Crystallinity, Crystalline Quality, and Microstructural Ordering inBoron Nitride Fibers

Samuel Bernard,*,w Fernand Chassagneux, Marie-Paule Berthet, David Cornu,* and Philippe Mielez

Laboratoire des Multimateriaux et Interfaces, UMR CNRS 5615, Universite Claude Bernard Lyon 1, F-69622Villeurbanne, France

A qualitative and quantitative analysis of the microtexture/microstructure in several experimental boron nitride fibers wasinvestigated by X-ray diffraction and electron microscopy ac-cording to their mechanical properties. Young’s modulus wascontrolled by grain orientation. The disordered microstructureof low modulus fibers consisted of amorphous regions and nano-meter-sized crystallites arranged in a wide range of orientation.In contrast, high modulus samples showed a coarse-grainedmicrotexture in which crystallites were ordered in a meso-hexagonal arrangement along the fiber-axis. Cracks were easilypropagated as crystallinity increased, leading to decreased fail-ure strain. The strength-microstructure relationship was ob-scured by the fiber porosity and defects.

I. Introduction

HEXAGONAL boron nitride (h-BN) is a ceramic material witha layer anisotropic structure, similarly to that of carbon

graphite. It offers some attractive properties such as high stiff-ness and toughness along the basal layers, a non-wettabilityagainst many metallic and silicate melts, a good oxidative re-sistance up to TB10001C, a good thermal stability up toTB25001C in a nitrogen atmosphere and a low coefficient ofthermal expansion (CTE) in the direction of basal layers.1–3

Based on these properties, h-BN has great potential for prepar-ing continuous fiber-reinforced ceramic–matrix composites(CFCCs) intended for high temperature applications. This newgeneration of BN/BN composites could favorably replace thetraditional carbon/carbon (C/C) composites, which are verysensitive to oxidative and hydrolytic environment above4001C. In addition, BN/BN composites could be used in radi-ation-transparent structures due to the low dielectric constant ofh-BN as well as in aerospace applications requiring light weightdue to its low density (2.27) compared with that of SiC- or ox-ide-based materials, for example.4 The development of mechan-ically reliable BN/BN composites requires the preparation ofBN fibers with optimal mechanical properties.

The polymer-derived ceramics (PDCs) route5–14 is particular-ly suited to the preparation of high-performance fibers as illus-trated with the worldwide works of Yajima et al.15 and Winteret al.16 dedicated to silicon-based fibers. Following this route, wehave investigated the preparation of BN fibers using poly[(al-kylamino)borazines] as starting preceramic polymers.17–21 Inparticular, the poly[(methylamino)borazine] (PolyMAB) is anappropriate candidate for providing continuous BN fibers with

high mechanical properties, fine diameters, and low density.20,21

In a previous paper, we showed that the high mechanical prop-erties (s5 2 GPa, E5 440 GPa) of polyMAB-derived BN fibersreflected a well-ordered and oriented BN microstructure.21

Advances in BN fibers are important in improving the per-formance of composite materials. With this aim in mind, signif-icant improvements in the preparation of polyMAB-derived BNfibers have been introduced. For instance, by varying the exper-imental conditions of melt spinning, different forms of micro-structure and microtexture affecting the mechanical properties(strength and flexibility) have been developed. Since the objec-tive fiber structure and texture providing such performances isnot clearly identified, the present paper aims at the characteri-zation of the crystallinity, crystalline quality, and microstruc-tural ordering of several BN fibers. In particular, the influence ofstructural parameters on mechanical properties is examined.

II. Experimental Procedure

Fibers were melt-spun in a nitrogen atmosphere using a Poly-MAB as polymeric precursor. The synthesis and spinning of thispolymer have been reported in the literature.22 Table I gives thespinning conditions (extrusion rate and wheel velocity duringthe drawing of the endless fiber emerging from the capillary) anddiameters of as-spun fibers. In particular, a die with seven cap-illaries was used to produce fibers A, B, and C whereas a singlecapillary was used to prepare fibers D, E, and F.

As-spun fibers were cured in an ammonia atmosphere up to4001C, then in ammonia (400–10001C) and nitrogen (1000–18001C) atmospheres.

Tensile tests were performed at room temperature on singlefilaments using an Adamel Lhomargy DY 22 apparatus (Ivry-sur-Seine, France) and diameters were measured by laser inter-ferometry as previously reported.24 Table II reports informationswhich includes average diameters and mechanical properties ofas-obtained BN fibers.

X-ray diffraction (XRD) measurements were performed usinga Philips diffractometer (Almelo, Netherlands) (CuKa radiation;l5 1.5406 A at 40 kV and 30 mA) with a PW 1830/40 generatorand a PC-APD (PW1877) software. XRD patterns were record-ed from as-obtained BN fibers in the 2y range 151–601, then inthe 2y range 151–901 from the same fibers, after grinding intofine powders to free from the grain orientation which occurs inas-obtained fibers. Intensities of (hkl) peaks were given in relativeunits, using the (002) peak as a reference (I(002)5 100%). Struc-tural properties were evaluated by matching experimental pat-terns against reference patterns from the JCPDS card data andby measuring some structural parameters such as:

(a) the d-spacing values, d002 and d10/100 from the (002) and(10)/(100) peaks using the Bragg formula,21,22

(b) the average stack height Lc and length La of crystallitesfrom the FWHM of the (002) and (10)/(100) peaks using theScherrer formula,21,22

(c) the average degree of orientation (o1, 0ro1r1) of the(002) layers along the fiber-axis using an empirical formula pre-viously reported.21,22

1607

JournalJ. Am. Ceram. Soc., 88 [6] 1607–1614 (2005)

DOI: 10.1111/j.1551-2916.2005.00318.x

R. S. Hay—contributing editor

We have the pleasure to dedicate with affection this article to Prof. Jean Bouix in rec-ognition of his outstanding contribution to material chemistry at the head of the Labo-ratoire des Multimateriaux et Interfaces.

*Member, American Ceramic Society.wAuthor to whom correspondence should be addressed. e-mail: Samuel.Bernard@

univ-lyon1.frzMember, IUF (Institut Universitaire de France).

Manuscript No. 11202. Received July 22, 2004; approved December 28, 2004.

Scanning electron microscopy (SEM, Hitachi S800, Tokyo,Japan) was used to observe the cross-sectional microtexture of fib-ers. Fibers were metallized with an Au/Pd film prior observation.

Transmission electron microscopy (TEM) was investigatedusing a Topcon 002B microscope (Yokohama, Japan). Sampleswere embedded in a resin and cut into thin foils with an ultra-microtome. Foils were then set on microgrids to observe themicrostructure/microtexture in both longitudinal as well ascross-sectional thin specimens. Analysis was carried out inbright-field (BF), high-resolution (002) lattice-fringes (HRTEM)and selected area electron diffraction (SAED) modes. A typicalSAED of BN fibers is presented in Fig. 1.

The SAED mode is used to estimate the local degree of orien-tation o2 (0ro2r1) of the (002) layers along the fiber-axis from

o2 ¼ ð180� zÞ=180 (1)

where z52w and w is the angle between the (002) layers directionand the fiber-axis.

III. Results

(1) XRD

XRD plots of as-obtained BN fibers are presented in Fig. 2. InXRD analysis, h-BN is identified by the presence of peaks at26.761 (002), 41.601 (100), 43.871 (101), 50.151 (102), 55.161(004), 75.931 (110), 82.181 (112), and 85.521 (105).25 In a gener-al way, two disordered BN phases, turbostratic-BN (t-BN) andamorphous-BN (a-BN), are found in addition to h-BN.26 t-BNshows a random stacking sequence of the (002) layers and adisorientation of these layers around the c-axis. a-BN representsa structure disordered at atomic level. XRD patterns of suchphases are significantly distinct from that of h-BN.

XRD of fibers show the principal reflections of h-BN withsignificant differences in peak intensity, position, and FWHM.For example, in A and B, the decrease in the profile intensityfrom 2y5 151 to 2yB201 and the presence of the broadened(002) peaks and diffuse (10)/(100) and (004) peaks which areshifted to the Bragg angles for h-BN are representative of a dis-ordered structure. In contrast, no decrease in the profile inten-sity up to 2yB201 is observed in other fibers. Moreover, theFWHM of the (002) and (10)/(100) peaks is narrower while thepeak position shifts toward 2y values of h-BN from C to F.These observations point to the fact that both crystallinity andcrystalline quality are improved from A to F. In addition, the

decrease in the relative intensity of the (10)/(100) peaks from Ato F is attributed to an enhancement in the grain orientationalong the fiber-axis.

Interestingly, due to the lack of grain orientation in ground-up fibers, peak intensity can be modified and additional peakscan appear in XRD patterns (Fig. 3).

For example, the relative intensity of the (10)/(100) peaks ishigher than that observed in as-obtained fibers (Fig. 2) while itpartially resolves into the (100) and (101) lines going from C toF. Additional (110), (112), and (105) peaks gradually appearfrom C to F. The appearance of additional peaks of h-BN isattributed to an enhancement in the structural ordering fromA to F.

The d-spacing and crystallite size data given in Table IIIhighly reflect previous observations.

Consistently with a poor crystallinity and a low crystallinequality, d002 values are high (d002B3.45 A) in A and B and suchfibers are also composed of small crystallites (Lc 5 31 A andLa5 86 A for A). Crystallinity and crystalline quality are morepronounced in other fibers. As a result, d002 values decrease upto 3.35 A in F, while Lc and La reach values as high as 117 and257 A, respectively. As observed in Fig. 2, the measured degree

Fig. 1. Typical selected area electron diffraction of boron nitride fiberswith (hkl) reflection indexing.

Table II. Average Diameters and Mechanical Properties ofAs-Obtained BN Fibers

As-obtained

BN fibers

Diameter

(mm)

Strength

s (GPa)

Modulus

E (GPa)

Strain

e (%)

A 7.8 1.5 105 1.30B 8.3 1.5 130 1.10C 9.8 1.0 235 0.43D 10.5 2.0 280 0.55E 9.3 1.6 315 0.46F 7.4 1.4 400 0.31

BN, boron nitride.

Table I. Melt-Spinning Conditions Used for the Preparationof As-Spun Fibers

As-spun fibers A B C D E F

Extrusion rate per capillary(m/min)

0.66 0.66 0.66 1.8 1.8 1.8

Wheel velocity (m/min) 97 91 69 151 187 300Drawdown ratiow 12.1 11.7 10.3 9.1 10.2 13Green fiber diameter (mm) 16.5 17 19.4 22 19.6 15.4

wDrawdown ratio is the diameter of the die capillary divided by the final fil-

ament diameter.23

Fig. 2. XRD patterns of as-obtained boron nitride fibers going fromA to F.

1608 Journal of the American Ceramic Society—Bernard et al. Vol. 88, No. 6

of orientation of the (002) layers along the fiber-axis increasesfrom A (o15 0.55) to F (o15 0.85).

(2) SEM Investigations

SEM observations can be easily correlated to XRD experiments.Figure 4 shows the SEM images of BN fibers starting from aglassy-like microtexture in A (Fig. 4(a)) to a coarse-grained mi-crotexture in F (Fig. 4(f)). In a general way, no grain orientationis shown in the cross-sectional microtexture of BN fibers but,with specific spinning parameters, a cross-sectional skin-coreheterogeneity can develop as shown with C (Fig. 4(c)).

Diameters are in accordance with those measured by laserinterferometry.

The glassy-like microtexture of A is similar to that observedin amorphous ceramic fibers which exhibit low apparent crys-tallite sizes (B10–20 A) and no orientation effects.12,27 The mi-crotextural definition slightly improves in B but the cross-sectional microtexture is much more coarse-grained in C, D,E, and F as crystalline materials. The transition observed be-tween A and D illustrated structural changes which occurs whenamorphous ceramic fibers are converted into crystalline materi-als at very high temperature.27

As mentioned before, the microtexture of C consists of a skin-core heterogeneity. Grains extend in a radial direction over theskin whereas they are randomly oriented in the core. Mostly,cracks running along radially oriented grains are propagatedthrough the whole thickness of the skin region while they extendover the full fiber length most probably due to differences in theCTE between their skin and core regions (Fig. 5).

(3) TEM Investigations

TEM is an ideal tool for describing the complex crystalline ar-rangement of BN fibers.

Fibers with low (A, 105 GPa), intermediate (D, 280 GPa),and high (F, 400 GPa) modulus are examined in the longitudinalsection. Since all fibers show a randomly grained cross-sectionalmicrotexture (see SEM analysis), only fibers D are studied in thismode with C.

(A) Longitudinal Sections: Figure 6 shows BF images ofthe longitudinal thin section samples A (a), D (b), and F (c).Corresponding SAED are inserted in each image. (hkl) reflectionindexing of SAED are reported in experimental section.

Consistently with SEM analysis, the BF TEM image of A isfeatureless and of low contrast, indicating a fine crystallinestructure whereas the grain growth is clear in D and F.

Data provided by SAED from such fibers highly reflect XRDresults. The amorphous halo imposed on the diffuse (002) arcsand the poorly resolved (10)/(100), (004), and (11)/(110) arcs

confirm the low level of crystallinity of A. In contrast, SAED ofD and F, which are composed of well resolved (002), (10)/(100),(004), and (11)/(110) arcs as well as (112) (D and F) and (105)(F) arcs, are representative of ordered materials. Furthermore,the decrease in the spread of the (002) arcs from A to F is at-tributed to an increase in the local degree of orientation o2 of the(002) layers from 0.60 to 0.85 (see experimental section for themeasurement of o2).

TEM observations are achieved with high-resolution tech-nique (HRTEM). HRTEM images of BN fibers are shown inFig. 7.

The HRTEM image of A shows that the long-range ordering,common in crystalline ceramics, is minimal and the crystallinityis inhomogeneous (Fig. 7(a)). In particular, the lattice imagehighlights the coexistence of amorphous regions and nanocrys-talline phases (basal structural units) of distinct shapes andorientations. These nanograins with thickness around five atom-ic basal planes and lengths o100 A are significantly buckled.Moreover, high tortuosities in their stacking sequences are fre-quently encountered similarly to those seen in t-BN.26 In con-trast, HRTEM images of fibers D as well as F show local regionswhich are composed of ordered and extended (002) layers withd0025 3.34 A (Fig. 7(b)), but these regions are frequently em-bedded in a t-BN matrix with d0025 3.37 A (Fig. 7(c)). The ori-entation of the (002) layers along the fiber-axis is highlighted inlattice images of fibers D and F.

(B) Cross Sections: Cross-sectional microtexture exam-ined by TEM analysis are well representative of SEM observa-tions.

BF TEM images show that the cross section of D consists of arandomly grained microtexture in accordance with SEM anal-ysis (Fig. 8). Since the crystals are randomly oriented in the fiber,the SAED shows (00l) rings.

The BF images of C (Fig. 9) and the corresponding SAED(Fig. 10) show that the cross-sectional microtexture consists of askin-core heterogeneity with radially oriented fibrils in the skinregion and randomly-oriented grains in the core region.

The SAED of the skin region (Fig. 10(a)) shows (002) arcsthat indicate a preferential orientation of the (002) layers alongthe fibrils. In contrast, the SAED of the core region shows (002)rings that reflect the random orientation of the (002) layers(Fig. 10(b)) as generally shown in the homogeneous cross-sec-tional microtexture of other fibers.

IV. Discussion

(1) Microstructural Ordering in BN Fibers

Previous results show that a relationship is established betweenthe microtexture/microstructure of BN fibers and their mechan-ical properties.

In XRD analysis, the lack of the main (hkl) reflections of h-BN and the only presence of reflections from the basal (00l) and(hk0) planes in fibers A and B show that the structure of lowmodulus fibers consists of a two-dimensional short-range order-ing. As a consequence, the (hk0) reflections can be abbreviated

Fig. 3. X-ray diffraction patterns of ground-up boron nitride fibers asfine powders going from A to F.

Table III. Structural Parameters of As-Obtained BN Fibers

As-obtained

BN fibers d002 (A) Lc (A) La (A) d10/100 (A) o1

A 3.45 31 86 2.13 0.55B 3.44 42 110 2.13 0.60C 3.37 100 280 2.14 n.o.w

D 3.38 109 280 2.15 0.78E 3.36 120 310 2.15 0.82F 3.35 117 357 2.16 0.85h-BN 3.33 2.17

wNot determined due to the heterogeneous cross-section of C (see SEM anal-

ysis). BN, boron nitride.

June 2005 Crystallinity, Crystalline Quality, and Microstructural Ordering in Boron Nitride Fibers 1609

Fig. 4. Scanning electron microscopy images of the cross-sectional microtexture in boron nitride fibers (a) A, (b) B, (c) C, (d) D, (e) E, and (f) F.

1610 Journal of the American Ceramic Society—Bernard et al. Vol. 88, No. 6

as (hk) reflections. Furthermore, the d002 values, which are4%–5% greater than that for h-BN, and the broadened peaksdue to small crystallites within the fibers are clearly representa-tive of a turbostatic-like microstructure. HRTEM investigationssupport these observations since they show that the microstruc-ture of low modulus fibers is composed of an alternation ofsmall crystals and amorphous BN regions.

In a general way, d002 values decrease from C to F and thelong-range ordering enhances as modulus increases. In particu-lar, the sharpening of the (00l) peaks and the appearance of the(100), (101), (112), and (105) reflections from C to F differen-tiates low modulus fibers (A and B) from high modulus fibers(C–F). However, beside these peaks, none of the other main re-flections of h-BN such as the (102) reflection at 2y5 50.1501 areobserved in XRD patterns, presumably because of the lack ofregular hexagonal three-dimensional ordering at very long dis-tance. This means that the structure of high modulus BN fiberstends to order as a transitional phase so-called ‘‘meso-hexago-nal’’ (ms-BN) which consists of hexagonal nanograins embed-ded in a turbostratic phase. The lack of strong (101) reflectionsoutside the (100) rings confirms the presence of ms-BN.

(2) Young’s Modulus

TEM analysis shows that BN fibers are made up of a series ofend-to-end linked crystallites, each one having a disorientationangle w which gradually decreases from A to F. XRD also showthat the grain orientation is closely related to modulus (TableII). For example, the most important feature of high modulusBN fibers is that their crystallites are nearly parallel to the fiber-axis. These results can be correlated to findings for high moduluscarbon fibers in which the (002) sheets are arranged nearlyparallel to the fiber-axis.28 Northolt et al.29 showed that therelationship between modulus and orientation distributionmeasured on graphite, in a direction making an angle y withthe c-axis, was given by:

ð1=EÞ ¼ ð1=EcÞ þ ð cos2 y� �

=GcÞ (2)

Fig. 5. Propagation of cracks through the whole thickness of the radialregion in boron nitride fibers C.

Fig. 6. Bright-field transmission electron microscopy images and asso-ciated selected area electron diffraction of the longitudinal sectionalmicrostructure in boron nitride fibers (a) A, (b) D, (c) F.

June 2005 Crystallinity, Crystalline Quality, and Microstructural Ordering in Boron Nitride Fibers 1611

with y5 (p/2)�w, Ec is the axial elastic modulus of the crystal,/cos2 yS is an orientation parameter describing the orientationof crystallites compared with the fiber-axis, and Gc is the axialshear modulus along the direction of crystallites.

Due to structural similarities between graphite and h-BN, theprevious model is applicable to BN fibers. The reciprocal of E is

Fig. 7. High-resolution transmission electron microscopy micrographsof the longitudinal sectional microstructure in boron nitride fibers (a) A,(b) D, (c) F.

Fig. 8. Bright-field transmission electron microscopy image and asso-ciated selected area electron diffraction of the cross-sectional micro-structure in boron nitride fibers D.

Fig. 9. Bright-field transmission electron microscopy image of thecross-sectional microstructure in boron nitride fibers C.

1612 Journal of the American Ceramic Society—Bernard et al. Vol. 88, No. 6

plotted versus the orientation parameter cos2 y calculated fromTEM data of fibers A, D, E, and F (Fig. 11).

Although some TEM data are missing in the range 105–280GPa (0.0035/1/ES0.0095) and despite of some uncertainties incalculating /cos2 yS, the linear increase of 1/E with cos2 y clear-ly clearly shows the relationship between the orientation of the(002) layers and the Young’s modulus. By extrapolating the lin-ear plot to a zero value of cos2 y, it is shown that BN fibers canreach high values of Young’s modulus when (002) layers areideally aligned along the fiber-axis. This can be correlated to theanisotropic configuration of h-BN which indicates a high in-plane average B–N bond energy (D(B–N)B500.75 kJ/mol).30

Furthermore, XRD results also show that the orientationlevel enhances as the grain size increases. As illustrated in TableII, large crystallites are better aligned along the fiber-axis thanthose having small size. This observation points to the fact thatend-to-end linked crystallites tend to coalesce and grow as theirorientation along the fiber-axis enhances leading to increasedmodulus.

(3) Failure Strain and Crack Propagation

XRD data show that the failure strain decreases as d002 valuesincrease (Fig. 12). Many other BN fibers produced in our lab-oratory successfully confirm this behavior.

The crack propagation between crystallites is probably thedominant failure mechanism which can explain this relationship.For example, the crack propagation is disturbed in low modulusfibers A and B by the high tortuosity of their microstructure.When tensile stress is applied to A and B, the good resistance ofsuch fibers to crack propagation gives high failure strain. Incontrast, HRTEM images show that D and F are composed ofextended (002) layers with a local flat layer microstructure nearlyoriented along the fiber-axis. The resistance to crack propaga-tion is low since the crystallite initiating the failure is sufficientlycontinuous with the neighboring crystallites to permit the crackto continue its propagation along the flat cleavage (002) layers.Therefore, elongation and failure strain of high modulus BNfibers decrease as crystallinity increases. In addition, it is alsoclear that the crack propagation highly depends on the type ofcross-sectional microtexture. In particular, crack propagation isfavored when the cross-sectional microtexture consists of a ra-dial orientation as shown with fibers C in which crack propa-gation occurs along the fibrils, which compose the skin region.

(4) Tensile Strength

The practical strength of BN fibers is always several orders ofmagnitude lower than the predicted tensile strength whichshould be proportional to modulus as following: s/EB0.1. Asa result, a s/E ratio of 0.014 is found for A whereas it decreasesto 0.0035 for F.

Fig. 10. Bright-field transmission electron microscopy images and as-sociated selected area electron diffraction of the skin (a) and core (b)regions of the cross-sectional microstructure in boron nitride fibers C.

Fig. 11. Plot of the reciprocal of modulus 1/E versus the orientationparameter cos2 y.

Fig. 12. Plot of the interlayer spacing d002 versus the failure strain e.

June 2005 Crystallinity, Crystalline Quality, and Microstructural Ordering in Boron Nitride Fibers 1613

Fibers A, C, and F are appropriate examples which reveal thefactors limiting their strength.

The microstructure and microtexture of A and F are stronglydifferent whereas their tensile strengths (sB1.5 GPa) and di-ameters (+B7.5 mm) are similar. The microstructure of A con-sists of a disordered arrangement of BN layers and its apparentdense microtexture is similar to that of glassy-like materials. Incontrast, the microstructure of F consists of a meso-hexagonalarrangement with large crystallites nearly oriented along thefiber-axis while its microtexture is coarse-grained. In a previouswork, we showed that a coarse-grained microtexture involved alow density (d5 1.85) and a high level of porosity.21 As a con-sequence, in contrast to low modulus fibers which probablyreach their expected strength, the practical strength of highmodulus fibers is probably reduced because of porosity. Asa consequence, the s/E ratio decreases as grains grow. Besidethese fibers, it is also shown that the presence of a skin-coreheterogeneity in the cross-sectional microtexture (fibers C) in-ducing cracks leads to decreased strengths (sB1 GPa). Theseresults show that an improved textural homogeneity free of sur-face defects is required to obtain high strengths. Obviously, thedependence of strength with critical defects through the Griffithequation is well-established in ceramic fibers.11 In addition, it isalso clear that the development of the structural ordering of BNlayers by controlling the porosity is another important require-ment to enhance the tensile strength of BN fibers.

V. Conclusion

The present study showed a relationship between the micro-structure/microtexture of BN fibers and their mechanical prop-erties. In particular, it was shown that the Young’s modulus wasrelated to the grain orientation. The microstructure of BN fiberswith low Young’s modulus consisted of a mixture of disorderednanocrystals and amorphous regions, whereas the large crystal-lites were stacked in a meso-hexagonal ordering and nearly ori-ented along the fiber-axis in high modulus fibers. In a generalway, the crystallinity enhanced as the microstructural orderingincreased. The development of a flat layer microtexture in-creased the crack propagation along the basal (002) layers caus-ing a general decrease in the fiber flexibility and failure strain.The influence of microstructure upon tensile strength of BNfibers was clear but the appearance of porosity with the growthof grains in high modulus fibers was the main factor in limitingthe tensile strength.

References

1R. T. Paine and C. K. Narula, ‘‘Synthetic Routes to Boron Nitride,’’ Chem.Rev., 90, 73–91 (1990).

2R. Haubner, M. Wilhelm, R. Weissenbacher, and B. Lux, ‘‘Boron Nitrides—Properties, Synthesis and Applications’’; pp. 1–45 inHigh Performance Non-OxideCeramic II, Structure Bonding, Vol. 102, Edited by M. Jansen. Springer-Verlag,Berlin, Heidelberg, 2002.

3J. Huang and Y. T. Zhu, ‘‘Advances in the Synthesis and Characterization ofBoron Nitride,’’ Defects Diffusion Forum, 186, 1–32 (2000).

4A. R. Bunsell and M.-H. Berger, ‘‘Fine Diameter Ceramic Fibers,’’ J. Eur.Ceram. Soc., 20, 2249–60 (2000).

5J. Bill and F. Aldinger, ‘‘Precursor-Derived Covalent Ceramics,’’ Adv. Mater.,7, 775–87 (1995).

6M. Peuckert, T. Vaahs, and M. Bruck, ‘‘Ceramics from Organometallic Pol-ymers,’’ Adv. Mater., 2, 398–404 (1990).

7M. Weinmann and F. Aldinger, ‘‘Precursor-Derived Ceramics’’; pp. 265–368in The Handbook of Advanced Ceramics, Edited by S. Somiya, F. Aldinger, N.Claussen, R. Spriggs, K. Uchino, K. Koumoto, and M. Kaneno. Elsevier,Amsterdam, 2003.

8K. J. Wynne and R. W. Rice, ‘‘Ceramics Via Polymer Pyrolysis’’, Ann. Rev.Mater. Sci., 14, 297–334 (1984).

9R. Riedel, ‘‘Advanced Ceramics from Inorganic Polymer’’; pp. 4–50 in A Com-prehensive Treatment, Materials Science and Technology, Vol. 17, Part B, Edited byR. W. Cahn, P. Haasen, and E. J. Kraner. Weinheim, Germany, 1996.

10P. Greil, ‘‘Polymer Derived Engineering Ceramics,’’ Adv. Eng. Mater., 2,339–48 (2000).

11J. Lipowitz, ‘‘Polymer-Derived Ceramic Fibers,’’ Am. Ceram. Soc. Bull., 70,1888–94 (1991).

12J. Lipowitz, ‘‘Fiber Synthesis Process’’; pp. 433–55 in Carbide, Nitride andBoride Materials Synthesis and Processing, Edited by A. W. Weimer. Chapman &Hall, London, 1997.

13V. S. R. Murphy, ‘‘Processing of Non-Oxide Ceramic Fibers,’’ Bull. Mater.Sci., 16, 87–108 (1993).

14K. Okamura, ‘‘Ceramic Fibres from Polymer Precursors,’’ Composites, 19,107–20 (1987).

15S. Yajima, J. Hayashi, M. Omori, and K. Okamura, ‘‘Development of a Sil-icon Carbide Fibre with High Tensile Strength,’’ Nature, 261, 683–5 (1976).

16G. Winter, W. Verbeek, and M. Mansmann, ‘‘Production of Shaped Articlesof Silicon Carbide and Silicon Nitride’’; US Patent 3 892 583, July 1, 1975.

17B. Toury, P. Miele, D. Cornu, H. Vincent, and J. Bouix, ‘‘Boron Nitride Fib-ers Prepared from Symmetric and Asymmetric Alkylaminoborazine,’’ Adv. Funct.Mater., 12, 228–34 (2002).

18B. Toury, S. Bernard, D. Cornu, J.-M. Letoffe, F. Chassagneux, and M. Mi-ele, ‘‘High-Performance Boron Nitride Fibers Obtained from Asymmetric Al-kylaminoborazine,’’ J. Mater. Chem., 13, 274–9 (2003).

19P. Miele, B. Toury, S. Bernard, D. Cornu, F. Chassagneux, and J. Bouix,‘‘Elaboration of High-Performance Boron Nitride Fibers from Tailored Borazine-Based Molecular and Polymeric Precursors’’; pp. 261–8 in The 10th InternationalCeramics Congress Proceedings, Vol. 31, Advances in Science and Technology, PartB, Edited by P. Vincenzini. Faenza, Italy, 2003.

20S. Bernard, D. Cornu, P. Miele, H. Vincent, and J. Bouix, ‘‘Pyrolysis ofPoly[2,4,6-tri(methylamino)borazine] and its Conversion into BN Fibers,’’ J. Or-ganomet. Chem., 657, 91–7 (2002).

21S. Bernard, F. Chassagneux, M. P. Berthet, H. Vincent, and J. Bouix, ‘‘Struc-tural and Mechanical Properties of a High-Performance BN Fibre,’’ J. Eur.Ceram. Soc., 22, 2047–59 (2002).

22S. Bernard, M. P. Berthet, D. Cornu, P. Miele, H. Vincent, and J. Bouix,‘‘Development of Boron Nitride Fibers with High Tensile Strength by the Prece-ramic Polymer Route’’; pp. 63–70 in Advanced Inorganic Fiber Composites forStructural Applications Proceedings, Vol. 40, Advances in Science and Technology,Edited by P. Vincenzini and C. Badini. Faenza, Italy, 2003.

23D. D. Edie and M. G. Dunham, ‘‘Melt-Spinning Pitch-Based Carbon Fibers,’’Carbon, 27, 647–55 (1989).

24S. Bernard, K. Ayadi, M.-P. Berthet, F. Chassagneux, D. Cornu, J.-M.Letoffe, and P. Miele, ‘‘Evolution of Structural Features and Mechanical Prop-erties During the Conversion of Poly[(methylamino)borazine] Fibers into BoronNitride Fibers,’’ J. Solid State Chem., 177, 1803–10 (2004).

25R. S. Pease, ‘‘An X-Ray Study of Boron Nitride,’’ Acta Crystallogr., 5, 356–61(1952).

26J. Thomas Jr., N. E. Weston, and T. E. O’Connor, ‘‘Turbostratic Boron Ni-tride, Thermal Transformation to Ordered-Layer-Lattice Boron Nitride,’’ J. Am.Chem. Soc., 84, 4619–22 (1963).

27Z.-Y. Chu, C.-X. Feng, and Y.-C. Song, ‘‘Development of a New Si–C–O–N–B Ceramic Fiber Using a Hybrid Precursor of Polycarbosilane and Poly-borosilazane,’’ J. Mater. Sci. Lett., 22, 725–8 (2003).

28M. Endo, ‘‘Structure of Mesophase Pitch-Based Carbon Fibres,’’ J. Mater.Sci., 23, 598–605 (1988).

29M. G. Northolt, L. H. Veldhuizen, and H. Jansen, ‘‘Tensile Deformation ofCarbon Fibers and the Relationship with the Modulus for Shear Between the Ba-sal Planes,’’ Carbon, 29, 1267–79 (1991).

30M. Cote, P. D. Haynes, and C. Molteni, ‘‘Boron Nitride Polymers: BuildingBlocks for Organic Electronic Devices,’’ Phys. Rev. B, 63, 125207-1-4 (2001). &

1614 Journal of the American Ceramic Society—Bernard et al. Vol. 88, No. 6