# Impact of fiber surface gas treatments on interfacial bonding in SiC/BN/SiC minicomposites

Charlotte Chanson, Stéphane Jouannigot, Sylvain Jacques

► To cite this version:

Charlotte Chanson, Stéphane Jouannigot, Sylvain Jacques. Impact of fiber surface gas treatments on interfacial bonding in SiC/BN/SiC minicomposites. International Journal of Applied Ceramic Technology, 2023, 21 (4), pp.2747-2756. 10.1111/ijac.14638 . hal-04774574

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# Impact of fiber surface gas treatments on interfacial bonding in SiC/BN/SiC minicomposites

C. Chanson, S. Jouannigot and S. Jacques*

LCTS UMR 5801, CNRS - University of Bordeaux - Safran - CEA, 3 allée de la Boétie, 33600, Pessac, France

*corresponding author

## Abstract

BN interphases with different microstructures and a structural gradient were prepared by chemical vapor deposition (CVD) within SiC/SiC minicomposites. The interfacial shear stress (ISS) measured by fiber push-out test and the debonding location in the fiber/matrix interface zone during mechanical loading depend on the BN interphase structure. The ISS is higher when the structural organization of the BN interphase is poor, in which case debonding occurs on the fiber side. The Hi-Nicalon S fiber surfaces were gas treated to strengthen the bond with the interphase. $BCl_3$-based reactive CVD treatments have produced a thin boron carbide layer on the fiber surface, resulting in a 30% increase in ISS. However, the minicomposites are brittle in tensile tests due to fiber bridging by parasitic boron clusters. Short time CVD of SiC increases the fiber surface roughness and doubles the ISS, but again the minicomposites are brittle, indicating the need to better modulate the resulting roughness and adjust the interphase. Finally, ammonia treatments have nitrided the fiber surface and increased the ISS by 12%. While the increase in interfacial bonding measured at the microscopic scale was small, this nitriding treatment did not embrittle the minicomposites at the macroscopic scale.
# 1 INTRODUCTION

The ceramic matrix composites (CMCs) developed for aeronautics applications are made of silicon carbide-based fibers and matrices. The reinforcement used in these SiC/SiC composites consists mainly of Hi-Nicalon™ type S fibers (NGS Advanced Fibers Co., Ltd.). The SiC matrix is commonly produced in the woven fiber preform by chemical vapor infiltration (CVI), a method derived from low pressure chemical vapor deposition (CVD). The interphase, located between the reinforcing fiber and the matrix, allows the composite to withstand mechanical damage without brittle failure. It is also synthesized by CVD/CVI. Due to its layered crystal structure, BN has become the main interphase material used for aeronautics applications.¹

The microstructural organization of layered interphase materials such as pyrocarbon or boron nitride, has an effect on interfacial debonding mechanisms and on the micro- and macromechanical properties of CMCs.²⁻⁴ Therefore, it is essential to control this microstructural organization. However, using Hi-Nicalon S fibers also means controlling their coupling to the interphase, as the bond between these two components must remain strong enough for the interphase to fully play its role as a mechanical fuse. This may involve modifying the surface chemistry of the fiber with specific treatments prior to interphase deposition.

Reactive CVD (RCVD) converts the surface of a substrate by direct reaction between the substrate and gases. Baklanova *et al*.⁵ and Piquero *et al*.⁶ obtained uniform and adherent carbide thin coatings on the surface of carbon fibers. In the latter work, B₄C was produced by reacting the carbon of the fiber with a gaseous mixture of BCl₃ and H₂. In another work, Baklanova *et al*. had also converted the Nicalon fiber surface to TaC by a gaseous route.⁷ Nitriding gas treatment is another option for modifying the surface of SiC-based fibers. Rebillat *et al*. obtained silicon nitride on the surface of SiC-based fibers by
high temperature treatment in NH₃.⁸ Interfacial coupling in SiC/SiC composites also depends on fiber surface roughness. The radial compressive stress and consequently the debond or interfacial shear stress increase with fiber surface roughness due to mechanical interlocking.², ⁹

The first aim of this study was to tune the microstructure and organization of the BN interphase to control debonding in the fiber/matrix interface zone of SiC/SiC composites. The second was to use fiber surface treatments to promote chemical bonding or mechanical interlocking of the interphase. In the former case, a thin interlayer was produced consisting of elements common to both the SiC fiber and the BN interphase. In the latter case, the surface roughness of the fiber was increased. These treatments were carried out using gaseous routes. This allows the fiber surface treatment and the interphase and matrix deposition to be performed in sequence in the same CVD reactor. The samples used in this study were minicomposites. They are unidirectional model composites consisting of one fiber tow composed of 500 SiC-based fibers, the BN interphase and the SiC matrix.¹⁰

# 2 EXPERIMENTAL PROCEDURE

## 2.1 Synthesis

The fibers (Hi-Nicalon S) were desized during a process comprising an oxidation step followed by a treatment step with an acid to remove the silica layer formed during the oxidation step.¹¹ As in the work of Tobin *et al*,¹² the goal was to remove both the PVA sizing coating and the excess surface carbon. The equipment used to gas treat the fiber surface and deposit the BN interphase followed by the SiC matrix on fiber tows was the same as that used by Carminati *et al*.¹³ It was a hot-wall CVD reactor working at low pressure. The ends of 70 mm long fiber tows were bonded with carbon paste to a silicon carbide frame placed in the center of the CVD reactor.
The BN interphases were deposited from NH₃ and BCl₃ with H₂ as a diluent gas using the same CVD conditions as Carminati *et al.*¹³ and Plaisantin *et al.*¹⁴ Three types of interphases, each 300 nm thick, were synthesized. The first two, referred to as BN-1 and BN-2 interphases, were synthesized at a constant temperature of 900 °C and 1200 °C for 5 min and 30 min, respectively (FIGURE 1A and FIGURE 1B). The third type, referred to as BN-3, was prepared by varying the temperature during the deposition process, while maintaining an uninterrupted precursor gas flow in the CVD reactor (FIGURE 1C). The temperature was kept constant at 900 °C for several minutes, the CVD reactor was then heated-up rapidly up to 1200 °C and cooled-down to 900 °C and the temperature was finally maintained constant at 900 °C for a few minutes. The SiC matrix was then deposited at low pressure in the same CVD apparatus from a mixture of hydrogen and methyltrichlorosilane (MTS) at a temperature of about 1000 °C for 5 hours. The matrix volume fraction of the minicomposites was measured to be 0.60 by weighing without considering the porosity.

<figure><img src="image_2.png" /><figcaption>**FIGURE 1** Temperature vs. time in the CVD reactor during the preparation of BN-1 (A), BN-2 (B), and BN-3 (C) interphases</figcaption></figure>

A number of fibers were subjected to a surface gas treatment in the same CVD reactor prior to the deposition of the BN interphase. Three types of surface gas treatments were used. The first was a RCVD treatment. The fibers were reacted with a mixture of BCl₃, H₂ and Ar flowing at low pressure in the
reactor at 1200°C for 10 min with a H₂/BCl₃ flow ratio = 1.5. The second was a short time, low pressure
CVD deposition of SiC. The deposition was performed at 940°C using MTS and H₂ for only 1.5 min. The
third was a nitriding treatment using a mixture of NH₃, H₂ and Ar flowing at low pressure at 1300°C for
60 min.

**2.2 Characterization**

A specific table-model testing machine was used to perform tensile tests to determine the fiber
strength at room temperature. Fiber failure stresses were obtained from 30 tests per batch with a
gauge length of 25 mm according to a single fiber test procedure.¹⁵ Each fiber was mounted on a paper
frame; the paper frame was aligned in the grips and then the frame sides were cut. The fiber was
loaded up to failure. The stresses at failure were determined by dividing the load at failure by the
section of the fiber. For each batch of fibers, the given value was an arithmetic mean obtained from
the 30 tested specimens.

Untreated and treated fiber tows (without interphase and matrix) and SiC/BN/SiC minicomposites
were tensile tested at room temperature using a machine (MTS Systems, Synergie 200, USA) equipped
with a 500 N load cell. 5 to 7 specimens per batch were tested. The minicomposite ends were glued
with an epoxy resin (3M, 9353 A/B) in metallic tubes separated by 25 mm that were then gripped into
the testing machine jaws. The crosshead speed was 0.5 mm/min. The displacement was measured by
two inductive sensors (to check system alignment) with an accuracy of 0.1 µm.

The stress, $\sigma_d$, required to initiate fiber-matrix debonding and sliding and the interfacial shear stress (ISS), $\tau$, were determined from fiber push-out tests using the same procedure as described by Delpouve *et al.*⁴

The fracture surface and the fiber pull-out of the minicomposite specimens and gas treated fibers were
observed after the mechanical tests by scanning electron microscopy (SEM, Hitachi S4500 FEG) using
a secondary electron (SE) detector or a backscattered electron (BSE) detector as in the work of Delpouve *et al.*⁴ Chemical analyses were also performed in the microscope by Energy Dispersive X-ray Spectroscopy (EDS). Interphases of as-processed and *post-mortem* minicomposites were also observed by transmission electron microscopy (TEM, Philips, CM30ST) using bright field (BF), high resolution (HR) and selected area electron diffraction (SAED) techniques according to the methods described in the work of Plaisantin *et al.*¹⁴ In particular, the BN structural anisotropy was quantified by measuring the orientation angle (OA) in SAED patterns, i.e. the azimuthal opening of the arcs associated with the first-order reflection of the sp²-hybridized atom basal layers.

Chemical analyses of RCVD-treated fibers were performed by Auger electron spectroscopy (AES-VG Microlab 310F spectrometer) with simultaneous Ar⁺ etching in order to obtain the concentration depth profiles from the surface. X-ray photoelectron spectroscopy (XPS) was used to determine the composition and the chemical bonds in which nitrogen is involved at the surface of a fiber tow after the nitriding treatment. The conditions of the AES and XPS analyses were similar to those of Jouanny *et al.*¹⁶

# 3 RESULTS AND DISCUSSION

## 3.1 BN interphase microstructures and influence on mechanical behaviors

The HRTEM observations of interphases in minicomposites are similar to those of Plaisantin *et al.* for BN interphases deposited under the same conditions in single 2D woven SiC fibre plies.¹⁴ For BN-1 interphase deposited at 900 °C, the images show coherent domains consisting of a stack of a few distorted basal atomic and domains limited in size up to 5 nm (FIGURE 2A). The crystallographic organization is very poor and the microstructure is quasi-isotropic. The arcs of the SAED pattern are diffuse and appear to coalesce to form almost a ring (OA =100°). BN-2 deposited at 1200 °C displays, by contrast, a high structural organization with sp²-BN basal atomic layers oriented preferably parallel
to the fiber surface and well aligned with respect to each other (FIGURE 2B). Rectilinear fringes can reach 10 nm in length and the coherent domains expand in places over 15 nm along stacking c axis. The SAED pattern is typical of a turbostratic material with an overall structural anisotropy illustrated by very distinct arcs (OA = 30°).

<figure><img src="image_3.png" /><figcaption>**FIGURE 2** HRTEM images and SAED patterns of BN-1 (A) and BN-2 (B) interphases</figcaption></figure>

The BF TEM observation of the BN-3 interphase clearly shows the sequence of three sublayers of type BN-1, BN-2 and BN-1 (Figure 3). While the heating and cooling times between the two deposition temperatures are long relative to the dwell time at 1200 °C (Figure 1C), the transitions between the textures are abrupt and the interfaces between the sublayers are rather sharp. This can be explained by the strong temperature dependence of the BN growth rate, which varies by an order of magnitude between 900 °C and 1200 °C.
<figure><img src="image_4.png" /><figcaption>**FIGURE 3** BF TEM image of BN-3 interphase deposited in a minicomposite on a BCl₃ RCVD treated fiber</figcaption></figure>

FIGURE 4 displays a typical minicomposite monotonic tensile stress-strain curve for each kind of interphase. In each case, the initial linear elastic domain is followed by an extended nonlinear domain, characteristic of a damageable behavior with matrix microcracking and fiber/matrix debonding, as previously observed for Hi-Nicalon S fiber-reinforced minicomposites, whether with a pyrocarbon or a BN interphase.¹⁷, ¹⁸ The tensile curves of the fiber-reinforced minicomposites without surface gas treatment show the same shape. The average strains and stresses at the onset of matrix cracking and minicomposite failure are also similar (standard deviation for average strength around 100 MPa). Unload/reload tensile tests show no differences in hysteresis loop widths and residual strains (not shown here). The trend is the same for all these minicomposites: no influence of the interphase BN textures on the mechanical behavior is observed from these tensile tests.
<figure><img src="image_5.png" /><figcaption>**FIGURE 4** Typical tensile stress-strain curves of minicomposites with different BN interphases.</figcaption></figure>

Various debonding areas were observed by SEM (supplemental material S1) on fracture surfaces of the minicomposites after tensile tests. For the minicomposites with BN-1 interphase, the interphase is retained on the inner surface of the holes left in the matrix by the fiber pull-out while the pulled-out fiber surfaces are smooth and free of any coating. Therefore, the debonding occurs preferentially on the fiber side (S1 A and B). Conversely, in the second minicomposite batch, the BN-2 interphase is present on the pulled-out fiber surfaces (S1 C and D). In this case, the debonding occurs preferentially on the matrix side. Finally, for minicomposites with BN-3 interphase, interphase remains both on the fiber surfaces and on the matrix. Thus, the debonding occurs within the BN-3 interphase (S1 E and F).

HRTEM observation of this last minicomposite confirms this feature, a large debonding separates the well organized BN basal atomic layers in the middle of the interphase in the part deposited at 1200°C (FIGURE 5).
<figure><img src="image_6.png" /><figcaption>**FIGURE 5** HRTEM observation of debonding in the BN-3 interphase in a tensile-tested minicomposite.</figcaption></figure>

FIGURE 6 shows typical curves of single-fiber push-out tests for SiC/SiC minicomposites with BN-1, BN-2 and BN-3 interphases reinforced with fibers without gas surface treatment prior to interphase deposition. The values of debonding stress $\sigma_d$ and interfacial shear stress $\tau$ derived from the push-out tests are shown in TABLE 1. $\sigma_d$ and $\tau$ for BN-1 interphase are about four and three times higher than those for BN-2 interphase, respectively. Values for BN-3 interphase are intermediate.
<figure><img src="image_7.png" /><figcaption>**FIGURE 6** Typical single-fiber push-out curves for mini composites with the 3 different BN interphases and with fibers without gas surface treatment.</figcaption></figure>

<table><caption>**TABLE 1** Tensile behaviors and push-out test results for SiC/SiC minicomposites with different BN interphases and additional fiber surface gas treatments. Values in brackets () are standard deviations.</caption><thead><tr><th rowspan="2">gas treatment</th><th rowspan="2">interphase</th><th rowspan="2">Tensile behavior</th><th colspan="2">Push-out</th></tr><tr><th>σ<sub>d</sub> (MPa)</th><th>τ (MPa)</th></tr></thead><tbody><tr><td colspan="5"><strong>No</strong></td></tr><tr><td>No treatment</td><td>BN-1</td><td>damageable</td><td>3300 (400)</td><td>76 (9)</td></tr><tr><td>No treatment</td><td>BN-2</td><td>damageable</td><td>750 (170)</td><td>23 (8)</td></tr><tr><td>No treatment</td><td>BN-3</td><td>damageable</td><td>1790 (700)</td><td>51 (9)</td></tr><tr><td colspan="5"><strong>B<sub>4</sub>C RCVD</strong></td></tr><tr><td>BCl<sub>3</sub> + H<sub>2</sub> + Ar</td><td>BN-1</td><td>brittle</td><td>3000 (600)</td><td>105 (14)</td></tr><tr><td>BCl<sub>3</sub> + H<sub>2</sub> + Ar</td><td>BN-3</td><td>brittle</td><td></td><td></td></tr><tr><td colspan="5"><strong>SiC CVD</strong></td></tr><tr><td>MTS + H<sub>2</sub></td><td>BN-1</td><td>brittle</td><td>3400 (300)</td><td>150 (10)</td></tr><tr><td>MTS + H<sub>2</sub></td><td>BN-2</td><td>brittle</td><td></td><td></td></tr><tr><td colspan="5"><strong>Nitridation</strong></td></tr><tr><td>NH<sub>3</sub> + H<sub>2</sub> + Ar</td><td>BN-1</td><td>damageable</td><td>3000 (400)</td><td>85 (6)</td></tr></tbody></table>
The debonding locations observed after fiber push-out (FIGURE 7) are consistent with those previously observed on the fracture surfaces of tensile-tested minicomposites. They occur preferentially on the fiber side, matrix side, and within the interphase for interphases BN-1, BN-2, and BN-3, respectively. However, in contrast to the mechanical test on minicomposites, push-out tests reveal differences in interfacial bonding with the different BN interphases through $σ_d$ and $τ$ values. Buet *et al.* had also shown that fiber push-out tests are more effective in highlighting the influence of interphase texture than tensile tests with SiC/PyC/SiC minicomposites.⁹

<figure><img src="image_8.png" /><figcaption>**FIGURE 7** BSE SEM observations of minicomposite sample surfaces after fiber push tests on the front side (A to C) and back side (D to F), and diagrams showing debonding locations (G to I) for interphase BN-1 (A, D and G), interphase BN-2 (B, E and H) and interphase BN-3 (C, F and I). (Fibers without surface gas treatment.)</figcaption></figure>
In the case of single-ply model composites, a higher ISS measured by fiber push-out test correspond to better tensile test behavior, such as stiffness and onset of the first matrix cracks.⁴ The consistency between macroscopic and microscopic behavior at room temperature can be explained by the fact that a significant proportion of the interphases - those located on the transverse fibers of single-ply composites - are stressed perpendicular to the main direction of the interfaces, making them sensitive to stress. This is not the case for unidirectional minicomposites where all the fibers are in the direction of stress. However, previous work on unidirectional model composites reinforced with previous generations of SiC-based fibers had shown significant correlations between ISS and tensile behavior.¹⁹⁻²¹ This is explained by the differences in stiffness between these fibers and the SiC matrix (reaching a factor of 2 for Nicalon fibers), resulting in a local shear effect in the interphase sufficient to stress the interphase and its interfaces during tensile loading. The similarity in stiffness between the High Nicalon S fibers used here and the SiC matrix prevents this phenomenon. In this case, only fiber push-out tests are effective enough to measure the mechanical effects of the fiber-matrix bond.

Although the ISS measured by fiber push-out is improved with the BN-1 interphase, the best $\tau$ value of 76 MPa obtained here with non-gas treated fibers is still below 100 MPa. This means that the BN interphase does not fully meet the earlier concept of a strong fiber-matrix interface with cohesive failure achieved within the BN coating in the case of treated Nicalon fibers.¹⁹

## 3.2 Fiber surface gas treatments and effect on mechanical behaviors

The AES atomic concentration profile of a fiber surface treated with $BCl_3$ by RCVD shows a strong presence of boron associated with carbon in smaller proportions and the absence of silicon (FIGURE 8). This profile reflects the presence of boron carbide. Silicon has certainly been removed during RCVD in the form of chloride gas ($SiCl_2$, $SiCl_4$ and $SiCl_3H$).
<figure><img src="image_9.png" /><figcaption>**FIGURE 8** AES atomic concentration profiles of the surface of a BCl₃ RCVD treated fiber (Sputter rate: 10 nm/min; ref. Ta₂O₅)</figcaption></figure>

Cross-sectional SEM observations (not shown) coupled with EDS analysis reveal the presence of a boron carbide layer between 70 and 100 nm thick on the fiber surface. In the case of minicomposites with fibers that have undergone this RCVD treatment prior to the deposition of the BN-3 interphase, TEM observation also shows the presence of this B₄C layer with a thickness of the order of a hundred nanometers (FIGURE 3). The nature of this phase was verified using the reciprocal lattice obtained by Fourier transform of the lattice fringes obtained in the HR mode.

According to the fiber push-out results, this treatment increases τ by 30% with interphase BN-1 (TABLE 1). However, the behavior of minicomposites with RCDV-treated fibers is brittle, both with interphase BN1 and even with interphase BN2, although the latter gives a weak fiber-matrix bond without prior gas treatment of the fiber surface. Fracture surfaces (not shown) confirm the mechanical results: there is little or no fiber pullout.
Single fiber tensile tests show that RCVD-treated fibers suffer little damage under these conditions. Their average mechanical strength is 2200 MPa, compared to 2800 MPa for untreated reference fibers (TABLE 2). This does not account for the very poor behavior of the minicomposites. On the other hand, the mechanical strength of the only tow that could be tested after this gas treatment was only 90 MPa, compared to 1870 MPa for the reference fiber tows, a drop of over 95%. (Other tows with the same treatment could not be tensile tested because they were difficult to handle due to their low mechanical strength).

<table><caption>**TABLE 2** Mechanical strength of single fibers and fiber tows, untreated and treated, determined by tensile testing. Values in brackets () are standard deviations.</caption><thead><tr><th rowspan="2">gas treatment</th><th colspan="2">Tensile strength (MPa)</th></tr><tr><th>Fiber</th><th>Tow</th></tr></thead><tbody><tr><td>No treatment</td><td>2800 (600)</td><td>1870 (200)</td></tr><tr><td>B<sub>4</sub>C RCVD</td><td>2200 (500)</td><td>90</td></tr><tr><td>SiC CVD</td><td></td><td>1770 (300)</td></tr></tbody></table>

RCVD-treated fiber tows were observed by SEM (FIGURE 9). Within a tow, some fibers are bridged with other fibers. This bridging effect is due to the presence of solid clusters containing, according to EDS analysis, 85-90% boron, possibly bonded with silicon. Thus, if a fiber has a critical flaw, its premature failure leads to the failure of the neighboring fibers to which it is bridged and to the precipitated overloading of other unbridged fibers.²², ²³ These local catastrophic failures and load sharing induce here a dramatic drop in tow strength.

The presence of boron indicates that the reactive gases no longer react with the carbon of the fiber. Several hypotheses are considered. Hydrogen may infiltrate more easily than BCl₃ into confined spaces (tight fibers), favoring the presence of boron-rich phases. In the absence of carbon and silicon, H₂ reduces BCl₃ to form B, especially when H₂ is present in large amounts. Initially, a layer of B₄C may form
on the fiber surface as expected, acting as a diffusion barrier and isolating the fiber from the gas phase.

Carbon and silicon are then locally absent from the chemical reaction system, allowing boron to form.

Finally, these confined spaces can make it difficult to evacuate the gases produced by RCVD treatments after the thin boron carbide layer has been formed. This would favor the formation of solid boron from $BCl_3$.

These bridged fibers should not have been sampled for the single fiber tensile tests because they are difficult to separate manually. Consequently, the single fiber tensile results could not account for this phenomenon since they were based on fibers free of boron cluster deposits, the only ones that could be easily extracted individually from the tow.

In conclusion, these interfiber bridges, present prior to interphase and matrix synthesis, can explain the brittle fracture of RCVD-treated fiber-reinforced minicomposites.

<figure><img src="image_10.png" /><figcaption>**FIGURE 9** SE SEM observation of a $BCl_3$ RCVD treated fiber tow</figcaption></figure>

Short time CVD of SiC on the fiber surfaces has little effect on the mechanical strength of the tows (TABLE 2). With this treatment, $\tau$ is multiplied by 2, indicating very strong fiber-matrix bonding in SiC/BN-1/SiC minicomposites (TABLE 1). SiC CVD does not delay the onset of debonding ($\sigma_d$ values
~unchanged), but induces significant resistance to fiber sliding. However, minicomposites with both BN-1 and BN-2 interphases are brittle under tensile stress.

To mitigate the interfacial friction, which appears to be too strong, the thickness of the BN-2 interphase was increased by doubling the deposition time (10 min to obtain a thickness close to 500 nm). The mechanical behaviors of some minicomposites of the new batch is still brittle. Nevertheless, of the seven minicomposites studied, three showed damageable behavior. This scattering within the same batch highlights the difficulty of infiltrating the thick BN interphase at 1200°C into the core of the minicomposites. The same issue has been reported in single ply composites by Delpouve et al.⁴ This feature is highlighted by SEM observations which show, in addition to interface roughness, the heterogeneity of deposition at the surface and core of the minicomposite (FIGURE 10). The infiltration of the tow by a thicker film of BN-2 is made difficult by the very high growth rate at this temperature.

The interphase studied does not appear to be suitable for smoothing friction phenomena, but SiC CVD deposition on the fiber appears to be a promising approach. Work on interphase optimization (temperature gradient infiltration, pressure, BN-3 type multilayer architecture) would prevent brittle fracture. Work also needs to be done to optimize SiC deposition to better control roughness by using different conditions to reduce growth rate and better control CVD initiation with longer times.
<figure><img src="image_11.png" /><figcaption>**FIGURE 10** SE SEM observations of minicomposites with short-time SiC CVD treatment on fibers and the BN2 interphase deposited for 10 min, core (A), surface (B).</figcaption></figure>

After the nitriding treatment, XPS analysis show that nitrogen is present at the extreme surface of the fibers with a content of 20% (FIGURE 11). It is mainly bound to silicon. The carbon content is lower than that of silicon, indicating that it has been partially consumed. The nitrogen content drops below 5% after etching between 50 and 100 nm depth (not shown).

<figure><img src="image_12.png" /><figcaption>**FIGURE 11** XPS analysis fiber tow surface after nitriding treatment (Etch Time = 2 s) – Chemical composition and N<sub>1s</sub> deconvoluted peak</figcaption></figure>
The mechanical behavior of the minicomposites with nitrided fibers and BN-1 interphase is very close to that of the reference without fiber surface gas treatment (FIGURE 4). Given the standard deviations of ~100 MPa, the average mechanical strength of the nitrided and reference batches can be considered similar. However, the shape of the tensile curves shows a more progressive transition between the linear and non-linear domains, with slightly higher stress levels for the same strain before ultimate failure. This may reflect a slightly stronger fiber-matrix bonding.

The nitriding treatment does not delay the onset of decohesion, since the $\sigma_d$ values obtained by fiber push-out are practically similar. Nevertheless, the average value of $\tau$ for the batch of nitrided fiber minicomposites is slightly higher than that of the reference, reaching 85 MPa (TABLE 1).

## CONCLUSION

In contrast to micromechanical fiber push-out tests, tensile tests on minicomposites allow only a simple distinction between brittle and damageable mechanical behavior, but they do not really account for the intensity of fiber-matrix interfacial bonding and its effect on the macromechanical behavior of the minicomposites. The minicomposite with the latest generation of SiC-based fibers, which are as stiff as the matrix, is probably too simplistic a model composite to characterize in detail the room temperature mechanical behavior of its interfaces and interphases at the macroscopic scale.

By adjusting the structural organization of the BN interphase and its gradient (via constant or varying temperatures from 900 to 1200 °C during the CVD process), it is possible to control the ISS and the location of debonding: inside, outside or within the interphase. The best $\tau$ value is obtained with the poorly organized interphase deposited at the lowest constant temperature but is remains below 100 MPa with non-gas treated fibers.
Gas treatments of fiber surfaces all result in an increase in fiber-matrix interfacial bonding. BCl₃-based reactive CVD and short time CVD of SiC treatments lead to brittle minicomposite. While the former treatment, with parasitic boron cluster deposition bridging the fibers, appears to be doomed to failure, the latter could be addressed by limiting the increase in fiber surface roughness and adjusting the BN interphase. Ammonia-based nitriding treatment was the only treatment capable of increasing the ISS without embrittling the macroscopic mechanical behavior of the minicomposites. However, since the increase in ISS obtained was modest, this treatment still needs to be improved.

## ACKNOWLEDGEMENTS

This work was supported by a grant from SAFRAN CERAMICS to C. CHANSON. The authors thank L. Lapuyade from LCTS and M. Lahaye and C. Labrugère from PLACAMAT (UMS 3626 CNRS University of Bordeaux) for AES and XPS analyses and data processing. The authors are indebted to A. Desenfant, who contributed to a number of sample treatments and characterizations. The authors pay tribute to the late P. Weisbecker, who participated in the TEM characterization at LCTS. They also pay tribute to the late E. Buet and thank his colleagues S. Loison and A. Delcamp from SAFRAN CERAMICS for their fruitful discussions.

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