofilament using a lab-scale spinning

44

apparatus set up in a glove-box filled with nitrogen. The as-spun fiber was transferred into a furnaceequipped with a silica glass tube directly connected to the glove-box in order to minimizecontamination by oxygen (air) and moisture. Curing and pyrolysis were achieved at atmosphericpressure under the tension imposed by the important shrinking effects which occur in the fiberswound on the drum. The importance of the tension was previously reported to improve tensilestrength and Young's modulus by producing straighter and stiffer fibers.6 The as-spun fiber wascured in an ammonia atmosphere (25°C-400°C, 0.8°C/min), then the as-cured fiber was treated in anammonia atmosphere to 1000°C (0.8°C/min) with a dwell time of Ih to remove the majority ofcarbon-based groups bearing by the polymer. After such treatments, the fiber was less sensitive tooxygen and moisture and could be handled and quickly transferred into a second furnace in the openair. An additional heat-treatment (10°C/min) was performed in a graphite furnace in a nitrogenatmosphere up to 1800°C for hold times of Ih. As-pyrolyzed BN fibers, 7.5 jim in diameter, werewhite colored and of flexible form. Their typical elemental composition was previously reported.8

RESULTS AND DISCUSSIONXRD investigations

Fig. 1. shows X-ray diffraction results taken on crushed fibers during the preparation of BNfibers through curing at 400°C (as-cured) then pyrolysis up to 1800°C (as-pyrolyzed).

Fig. 1. X-ray diffraction patterns of polymer (polyMAB) fibers during their conversion into ceramic(BN) fibers up to 1800°C. All reflections can be assigned to /z-BN.

As-cured fibers exhibited every broad X-ray reflections of /z-BN with low relative peak intensity.XRD patterns did not change in appearance on further heating in an ammonia atmosphere at 1000°Cand the amorphous nature of the BN phase was even preserved after pyrolysis to 1400°C. Onlyabove 1400°C was crystallization detectable by XRD. An increase in the intensities of the (001)reflections, which means an ordering of the structure in basal planes, accompanied by a sharpeningof these reflections were detected at 1500°C. On further heating to 1800°C, fibers exhibited themajority of /z-BN reflections. In particular, the diffraction from the (100) and (101) peaks wasobserved, which should indicate the formation of a three-dimensional crystal structure. However, the

45

crystallization was not fully completed in as-pyrolyzed BN fibers as indicated by the lack of (hOl)type of reflection in the corresponding XRD pattern (1800°C). The results indicated a meso-hexagonal structure with an incomplete three-dimensional ordering.

In support of the amorphous-to-crystalline transition was the observations that the graingrowth started from 1400°C to form a stabilized crystallization at 1800°C (Fig. 2).

L (nm)La(nm)

1000 1200 1400Temperature/°C

Fig. 2. Dependence of the average crystallite size of fibers during their conversion up to 1800°C.

The crystallites sizes at different temperatures were estimated on the basis of the FWHM of the (002)(29 = 26.76°) and (10) (29 = 41.60°) peaks using the Scherrer equation (Eq. I)14:

B = KV(Lcos9) (Eq. 1)

,where K = 0.9 (Lc) and 1.84 (La), X, the wavelength of the X-rays ((X= 0.15418 nm), Lrepresents the average stack height (Lc) and the average length (La) of the crystallites, 9 is thediffraction angle of the corresponding reflections and B is the FWHM of the peaks.For example, the average value of the particle size (La) increased from <6 nm at 1400°C to ~16 nmat 1500°C and reached -25 nm after pyrolysis at 1800°C.

SEM observationsFig. 3 shows the microtextural changes during the preparation of BN fibers. SEM

observations are entirely consistent with XRD data. The cross-section of the as-spun fibers iscommon to that of organic polymer fibers. When fibers were cured at 400°C, then pyrolyzed at1000°C and subsequently at 1400°C, their microtexture remained featureless reflecting theamorphous nature of the material below 1400°C. The amorphous-to-crystalline transition occurredthrough a featureless-to-granular microtexture transformation as fibers were heated from 1400 to1500°C. The granular microtexture reflected the polycristalline nature of the BN fibers after suchheat-treatments. Upon further heating at 1800°C, the coarsening of the grains through the cross-section of the as-pyrolyzed BN fibers was more pronounced in relation with the improvement of thestructural ordering of the BN phase.

46

(d) (e) (f)Fig. 3. SEM micrographs of the polymer fibers (a) and fibers prepared at 400°C (as-cured fibers) (b),1000°C (c), 1400°C (d), 1500°C (e) and 1800°C (as-pyrolyzed fibers) (f).

TEM observationsIn a previous work, TEM studies of BN fibers prepared at 1800°C indicated that the large

crystallites were stacked in a mesohexagonal BN ordering and nearly oriented along the fiber-axis.11

In order to complete TEM investigations, microstructures of intermediate fibers heated at 1000°C,1400°C and 1500°C were observed in the longitudinal section.Consistently with XRD results, the SAED of the fibers heated at 1000°C showed an amorphous haloimposed on the diffuse (002) arcs and poorly-resolved (10)/(100), (004) and (11)7(110) arcs whichindicate the presence of a poorly-ordered structure (Fig. 4).

Fig. 4. BF TEM image of the longitudinal thin sectional fibers prepared at 1000°C.

47

In addition, the BF TEM image of such fibers was featureless and of low contrast and, in particular,it showed that the microstructure of the fibers prepared at 1000°C was damaged during the specimenpreparation when the Leitz diamond knife cut the fibers into thin foils (Fig. 4). As an explanation, theknife moving along the fiber-axis damaged the poor microstructural cohesion of the fibers byseparating the matter into elongated fibrils. This phenomena, which was also observed in fibersprepared at 1400°C, results from the lack of regular crystalline BN planes in the fibers heated below1400°C as illustrated in Fig. 5.

Fig. 5. BF TEM image of the longitudinal thin sectional fibers prepared at 1400°C.

In accordance with BF images, the HRTEM image of fibers prepared at 1000°C presented in Fig. 6exhibited a completely disordered microstructure with nanograins of varying size (less than 5 basalplane in thickness, doo2 values around 0.350 nm) and orientation arising from an amorphous phase.

Fig. 6. HRTEM image of the longitudinal thin sectional fibers prepared at 1000°C.

It should be mentioned that a similar nanostructure was also detected in fibers prepared at 1400°C(doo2 -0.347 nm). As a consequence, we can postulate that the poor microstructural ordering and thelow level of orientation which compose the fibers prepared below 1400°C result in poor mechanicalproperties (strength and modulus) as reported in Table 1. In contrast, in such poorly crystallizedfibers, the slow crack propagation, due to the lack of well-defined crystalline planes in the structure,result in a high fiber flexibility and therefore, in a high strain (Table 1).As observed in XRD and SEM analysis, the amorphous-to-crystalline transition was also reflected inTEM images when fibers were heated from 1400 to 1500°C. Figure 7 shows the BF TEM images ofthin sectional fibers prepared at 1500°C.

48

Table I. Changes in the mechanical properties during the preparation of BN fibers.As-pyrolyzed fibers (1800°C)

1.4Temperature/°C

Average strength/GPa360 Average modulus/GPa0.35 Average strain/%

400 (as-cured)0.2171.1

10000.3231.2

14000.4400.9

15000.71500.5

Fig. 7. BF TEM images of the longitudinal thin sectional fibers prepared at 1500°C.

Consistent with SEM observations, a granular microtexture is clearly seen in such fibers. No damageof fibers was observed which means that the microstructural cohesion of the BN network is strongerin accordance with the higher values of mechanical properties compared to those measured afterthermal treatment at 1000 and 1400°C (Table 1). In contrast, as the microtexture becomes moregrained, cracks proceed along the well-defined crystal planes, i.e., along the basal (002) sheets. Suchfibers are therefore accompanied by a higher brittleness and, therefore, a lower flexibility. SAED,which are composed of well-resolved (002), (10)/(100), (004) and (110) arcs, were alsorepresentative of a better ordered material. Furthermore, the HRTEM image presented in Fig. 8 well-reflected the BN phase evolution between the fibers prepared at 1000°C (Fig. 6), with a highproportion of 0-BN and those heated at 1500°C exhibiting better ordered and extended (002) layerswithdoo2 =

Fig. 8. HRTEM image of the longitudinal thin sectional fibers prepared at 1500°C.

49

The decrease in the doo2 values from 1400 (doo2 = 0.347 nm) to 1500°C (doo2 = 0.340 nm) which werehigher than the ideal value of /z-BN (doo2 = 0.333 nm) were assumed as an indicator in assigning theturbostratic structure to the fibers were prepared at 1500°C. On further heating to 1800°C, as-pyrolyzed fibers exhibited a well-ordered homogeneous nanostructure and basal layers nearlyoriented along the fiber-axis in good agreement with their high mechanical properties as shown in aprevious work.11 doo2 values decreased up to 3.35 nm indicating a meso-hexagonal structure.

CONCLUSIONThe present study reported the microtextural/microstructural changes which occurred during

the preparation of BN fibers from a B-aminoborazine-based polymer according to the Polymer-Derived Ceramic (PDC) route. In particular, it was shown that the microstructure of fibers preparedbelow 1400°C consisted of a mixture of disordered nanocrystals and amorphous regions, whereas thelarge crystallites tended to stack in a meso-hexagonal ordering and to nearly align along the fiber-axis in fibers prepared above 1400°C. The amorphous-to-crystalline transition occurred with theincrease in the mechanical properties.

REFERENCES1M. Peuckert, T. Vaahs, and M. Briick, "Ceramics from Organometallic Polymers," Adv.

Mater., 2, 398-404 (1990).2J. Bill, and F. Aldinger, "Precursor-Derived Covalent Ceramics," Adv. Mater., 7, 775-87

(1995).3J. R. Strife, J. J. Brennan, and K. M. Prewo, "Status of Continuous Fiber-Reinforced

Ceramic Matrix Composite Processing Technology," Ceram. Eng. Sci. Proc., 11, 871-919 (1990).4S. Yajima, J. Hayashi, M. Omari, and K. Okamura, "Development of a SiC Fibre with High

Tensile Strength," Nature, 261, 683-85 (1976).5E. Fitzer, and M. Heine, "Carbon Fiber Manufacture and Surface Treatment," In Fiber

Reinforcements for Composite Materials, A. R. Bunsell (eds), Elsevier, New York, pp. 79-148(1988).

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

7B. Toury, S. Bernard, D. Cornu, F. Chassagneux, J.-M. Letoffe, and P. Miele, "High-Performances Boron Nitride Fibers Obtained from Asymmetric Alkylaminoborazine," /. Mater.Chem., 13, 274-279 (2003).

8S. 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," /. Organomet. Chem., 657, 91-97(2002).

9R. S. Pease, "Crystal Structure of Boron Nitride," Nature, 5, 722-723 (1950).10K. Goda, and H. Fukunaga, "The Evaluation of the Strength Distribution of Silicon Carbide

and Alumina Fibers by a Multi-Modal Weibull Distribution," /. Mater. Sci., 21, 4475-80 (1986).nS. Bernard, F. Chassagneux, M. P. Berthet, H. Vincent, and J. Bouix, "Structural and

Mechanical Properties of a High-Performance BN Fibre," /. Eur. Ceram. Soc., 22, 2047-59 (2002).

50