The high-temperature corrosion behavior of SiBCN fibers for

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The high-temperature corrosion behavior of SiBCN fibers for aerospace applications Xiaoyu Ji, Shanshan Wang, Changwei Shao, and Hao Wang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b04497 • Publication Date (Web): 16 May 2018 Downloaded from http://pubs.acs.org on May 16, 2018

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ACS Applied Materials & Interfaces

The High-Temperature Corrosion Behavior of SiBCN Fibers for Aerospace Applications Xiaoyu Ji, Shanshan Wang, Changwei Shao*, Hao Wang*

Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, PR China

KEY WORDS: SiBCN fiber, oxidation resistivity, corrosion, combustion, aerospace applications, mechanism

ABSTRACT

Amorphous SiBCN fibers possessing superior stability against oxidation have become a desirable candidate for the high-temperature aerospace applications. Currently, investigations on the high-temperature corrosion behavior of these fibers for the application in high-heat engines are insufficient. Here, our polymer-derived SiBCN fibers were corroded at 1400 °C in air and simulated combustion environments, respectively. The fibers’ structural evolution after corrosion in two different conditions and the potential mechanisms are investigated. It shows that the as-prepared SiBCN fibers mainly consist of amorphous networks of SiN3C, SiN4, B-N hexatomic rings, free carbon clusters and BN2C units. HRTEM cross-section observations combined with EDS/EELS analysis exhibit a tri-layer structure with no detectable cracks for fibers after corrosion, including the outermost SiO2 layer, the h-BN grain-contained interlayer, and the uncorroded fiber core. A high percentage of water vapor contained in the simulated combustion environment triggers the formation of abundant α-cristobalite nanoparticles dispersing in the amorphous SiO2 phase, which are absent in fibers corroded in air. The formation of h-BN grains in the interlayer could be ascribed to the sacrificial effects of free carbon clusters, Si-C and Si-N units reacting with oxygen diffusing inward, which protects ACS Paragon Plus Environment

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h-BN grains formed by networks of B-N hexatomic rings in original SiBCN fibers. These results improve our understanding of the corrosion process of SiBCN fibers in high temperature oxygen- and water-rich atmosphere.

1. Introduction

Humankind’s aerospace aspirations are placing increased demand for highly-efficient and robust aircraft and spacecraft propulsion systems, which can function properly in more severe operating conditions, such as high temperature and oxidative environments. In this context, the availability of high-temperature structural materials plays a key role in improving the efficiency of energy transformation in turbines and combustors.1-3 Metallic materials, such as superalloys, are one of the primary choices for gas and jet turbine engines. However, they are limited by the intrinsic weakness in the resistance to creep deformation and corrosion at elevated operating temperatures expected in the future aerospace transportation.4-6 Fortunately, ceramic materials, especially fiber reinforced ceramic matrix composites, have a great potential to fulfill such demand. They are gradually being employed in new-generation high-heat engines.1 The thermal stability advantage of advanced structural ceramics over superalloys could reduce fuel consumption by 6 to 8% in aircraft engines and by 10 to 15% in gas turbines.6

SiC fiber reinforced SiC matrix composites (SiCf/SiC CMCs), as a representative high-performance ceramic system, have been significantly invested, intensively researched and successfully used in hot-section components of high-heat engines.7 Acceptable high-temperature endurance of CMCs is highly dependent on the judicious selection of ceramic fiber reinforcement with proper chemical, physical and mechanical properties. Compared to crystalline SiC fibers,8 SiBCN fibers composed of silicon, boron, carbon, and nitrogen inorganic networks perform superior high-temperature creep resistance, which can be attributed to a grain boundary-free amorphous structure and are considered as promising candidates for the application under mechanical loading conditions in high-heat engines.9 Moreover, due to the presence of BN(C) layers ACS Paragon Plus Environment

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which could encapsulate the Si3N4 nanodomains and prevent carbothermal reduction, SiBCN bulk ceramics and fibers perform unprecedented high-temperature stability.6, 10-12

However, detailed investigations of high-temperature corrosion behaviors of SiBCN fibers under combustion and air environments are still lacking, which is of great importance for material’s realistic application in aircraft and spacecraft propulsion systems. Although the remarkable oxidation resistance of the SiBCN bulk ceramics has been confirmed by several reports,13-14 to date, only two repots focus on the stability against oxidation of SiBCN fibers and provide discrepant results. One report suggested that SiBCN fibers exhibited outstanding oxidation resistance due to a thin, crack-free, amorphous SiO2/BN two-layer structure.6 In contrast, using the fibers produced in the same year, another report indicated the formation of a more complex cristobalite/SiO2/SiBCNO (BN) tri-layer scale15 and the studies on bulk SiBCN ceramics also pointed out the absence of the SiO2/BN double layer.14,

16

Moreover, the study on high-temperature

corrosion behavior of SiBCN fibers in water-contained environment, which is closer to the real combustion in high-heat engines, is nearly blank.

In this work, the oxidization resistance and corrosion behavior of SiBCN fibers in simulated combustion environment is reported for the first time. By means of FIB/SEM-Slice and TEM/EELS techniques, the complex layered structures radially from surface to core of corroded fibers were captured and analyzed on a nanometer scale. Meanwhile, the oxidization behavior and structure evolution under air atmosphere were also investigated in detail.

2. Experimental

2.1 Preparation of SiBCN fibers

The SiBCN fibers were fabricated via polymer-derived route. The polyborosilazane precursor was synthesised

by

reacting

boron

trichloride

and

dichloromethylsilane

simultaneously

with

hexamethyldisilazane in one pot.17 Detailed information of the precursor was shown in Supporting ACS Paragon Plus Environment


Information. Green fibers were then prepared using a lab-scale melt-spinning system. Electron beam irradiation was employed to render the as-spun fibers infusible for the first time.18 The continuous green fibers with graphite spool were set on a rotating stain-less-steel cylinder in a helium filled irradiation chamber. The electron beam was then generated under the condition of a 2 MeV acceleration voltage, 2-15 kGy/s dose rate and 17 MGy dose. Finally, the cured fibers were converted into ceramic fibers by pyrolysis at 1300 °C and sintering at 1500 °C in a nitrogen atmosphere. The chemical composition of the SiBCN fibers was shown in Table S1 and the calculated formula was SiB0.9C1.6N2.4.

2.2 Corrosion of SiBCN fibers in air

The as-received fibers were cut into ~ 5 cm in length and put into an alumina crucible. The alumina crucible with the fibers was then placed in a muffle furnace and heated in static laboratory air. The temperature was progressively increased to 1400 °C at a rate of 5 °C/min and was then held at this temperature for two hours. Finally, the specimens were cooled from 1400 °C to ambient temperature in a natural process and named as ACF (air-corroded fiber).

2.3 Corrosion of SiBCN fibers in simulated combustion environment

A high temperature tubular furnace associated with a steam generator was fabricated to simulate the combustion environment. The dry N2-O2 mixture gas flows through a heated water device in order to be saturated in steam before its introduction into the alumina tube of the furnace. The accurate control of the gas ratio was realized by adjusting the amount and the temperature of water and the flow rate of the gas mixture, which could be monitored using the water vapor detector and the oxygen analyser. Calculation has shown that 5-10% of water vapor is produced under an equilibrium condition in the combustion process. In the current study, a gas mixture with PH2O: PO2: PN2 = 14: 8: 78 were employed based on previous studies.19-21 In this corrosion test, the same SiBCN fiber specimens were firstly heated in N2 atmosphere with a ramping rate of 5 °C/min to 1400 °C. Subsequently, the corrosive gases instead of N2 were introduced to the furnace ACS Paragon Plus Environment

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at a flow rate of 200 ml/min and the temperature was held at 1400 °C for two hours. The specimens corroded in this combustion environment was named as CCF (combustion-corroded fiber).

2.4 Characterization

The Auger electron spectroscopy (AES) depth profile of the fiber surface was conducted using an Auger electron spectroscope (ULVAC PHI-700, Japan) coupled with Ar+ etching. The fibers were milled into powders for XPS, Raman and MAS NMR experiments. The X-ray photoelectron spectroscopy (XPS) measurements were recorded using Al Kα as the excitation source (Thermo Scientific Escalab 250Xi, USA) to detect the surface compositions and chemical bonding states within a few atomic layers. The Raman spectrum was obtained using JY HR-800 Raman microscope (Horiba, France). The excitation source used was the 514 nm line of an Ar-ion laser with a beam size of ~ 2 µm. Solid state 29Si, 11B and 13C MAS NMR experiments were carried out on a JNM-ECZ600R spectrometer (Jeol, Japan) by spinning powdered samples packed into zirconia rotors at the field strength of 14.1T (600 MHz) and the max frequency of 15 kHz. The relaxation delays of the

29

Si,

11

B and

13

C MAS NMR are 5s, 3s and 3s, respectively. For

transmission electron microscopy (TEM) studies of the original SiBCN fibers, powdered samples were dispersed in alcohol using 10 min ultrasonication, and a drop of the dispersed solution was deposited on a 3 mm carbon-coated copper grid. High resolution TEM (HRTEM) images were collected using a Titan G2 60-300 (FEI, USA) instrument with an accelerating voltage of 300 kV with finely pulverized samples.

The morphologies of SiBCN fibers before and after corroded in air and simulated combustion environment were examined with a scanning electron microscope (SEM; FEI Helios Nanolab 600i, Japan). Cross section of the corroded fibers was sliced into thin films by focused ion beam (FIB; Zeiss Auriga, Germany). A rough FIB milling in a current range of 1-20 nA with a beam energy of 30 keV was used to obtain the thin slice containing the whole fiber cross section. Partial cross-section face was then polished with a lower beam current in the range of 300-1000 pA. The FIB cross-section face of the corroded fibers was further ACS Paragon Plus Environment


characterized by transmission electron microscopy (TEM; Tecnai F20st, USA), operating at 200kV. The elemental analysis of the cross section was carried out using an energy dispersive spectrometer (EDS; Hitachi S-4800, Japan). The local chemical compositions radially from the surface to the core were examined by electron energy-loss spectroscopy (EELS; Gatan GIF Quantum 965). Silicon L-edge, boron K-edge, carbon K-edge, nitrogen K-edge and oxygen K-edge were considered.

3. Results and discussion

3.1 Composition and microstructure of the SiBCN fibers

Figure 1．(a＆b) SEM images showing the as-synthesized SiBCN fibers. Inset of panel b is the EDS elemental mapping of Si, B, N, and C taken from the region marked in the red rectangle box in panel b. (c) HRTEM images of the as-synthesized SiBCN fibers. Insets of panel c are the SEAD pattern and the magnified image taken from the region marked in the red square box in panel c. (d) AES depth profile showing the concentration change of elements in the surface region of the fiber as a function of the sputtering depth. ACS Paragon Plus Environment

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As shown in Figure 1a＆b, the fibers exhibit densified structure with smooth surfaces and uniform diameter of ~ 11.5 µm. HRTEM is also employed to further study the microstructure of the SiBCN fibers (Figure 1c). Both the HRTEM images and the inset SEAD pattern illustrate that the as-received fibers are amorphous on a nanometer scale. The elemental mappings of the fiber cross section by EDS illustrate that silicon (Si), boron (B), carbon (C), and nitrogen (N) have an approximately homogeneous distribution from the surface to the core. Interestingly, the detailed composition within a very short distance of the fiber surface varies. Figure 1d is the AES depth profile showing the atomic concentration variations of five primary elements of the fiber, including C, Si, N, B and O, as a function of the drilling depth by the Ar ion. Distinct from the cases of Si, N, and B, whose atomic concentrations increase steadily within ~ 40 nm from the fiber surface followed by a level-off, the concentration of C reaches its maximum at 8 nm with a value of 89.1 at % followed by a sharp decrease to 33.2 at % when the etching deepens down to ~ 40 nm. It indicates the presence of a carbon-rich layer less than 40 nm at the fiber surface, which might derive from the deposition of pyrolytic carbon produced by CH4 released in the pyrolysis process.22 The content of O shows a slow decrease as the depth enhances. At the depth of 160 nm, the concentrations of Si, B, C, N, and O cease to change further and stabilizes.

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Figure 2. (a-d) XPS fine spectra showing the chemical bonding of Si2p, B1s, N1s and C1s, respectively, at the depth of 200nm from the SiBCN fiber surface. (e) Raman spectrum showing the structure of free carbon in as-synthesized SiBCN fibers. (f-h) MAS NMR spectra showing the

29

Si,

11

B and

13

C MAS NMR of

as-synthesized SiBCN fibers, respectively.

The chemical bonding of the SiBCN fibers was analyzed by XPS. As shown in Figure 2a, the Si2p signal can be deconvoluted into two peaks. The stronger one at 101.7 eV corresponds to the Si-N bond, while the weaker one at 100.6 eV is assigned to the Si-C bond. The B1s peak (Figure 2b) can also be fitted into a strong peak at 190.6 eV and a weak one at 189.0 eV, which correlate to two types of boron atoms bonded to nitrogen and carbon atoms, respectively. The N1s peak is associated with N-Si bonds locating at 397.8 eV and N-B bonds sitting at 398.4 eV, respectively (Figure 2c). The fitting results indicate that Si, N and B ACS Paragon Plus Environment

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atoms are mainly connected in the form of Si-N bonds, B-N bonds and a small amount of B-C and Si-C units with a mole ratio of 1.0: 1.9: 0.6: 0.2 calculated by the integration of the peak area. For C1s (Figure 2d), two weak peaks at 283.0 and 285.5 eV arise from Si-C bonds and sp3 hybridized C-C bonds, respectively, whereas the strongest peak at 284.3 eV can be attributed to C-B bonds or sp2 hybridized C-C bonds. However, by means of XPS, it is difficult to determine whether the graphite-like carbon exist due to the overlap of B-C signal and sp2 hybridized C-C signal.23-26

Further characterization was carried out to confirm the structure of carbon in SiBCN fibers using Raman spectroscopy, which is a sensitive nondestructive tool to monitor the structure and bonding manner of carbon atoms.27-28 As shown in Figure 2e, two Raman bands were observed at ∼ 1350 and 1582 cm-1. The first one can be assigned to the defect activated D band, representing a breathing mode of sp2 carbon rings adjacent to graphite edges or defects, while the second peak is associated with the in-plane vibration of sp2 hybridized carbon atoms, denoted as G band. The presence of G band suggests the existence of graphitic structures, while the broad and strong D bond indicates a large amount of disordered structure of carbon atoms with diversified defective structures. The disorder and defects in the carbon phase could be primary induced by the presence of edges in the graphite layers, by the deviation from planarity of graphite layers and also by the presence of carbon atoms in sp3 hybridization state.29 The intensity ratio of the D and G modes, ID / IG, enables the evaluation of the carbon-cluster size by using the formula reported by Ferrari and Robertson.30 It can be expressed as ID / IG = C’(λ) La2, where La is the size of carbon domains along the six fold ring plane (lateral size), and C’ is a coefficient that depends on the excitation wavelength of the laser. The value of the coefficient C’ for the wavelength of 514.5 nm of the Ar-ion laser employed here is 0.0055 Å−2. Lorentz curve fitting of the D band and Breit-Wigner-Fano (BWF) curve fitting of the G band were performed in order to extract the ID / IG intensity ratios and the size of the free carbon cluster formed in the ceramics is determined as 1.71 nm.

The disordered free carbon clusters on a nanometer scale are fixed in

the amorphous Si-B-C-N network with linkages of Si-C and B-C, which could be much less susceptible to ACS Paragon Plus Environment


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oxidation.24

The structure of the SiBCN fibers could be eventually verified by MAS NMR experiments. As shown in Figure 2f, the 29Si MAS NMR signal is characterized by overlapping resonances without any fine structure, which is centered at -44.5ppm with a FWHM of ~ 42 ppm. Individual NMR signals arising from distinct structural units thus cannot be resolved. The appearance of such significantly broadened signal is caused by overlapping resonances of SiCN3, SiN4 and SiN3O sites, which typically range from -15 to -35 ppm, -35 to -55 ppm and -55 to -70 ppm.31-32 The weak SiN3O signal illustrates that the trace of O in fibers exists in the form of Si-O-Si units. The

11

B MAS spectrum (Figure 2g) shows a sharp isotropic chemical shift value of

24.5 ppm. In accordance with the

11

B MAS NMR spectra of known SiBCN ceramics and h-BN, the data

confirms that boron is trigonally planar coordinated by three nitrogen atoms or in the form of B-N hexatomic rings.33-34 A weak deconvoluted signal at ~ 35.1 ppm indicates the presence of BN2C coordination sphere, which is in satisfactory accordance with the result of XPS fine spectrum of B1s. As shown in Figure 2h, the carbon is found to exist in two main fractions. The broad spectral component ranging from 150 to 100 ppm is strongest at 130 ppm, which can be assigned to sp2 hybridized free carbon.35 In addition, a shoulder peak locating in a range between 186 and 167 ppm might be considered as partial free carbon bonding with the boron atom of B-N hexatomic rings. The other main resonance peak at 10-50 ppm was attributed to the unit of CSix (x=1 or 2) units in a Si-C-N matrix.36 Consequently, the as-received SiBCN fibers consist of the amorphous network of silicon, boron, carbon and nitrogen, which mainly exist in the form of SiN3C, SiN4, BN3 (B-N hexatomic rings) units and free carbon clusters, accompanied with a small number of BN2C and Si-O-Si units.

3.2 Corrosion behavior in air and combustion environments.


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Figure 3. SEM images showing the surface (a＆b) and cross section morphologies (c＆d) of SiBCN fibers annealing at 1400 °C for 2 h in air.

Figure 3 exhibits the morphologies of the SiBCN fiber after annealing at 1400 °C for 2 h in air (ACF). It shows a slight increase in diameter from ~ 11.5 µm (unoxidized fiber) to ~ 12.1 µm, which could be ascribed to the absorption of O2 and the formation of oxide layer. The surface morphology of the ACF are shown in Figure 3 a＆b, where is smooth and free of pores and cracks. As shown in Figure 3c＆d, the cross section of the ACF performs complicated microstructure which could be divided into at least two layers radially from outside to inside. The interfaces among these layers show good adhesion to each other and no cracks are observed on the cross section of the fiber.



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Figure 4. (a) TEM images showing the cross section of the fibers annealing at 1400 °C in the static air for 2h. (b-d) High magnification TEM images corresponding to the zones marked by B, C and D in (a), respectively. Insets of panel c and d are SEAD patterns of the region marked in the red circles. (e-g) HRTEM images taken from the glass-like area in (b), (c) and the amorphous area in (d). Insets of panel e, f and g are their corresponding SEAD and FFT patterns, respectively. (h) EDS spectra (I), (II) and (III) corresponding to the zones marked by EDS-I, II and III in (a), respectively. (i) EELS spectra with the distance of 0.25, 0.75, 1.50 and 2.50 µm from the fiber surface in (a).

To further understand the detailed composition and structure of ACF, FIB/SEM-Slice was employed to slice the oxidized fiber into a thin film and the as-sliced fiber cross section was systematically characterized by TEM, EDS and EELS analysis. Figure 4a. exhibits an integral TEM image of ACF containing three distinct zones with different features: (i) the outermost layer of glass-like structure; (ii) the intermediate layer with ACS Paragon Plus Environment

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dispersive nanocrystals; and (iii) the innermost homogeneous and amorphous fiber core. All the three zones are dense and intact with no clear boundaries.

As shown in Figure 4b, the fiber surface is a glass-like phase with approximately 470 nm thick. EDS-I spectrum (Figure 4h) suggests that this layer consists of only Si and O elements, containing no detectable amounts of B, C and N. The HRTEM image and SAED analysis (Figure 4e) fails to locate any crystalline phases, which suggests that the surface of the oxidized fiber is the amorphous silicon oxide. Beneath this layer is a crystallized interlayer with approximately 1.25 µm thick, where both B and N are detected by EDS-II, with trace of C as well (Figure 4h). As shown in Figure 4c and Figure 4d, crystal particles of ≤30 nm in size are spread all over an amorphous matrix. Based on the HRTEM image of crystal particle in Figure 4c and the corresponding FFT data (Figure 4f), the lattice fringes with a spacing of 0.33 nm correspond to the (002) lattice plane of h-BN, which implies that the particle could be identified as h-BN nanocrystal. The polycrystal SAED patterns in Figure 4c and Figure 4d also provides evidence for the formation of h-BN nanocrystals. In addition, a gradient distribution in size of the nanocrystals across the intermediate layer can be observed, which varies from 20-30 nanometers at the top of the interlayer to a couple of nanometers close to the indistinct boundary with the inner layer. With the depth further increasing, Figure 4d and Figure 4g exhibit the boundary between the crystallization interlayer and the inner amorphous core and the HRTEM image with SAED pattern of the amorphous area, where no crystalline phases have been detected. Compared with the oxidized crystallization layer, C content obviously increases, whereas O content decreases dramatically to 6.88 at% in this area (Figure 4h EDS-III), which is similar to the AES result of the as-received SiBCN fibers. Therefore, the third layer from the surface could be deduced as unoxidized SiBCN fiber core.

Due to the difficulties to differentiate the light elements by means of EDS, EELS, which is more sensitive to B, C and N,37 was also employed to further judge the composition and structure variation of SiBCN fibers. EELS data measured on materials with distances of 0.25, 0.75, 1.50 and 2.50 µm from the surface, ACS Paragon Plus Environment


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respectively, are showed in Figure 4i. For the material locating within 0.25 µm from the surface (in the outermost layer), the silicon L-edge (Si-L2,3) at 107 eV and oxygen K-edge (O-K) at 532 eV are detected. The two sharp peaks labelled as Si-L2 and Si-L3 and the following strong broad peak are the characteristic peaks of SiO2.38 For the material within the range of 0.75 and 1.50 µm from the surface, respectively (the crystallized interlayer), the EELS clearly show boron K-edge (B-K) at 188 eV and nitrogen K-edge (N-K) at 401 eV. Each of these edges first shows a sharp pre-peak due to the transition from 1s to π*, followed by a second one, due to the transition from 1s to σ* (Figure 4i). This fine structure of the B-K and N-K presents the classical features that have been reported for h-BN or BN compounds involving sp2 bonding.39-41 It is consistent with the result of HRTEM image and demonstrates that the grains dispersed in the interlayer are h-BN nanocrystals. In addition, radially from the surface to the core in this layer, the matrix that h-BN crystalline grains precipitate in contains varying compositions. Compared to the material in the position of 0.75 µm (below fully oxidized silica layer), the material in the position of 1.50 µm (closest to the amorphous core layer) shows a weaker intensity of Si-L2 edges and O-K edge, which indicates a decrease of the SiO2 phase. In addition, C-K edge is not observed at 0.75 µm but shows obviously π* and σ* edge at 1.50 µm, which illustrate the existence of sp2-bonded C atoms in the position closest to the amorphous core layer. When the position move closer to the amorphous core layer, the π* and σ* N-K edges are not as obvious as those of the position below fully oxidized silica layer. It could be explained that the N content increases and exist not only in the form of h-BN, but also in the amorphous SiBCN(O) matrix. The EELS of the innermost core (Figure 4d) shows the disappearance of O-K edge. Only a single Si-L edge and a broad single N-K edge were present, which could be judged as the structure of amorphous SiBCN ceramics.42


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Figure 5. SEM images showing the surface (a＆b) and cross section morphologies (c＆d) of SiBCN fibers annealing at 1400 °C for 2 h in simulated combustion environment.

In the combustion process, substantial amounts of water vapor are produced by burning hydrocarbon fuels in air. Therefore, combustion environment is simulated with the partial pressure ratio of H2O, O2 and N2 being 14%:8%:78%, which has been applied to simulated combustion environment to corrode SiC fibers.19-20 Figure 5 exhibits the morphologies of the SiBCN fiber after annealing at 1400 °C for 2 h in a simulated combustion environment (CCF). Similar morphologies are observed in SEM graphs with ACF: smooth surface with no detectable pores and cracks and complex cross section which exhibits multilayer oxide structure with good adhesion to each other.



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Figure 6. (a) TEM images showing the cross section of the fibers annealing at 1400 °C in simulated combustion environment for 2h. (b-d) High magnification TEM images corresponding to the zones marked by B, C and D in (a), respectively. (e-g) HRTEM images taken from regions marked in the red rectangle box in panel b, c and d, respectively. Insets of panel e, f and g are their corresponding SEAD and FFT patterns, respectively. (h) EDS spectra (I), (II), (III) and (IV) corresponding to the zones marked by EDS-I, II, III and IV in (a), respectively.

Systematical composition and structure analysis of CCF is also carried out. As shown in Figure 6a-d, CCF exhibits three distinct zones: two-layer oxide scale and unoxidized fiber core, which are similar to ACF. EDS analysis of the outermost layer shows that only Si and O elements exist in this area (EDS-I shown in Figure 6h), which illustrate that this layer mainly consist of SiO2. However, as marked in Figure 6b, the ACS Paragon Plus Environment

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thickness of outmost layer of CCF is ~ 160 nm, which is 300 nm thinner than that of the ACF (~ 470 nm). Moreover, abundant freckle-like particles spread around the outermost SiO2 layer and extend to the exterior of the crystallized interlayer. The HRTEM image of the outermost layer (Figure 6e) exhibits that the size of these particles is less than 10 nm and the inset FFT data provides evidence for the lattice fringes with a spacing of 0.23 nm, which can be assigned to the (201) lattice plane of α-cristobalites. It reveals that crystallization of the amorphous SiO2 occurs when corroded in the combustion environment.

Beneath the cristobalite-contained silica layer, the variation of elements performs the same tendency with that of ACF. EDS-II, III and IV indicate that N, B, and C are detected in the corresponding areas and, with the depth increasing from EDS-II to EDS-IV, the contents of N, B, and C rise whereas the O content reduces. However, the intermediate layer of CCF exhibits more complex structure than that of ACF. As shown in Figure 6a＆c, freckle-like α-cristobalites are still dispersed in the upper position of this layer. However, with the depth increasing, the size of α-cristobalites starts to reduce and disappear. Figure 6f, the HRTEM image taken in the red-marked area of Figure 6c, shows two forms of crystal structure. The inset SAED pattern exhibits (201), (113) and (332) lattice planes of α-cristobalites (labeled as 2, 3 and 5) and the (002) and (104) lattice planes of h-BN (labeled as 1 and 4), which verify the existence of both α-cristobalites and h-BN. As shown in the HRTEM image, the size of α-cristobalites has decreased to 1-2 nm and tends to disappear, whereas the h-BN nanocrystals still exist. In addition, the size of h-BN nanocrystals at the top of the interlayer is 10-15 nm, which is smaller than those of ACF (20-30 nm). Figure 6d and g exhibits an obscure boundary of crystallized and amorphous section. Near the line, the h-BN are turbostratic and the size is restricted to 1-2 nm. Beneath the line, the amorphous area can be considered as the uncorroded fiber core, which is consistent with the results of SAED pattern in Figure 6a and EDS-IV in Figure 6h.

3.3 Corrosion mechanism in air and combustion environments.

Different from the amorphous SiO2/BN two-layer scale and the complex α-cristobalite/SiO2/SiBCNO(BN) ACS Paragon Plus Environment


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tri-layer scale reported before,6, 15 the fibers corroded in air and simulated combustion environments in this study shows such two reaction layers: SiO2 outer layer and BN grain-contained interlayer. A scheme summarizing the microstructure of SiBCN fibers before and after annealing in air and simulated combustion environment at 1400 °C and the corresponding potential corrosion mechanisms is illustrated in Figure 7.

Figure 7. Schematic representation of the microstructure of the original SiBCN fiber (a), ACF (b) and CCF (c) and the corresponding corrosion mechanisms.

The formation of the distinct two-layer structure in ACF is mainly attributed to the oxygen inward diffusion and the oxygen pressure gradient from the fiber surface to the core developed in this gas-solid interaction (Figure 7b). Although SiBCN fibers are not a mechanical mixture of Si3N4, BN, C and SiC materials, as shown in Figure 7a, the main structural units of SiBCN fibers are Si-N, B-N, sp2 hybridized C-C and Si-C ACS Paragon Plus Environment

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bonds. Moreover, the SiN3C, SiN4 units can be regarded as a composite structure of Si3N4 and SiC, and the B-N hexatomic rings can be considered as the microstructure of BN. Figure 8 provides the oxidation reaction of C, SiC, Si3N4 and BN, and the corresponding thermodynamic results calculated by the Reaction Web module of FactSage software. Within the temperatures of 200-2000 °C, it can be found that all the standard Gibbs free energy changes of the Reactions (1)-(4) are negative, which indicates the reactions can occur during oxidation. At the interface of the ambient atmosphere and the fiber surface, the partial pressure of O2 (P(O2)) is ~ 2×104 Pa, which is sufficient for the oxidation of all components in the Si-B-C-N network. At 1400 °C, free carbon clusters, Si-N, Si-C and B-N units are oxidized simultaneously to form CO (g), N2 (g), SiO2 (s) and B2O3 (l). Gaseous CO and N2 escape, liquid B2O3 volatilizes rapidly whereas only the solid SiO2 is stable and left to form the amorphous silica outer layer. Previously, Butchereit and Nickel reported a heterogeneous formation of bubbles along with crystallization on the edges of sponge-like SiBCN ceramics after longer oxidation time.14, 43-44 In contrast, the outmost layer of SiBCN fibers in this study is free of either the bubbles or the cristobalites. This could be attributed to the higher specific surface area of fibers which accelerates the volatilization of B2O3 and decrease the flowability of the SiO2 layer. In addition, the factors of different operation conditions between our studies also cannot be excluded.

The formation of h-BN nanocrystals beneath the silica layer could be due to the decrease of P(O2) and the selectivity of oxidation. With the thickness of silica layer increasing, oxygen diffusion inward leads to a gradient of P(O2) from the outer surface, thus, causing a low P(O2) beneath the silica layer. Previously, volatility diagrams of Si-N-O and Si-C-O systems have been compiled by Heuer et al. and have demonstrated that the lowest P(O2) for oxidation of Si3N4 at 1400 °C is ~ 10-15 Pa, and that for SiC is lower.45 Moreover, thermodynamic calculations also indicate that SiC and Si3N4 oxidizes at a lower oxygen potential than BN.46-47 When the supply of oxygen is limited, SiC and Si3N4 could effectively getter O2 so that BN does not get oxidized. Correspondingly, at low P(O2), free carbon clusters, Si-C and Si-N units of SiBCN fibers preferentially getter O2 and are oxidized to yield SiO2, whereas B-N hexatomic rings remain. ACS Paragon Plus Environment


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At high temperature, benefited from the oxidative elimination of the carbon in BN2C units, the B-N hexatomic rings move together aided by liquid SiO2 and transform to h-BN crystalline grains with different sizes. As mentioned before, the size of h-BN grains and composition of the matrix they precipitate are gradient changed from the outside of this layer radially, which are relative and attributed to the oxygen inward diffusion mechanism as well. The outer zone of the intermediate layer where contains higher P(O2) and SiO2 concentration leads to a lower viscosity of matrix. Therefore, small BN crystalline grains could more easily gather and grow up to larger grains in the outside area. With the decrease of P(O2), most of the small h-BN grains are fixed and embedded in the amorphous SiBCNO network.

2C + O2 = 2CO (g)

(1)

2/3SiC + O2 = 2/3SiO2 (s) + 2/3CO (g)

(2)

1/3Si3N4 + O2 = SiO2 (s) + 2/3N2 (g)

(3)

2/3BN + O2 = 1/3B2O3 (l) + 1/3N2 (g)

(4)

Figure 8. Standard Gibbs free energy changes of Reactions (1)-(4) as a function of the temperature during oxidation.11, 48-49 The standard state of the gases is a pressure of 1 atmosphere.

As shown in Figure 7c, the CCF performs similar two-layer scale but more complex structure than that of the ACF. With the function of oxygen, the CCF generates the SiO2 out layer and the h-BN nanocrystal-contained interlayer as well. However, the thickness of SiO2 layer of the CCF decreases visually when compared with the ACF. Due to the introduction of H2O in the simulated combustion environment, the oxidation product, SiO2, reacts with the water vapor and yields gaseous Si(OH)4,50-51 which subsequently escapes from the SiO2 layer, thus resulting in the decrease of the thickness. The thermogravimetric analysis (Figure S2) of the SiBCN fibers via heat treatment up to 1400 °C under air and simulated combustion also provides indirect evidence for the difference between ACF and CCF. No detectable mass change could be ACS Paragon Plus Environment

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observed under air atmosphere whereas the fibers in simulated combustion environment exhibit approximately 3 wt% weight loss from 1200 to 1400 °C. The zero-mass change for ACF could be attributed to a balance of mass gain caused by the formation of SiO2 and the mass loss by the escape of N2, CO and B2O3. On this basis, the mass loss of CCF under water-contained environment might result from the volatilization of silica as reaction (5).

SiO2 + 2H2O (g) = Si(OH)4 (g)

(5)

Moreover, water vapor in the mixed atmosphere also plays an important role in the structure evolution for the formation of α-cristobalites in the amorphous silica matrix, which can be attributed to the catalytic action of water. At high temperature, water vapor can weaken the silica structure through the formation of hydroxyls, where the strong Si-O-Si bond is replaced by Si-OH group and produce the much weaker [Si-OH HO-Si] intermediate. The increasing [Si-OH HO-Si] content of the amorphous silica reduces the activation energy for viscous flow and promotes the conversion from glass to cristobalite.52 Moreover, H2O can also accelerate the volatilization of B2O3 at high temperature, which gives rise to the local viscosity and induces devitrification of silica into cristobalite.44

The size and distribution of α-cristobalites are influenced by the H2O inward diffusion. In the SiO2 layer, α-cristobalites spread all over the amorphous matrix with mean size of 10nm due to a higher P(H2O). When the position comes to the h-BN-contained interlayer, the size of α-cristobalites reduces and disappears in the middle region of this layer, where the h-BN nanocrystals still exist. This is attributed to the low diffusion coefficient of the water vapor in silica, which is ~ 100 times lower than oxygen.53 In the middle of the interlayer, P(H2O) has been low enough and is difficult to render crystallization of silica. In addition, it could also be observed that the size of h-BN nanocrystals at the top of the interlayer is ~ 15 nm, which is smaller than those of ACF (20-30 nm). This could be explained by the crystallization of amorphous silica to α-cristobalite grains, which increases the viscosity (decrease the liquidity) of the SiO2-contained matrix, thus ACS Paragon Plus Environment


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suppressing gathering and growing up of small BN crystalline grains.

4. Conclusion

In summary, this work has investigated the composition and structure of polymer-derived SiBCN fibers before and after corrosive tests at 1400 °C in air and simulated combustion environments, respectively. The as-prepared amorphous SiBCN fibers mainly consist of SiN3C, SiN4, B-N hexatomic rings, free carbon clusters and BN2C units. The fibers separate to three layers with different compositions, mainly having SiO2, h-BN grains and unaffected SiBCN, respectively, from the edge to the center after the corrosive tests. Three portions show good adhesion to each other with no detectable cracks. The prominent difference for fibers corroded in the simulated combustion condition compared with those annealed in air is the formation of α-cristobalite nanocrystals in the amorphous SiO2 matrix rather the pure amorphous SiO2 itself, which might be induced by the presence of water vapor in the high temperature oxidation reaction. These α-cristobalite nanoparticles hinder the aggregation of small h-BN grains in the second layer. In addition, the water vapor also triggers the conversion of SiO2 to the gaseous Si(OH)4, leading to the thinning of the outermost SiO2 layer. We also propose a potential mechanism for the preservation of h-BN grains in the interlayer in such a high temperature environment with oxygen. It could be due to the sacrificial effect of free carbon clusters, Si-C and Si-N units, which react with oxygen prior to B-N hexatomic rings, thus effectively decreasing the amount of oxygen diffusing inward and allowing the formation of h-BN nanoparticles. These results improve our understanding of the corrosion process of SiBCN fibers and more experiments on the structure and properties of these fibers will be done for the applications in aerospace in the future.

AUTHOR INFORMATION Corresponding Authors *

Changwei Shao, [email protected]


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*

Hao Wang, [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT The financial support by the National Natural Science Foundation of China (51203184, 51772327) is gratefully acknowledged. Valuable time provided by Dr. Xinguo Chen in the Ningbo Institute of Material Technology and Engineering for FIB, TEM and EELS investigations is also appreciated.

5. References (1) Padture, N. P. Advanced Structural Ceramics in Aerospace Propulsion. Nature Materials 2016, 15 (8), 804-809. (2) Council, N. R. Materials Needs and R&D Strategy for Future Military Aerospace Propulsion Systems, National Academies Press: 2011. (3) Dever, J. A.; Nathal, M. V.; DiCarlo, J. A. Research on High-Temperature Aerospace Materials at NASA Glenn Research Center. Journal of Aerospace Engineering 2013, 26 (2), 500-514. (4) Pope, D. P.; Darolia, R. High-Temperature Applications of Intermetallic Compounds. MRS Bulletin 1996, 21 (5), 30-36. (5) Bochenek, K.; Basista, M. Advances in Processing of NiAl Intermetallic Alloys and Composites for High Temperature Aerospace Applications. Progress in Aerospace Sciences 2015, 79, 136-146. (6) Baldus, P.; Jansen, M.; Sporn, D. Ceramic Fibers for Matrix Composites in High-Temperature Engine Applications. Science 1999, 285 (5428), 699-703. (7) Flores, O.; Bordia, R. K.; Nestler, D.; Krenkel, W.; Motz, G. Ceramic Fibers Based on SiC and SiCN Systems: Current Research, Development, and Commercial Status. Advanced Engineering Materials 2014, 16 (6), 621-636. (8) Hart, A. H.; Koizumi, R.; Hamel, J.; Owuor, P. S.; Ito, Y.; Ozden, S.; Bhowmick, S.; Syed Amanulla, S. A.; Tsafack, T.; Keyshar, K. Velcro-Inspired SiC Fuzzy Fibers for Aerospace Applications. ACS Applied Materials & Interfaces 2017, 9 (15), 13742-13750. (9) Aldinger, F.; Weinmann, M.; Bill, J. Precursor-Derived Si-BCN Ceramics. Pure and Applied Chemistry 1998, 70 (2), 439-448. (10) Spanio, H. Novel SiBNC Ceramic Fibers for Production, Energy and Traffic Technology.; Fraunhofer-Institutfür Silicatforschung: Würzburg, 2009; pp 68-71. (11) Weinmann, M.; Schuhmacher, J.; Kummer, H.; Prinz, S.; Peng, J.; Seifert, H. J.; Christ, M.; Müller, K.; Bill, J.; Aldinger, F. Synthesis and Thermal Behavior of Novel Si−B−C−N Ceramic Precursors. Chemistry of Materials 2000, 12 (3), 623-632. (12) Riedel, R.; Kienzle, A.; Dressler, W.; Ruwisch, L.; Bill, J.; Aldinger, F. A Silicoboron Carbonitride Ceramic Stable to 2,000 °C. Nature 1996, 382 (6594), 796-798. (13) Baldus, H.; Passing, G.; Sporn, D.; Thieraf, A. Si-B-(N, C) a New Ceramic Material for High Performance Applications; American Ceramic Society, Westerville, OH (United States): 1995. (14) Butchereit, E.; Nickel, K. G.; Müller, A. Precursor-Derived Si−B−C−N Ceramics: Oxidation Kinetics. Journal of the American Ceramic Society 2001, 84 (10), 2184-2188. (15) Cinibulk, M. K.; Parthasarathy, T. A. Characterization of Oxidized Polymer-Derived SiBCN Fibers. Journal of the American Ceramic Society 2001, 84 (10), 2197-2202. (16) Butchereit, E.; Nickel, K. In Oxidation Behavior of Precursor Dervied Ceramics in the System Si-(B)-CN, Key Engineering Materials, Trans Tech Publ: 2000; pp 69-78. (17) Ji, X.; Shao, C.; Wang, H.; Wang, J.; Cheng, J.; Long, X.; Mao, T. A Simple and Efficient Method for the Synthesis of SiBNC



Page 24 of 34

Ceramics with Different Si/B Atomic Ratios. Ceramics International 2017, 43 (10), 7469-7476. (18) Ji, X.; Shao, C.; Wang, H.; Cheng, J.; Li, H. Curing Green Fibres Infusible by Electron Beam Irradiation for the Preparation of SiBNC Ceramic Fibres. Ceramics International 2017, 43 (14), 11218-11224. (19) Yao, R.; Feng, Z.; Chen, L.; Zhang, Y.; Zhang, B. Oxidation Behavior of Hi-Nicalon SiC Monofilament Fibres in Air and O2–H2O– Ar Atmospheres. Corrosion Science 2012, 57, 182-191. (20) Li, S.; Feng, Z.; Mei, H.; Zhang, L. Mechanical and Microstructural Evolution of Hi-Nicalon Trade Mark SiC Fibers Annealed in O2–H2O–Ar Atmospheres. Materials Science and Engineering: A 2008, 487 (1-2), 424-430. (21) Jacobson, N. S. Corrosion of Silicon-Based Ceramics in Combustion Environments. Journal of the American Ceramic Society 1993, 76 (1), 3-28. (22) Hu, T.; Li, X.; Li, G.; Wang, J. SiC Fibers with Controllable Thickness of Carbon Layer Prepared Directly by Preceramic Polymer Pyrolysis Routes. Materials Science and Engineering: B 2011, 176 (9), 706-710. (23) Kong, J.; Wang, M.; Zou, J.; An, L. Soluble and Meltable Hyperbranched Polyborosilazanes toward High-Temperature Stable SiBCN Ceramics. ACS Applied Materials & Interfaces 2015, 7 (12), 6733-6744. (24) Zhang, Q.; Yang, Z.; Jia, D.; Chen, Q.; Zhou, Y. Synthesis and Structural Evolution of Dual-Boron-Source-Modified Polysilazane Derived SiBCN Ceramics. New Journal of Chemistry 2016, 40 (8), 7034-7042. (25) Song, Y.; He, L.; Zhang, X.; Liu, F.; Tian, N.; Tang, Y.; Kong, J. Highly Efficient Electromagnetic Wave Absorbing Metal-Free and Carbon-Rich Ceramics Derived from Hyperbranched Polycarbosilazanes. The Journal of Physical Chemistry C 2017, 121 (44), 24774-24785. (26) Kong, J.; Kong, M.; Zhang, X.; Chen, L.; An, L. Magnetoceramics from the Bulk Pyrolysis of Polysilazane Cross-Linked by Polyferrocenylcarbosilanes with Hyperbranched Topology. ACS Applied Materials & Interfaces 2013, 5 (20), 10367-10375. (27) Wall, M. The Raman Spectroscopy of Graphene and the Determination of Layer Thickness. Thermo Scientific-Application Note 2011, 52252. (28) Sarkar, S.; Gan, Z.; An, L.; Zhai, L. Structural Evolution of Polymer-Derived Amorphous SiBCN Ceramics at High Temperature. The Journal of Physical Chemistry C 2011, 115 (50), 24993-25000. (29) Gao, Y.; Mera, G.; Nguyen, H.; Morita, K.; Kleebe, H.-J.; Riedel, R. Processing Route Dramatically Influencing the Nanostructure of Carbon-Rich SiCN and SiBCN Polymer-Derived Ceramics. Part I: Low Temperature Thermal Transformation. Journal of the European Ceramic Society 2012, 32 (9), 1857-1866. (30) Ferrari, A. C.; Robertson, J. Interpretation of Raman Spectra of Disordered and Amorphous Carbon. Physical review B 2000, 61 (20), 14095. (31) Ionescu, E.; Papendorf, B.; Kleebe, H.-J.; Breitzke, H.; Nonnenmacher, K.; Buntkowsky, G.; Riedel, R. Phase Separation of a Hafnium Alkoxide-Modified Polysilazane upon Polymer-to-Ceramic Transformation—A Case Study. Journal of the European Ceramic Society 2012, 32 (9), 1873-1881. (32) Ruwisch, L. M.; Dürichen, P.; Riedel, R. Synthesis of Silyl Substituted Organoboranes by Hydroboration of Vinylsilanes. Polyhedron 2000, 19 (3), 323-330. (33) Gervais, C.; Maquet, J.; Babonneau, F.; Duriez, C.; Framery, E.; Vaultier, M.; Florian, P.; Massiot, D. Chemically Derived BN Ceramics: Extensive

11

B and

15

N Solid-State NMR Study of a Preceramic Polyborazilene. Chemistry of Materials 2001, 13 (5),

1700-1707. (34) Schuhmacher, J.; Berger, F.; Weinmann, M.; Bill, J.; Aldinger, F.; Müller, K. Solid-State NMR and FT-IR Studies of the Preparation of Si−B−C−N Ceramics from Boron-Modified Polysilazanes. Applied Organometallic Chemistry 2001, 15 (10), 809-819. (35) Gottardo, L.; Bernard, S.; Gervais, C.; Weinmann, M.; Miele, P. Study of the Intermediate Pyrolysis Steps and Mechanism Identification of Polymer-Derived SiBCN Ceramics. Journal of Materials Chemistry 2012, 22 (34), 17923-17933. (36) Zhao, H.; Chen, L.; Luan, X.; Zhang, X.; Yun, J.; Xu, T. Synthesis, Pyrolysis of a Novel Liquid SiBCN Ceramic Precursor and its Application in Ceramic Matrix Composites. Journal of the European Ceramic Society 2017, 37 (4), 1321-1329. (37) Egerton, R. F. Electron Energy-Loss Spectroscopy in the Electron Microscope, Springer Science & Business Media: 2011. (38) Gregori, G.; Kleebe, H.-J.; Blum, Y. D.; Babonneau, F. Evolution of C-Rich SiOC Ceramics: Part II. Characterization by High Lateral Resolution Techniques: Electron Energy-Loss Spectroscopy, High-Resolution TEM and Energy-Filtered TEM: Dedicated to Professor Dr. Fritz Aldinger on the Occasion of his 65th Birthday. Zeitschrift für Metallkunde 2006, 97 (6), 710-720. (39) Serin, V.; Brydson, R.; Scott, A.; Kihn, Y.; Abidate, O.; Maquin, B.; Derré, A. Evidence for the Solubility of Boron in Graphite by Electron Energy Loss Spectroscopy. Carbon 2000, 38 (4), 547-554.


Page 25 of 34 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(40) Wang, X.; Weng, Q.; Wang, X.; Li, X.; Zhang, J.; Liu, F.; Jiang, X.; Guo, H.; Xu, N.; Golberg, D. Biomass-Directed Synthesis of 20 g High-Quality Boron Nitride Nanosheets for Thermoconductive Polymeric Composites. ACS Nano 2014, 8 (9), 9081-9088. (41) Yang, X.; Guan, Z.; Zeng, M.; Wei, J.; Wang, W.; Bai, X. Facile Synthesis of Large-Area Ultrathin Hexagonal BN Films via Self-Limiting Growth at the Molten B2O3 Surface. Small 2013, 9 (8), 1353-1358. (42) Müller, A.; Zern, A.; Gerstel, P.; Bill, J.; Aldinger, F. Boron-Modified Poly (propenylsilazane)-Derived Si–B–C–N Ceramics: Preparation and High Temperature Properties. Journal of the European Ceramic Society 2002, 22 (9-10), 1631-1643. (43) Butchereit, E.; Nickel, K. G. In Comparison of the Oxidation Kinetics of Boron Free and Boron Containing Precursor Derived Ceramics, High Temperature Corrosion and Materials Chemistry III: Proceedings of the International Symposium, The Electrochemical Society: 2001; p 112. (44) Chollon, G., Oxidation Behaviour of Polymer-Derived Ceramics. DEStech Publications: 2010. (45) Heuer, A. H.; Lou, V. L. K. Volatility Diagrams for Silica, Silicon Nitride, and Silicon Carbide and their Application to High-Temperature Decomposition and Oxidation. Journal of the American Ceramic Society 1990, 73 (10), 2789-2803. (46) Jacobson, N. S.; Morscher, G. N.; Bryant, D. R.; Tressler, R. E. High-Temperature Oxidation of Boron Mitride: II, Boron Nitride Layers in Composites. Journal of the American Ceramic Society 2010, 82 (6), 1473-1482. (47) Wakasugi, T.; Tsukihashi, F.; Sano, N. Thermodynamics of Nitrogen in B2O3, B2O3-SiO2, and B2O3-CaO Systems. Journal of the American Ceramic Society 1991, 74 (7), 1650-1653. (48) Liu, Z.; Gong, Y.; Zhou, W.; Ma, L.; Yu, J.; Idrobo, J. C.; Jung, J.; MacDonald, A. H.; Vajtai, R.; Lou, J. Ultrathin High-Temperature Oxidation-Resistant Coatings of Hexagonal Boron Nitride. Nature Communications 2013, 4, 2541-2548. (49) Bernard, S.; Weinmann, M.; Gerstel, P.; Miele, P.; Aldinger, F. Boron-Modified Polysilazane as a Novel Single-Source Precursor for SiBCN Ceramic Fibers: Synthesis, Melt-Spinning, Curing and Ceramic Conversion. Journal of Materials Chemistry 2005, 15 (2), 289-299. (50) Wu, S.; Cheng, L.; Zhang, L.; Xu, Y.; Zhang, J.; Mei, H. Wet Oxidation Behaviors of Hi-Nicalon Fibers. Applied Surface Science 2006, 253 (3), 1447-1450. (51) Jacobson, N. S.; Opila, E. J.; Lee, K. N. Oxidation and Corrosion of Ceramics and Ceramic Matrix Composites. Current Opinion in Solid State and Materials Science 2001, 5 (4), 301-309. (52) Wagstaff, F.; Richards, K. Kinetics of Crystallization of Stoichiometric SiO2 Glass in H2O Atmospheres. Journal of the American Ceramic Society 1966, 49 (3), 118-121. (53) Opila, E. J. Variation of the Oxidation Rate of Silicon Carbide with Water-Vapor Pressure. Journal of the American Ceramic Society 1999, 82 (3), 625-636.

Table of Contents (TOC)




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Figure 1．(a＆b) SEM images showing the as-synthesized SiBCN fibers. Inset of panel b is the EDS elemental mapping of Si, B, N, and C taken from the region marked in the red rectangle box in panel b. (c) HRTEM images of the as-synthesized SiBCN fibers. Insets of panel c are the SEAD pattern and the magnified image taken from the region marked in the red square box in panel c. (d) AES depth profile showing the concentration change of elements in the surface region of the fiber as a function of the sputtering depth. 150x110mm (300 x 300 DPI)



Figure 3. SEM images showing the surface (a＆b) and cross section morphologies (c＆d) of SiBCN fibers annealing at 1400 °C for 2 h in air. 150x129mm (300 x 300 DPI)


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Figure 4. (a) TEM images showing the cross section of the fibers annealing at 1400 °C in the static air for 2h. (b-d) High magnification TEM images corresponding to the zones marked by B, C and D in (a), respectively. Insets of panel c and d are SEAD patterns of the region marked in the red circles. (e-g) HRTEM images taken from the glass-like area in (b), (c) and the amorphous area in (d). Insets of panel e, f and g are their corresponding SEAD and FFT patterns, respectively. (h) EDS spectra (I), (II) and (III) corresponding to the zones marked by EDS-I, II and III in (a), respectively. (i) EELS spectra with the distance of 0.25, 0.75, 1.50 and 2.50 µm from the fiber surface in (a). 184x135mm (300 x 300 DPI)



Figure 5. SEM images showing the surface (a＆b) and cross section morphologies (c＆d) of SiBCN fibers annealing at 1400 °C for 2 h in simulated combustion environment. 150x120mm (300 x 300 DPI)


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Figure 6. (a) TEM images showing the cross section of the fibers annealing at 1400 °C in simulated combustion environment for 2h. (b-d) High magnification TEM images corresponding to the zones marked by B, C and D in (a), respectively. (e-g) HRTEM images taken from regions marked in the red rectangle box in panel b, c and d, respectively. Insets of panel e, f and g are their corresponding SEAD and FFT patterns., respectively. (h) EDS spectra (I), (II), (III) and (IV) corresponding to the zones marked by EDS-I, II, III and IV in (a), respectively. 176x145mm (300 x 300 DPI)



Figure 7. Schematic representation of the microstructure of the original SiBCN fiber (a), ACF (b) and CCF (c) and the corresponding corrosion mechanisms. 279x235mm (300 x 300 DPI)


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Figure 8. Standard Gibbs free energy changes of Reactions (1)-(4) as a function of the temperature during oxidation.11, 48-49 The standard state of the gases is a pressure of 1 atmosphere. 209x148mm (300 x 300 DPI)



TOC 99x101mm (300 x 300 DPI)


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The high-temperature corrosion behavior of SiBCN fibers for

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