Structural and mechanical behavior of boron nitride fibers derived from poly[(methylamino)borazine] precursors: Optimization of the curing and pyrolysis procedures

Article abstract of DOI:10.1111/j.1551-2916.2005.00685.x

A series of boron nitride fibers were prepared by varying the tension applied on a same lot of poly(methylamino)borazine fibers during their pyrolytic conversion at 1800°C. Their microstructure/microtexture was studied by X-ray diffraction, scanning electron microscopy, and transmission electron microscopy investigations. Such analyses showed that ceramic fibers presented good crystallinity, but crystallites oriented along the fiber axis with fiber stress. It is most interesting that stretching was essentially effective during the pyrolysis to 1000°C, preventing fiber crimping, and decreasing the fiber diameter. Therefore, the fiber strength increased, while the fiber modulus also increased because of an improvement of the basal layer orientation along the fiber axis.

Full text of DOI:10.1111/j.1551-2916.2005.00685.x

J. Am. Ceram. Soc., 89 [1] 42–49 (2006)
DOI: 10.1111/j.1551-2916.2005.00685.x
r 2005 The American Ceramic Society
ournal
J
Structural and Mechanical Behavior of Boron Nitride Fibers Derived
from Poly[(Methylamino)Borazine] Precursors: Optimization of the
Curing and Pyrolysis Procedures
Pascaline Toutois, Philippe Miele,w,z Sylvain Jacques, David Cornu,* and Samuel Bernard*
Laboratoire des Multimateriaux et Interfaces, UMR CNRS 5615, Universite
´ ´
Claude Bernard–Lyon 1, F-69622
Villeurbanne, France
A series of boron nitride fibers were prepared by varying the
tension applied on a same lot of poly(methylamino)borazine
fibers during their pyrolytic conversion at 18001C. Their
microstructure/microtexture was studied by X-ray diffraction,
scanning electron microscopy, and transmission electron micros-
copy investigations. Such analyses showed that ceramic fibers
presented good crystallinity, but crystallites oriented along the
fiber axis with fiber stress. It is most interesting that stretching
was essentially effective during the pyrolysis to 10001C, pre-
venting fiber crimping, and decreasing the fiber diameter. There-
fore, the fiber strength increased, while the fiber modulus also
increased because of an improvement of the basal layer orien-
tation along the fiber axis.
shocks up to 20001C in an inert atmosphere, high mechanical
properties along the (002) layers, an excellent chemical stability,
and a very good resistance toward oxidation at a high temper-
ature (upon B9001–10001C).7 Above all, it exhibits a low coef-
ficient of thermal expansion. Therefore, h-BN-based fibers (BN
fibers) should be ideally suited for the reinforcement of ceramic
matrix composites (CMCs) instead of carbon fibers in thermo-
structural applications, insofar as a method allows their prepa-
ration.
The fabrication of BN materials with complex shapes is dif-
ficult to achieve using conventional powder routes, particularly
in the case of fibers. To prepare fibers, the pyrolysis of suitable
molecular and polymeric precursors, usually referred to as the
polymer-derived ceramics (PDCs) technique, appears to be the
most appropriate.8–10 This chemical approach is composed of
several complex sequences summarized as follows: (1) design
and synthesis of a molecular precursor, (2) its polymerization
into a preceramic polymer composed of an inorganic backbone
with pendent organic groups, (3) preparation of a covalent and
amorphous ceramics at an intermediate temperature (with the
aid of a chemical treatment if necessary), and (4) conversion into
ultimate crystalline ceramics by a high-temperature thermal
treatment. This route, in which a shaping step may be inserted
between steps (2) and (3), was first developed for the fabrication
of SiC and Si/C/N fibers.11–13 Later it was extensively studied
for other types of non-oxide ceramic fibers such as BN14–18 and
SiBCN.19–21
Most spinnable polymers are based on the borazine molecule,
which already foreshadows the hexagonal structure of h-BN
and its B:N stoichiometry.22–25 These studies showed that the
presence of alkyl groups in the pendent borazine units and the
type of linkages between the borazine rings in the polymer
structure were prerequisites for preparing BN fibers. These
general trends were confirmed in recent works published by
our group.14–16,25–28 In these studies, we showed that tri(alkyl-
amino)borazines are excellent molecular precursors for devel-
oping polymers with controlled melt spinnability and thermal
reactivity. Particularly, we have devoted a large part of our
efforts toward poly[tri(methylamino)borazine] (polyMAB), be-
cause this polymer can exhibit a set of physicochemical proper-
ties of particular interest for the preparation of fine-diameter BN
fibers with controllable mechanical and structural properties us-
ing spinning, curing, and pyrolysis procedures.16,25–26 In previ-
ous works, we reported the chemical modifications, and the
changes in the structural and mechanical features of fibers (crys-
tallite sizes, crystallization degree, fiber strength, etc.) occurring
during the polymer-to-ceramic conversion.16,26 It was demon-
strated that the curing of green fibers at 4001C in an ammonia
atmosphere followed by the pyrolysis of as-cured fibers in am-
monia (10001C), then nitrogen (18001C) atmospheres, yielded
purely BN fibers with high mechanical properties.
I. Introduction
HE existence of ceramics has been known for a long time,
and they are fascinating materials, resulting in them being
T
present in a wide range of domains ranging from day-to-day life
utilization to specific high-tech applications. However, a com-
mon drawback of ceramics is their intrinsic brittleness, which
limits their utilization, particularly as thermostructural materi-
als. To overcome this problem, it was demonstrated that the
fabrication of continuous fiber-reinforced ceramic composites
(CFCCs) is the most effective alternative for development of
nonbrittle ceramics.1–3 ‘‘Graphitized’’ carbon fibers have been
the most extensively studied reinforcing agent, resulting in
numerous commercial developments of C/C composites.4
Nevertheless, the low resistance of graphite toward oxidation
(Tonset ꢀ 4501C) excludes the utilization of such fibers in a high
temperature and a harsh environment. On looking at the main
properties of several non-oxide ceramics one can notice that
boron nitride (BN) could represent an interesting alternative to
carbon for high-temperature structural applications owing to
the isoelectronic character of the C–C and B–N bonds and the
structural similarity of graphite and the hexagonal form of BN
(h-BN), which is the most thermodynamically stable phase ob-
tained under standard conditions.5,6 Its lamellar structure is
composed of basal layers (planes of hexagons), in which boron
and nitrogen atoms alternate. Within the basal (002) layers (a-
axis), the boron and nitrogen atoms are covalently bonded,
whereas they are bonded through weak van der Waals interac-
tions (c-axis) between subsequent stacked (002) layers. This
structure results in unique properties, mostly anisotropic, which
are of great interest for advanced technological applications.
The most attractive properties are a high resistance to thermal
D. Green—contributing editor
In the present work, we focus our attention on the influence
of some crucial parameters of the polymer-to-ceramic conver-
sion on optimization of this conversion, and therefore control-
ling the mechanical properties of the final ceramic fibers. In
Manuscript No. 20576. Received May 18, 2005; approved July 18, 2005.
*Member, American Ceramic Society.
wAuthor to whom correspondence should be addressed. e-mail: Philippe.Miele@
univ-lyon1.fr
zMember, Institut Universitaire de France (IUF).
42

January 2006
Structural and Mechanical Behavior of Boron Nitride Fibers
43
particular, the present paper is devoted to both the study of the
fiber stress that occurs during the pyrolytic conversion, and the
effect of fiber tensioning on the mechanical and structural prop-
erties of BN fibers.
H
N
H
N
H
N
H C
3
N
N
B
N
B
N
CH
H C
B
N
B
N
3
3
∆
H
H
B
N
H
B
N
II. Experimental Procedure
H
(1) Characterization Techniques
H
H C
3
CH
n
3
All reactions were performed in an argon atmosphere and
anhydrous conditions using a standard vacuum line, Schlenk
techniques, and an efficient dry glove-box. The solvents were pu-
rified over appropriate drying agents by standard methods. The
starting monomer, tri(methylamino)borazine (MAB), and the
resulting polymer (polyMAB) were prepared following a proce-
dure that has been described previously in the literature.16–17
The differential scanning calorimetry analysis was performed on
a ca. 15.0 mg sample using a TA8000 Mettler-Toledo apparatus
(Viroflay, France) between ꢁ101 and 1701C at a heating rate of
101C/min in an argon atmosphere. The 11B NMR spectrum was
recorded on a Bruker AM300 spectrometer (Attlingen, Germa-
ny) at 96.28 MHz with Et2O ꢂ BF3 as an external reference (pos-
itive values downfield) in CH2Cl2 solutions containing C6D6 as
an internal reference. 1H and 13C spectra were, respectively, ob-
tained at 300 and 75 MHz in CD2Cl2. Size-exclusion chroma-
tography analysis was performed in CH2Cl2 using a Gilson 118
UV detector (Villiers le Bel, France) at 240 nm and water
styragel columns. Thermogravimetric analysis (TGA) was
performed on a TGA 92 16–18 Setaram (Caluire, France)
instrument at 15001C (10001C (ammonia atmosphere) 1
15001C (nitrogen atmosphere) at 0.81C/min). Chemical compo-
sitions of the polymer were obtained from the Service Central de
Microanalyses du C.N.R.S (Vernaison, France). X-ray diffrac-
tion (XRD) patterns were recorded between 2y 5 201 and 451
using a Philips diffractometer (Almelo, the Netherlands) and
CuKa1 radiation (l 5 0.1540 nm at 40 kV and 30 mA) from as-
obtained fibers placed in a position so that the fiber axis is per-
pendicular to the incidence plane. The X-ray patterns obtained
from the BN fibers exhibited some peak characteristics of h-BN
by matching experimental patterns against reference patterns
from the Joint Committee on Powder Diffraction Standards
card data.29 Intensities of the (hkl) peaks were given in relative
units using the (002) peaks as a reference (I(002) 5 100%).27
Transmission electronic microscopy (TEM) micrographs were
obtained using a Topcon EMB-002B microscope (Yokohama,
Japan) in both imaging and diffraction modes (diffraction pat-
terns were obtained on 0.5 mm selected-area electron diffraction
(SAED) patterns) on fragments of fibers embedded in an epoxy
resin and then cut with an ultramicrotom (LEICA S, Vienna,
Austria) as fully described in Chassagneux et al.28 The SAED
mode was also used to estimate the local degree of orientation
o2 of the basal layers along the fiber axis as described in Bernard
et al.27 Fibers were tensile tested at room temperature using an
MTS-Adamel DY-22 tensile tester (Adamel Lhomargy, Roissy
en Brie, France) equipped with a 5 N load cell, with the cross-
head speed being 0.1 mm/min. The Weibull statistic was used to
estimate the averaged tensile strength of fibers for a failure
probability of 0.632 from a series of 50 fibers with a gauge length
of 10 mm. The average Young’s modulus was evaluated from
the strain–strength curves.
MAB
PolyMAB
Scheme 1. Schematic representation of the poly[tri(methylamino)
borazine] (polyMAB) structure.
found to be lower than 2 wt%, and was therefore neglected. The
other values were as follows: 13C NMR (CD2Cl2) d (ppm): 30.7
(NCH3 bridges), 27.68 (NHCH3);. 11B NMR d (ppm): 25.96
(br); H NMR (CD2Cl2) d (ppm): 1.89 br (NHCH3), 3.39 br
(NH borazine), 2.58 vbr (NHCH3, NCH3); and FT-IR/KBr
(cmꢁ1): 3447 (nNꢁH), 2948, 2894 (nCꢁH), 1478 ðdNðHÞCH þ nBꢁN
1
3
þdasCH Þ, 1161 (nCꢁN), and 609 (dBN, out-of-plane).
3
(3) Fiber Preparation
Poly(methylamino)borazine fibers (green fibers) were obtained
using the general melt-spinning process (nitrogen atmosphere)
described previously26 from a same lot of polymer. In particular,
the polymer was spun at 1801C, and emerged as an endless green
fiber from the spinneret with a capillary of 200 mm diameter. The
as-spun fiber was immediately stretched into a finer-diameter
green fiber by a winding graphite spool that exhibited a negli-
gible coefficient of thermal expansion (Carbone Lorraine, Paris,
France). The polyMAB fiber emerged as continuous green fibers
wound on the spool.
After the melt-spinning operation, a series of three bundles of
the green fiber, labeled A, with an initial equivalent length (15.70
cm) and diameter (14.9 mm), was isolated to undergo pyrolytic
conversion (in an ammonia atmosphere up to 10001C (0.81C/
min, atmospheric pressure), and then in a nitrogen atmosphere
up to 18001C (0.81C/min, atmospheric pressure)) as follows:
(1) The first bundle of A was cut along the spool after the
melt spinning, which resulted in a lack of any tension during the
subsequent pyrolytic conversion into BN fibers A-1.
(2) The second bundle of A was maintained around the
spool to undergo curing (4001C; ammonia atmosphere) and
mineralization (10001C; ammonia atmosphere) steps under ten-
sion. After such heat treatments at 10001C, as-mineralized fibers
were cut along the spool to be transformed without tension into
BN fibers A-2 between 10001 and 18001C in a nitrogen atmos-
phere.
(3) A third bundle of A was maintained around the spool
during the whole classical pyrolytic conversion at 18001C into
BN fibers A-3. It is therefore considered to be under tension
during the whole pyrolytic conversion.
III. Results
(1) Scanning Electron Microscope (SEM) Observations
The green fibers A, 15.70 cm in length and 14.9 mm in diameter,
were prepared in the same batch, and thus pyrolyzed together up
to 18001C with different applied tensions as mentioned in the
experimental procedure. Figure 1 shows the SEM image of a
typical green fiber A after a melt-spinning operation. Its fea-
tureless cross-sectional microtexture is similar to that of com-
mon organic polymer fibers. Defects appeared neither in the
volume nor on the fiber surface.
(2) Precursor Synthesis
114.7 g (683 mmol) of MAB was heated up to 2001C following a
nine-step procedure: 1 h at 1001C, 1 h at 1201C, 2 h at 1401C, 20
h at 1501C, 7 h at 1601C, 17 h at 1701C, 24 h at 1801C, 5 h at
1901C, and 2 h at 2001C via major condensation of –N(H)CH3
units as illustrated in Scheme 1.
A vitreous yellow air- and moisture-sensitive solid, po-
ly[(trimetylamino)borazine] (89.1 g) was obtained. Its molecu-
lar weight was Mw 5 910; and glass-transition temperature,
Tg 5 861C. The ceramic yield was a 5 53.9%. Anal. found
(wt%): C, 19.6; H, 7.9; B, 24.1; N, 48.4. Oxygen content was
The green fiber was cured in an ammonia atmosphere
(4001C), and then pyrolyzed in ammonia (10001C) and nitro-
gen (18001C) atmospheres yielding BN fibers A-1 (without
applied tension), A-2 (green fiber maintained under tension up
to 10001C), and A-3 (green fiber maintained under tension up to

44
Journal of the American Ceramic Society—Toutois et al.
Vol. 89, No. 1
tural differences were observed from SEM investigations as
illustrated in Fig. 2.
In particular, the three types of fibers displayed a circular
cross section similar to that of the original green fiber A,
which means that the polymer possessed a reactivity tailored
for the curing and pyrolysis procedures, and is thus ideal for
generating materials without loss in fiber integrity. Fibers
were smooth and free of surface defects, and their cross-section-
al microtexture did not exhibit marked differences. Observa-
tions of these cross-sections revealed a similar grainy microtex-
ture as opposed to the featureless cross-sectional microtexture of
the green fiber (Fig. 1). This coarse-grained microtexture is
characteristic of sintered crystalline ceramic materials. Further-
more, considering the findings in SEM and laser interferome-
try investigations (Table I), the fiber diameter decreased in
the sequence A-1 (10.7 mm, Df/f0 5 28%)4A-2 (9.4 mm,
Df/f0 5 37%)4A-3 (8.4 mm, Df/f0 5 43.6%). Besides, the
original length of the green fiber A (l 5 15.70 cm) was not
retained in the resulting BN fibers A-1 (lB12 cm), and A-2
(lB14.6 cm). In contrast, it was slightly increased in BN fibers
A-3 (lB16.20 cm).
The observed differences in the measured values of dia-
meters and lengths could be easily correlated to the shrink-
age that the green fiber A underwent during its pyrolytic con-
version.
As the same green fiber was used as precursor for BN fibers,
and the same temperature and heating rate conditions were ap-
plied, Table I shows that the tensile strength of BN fibers is in-
directly related to the fiber stress, as it was shown that the
strength of fibers A-1 was about 20% lower than those of BN
fibers A-2 and A-3. The strengths of fibers A-2 and A-3 are sim-
ilar, and as-obtained strengths for such fibers heated under
Fig. 1. Scanning electrom microscope image of the cross-sectional mi-
crotexture of the green fiber A.
18001C). After the pyrolytic conversion, as-obtained ceramic
fibers were white in color and of flexible form, but it was
observed that fibers A-1 were considerably crimped. Neverthe-
less, despite the different morphologies, no significant microtex-
Fig. 2. Scanning electron microscope images of the cross-sectional microtexture of boron nitride fibers A-1, A-2, and A-3.

January 2006
Structural and Mechanical Behavior of Boron Nitride Fibers
45
Table I. Physical Properties of the Three Types of BN Fibers
Fiber type
Diameter (mm)
Strength s (GPa)
Modulus E (GPa)
Strain e (%)
Lc (nm)
La (nm)
d002 (nm)
o
o2
A-1
A-2
A-3
10.7
9.4
8.4
0.9
1.1
1.2
110
285
285
0.72
0.36
0.44
9.9
9.0
8.6
21.3
21.9
18.2
0.339
0.336
0.337
0.37
0.85
0.85
0.3
0.9
0.9
BN, boron nitride.
tension were typically those found for high-performance BN
fibers studied in previous papers.26,27
patterns (o1). Both values are close to that reported by Okano
and Yamashita30 in a patent describing the fabrication of high-
performance BN fibers with highly oriented crystallites. The
comparison between (002) dark-field (DF) mode and the BF
normal mode (Fig. 5) obtained at higher magnification of the
same part of the fibers allows positive identification of the re-
gions in the specimen that contributed to a particular reflection,
i.e., (002) reflection, in the diffraction pattern. By selecting only
the (002) beam to form the DF image, the regions of the fiber
having scattered, this beam appeared bright on a dark back-
ground, so that the image is formed only from those crystallites
diffracting electrons into the (002) beam. Each bright domain
images a single crystallite (circled in Fig. 5), and the extent of the
network is that of a local crystallite orientation. The dark areas
in the DF image correspond to other crystallite orientations, ei-
ther twisted or tilted considerably to allow their (002) beams to
pass through the aperture.
The presence of some rhombohedral BN crystals generally
found in the close vicinity of internal pores in high-resolution
TEM (HRTEM) images (Fig. 6), as it was recently described, is
also remarkable.28 Nevertheless, HRTEM investigations of such
fibers frequently show local hexagonal segregation embedded in
a turbostratic BN phase. The turbostratic BN phase (t-BN) has
already been described elsewhere.31
(2) XRD Investigations
The XRD patterns of the three types of as-obtained BN fibers
were nearly identical (in terms of definition and broadening of
the peaks), and revealed a polycrystalline nature (Fig. 3), despite
a slight shift of the (002) peaks to higher Bragg’s angles from
A-1 (2y(002) 5 26.271) to A-3 (2y(002) 5 26.431), approaching clos-
er to the 2y(002) position found in h-BN crystals (2y(002) 5 26.761).
As an illustration, XRD data are reported in Table I. Table I
shows that significant differences were found neither in the
average crystallite sizes nor in interlayer spacing. However,
although the crystallite sizes are similar, irrespective of the
kind of ceramic fibers, their orientation along the fiber axis dif-
fers significantly. This characteristic is immediately clear in
XRD patterns reported in Fig. 3. We recently reported that an
average degree of orientation o1 of the (002) layers along the
fiber axis could be estimated by comparing the intensity ratio of
the above-mentioned diffraction peaks in both XRD patterns
obtained from as-pyrolyzed fibers (Fig. 3) and those after grind-
ing (groundup samples, not shown in the present paper).27 Ac-
cordingly, the higher the I(002)/I(10) ratio in as-obtained BN fibers
(Fig. 3), the higher the average orientation degree o1. This ratio
theoretically increases from 0 for polycrystalline h-BN without
orientation effects of grains to 1 for perfectly oriented BN
A similar analysis was performed on fibers A-2 and the same
conclusions concerning the orientation of the (002) layers were
drawn. For example, Fig. 7 shows the BF TEM image of a lon-
gitudinal section of such fibers, which is similar to that of the BN
fibers A-3.
grains. For fibers A-2 and A-3, o1 is around 0.85 (high I(002)
/
I(10) ratio, Fig. 3), which suggests the same degree of orientation
of the (002) planes compared with the fiber axis, whereas fibers
A-1 display a poor orientation level (low I(002)/I(10) ratio, Fig. 3).
According to the modulus data given in Table I, the increased
orientation of the basal (002) layers from A-1 to A-3 is clearly
accompanied by a significant increase in the fiber modulus of the
corresponding fibers, as generally observed in BN fibers.27 As
the modulus E emerged stronger than the fiber strength s with
the applied fiber stress, it was observed that the failure strain e
(e 5 s/E) significantly decreased with the applied fiber stress.
The associated SAED pattern could consistently be indexed
as previously. SAED of fibers A-2 is also representative of an
ordered material. A good orientation can immediately be ob-
served in the low spread of the (002) spot (o2B0.9). Similarly,
the orientation of the (002) planes was consistent with the per-
fect superimposition of the BF and (002) DF images (not shown
in the present paper) taken at a higher magnification.
The BN fiber A-1 was analyzed in the same manner. Figure 8
shows the BF TEM micrograph in which the crystal growth is
clear as observed previously. This observation reflected XRD
data and highlighted the similarity in the crystallinity level
between the three types of BN fibers.
(3) TEM Studies
In order to confirm the structural/textural/mechanical proper-
ties’ relationship, the microstructure and the texture of BN fibers
were investigated by TEM. For example, a bright-field (BF) mi-
crograph of a longitudinal thin section of fibers A-3, in which
the corresponding SAED pattern is inserted, is shown in Fig. 4.
Consistent with SEM analysis, the BF TEM image of A-3 rep-
resents a coarse-grained microtexture indicating a well-ordered
and crystallized material. Data provided by SAED of such fibers
reflected the XRD results considerably. In particular, the SAED
pattern displayed rings of (002), (10), (004), (110), and (112)
planes of h-BN.29 Moreover, the diffracted (00l) rings were dis-
continuous as arcs, which demonstrated a preferred orientation
of these planes compared with the fiber axis.y
(002)
(10)
BN Fibers A-3
BN Fibers A-2
BN Fibers A-1
The estimation of the alignment of (002) layers in a direction
parallel to the fiber axis, as quoted above from the XRD data,
was also possible from SAED patterns by calculating the degree
of orientation o2. In fibers A-3, it was found to be around 0.9, in
the same order of magnitude as that obtained from XRD
20
25
30
35
40
45
yIn the case of randomly oriented planes of hexagons, the diffraction spot would be a
circle of regular intensity and an infinitely small point for a perfectly oriented crystal. It
should also be noted that there is a shift of angle between the image mode and the diffrac-
tion mode, resulting in apparent different directions between the image and the diffraction
pattern.
2 theta [°]
Fig. 3. X-ray diffraction patterns of as-obtained boron nitride (BN)
fibers, from A-1 to A-3.

46
Journal of the American Ceramic Society—Toutois et al.
Vol. 89, No. 1
Fig. 6. High-resolution transmission electron microscopy micrographs
of the longitudinal sectional microstructure in boron nitride fibers A-3.
microstructural and microtextural state of the BN single phase.
In particular, the fibers must undoubtedly report a high level of
crystallinity and microstructural ordering of the hexagonal po-
lymorph, while the average preferred orientation of the layers of
hexagons (002 planes) also appears to be of particular impor-
tance.26 Indeed, with the BN fibers investigated in the present
paper, quite independent of their origin, a well-established cor-
relation exists between the preferred orientation of the (002)
layers and Young’s modulus: the more highly oriented the (002)
layers, the higher the fiber modulus.
Fig. 4. Bright-field transmission electron microscope image and the as-
sociated selected-area electron diffraction of the longitudinal sectional
microstructure of boron nitride fibers A-3.
In contrast, some differences appeared in the associated
SAED pattern. Although the microstructural ordering of such
fibers seems to be similar to those of fibers A-3 and A-2, the
SAED pattern is composed of almost (002) circles, demonstrat-
ing that the (002) layers are randomly distributed compared with
the fiber axis. The low calculated degree of orientation o2B0.3
confirmed such observations.
(1) Shrinking Effects
The use of tension during the heat treatment of green fibers was
investigated as a possible means of increasing the mechanical
properties of final BN fibers.
BN fibers A-1, A-2, and A-3 were prepared from the same
batch of polymer using the same spinning parameters, and with
identical curing and pyrolysis procedure conditions (tempera-
ture, heating rate, and atmosphere). In order to evaluate the role
of the tension applied on the green fiber A because of the com-
IV. Discussion
Previous results demonstrated that the control of the green fiber
stress that occurs during the whole pyrolytic conversion is an
important requirement to maintain the original fiber morphol-
ogy and improve the mechanical properties of BN fibers. As
previously reported by our group, it was demonstrated that
room temperature strength and toughness were correlated to the
Fig. 5. Imaging of the preferred orientation of (002) planes along the fiber axis.

January 2006
Structural and Mechanical Behavior of Boron Nitride Fibers
47
The sintering mechanism by which green fibers provided ce-
ramic fibers could be explained based on the mass changes and
density increase data. The shrinkage, which occurred during the
curing and pyrolysis procedures of the green fibers A into ce-
ramic fibers A-(1-3) inducing loss of gaseous byproducts (me-
thylamine and ammonia16), occurred according to the following
equation:
As ! A ꢁ ð1 ꢁ 3Þs þ gaseous byproducts
(1)
The fiber shrinkage depends on changes in material mass and
density increase. In this context, taking into account the circular
cross-sectional morphology of fibers, the relative change in vol-
ume (DV) could be expressed in terms of the ceramic yield (a)
and the density (d) as follows:
100m0ðdc ꢁ dpÞ 100ðdc ꢁ adpÞ
DVð%Þ ¼
¼
(2)
dpdcVp
dc
with dc and dp being the density of the ceramic fibers (dc 5 1.926)
and that of green fibers (dp 5 1.2), respectively. m0 is the initial
mass of the polymer, and Vp and Vc are the polymer and ce-
ramic volumes, respectively.
Fig. 7. Bright-field transmission electron microscope image and the as-
sociated selected-area electron diffraction of the longitudinal sectional
microstructure of boron nitride fibers A-2.
According to the value of the final mass loss (Dm/
m0 5 46.1%, a 5 53.9%) after the curing and pyrolysis proce-
dures of the green fiber to form BN fibers, a shrinkage in volume
by 66% was estimated. Considering this estimated value, it is
shown that the empirical volume loss associated with the A-to-
A-1 conversion is only B58%. This lower value suggests that
the amount of stress to which the green fiber was subjected dur-
ing the pyrolytic conversion also affected the fiber density. It can
be estimated at 1.65 in the present case. It is therefore probable
that the fiber stress causes the internal pores (owing to the
spinning and pyrolytic conversion procedures) to collapse.
Therefore, the considerable reduction in length during the trans-
formation of the green fiber A into ceramic fibers A-1 occurred
to the detriment of that of diameter, and yielded crimped BN
fibers. In contrast, fibers A-2 and A-3 were maintained around
the spool either during a part (10001C, fibers A-2) or the whole
(18001C, fibers A-3) pyrolytic conversion, so that the longitudi-
nal shrinkage was significantly reduced, especially in the case of
fibers A-3. The latter were even slightly slackened around the
spool, resulting in an increased length. Such phenomena could
be related to the slight unwinding of the lubricating continuous
BN fiber around the spool at high temperature. Therefore, the
fiber crimping was limited, and most of the volume change was
associated with diameter loss. It was interesting to note that the
longitudinal shrinkage occurred in an ammonia atmosphere up
to 10001C as well as in a nitrogen atmosphere from 10001 to
18001C. The concomitant decrease in the practical strengths
with increased diameters showed that tension was an essential
requirement for enhancing the strength of BN fibers.
bination of the shrinking effects of the material and the use of a
spool during the pyrolytic conversion, two types of fibers were
cut at different stages of the transformation: (1) before the pyro-
lytic conversion leading to BN fibers A-1 and (2) after the ther-
mal treatment in an ammonia atmosphere at 10001C providing
BN fibers A-2. In contrast, BN fibers A-3 were produced
through the classical pyrolytic conversion of the preceramic
green fiber A wound on the spool.
The first consequence of these changes in tension was the in-
creased diameter of ceramic fibers, and the concomitant de-
creased length when the global longitudinal fiber stress
decreased. In particular, the green fiber A significantly shrank
in length owing to the large two-step weight loss (Dm/
mo 5 43.9%) and the density increase (Dd/d0 5 B30%26) that
occurred during the pyrolytic conversion up to 10001C in an
ammonia atmosphere. Upon heat treatment in a nitrogen at-
mosphere in the temperature range 10001–18001C, minor chang-
es in fiber lengths were observed, as the total ceramic yield was
53.9%, while a fiber density of 1.9 was generally obtained.26
(2) Effect of Tension on the Fiber Properties
SEM and XRD of the three fibers (see Fig. 2) reflected a good
level in the crystallinity and microstructural ordering. The pat-
terns displayed most of the peaks corresponding to h-BN, de-
spite collecting the (100) and (101) reflections in a single broad
(10) peak, which is characteristic of a two-dimensional ordering
as a turbostratic boron nitride (t-BN).31 Consequently, the phas-
es displayed only a partial three-dimensional order as a meso-
hexagonal BN phase.27 The comparatively high crystallinity was
reflected in the generally sharp full-width at half-maximum of
the diffraction peaks, an effect usually thought to be because of
a high crystallite size as calculated. The values were in the order
of 9 and 20 nm for Lc and La, respectively, which showed that
the BN fibers were relatively well-ordered, from a structural
point of view, and that the grain growth seemed to be inde-
pendent of the fiber stress during the pyrolytic conversion. In
particular, it is interesting to note that the crystallite stack height
Fig. 8. Bright-field transmission electron microscope image and the as-
sociated selected-area electron diffraction of the longitudinal sectional
microstructure of boron nitride fibers A-1.

48
Journal of the American Ceramic Society—Toutois et al.
Vol. 89, No. 1
and length do not increase for crystallites oriented at an in-
creased angle to the fiber axis.
was shown that tensioning was already effective below 10001C,
improving the degree of orientation of the basal layers along the
fiber axis, and therefore, the modulus in the final BN fibers. In
addition, tensioning was also effective in decreasing the fiber
diameter, and therefore increasing the strength of the resulting
final BN fibers produced at 18001C. Further studies are in
progress to quantify and optimize (1) the stress that has to be
applied on green fibers in order to prevent fiber breakage and (2)
the restricted temperature range within which the fibers must be
stressed.
XRD and TEM data highlighted the correlation between the
fiber modulus and the microstructural and microtextural prop-
erties. In particular, the present paper shows that the Young’s
modulus of BN fibers was closely related to the orientation of
the basal (002) sheets along the fiber axis. It was observed that
the fiber stress appears to be one procedure that may yield con-
siderably increased fiber modulus. This trend has been clearly
established in the case of carbon fibers,32 and could be logically
expected for BN fibers considering their structural similarity and
the high in-plane average B–N bond energy (D(B–N)B500.75
kJ/mol33). The Young’s modulus was significantly lower for BN
fibers A-1 than those determined for fibers A-2 and A-3, in ac-
cordance with the values found in the orientation degrees o1 and
o2. However, considering the identical values in modulus and
orientation degree for fibers A-2 and A-3, it was surprising to
note that the first step of the pyrolysis occurring in an ammonia
atmosphere up to 10001C under tension of the green fiber A was
an important requirement for increasing such values. These ob-
servations can be related to those in carbon fibers, which showed
that the application of tension during curing is important to
prevent polymer chains from relaxing and losing their orienta-
tion.34 In contrast, the applied tensioning was observed to be
essentially ineffective during the heat-treatment region from
10001 to 18001C, and in particular, during the crystallization
process (14001–18001C26), despite the importance of fiber stress
as discussed above. Such results are in contradiction with those
of carbon fibers for which a tensioning at high temperature
should be maintained to improve the mechanical properties of
fibers.34 However, it should be mentioned that carbon fibers are
derived from polyacrilonitrile, which displays a noncomparable
higher molecular weight than the poly[(methylamino)borazine].
Nevertheless, the reason for this ineffectiveness at very high
temperature has not yet been understood, but it can be viewed
that a higher mechanical tension should be applied during the
ceramic stage to improve orientation.
Acknowledgments
We wish to thank Marie-Paule Berthet (tensile tests) and Dr Fernand Chas-
sagneux (TEM investigations) for their contributions to sample characterization,
and their exciting and informative discussions.
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