Effect of carbon nanotube on processing, microstructural, mechanical and ablation behavior of ZrB2-20SiC based ultra-high temperature ceramic composites

Accepted Manuscript Effect of carbon nanotube on processing, microstructural, mechanical and ablation behavior of ZrB2-20SiC based ultra-high temperature ceramic composites Ambreen Nisar, S. Ariharan, T. Venkateswaran, N. Sreenivas, Kantesh Balani PII:

S0008-6223(16)30854-5

DOI:

10.1016/j.carbon.2016.10.002

Reference:

CARBON 11361

To appear in:

Carbon

Received Date: 15 July 2016 Revised Date:

30 September 2016

Accepted Date: 1 October 2016

Please cite this article as: A. Nisar, S. Ariharan, T. Venkateswaran, N. Sreenivas, K. Balani, Effect of carbon nanotube on processing, microstructural, mechanical and ablation behavior of ZrB2-20SiC based ultra-high temperature ceramic composites, Carbon (2016), doi: 10.1016/j.carbon.2016.10.002. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Effect of Carbon Nanotube on Processing, Microstructural, Mechanical and Ablation Behavior of ZrB2-20SiC Based Ultra-High Temperature Ceramic

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Composites Ambreen Nisar1, S. Ariharan1, T. Venkateswaran2, N. Sreenivas2 and Kantesh Balani1,* 1

Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kanpur-208016, India 2

Vikram Sarabhai Space Centre, Indian Space Research Organization, Trivandrum- 695022, India

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Abstract

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Herein, ZrB2 is reinforced with silicon carbide (SiC) and carbon nanotube (CNT) to provide enhanced structural stability and oxidation protection against extreme thermal (> 2400 °C) and oxidative environments. The ablation resistance of ZrB2-based composites was evaluated using plasma arc-jet with a heat-flux of 2.5 MW/m2 for 30 s, and the decreased oxidation-rate

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(from 0.77 µm/s to 0.44 µm/s) is attributed to enhanced thermal conductivity (42.3 to 52.3 W/mK at 1200 °C) with synergistic reinforcement of SiC and CNT.

The increased onset

temperature (from 679 °C to 706 °C) and decreased enthalpy of oxide formation (from 1.6 to 0.6

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kJ/g), insinuates an increase in thermal stability and oxidation resistance with the synergistic addition of both SiC and CNT in ZrB2. The increase in the hardness of ZrB2 in the as-processed

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composites (up to 1.6 times) as well as after plasma arc jet exposure (up to 2.1 times) with synergistic reinforcement of SiC and CNT has shown to suppress crack-formation and restrict oxidation. The reduction in the analytically evaluated tensile interfacial residual stress indicates enhanced structural integrity of ZrB2-SiC-CNT composites, mandatory for aerospace applications.

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Corresponding Author. Email: [email protected], Tel: +91-512-259-6194 (Kantesh Balani)

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Keywords: ZrB2-20SiC ceramic; carbon nanotube (CNT); spark plasma sintering (SPS); high temperature oxidation; TEM.

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1. Introduction: Zirconium diboride (ZrB2), an ultra-high temperature ceramics (UHTC) possesses low density (6.09 g/cc) [1], high melting temperature (Tm = 3245 °C) [2] and high thermal

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conductivity (60 W/mK) [3], which makes it attractive for use in thermal protection system and scramjet engine components for hypersonic vehicles. The advantages of the ZrB2-based UHTCs

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is not only from their high-temperature stability, but also from their high thermal conductivity at high temperatures [3], which ensure its utilization as sharp leading edges for re-entry space vehicles. The superior high temperature stability is primarily due to strong covalent bonding in ZrB2 structure [3], however, strong bonding and low diffusion coefficient make the processing of

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ZrB2 a challenge. Reinforcements are often added to ZrB2 to improve densification, mechanical (hardness, elastic modulus, and fracture toughness), thermal and oxidative properties [4-6]. Several investigations indicated that SiC reinforcement not only improves flexural strength,

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fracture toughness but also improves oxidation properties by the formation of passive SiO2 layer which protects ZrB2 at elevated temperature [6-11]. It has been established that ZrB2-20SiC

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(vol%) composites is optimal for high-temperature thermomechanical applications [9, 12]. On one hand, where the thermal resistance of SiC reinforcement is well established, reinforcement of thermally conductive CNT (~3000 W/mK) [13, 14] may provide enhanced resistance to thermal damage of ZrB2-composites [6]. Recently, carbon nanotube (CNT), graphite and carbon fibers have also been used as nano-scale filler in the ZrB2 ceramics [6, 15, 16], resulting in nearly complete densification [6, 16]. Tian et al. [15] investigated the effect of CNT (2 wt%) on the properties of ZrB2-SiC ceramics, showing that the fracture toughness increased about 15% (i.e. 2

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from 4 MPam1/2 to 4.6 MPam1/2), but hardness (15.8 GPa to 15.5 GPa), and thermal conductivity (at 400 °C, 72.5 W/mK to 75.4 W/mK) did not vary significantly.

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Plasma-arc jet testing closely approximates the intense aerothermal conditions [17] which typically exceeds 2000 °C during the test and provide high shear stresses (due to the flux of 1-3 MW/m2) that can influence the microstructural evolution (enough to melt most UHTC oxides)

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leading to mechanical failure. While improving mechanical properties of ZrB2 is the crucial step in making it a viable material, it is not sufficient if the oxidation properties are not improved

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upon. While evaluating the oxidation behavior of these UHTCs, it is necessary to evaluate the materials performance by simulating environmental conditions experienced during hypersonic flight.

Till date, the oxidation mechanism of only monolithic ZrB2 is available [18], but, ZrB2

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reinforced with 20 vol% SiC (ZrB2-20SiC) is established as the baseline material in literature so, a comprehensive understanding on the effect of microstructural evolution on thermo-mechanical performance (of UHTC) is required for use of these materials in hypersonic and propulsion

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applications. In accordance with the recent study on TaC-based composites [19, 20], the synergy of SiC and CNT reinforcement has been extended to ZrB2-based composites in the current work.

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The scope of this work is to compare the effect of CNT on the oxidation behavior of ZrB2-20SiC under plasma arc jet exposure. In agreement with the previous study in literature [6, 15], where up to 6 vol% of CNT is added in ZrB2 and ZrB2-20SiC), these authors have observed an increasing trend of fracture toughness even with the maximum CNT reinforcement of 6 vol. % (i.e. 3.53 MPam1/2 in ZrB2, and 4.6 MPam1/2 in ZrB2-20SiC composite). Thus, the aspect of agglomeration may not have occurred in these composites. On that basis, 10 vol% (which corresponds to 4.1 wt %) of CNT was reinforced in order to improve densification behavior, mechanical properties, 3

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thermal conductivity, and oxidation resistance of ZrB2-20SiC. Enhanced content of CNT may protect the base ZrB2 material via sacrificial oxidation at high temperatures. Thus, the consequence of CNT addition on the the high temperature thermal conductivity is

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correlated with the oxidation behaviour of the ZrB2-based composites. The composition of the surface layer and variation in the thickness of oxide scale (after plasma arc jet exposure) is also investigated to analyse the oxidation resistance of ZrB2-SiC-CNT UHTC composites. The

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oxidation mechanism has been proposed according to the experimental results and thermodynamic considerations, based on which sample ranking has been made against the

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performance upon plasma arc jet exposure. Furthermore, the thermogravimetric analysis is carried out at 1500 °C (for 30 min.) in an oxygen environment, to complement the results obtained from plasma arc jet exposure. Through such investigation, the purpose is to correlate the effect of microstructural evolution on the structural integrity (with the synergistic addition of SiC and

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CNT), during processing, which later governs the oxidative performance of ZrB2-based UHTC against plasma arc jet testing.

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2. Materials and Method:

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2.1. Materials processing:

Commercial powders of ZrB2 (H.C. Starck, Germany, 99.9% pure, particle size < 2µm, Fig. 1a), SiC (H.C. Starck, Germany, 99.9% pure, particle size < 1µm, Fig. 1b) and multi-walled carbon nanotubes, (CNT, Nanostructured and Amorphous Materials Inc., TX, USA, 94% pure, outer diameter of 50 nm, inner diameter of 30 nm, and 1-2 µm long, Fig. 1c) were used as the starting materials. The powder mixtures: ZrB2 mixed with 20 vol% SiC (referred to as Z20S), ZrB2 with 10 vol % CNT (referred to as Z10C), and ZrB2 with 20 vol% SiC and 10 vol% CNT 4

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(referred to as Z20S10C) were dry ball milled with ball to powder ratio of 5:2 for 8 min at 500 rpm using tungsten carbide jar and tungsten carbide balls. Bright-field TEM micrograph of

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in ZrB2 matrix and the retention of CNT after ball-milling.

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Z20S10C (Fig. 1d) illustrates the distribution of secondary phase particles (SiC as well as CNT)

Figure 1: TEM micrographs of staring powder (a) ZrB2 (b) SiC (c) CNT and (d) Z20S10C composite powder.

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The ball-milled composite powders were processed via spark plasma sintering (SPS, Dr. Sinter, SPS-515S, Japan) for resultant pellet of 15 mm diameter and 2-3 mm thickness using graphite die and punches with heating rate of 150 ºC/min by holding 5 min at each step (1300 °C

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(0.4 Tm), 1500 °C (0.45 Tm) and 1650 °C (0.5 Tm)) at 40 MPa in argon (Ar) atmosphere. After final stage of holding at final temperature, the sample was allowed to cool naturally in the presence of Ar gas. The theoretical density of these composites were calculated using rule of

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water as immersion medium) on a hydrostatic balance.

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mixture (ROM), however, the final densities were measured by Archimedes method (using distil

2.2.Oxidation Test:

SPS samples of 15 mm diameter and 3 mm thickness, encased in carbon phenolic guard, were plasma arc-jet exposed under heat flux of 2.5 MW/m2 for 30 s. The heat flux (2.5 MW/m2)

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is directly measured using water cooled thin foil heat flux transducers (HFT) which has an uncertainty of ±5%. The HFT is placed at a desired axial location from nozzle exit and power and gas flow to plasma generator is adjusted to get this heat flux. To monitor back face

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temperature, a K-type thermocouple is attached at the back face using high temperature cement [19]. Location of the required heat flux and operating conditions of the plasma generator with

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argon calibrated using water cooled Gardon gauges. The linear oxidation rate (R) is calculated using the formula:


=

    

(1)

where, doxide is the oxide scale thickness after plasma arc jet exposure and texposure is the exposure time (30 s). Complimentarily, simultaneous thermal analysis (STA 8000, Perkin Elmer, USA) i.e. thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) is used to 6

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evince oxidation mechanism and thermal stability. Thermal studies were performed on samples with an initial mass of 50-100 mg from room temperature to 1500 °C at a rate of 20˚C/min in an oxygen atmosphere for 30 min holding at maximum temperature. The accuracy of the equipment

reported value is an average of at least 5 experiments.

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2.3. Thermal conductivity measurements:

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is 0.2% with the thermal stability of furnace within ±0.5 °C. Under similar conditions, the

The thermal conductivity of the investigated samples was measured by laser flash diffusivity

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technique (Flashline, thermal properties analyzer, Anter Corporation, USA) using pulsed Nd:YAG laser for ~300 µs in Ar atmosphere. The back face temperature rise of the samples (diameter of 14 mm and thickness of 3mm) was recorded by In-Sb photo detector. The thermal diffusivity, D is calculated using relation (Eq. 2). 0.13879  /

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 =

(2)

where, L is specimen thickness and t1/2 is the half time required for initiation of the pulse for

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back face temperature to reach half of the maximum rise in temperature. Thermal conductivity, κ

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is calculated using the relation (Eq. 3) [21]: κ =  



(3)

where, ρ is the measured density, Cp is the measured specific heat (simultaneously measured using laser flash diffusivity technique with an accuracy of ±4 %) as per Kopp-Neumann rule. It may be indicated that the effect of CNT agglomeration is not accounted for in the estimation of thermal conductivity measurements. The thermal diffusivity, specific heat, and thermal

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conductivity of these materials were analyzed under inert atmosphere in order to better understand the influence of porosity and composition. The true thermal diffusivity values are

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estimated to be within an error of ± 1.44 % with 95 % confidence limit. 2.4. Phase, microstructural and mechanical characterizations:

Phase was carried out using Rich-Seifert, 2000D diffractometer operated at 25 kV and 15

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mA (Cu Kα, λ = 1.54 Å at a scan speed of 0.5 s/step and a step size of 0.02°) in the 2θ range from 30° to 90°. The quantification of retained ZrB2 phase after plasma arc jet test was

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calculated as an average of area fraction of (001), (100) and (101) XRD peaks. The crystallite size (t) of ZrB2-based composites before and after plasma arc-jet test was evaluated using Scherrer’s equation: 0.9 ! " cos θ

(4)

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 =

where, β = (β)  − β+  , βm and βs are full-width half maximum (FWHM) of the most intense peak of the sample and that of the standard Al2O3 sample at same θ (Bragg peak position in

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radians), respectively, λ is the wavelength of incident X-ray, and θ is Bragg’s diffraction angle. Further, to study the damage of CNT after the plasma arc jet exposure, micro-Raman

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spectroscopy (Princeton Instruments, STR Raman, TE-PMT detector) was carried out using NdYAG green laser (λ= 532 nm) with laser power of 12.5 mW, an objective of 50X and spatial resolution of 0.5 cm-1 in the backscattered mode. The surface morphology of ZrB2-SiC-CNT composites (SPS and after plasma exposure) and profile of oxide layer after plasma exposure along the cross section were imaged using W-SEM (JEOL JSM-6010LA and Zeiss Model EVO

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50). Transmission electron microscopy (TEM, FEI UT 20) of the oxidized Z20S10C sample was performed after scratching the surface oxide for analysis at an operating voltage of 200 kV. Hardness and elastic modulus of the as-sintered compacts and on the cross-section of the

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oxide layer were measured by using instrumented micro indenter (MHT, CSM instruments, Switzerland) on the polished samples. The load function comprises loading to 2N (at rate of 4 N/min and holding at maximum load for 10 s followed by unloading at the same rate using

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Vickers’ indenter of type V-I 51. Diametrical compression test was carried out using 100kN Universal Testing Machine (UTM, Model: BiSS Ltd.) at a strain rate of 0.05 mm/min for all the

23]: ,- =

2 / 0ℎ

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processed composites. The fracture strength (σf) of the material is calculated using formula [22,

(5)

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where, P is the maximum load at which fracture occurs, h is the thickness and d is the diameter of the sample. Three samples were tested and an average value (obtained within 15% deviation) is reported.

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3. Results and discussion

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3.1. Densification of ZrB2-SiC-CNT composites: The densification of ZrB2-based composites increased from 93% for pure ZrB2 to ~100% with synergistic reinforcement of both SiC and CNT in ZrB2 under similar SPS processing conditions, presented in Fig. 2 (also see Table 1).

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Table 1: Nomenclature, theoretical and experimental densities, and % relative densification of ZrB2-SiC-CNT composites. Composition

ρtheo (g/cc)

ρexpt (g/cc)

% densification

ZrB2

ZrB2

6.1

5.7

93.1

Z20S

ZrB2+20vol%SiC

5.5

5.2

95.0

Z10C

ZrB2+10vol%CNT

5.7

5.5

95.9

Z20S10C

ZrB2+20vol%SiC+10vol%CNT

5.1

5.0

99.7

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Sample ID

Figure 2: Plot showing the variation of instantaneous densification with time processed in subsequent stages during SPS processing of ZrB2-based composites. 10

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During SPS, the recorded punch displacement profile reveals three distinct regions, i.e. initial compaction (step I) due to the rearrangement of powder particles; intermediate step (step II)

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corresponding to thermal expansion, and shrinkage (step III) corresponding to final densification as shown in Fig. 2. Since, step III corresponds to shrinkage, is a crucial step for the final densification, which is observed to increase with SiC and CNT reinforcement [20].

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3.2. Phase retention in spark plasma sintering of ZrB2-SiC-CNT composites:

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X-ray diffraction (XRD) patterns for powder mixtures as well as SPS sintered ZrB2-based composites are presented in Fig. 3a. The presence of characteristic peaks of ZrB2 and SiC elicit retention of phases without any undesirable interfacial reactions after SPS processing. However, the intensity of SiC phase was low due to poor X-ray reflecting the ability of SiC [20]. No peaks corresponding to CNT were detected in XRD, whereas, Raman spectra (Fig. 3b) elucidate D- and

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G-peaks confirming the retention of CNT even after SPS processing. The shift in the D- and Gpeaks in the SPS samples when compared to pristine CNT (D-peak at 1351.8 cm-1 and G-peak at

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1584.4 cm-1) elicits that the compressive stress was introduced after SPS processing, possibly due to thermal contraction of the ceramic matrix during sintering [6, 20, 24]. The ratio of the

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intensity of D- to G-peak (ID/IG) increased from 0.97 for pristine CNT powder to 0.99 (for Z10C) and 1.03 (Z20S10C) indicating damage of CNT structure (bending and buckling in CNT, shown later in Fig. 4).

3.3 Structural stability of CNT during SPS processing: Structural stability of CNT is essential in rendering enhanced toughness, so it is imperative to investigate the interfaces and structure of various phases of the SPS composite pellets (see Figs.

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4a-d). The absence of impurities (oxide, glassy phase etc.) at triple junctions and presence of dislocation tangles even in pure ZrB2 (Fig. 4a) suggests removal of oxide impurities during SPS,

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CNT are tightly encapsulated within ZrB2 matrix.

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which enables solid-state sintering. It can be well accentuated from Figs. 4b and c that SiC and

Figure 3 (a): XRD pattern of initial composite powders as well as ZrB2-SiC-CNT composite pellets, and (b) Raman spectra showing D- and G-peaks in pristine CNT powder as well as SPS ZrB2-SiC-CNT pellets. 12

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Figure 4: Bright field TEM micrographs of the ZrB2-based composites (a) pure ZrB2 (b) Z20S (c)

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Z10C (inset shows the buckling and bending of CNT after SPS processing) and (d) Z20S10C showing SiC and CNT sits at the inter-granular junctions. Selected area diffraction (SAD) pattern

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showing (e) hexagonal ZrB2 (f) hexagonal SiC and (g) nanocrystalline CNT.

However, the bending and buckling of CNT can also be seen in the inset of Fig. 4c which microstructurally supports the increment in ID/IG ratio obtained during Raman analysis (Fig. 3b). Again, along with the presence of anisotropic ZrB2 grains, Fig. 4d shows the presence of both SiC and CNT, which is further confirmed by the selected area diffraction (SAD) patterns as 13

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shown in Figs. 4e-g, respectively. The SAD pattern for hexagonal ZrB2 (Fig. 4e) and 6H-SiC (Fig. 4f) indexed corresponding to the zone axis [111]. The nano-crystalline nature of the CNT is

3.4.

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evident in Fig. 4g. Phase and microstructural characterization after plasma arc-jet testing:

Pre- and post-oxidation investigation of ZrB2-SiC-CNT composites showed no erosion on the

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exposed face, which elicits that all composites survived the flux of 2.5 MW/m2 for 30 s, see Fig. 5. The test (with flux of 2.5 MW/m2) was enough to simulate thermal shock, which apparently is

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withstood by all the samples as the oxide scale was adherent (no occurrence of spallation or micro-cracks) for all the composites (see Fig. 5a). Apparently, surfaces of ZrB2 and Z10C have turned white due to the oxidation at high temperature, however, no considerable changes were visible in Z20S and Z20S10C. The back face rise in temperature for all plasma arc-jet tested

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UHTC composites (shown in Fig. 5a) is observed to decrease with SiC addition and increases with CNT reinforcement in the ZrB2 matrix.

Oxidation of ZrB2-SiC-CNT composites during plasma arc jet exposure has been confirmed

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from the subsequent appearance of monoclinic ZrO2 and SiO2 peaks obtained in XRD spectra (Fig. 6a), however retention of ZrB2 phase was observed and further quantified (on the surface

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oxide) as an average of area fraction of ZrB2 under XRD peaks, see Table 2. It is elucidated that the formation of surface oxide was suppressed with SiC addition as the amount of retained ZrB2 phase was increased from ~30 % (for ZrB2) to 47 % (for Z20S), 38 % (for Z10C) and 42 % (Z20S10C). The crystallite size (t) evaluated using Scherrer’s formula (Eq. 4) for all as-prepared samples are of similar order (~39-44 nm), however, the crystallite size noticeably increased after plasma arc-jet exposure (~53-62 nm), shown in Table 2.

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Table 2: Crystallite size, retention of ZrB2 phase, the thickness of the oxide layer, and oxidation rate after plasma arc-jet exposure of ZrB2-SiC-CNT composites.

Crystallite size (t) of ZrB2 in SPS pellet (nm)

Retention of

Thickness of

Linear oxidation

before plasma arc jet

after plasma arc jet

after exposure

layer (µ µm)

rate (µ µm/s)

ZrB2

44 ± 2

57 ± 2

~30%

22.9 ± 0.3

0.77

Z20S

39 ± 1

62 ± 2

~47%

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Sample

10.1 ± 0.1

0.34

Z10C

40 ± 2

53 ± 2

~38%

16.8 ± 0.3

0.56

Z20S10C

40 ± 1

55 ± 1

~42%

13.2 ± 0.3

0.44

the oxide

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ZrB2 phase

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Raman spectrum showed two distinctive peaks at 1348.3 and 1573.6 cm-1 for Z10C and 1349.2 and 1588.6 cm-1 for Z20S10C (Fig. 6b), which corresponds to D- and G-peaks of carbon, respectively [20]. Thus, the retention of CNT is confirmed in SPS processed ZrB2-SiC-CNT

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composites even after plasma arc jet exposure. Cross-sectional Raman line analysis across the oxide scale shown in Fig. 6c (along a line with distinct point, each at a distance of 5 µm) of Z20S10C samples is presented in Fig. 6d. It is elicited from the graph (Fig. 6d) that ZrB2, being

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Raman in-active, showed no peaks. However, D- and G-peaks were observed wherever CNT was present both in the oxide scale as well as the unaffected sample. Correspondingly, the presence

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of oxide bands at 630.9 cm-1, 959.1 cm-1 and 1011.4 cm-1 is mainly due to a vibrational mode involving Si-O-Zr linkages were also confirmed in the samples after plasma arc jet exposure [25]. Among these, the Raman peak at 959.1 cm-1 corresponds to the break-up of Si-O-Zr linkages due to the exsolution of ZrO2 as tetragonal zirconia [25]. Of the various reaction products mentioned in the reaction schemes (Eqs. 6-10), ZrO2 is stable at > 2000 °C with a melting point of ~2715 °C [26], also revealed from the XRD results. 15

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B2O3 has a low melting point (450 °C) and high vapor pressure and hence it vaporizes at temperatures above 1100 °C as per the below mentioned reactions [27]. 



345 678 + : 6;8 → 34: 678 + 5 := 6>8 2

(6)

2

5 := 6> 8 → 5 := 6;8 34: 678 + 5 := 6> 8 + 5 678 → 345 678 + 5 : 6;8 2 2 2 @A 678 + : 6;8 → @A: 678 + :6;8 3 3 3 @A: 678 + 3 678 → @A 678 + 2 : 6;8

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2

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(7) (8) (9)

(10)

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The plot of standard free energy change against absolute temperature (using FactSage [28]) for reactants and products in the aforementioned reactions (see Fig. 7), from which it is inferred that Eqs. (6–10) are feasible at temperatures ≥ 1196.5 °C, 575.8 °C, 1499.4 °C, 1222.6 °C and 1452.4 °C respectively under the pressure of BCD 10-15 atm (during re-entry in real-life application [29]).

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It is anticipated that the oxidation of ZrB2 to ZrO2 (Eq. 6, exothermic) and SiC to SiO2 (Eq. 9, exothermic) are responsible for the net weight gain (later explained in section 3.5.2) while Eq. 7 (endothermic) result as the weight loss. However, from Eqs. 8 and 10 (endothermic), it is elicited

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that the addition of CNT suppresses the formation of oxide scales (i.e. ZrO2 and SiO2), evident

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from the XRD analysis (Table 2).

Typical SEM images depicting composites before and after oxidation, are shown in Fig.

8. Figure 8a shows surface morphology of as processed ZrB2 with porosity (~7%) while the presence of SiC and CNT (darker spots, Figs. 8b-d) appears to be uniformly distributed in ZrB2 matrix. Formation of multiple cracks and deep pits seen on ZrB2 sample (Fig. 8e) are attributed to the escape of B2O3 (g) phase (see Eq. 7), resulting in pore formation which accelerates oxidation diffusion. In contrast, a few traces of ZrO2 with the absence of cracks was observed in 16

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Z20S (Fig. 8f). Z10C and Z20S10C samples in Figs. 8g and h showed fewer cracks and small pits throughout the sample.

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The cross-sectional SEM micrographs of plasma exposed samples elicit the thickness (presented in Table 2), morphology and composition (see Appendix 1) of oxide scale (ZrO2), shown in Figs. 9a-d and Figs. 9e-h respectively, however, the composition of oxide scale is presented in . It is to be mentioned here that the oxide scale spalled off from ZrB2 (Fig. 9a), has

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occurred during preparation of cross-sectional sample (sectioning followed by hot mounting and ultra-sonication). The thickness of oxide scale decreased with SiC and CNT reinforcement from

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23 µm (in ZrB2, Fig. 9a) to 10 µm (in Z20S, Fig. 9b), 17 µm (in Z10C, Fig. 9c) and 13 µm (in Z20S10C, Fig. 9d). It is evident from the microstructure (Figs. 9e-h) that the temperature experienced by the composites is in excess of 2715 °C, i.e. the melting temperature of ZrO2 [26]. The linear oxidation rate (calculated using Eq. 1), presented in Table 2 elicits that the oxidation

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rate decreases from 0.77 µm/s for pure ZrB2 to 0.34 µm/s (for Z20S), 0.56 µm/s (for Z10C) and 0.44 µm/s (for Z20S10C) has been attributed to the high thermal conductivity (discussed in the

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following section) with SiC and CNT reinforcement. TEM of plasma arc jet exposed Z20S10C sample showed the presence of CNT and

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transformed CNT as shown in Fig. 10a, which supplement the results obtained in Raman spectroscopic analysis (Figs. 6b-d) and microstructure obtained before plasma exposure (Figs. 8c and d). The morphology of the oxide scales of SiO2 and ZrO2 is elucidated from the Figs. 10b and c. The SAD pattern in Figs. 10d-f corresponds to the zone axis [111] for monoclinic ZrO2, triclinic SiO2, and CNT, respectively. The survival of CNT (along with damaged CNT), even after plasma arc jet exposure, insinuates the plausibility of CNT in resisting oxidation of

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ZrB2 by both physical (Fig. 10 a) and chemical means (Eqs. 8 and 10) as explained in the

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following section.

Figure 5: (a) Photographs of the SPS pellets before and after plasma-arc jet testing and (b) Back face temperature profiles of ZrB2-SiC-CNT composites with time when subjected to plasma arc jet for 30 s.

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Figure 6 (a): XRD spectra of plasma arc jet exposed ZrB2-SiC-CNT composites pellet (b) Top surface Raman spectra of Z20S and Z20S10C samples, (c-d) Optical image and corresponding cross-sectional Raman spectra of Z20S10C sample showing the presence of graphitic peaks even after plasma arc jet exposure. Orange coloured line indicates the direction in which Raman line analysis has been carried out.

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Figure 7: Plot of standard free energy change vs absolute temperature using FactSage [28] for

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the reactant and product.

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Figure 8: SEM micrographs of SPS processed ZrB2-based composites (a) ZrB2, (b) Z20S, (c) Z10C and (d) Z20S10C and (e-h) display cracks, pits and oxide scale obtained after plasma arc jet exposure. 21

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Figure 9: Back-scattered SEM micrographs showing the thickness of the oxide scale of (a) ZrB2, (b) Z20S, (c) Z10C and (d) Z20S10C and (e-h) are details of the oxide scale (secondary electron mode) in the respective composite. Double-headed arrow (in Figs. a-d) shows the oxide scale thickness, while single-headed arrows (in Figs. e-h) represent the melting and re-solidified grains.

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Figure 10: TEM micrographs of plasma exposed Z20S10C sample showing (a) oxide scale and the presence of CNT(represented by single-headed arrows), (b) SiO2 particle (inside feature represents twin boundary), (c) ZrO2 particle and (d-f) SAD pattern confirming the crystal structure of the ZrO2, SiO2 and retained CNT after plasma arc jet exposure.

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3.5.

Oxidation Behavior:

3.5.1. Thermal conductivity:

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The thermal conductivity of ZrB2-SiC-CNT based UHTCs was determined based on the values calculated from the correlation function of thermal diffusivity, specific heat capacity and density (Eqs. 2 and 3) are presented in Table 3. Thermal conductivity for pure ZrB2 (~48.9

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W/mK) is well in accordance with that reported in the literature [8], and is observed to increase with synergistic reinforcement of SiC and CNT (Z20S10C) to 61.8 W/mK. As the temperature

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increases (from 50-1200 °C), defect concentration in the material increases, which leads to phonon scattering (known as Umklapp scattering) and, thus, a decrease in the thermal conductivity. The addition of SiC and CNT increases the thermal conductivity based on the dispersed phase models [14, 19], providing a conductance path (along the grain boundary, see Fig. 6) for heat dissipation, thereby, providing improved oxidation resistance (the high bow

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shock associated with plasma arc-jet is enough to generate thermal shock). However, a little decrease in thermal conductivity values, in case of CNT reinforced composites (Z20S10C, when compared with Z20S), is attributed to the distribution and alignment of CNT because CNT

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reinforcement shows higher thermal conductivity when it is aligned in one direction) [18]. The

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measured thermal conductivity values are higher than for some alloys and ceramic matrix composites (CMCs, for example: at 300K, λ ~ 13.6 W/mK for Inconel 617 and λ ~ 14.5 W/mK for a standard C/C-SiC [30].

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Table 3: Thermal diffusivity, specific heat capacity, thermal conductivity and back face rise in temperature of ZrB2-SiC-CNT based composites.

Sample

50 °C D

Cp

300 °C D

κ

Cp

600 °C κ

D

Cp

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Thermal properties of ZrB2-based composites at different temperatures 900 °C κ

D

Cp

1200 °C

κ

D

Cp

Tb (°C) κ

0.236 340.4 48.9 0.180 415.9 45.4 0.147 512.7 45.7 0.131 536.7 42.7 0.125 559.7 42.3

Z20S

0.276 446.0 67.6 0.197 623.5 67.6 0.173 668.5 65.2 0.151 732.8 60.5 0.121 761.2 50.6

853.2

Z10C

0.204 415.0 48.9 0.159 549.7 49.8 0.133 744.8 56.2 0.130 632.6 46.6 0.130 671.3 49.5

1081.3

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ZrB2

Z20S10C 0.240 485.2 61.8 0.177 624.3 58.7 0.160 799.0 67.4 0.127 802.4 51.4 0.121 816.7 52.3 where, D in mm2s-1, Cp in Jkg-1K-1, κ in W/mK and Tb = back-face temperature. The rise in back face temperature (see Table 3) is attributed to the thermal conductivity,

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porosity content, and emissivity of the material. Rise in temperature (Tb) of ZrB2 (1034 °C) in spite of lower thermal conductivity as compared to reinforced composites is due to high porosity content (~7%) in the material as per the effect of phonon-pore scattering on thermal conductivity

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[18]. Conversely, the maximum rise in temperature in Z10C (1081 °C) is due to the high thermal conductivity of the CNT, however, composites with SiC (i.e., Z20S (853.2 °C) and Z20S10C

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(842.3 °C)) show the least rise in temperature due to their higher surface emissivity (~0.98) [31]. Also, composites with SiC reinforcement forms oxide scale of SiO2-ZrO2 on the surface (passive oxidation), which hinders the rise in back face temperature (Tb) due to the poor thermal conductivity of oxide.

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842.3

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3.5.2. Thermal studies using TGA/DSC: Thermal studies using TGA (although, conditions are not as severe as that of plasma arc jet

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exposure) has been carried out to provide an insight to the oxidation mechanisms occurring in ZrB2-SiC-CNT composites. The final increase in the weight for the composites are: 24%, 5%, 18% and 2 % respectively for ZrB2, Z20S, Z10C and Z20S10C (see Table 4) illustrating that the

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resistance to oxidation increases with SiC and CNT addition (see Fig. 11 a), due to the formation of SiO2 (Eq. 9) and sealing mechanism offered by CNT [19]. Figure 11a elicits the shift to higher

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onset temperature of oxide formation after which a rapid weight change is observed (Table 4), i.e. 718.5 °C, 697.1 °C and 706.3 °C for Z20S, Z10C, and Z20S10C, respectively, when compared to that of pure ZrB2 (679.3 °C). This oxidation products (in liquid form) formed during the initial stage can penetrate into the pellet through surface pores [32] and offers good ablation

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resistance [32, 33] when compared to that of pure ZrB2.

Table 4: Thermal studies of ZrB2-SiC-CNT composites up to 1500 °C in oxygen atmosphere % wt.

Onset rapid weight

Peak temperature of

Enthalpy of oxide

gain

change temperature (°C)

oxide formation (°C)

formation ∆H (kJ/g)

679.3 ± 6.4

652.2 ± 2.8

-1.6 ± 0.1

718.5 ± 3.8

661.9 ± 4.4

-0.3 ± 0.1

24

Z20S

5

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ZrB2

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Composition

Z10C

18

697.1 ± 3.7

676.1 ± 2.9

-0.8± 0.2

Z20S10C

2

706.3 ± 2.1

690.5 ± 3.1

-0.6 ± 0.1

Figure 11b showed an exothermic peak at ~652.2 °C for pure ZrB2, which shift towards higher temperature i.e., 661.9 °C, 676.1 °C and 690.5 °C respectively for Z20S, Z10C, and Z20S10C with SiC and CNT addition. The presence of an endothermic peak at around 1350 °C, observed in case of ZrB2 and Z10C suggest a possible transformation of m
t ZrO2 [10]. Based 26

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on this, the efficacy of SiC reinforcement in protecting ZrB2 from oxidation and the absence of transformation of m
t ZrO2 is well elucidated. The enthalpy (∆H) of formation of oxide decreases from 1.6 kJ/g for pure ZrB2 to 0.3 kJ/g, 0.8 kJ/g and 0.6 kJ/g for Z20S, Z10C and

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Z20S10C respectively, insinuates an increase in the thermal stability with SiC and CNT

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reinforcement.

Figure 11: Thermal analysis at 1500 °C on ZrB2-SiC-CNT based composites: (a) TGA and (b) DSC.

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3.5.3. Effect of SiC and CNT reinforcement on oxidation mechanism: The schematic illustration of the factors that govern the oxygen transport mechanism (SiC

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and CNT reinforcement; processing induced defects) is shown in Figs. 12a-c. The presence of dislocation network in ZrB2 grain (see TEM image, Fig. 12d) act as an active site, which triggers the diffusion of oxygen via grain and grain boundaries (GBs) [34]. In the case of SiC reinforced

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composites, oxygen can transport via ZrB2, SiC particles, and GBs (see Fig. 12c), however, transport through GBs is favored over particles [19]. Also, the rapid oxidation of SiC is usually

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observed above 1700 °C, therefore, it protects ZrB2 by the formation of passive SiO2 (see Fig. 12b), further explored by the negligible oxidation observed during thermal studies till 1500 °C (see Fig. 11a). Moreover, the presence of low angle grain boundaries (LAGBs) at the interface of ZrB2 and SiC (TEM image, Fig. 12e) act as dormant site (having low energy) and thus provides oxidation resistance [35]. Similarly, the survival of CNT from harsh plasma jet exposure (shown

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earlier in Figs. 6 and 10) establish restriction of oxidation of ZrB2 via grain sealing mechanism (Fig. 12c and f) [19]. Oxygen diffusion through interfaces may be preferred, thus, CNTs, sealing the grains, may be sacrificially oxidized and protect the underlying ZrB2 and SiC (see Eqs. 8 and

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10). Succinctly, the synergy of SiC and CNT reinforcement ascertain ZrB2 as a potential

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structural material with high damage tolerance.

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Figure 12: Schematic illustration of

oxygen transport in (a) ZrB2 matrix, (b) with SiC

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reinforcement (c) grain sealing hindering diffusion of oxygen with CNT addition and (d-f) corresponding TEM image eliciting the microstructural features (after SPS processing) affecting

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oxidation kinetics. Here, GB = grain boundary (in Figs. a-c) and LAGBs = low angle grain boundaries (in Fig. e).

3.5.4. Evaluating mechanical performance of ZrB2-SiC-CNT composites: To further investigate the structural integrity of ZrB2-based composites, the mechanical properties of the as-processed and after plasma arc-jet tested ZrB2-SiC-CNT composites were 29

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evaluated using instrumented indentation technique. The hardness and Young’s modulus, tabulated in the Table 5 are calculated from the load vs displacement graph (shown in Fig. 13) using Oliver-Pharr indentation method for ZrB2-based composites before and after plasma arc jet

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exposure. Vickers’ hardness for as-processed ZrB2 is increased from 13.2 GPa to 19.3 GPa for Z20S, 18.5 GPa for Z10C and 21.0 GPa for Z20S10C, attributed to the higher relative densities of composites after SPS processing. It may be noted that plasma arc jet exposed samples showed

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lower hardness than that of unexposed samples. However, a similar trend of hardness was maintained after plasma exposure where it increases from 6.2 GPa for pure ZrB2 to 12.0 GPa (for

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Z20S), 10.2 GPa (for Z10C) and 12.8 GPa (for Z20S10C). Z20S10C showed the highest hardness values elucidating the synergy of SiC and CNT reinforcement in ZrB2 matrix before and after plasma arc jet exposure.

The measured Young's modulus for as-processed ZrB2 (~93% dense) ∼378 GPa is in

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agreement with the reported value [36].This value increases to 405 GPa for Z20S, 401 GPa for Z10C and 411 GPa for Z20S10C when compared with pure ZrB2. Correspondingly, the modulus decreases to 212 GPa for monolithic ZrB2 (after plasma arc jet exposure) which then increases

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with SiC and CNT reinforcement to 331 GPa (for Z20S), 288 GPa (for Z10C) and 342 GPa (for

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Z20S10C) after plasma arc jet exposure. The H3/E2 ratio determined experimentally by instrumented indentation (Table 5) confirms an increase in resistance to plastic deformation with SiC and CNT reinforcement in ZrB2 before and after plasma arc jet exposure. Based on the hardness and modulus of these composites, the H3/E2 ratio was determined

(see Table 5), which confirm a higher resistance to plastic deformation with SiC and CNT reinforcement in ZrB2 before and after plasma arc jet exposure when compared to that of ZrB2.

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Figure 13: Load-displacement curves for ZrB2-SiC-CNT composite before and after plasma arc jet exposure.

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Table 5: Comparison of mechanical properties of as-prepared and plasma arc jet exposed ZrB2-

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SiC-CNT composites using instrumented indentation technique. Before oxidation

Hv

E

H3/E2

(GPa)

(GPa)

(MPa)

ZrB2

13.2±1.0

378±5

16.1

Z20S

19.3±0.6

405±6

Z10C

18.6±0.8

Z20S10C

21.0±0.7

Sample ID

After Oxidation Hv

E

H3/E2

(GPa)

(GPa)

(MPa)

32.6

6.2±1.4

232±11

3.1

43.8

78.0

12.0±0.6

331±9

15.8

401±10

40.1

54.3

10.2±0.9

288±8

12.8

411± 7

54.8

81.1

12.8±0.5

342±2

17.9

σf (MPa)

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The load versus displacement relationship during the diametrical compression test (see Appendix 2) is used to calculate fracture strength of as processed ZrB2-based composites (see Table 5). The fracture strength obtained for pure ZrB2 (32.6 MPa) is in agreement with the

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reported value [8]. The fracture strength increased to 78 MPa (2.5 times for Z20S), 54.3 MPa (1.8 times for Z10C) and 81.1 MPa (2.7 times for Z20S10C) when compared to pure ZrB2. The increased strength with reinforcement indicates higher load bearing capacity of composites via

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CNT bridging and interfacial bonding of matrix with reinforcement [20].

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3.5.5. Effect of residual stress on structural stability of ZrB2-SiC-CNT composites: There are several strengthening mechanisms that occur in composite materials: (i) crack bridging (ii) micro-cracking when the misfit between CTE introduces a stress field, and (iii) residual stress toughening due to difference in thermal expansion coefficient (CTE) of matrix and reinforced particles, which creates a local compressive stress field in the matrix, thus,

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decreasing the stress intensity factor. Herein, authors have isolated the contributions of thermal residual stresses in composites, which arise due to the coupling of different phases with different thermo-elastic properties. Hence, the strength of the ZrB2 is governed by the reinforcement of

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SiC and CNT (ignoring the effect of distribution and size) and the generation of interfacial

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residual stresses (computed analytically) in the material. due to mismatch in the CTE of matrix and reinforced particles. It may be mentioned that the extrinsic effects (such as glassy phase content, micro-cracking, porosity, etc.) have not been accounted for in estimating residual stresses at the interface. The linear thermal expansion coefficient CTE of a two-phase composite (assuming that each phase is isotropic), computed by Levin [37], Rosen and Hashin [38], is provided as:

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6FGHH 8I = HJ FJ + H4 F4 +

(F J −F 4 ) 1 KJ

1 −K 4

L

1

KGHH

−

1

KJ

−

1

K4

M

(13)

Considering that K ≤ K -- ≤ K , the upper and lower bounds of K can be obtained

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from Hashin-Shtrikman bounds [39]. Under such conditions, the above Eq. (13) is further modified to calculate both the upper and lower bounds of α in the following manner: K4 (3KJ +4 PJ ) KJ (3K4 +4 PJ )+4 H4 PJ (K4 − KJ )

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(FGHH )I = FJ − H4 (FJ − F4 )

KJ (3K4 +4 P4 ) (3K 4 J +4 P4 )+4 HJ P4 (KJ − K4 )

(15)

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(FGHH )> = F4 + HJ (F4 − FJ ) K

(14)

where, (F -- ) and (F -- ) are the upper and lower bounds of CTE respectively of a given composite whereas α, K, P and H are CTE, bulk modulus, shear modulus and volume fraction with subscript “m” and “r” set for matrix and reinforcement respectively. The values of ν (for

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ZrB2 = 0.17, SiC = 0.14 and CNT = 0.17, calculated using ROM), α ( for ZrB2 = 6.8 x 10-6 K-1, SiC = 3.5 x 10-6 K-1 and CNT = 2.5 x 10-6 K-1), K (for ZrB2 = 229 GPa, SiC = 234 GPa and CNT = 190 GPa) and P (for ZrB2 = 211 GPa, SiC = 41 GPa and CNT = 150 GPa) used for

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calculations (are taken from the literature [19, 20, 40].

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The quantification of such thermal residual stresses in composites becomes primarily important, since processing has been carried out at a very high temperature where these stresses develop (ΔT). Since, the composites are being cooled from the final densification temperature (1650 °C) to room temperature; ZrB2 (6.8×10-6 K-1) tends to shrink faster than the reinforced particles (SiC, 3.5×10-6 K-1 and CNT, 2.5×10-6 K-1) leading to tensile stress state in the matrix, σm and corresponding compressive stress in the reinforced particles, σr as evaluated using Taya’s model [41], tabulated in Table 6. 33

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, Q −

6R-8 S T -

,U

(16) -WS T

XY

\

" = RYZ \ [ [

(17)

Z 8 X =W- 6RYZ 8

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,U Q + VU 6R-86WX86XY

Z

] ^ = 6F − FU 8 _`

(18)

(19)

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where, V , VU , a and aU are the Young’s modulus and Poisson’s ratio (for ZrB2 = 0.17, SiC =

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0.14 and CNT = 0.17, calculated using ROM [19, 20, 40]) of the reinforcement and matrix respectively; ] ^ is the thermal expansion misfit strain and _` is the temperature at which stresses begin to accumulate (set as 1400 °C) [42, 43]. Young’s modulus for the reinforced phase, Er is estimated by ROM, considering modulus of matrix (ZrB2) as Em = 378 GPa and modulus Ec for

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each composite is both from values presented in Table 4.

Correspondingly, the developed biaxial residual stress (σ) is obtained from the relation: VJ M ] ′ 61 − aJ 8

(20)

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, = L

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Table 6: Theoretical calculation of the coefficient of thermal expansion (CTE), modulus and residual stresses. CTE (×10-6 /K) Rule of Mixture

Hashin-Shtrikman model

ID

Taya’s model

Lower bound

ZrB2

N.A.

N.A.

N.A.

Z20S

6.14

6.15

6.13

Z10C

6.37

6.35

6.30

Z20S10C

5.71

5.69

where, N.A= not applicable

5.62

Er

Biaxial

σm

σr

N.A.

N.A.

N.A.

N.A.

378 405

513

630.1

-34.3

520.1

401

608

352.4

-43.1

347.6

488

428.0

-39.4

481.7

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Upper bound

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Ec (ROM)

Residual stress (MPa)

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Modulus (GPa)

Sample

411

The stability of the oxide scale formed also depends on internal stresses developed during

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plasma arc jet exposure. After the plasma arc jet exposure, both the oxide and the remaining composite contract upon cooling [37], leading to compressive strain in oxide scale and tensile strain in the remaining composite. So, the oxide scale relaxes (as the cooling is very rapid) by

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generating cracks on the surface as observed in Fig. 8. It is elucidated from Table 6 that the addition of secondary phase (SiC and CNT) introduce an interfacial stress (compressive in the

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reinforcement and tensile in the matrix), which provides enhanced structural integrity and toughness [20]. The values of compressive stress ranges from 34.3 MPa (for Z20S) to 43.1 MPa (for Z10C) and 39.4 MPa (for Z20S10C), but are only ~5% of that tensile stress in the matrix. It is observed that the residual tensile stress in the matrix decreases with reinforcement of SiC and CNT (Table 6, Taya’s model). The biaxial model follows the same trend of residual stress with SiC and CNT reinforcement. Thus, CNT-reinforced composites (Z10C and Z20S10C) showed structural stability (due to a reduction in residual tensile-stresses) along with the thermal stability 35

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established by thermal studies (discussed earlier in section 3.5.2.). Hence, enhancement in the interfacial residual compressive stresses with the synergistic addition of SiC and CNT make an

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assessment of structural integrity which is mandatory for aerospace applications. Conclusions:

ZrB2-based UHTC composites were spark plasma sintered achieving enhanced densification

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from 93% to ~99.8% with synergistic reinforcement of SiC and CNT. Oxidation behaviour of ZrB2-SiC-CNT composites was evaluated through plasma arc jet with 30 s exposure under heat

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flux of 2.5 MW/m2. The decrease in the linear oxidation rate from 0.77 µm/s (for ZrB2) to 0.34 µm/s (for Z20S), 0.56 µm/s (for Z10C) and 0.44 µm/s (for Z20S10C) is attributed to the enhanced thermal conductivity with SiC and CNT addition. The structural integrity of these UHTC composites was measured in terms of hardness which increases up to 1.6 times (13.2 to

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21.0 GPa) before and up to 2.1 times (6.2 to 12.8 GPa) with synergistic reinforcement of SiC and CNT in ZrB2 after exposure to harsh aerothermal conditions. Raman spectroscopic analysis revealed that CNT survived the extreme oxidising environment, further supplemented by TEM

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analysis of Z20S10C. In addition, enhanced thermal conductivity with SiC and CNT addition in ZrB2 indicates an improved oxidation resistance via heat dissipation mechanism. The high

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temperature shift of the oxide formation (from 679 °C to 706 °C) and decrease in the enthalpy of oxide formation (from 1.6 to 0.6 kJ/g), elicits enhanced thermal stability and oxidation resistance with synergistic reinforcement of SiC and CNT in ZrB2. The synergy of SiC and CNT in providing oxidation resistance (protective SiO2 formation by SiC and grain sealing by CNT) and enhanced thermal conductivity (from 42.3 W/mK to 52.3 W/mK at 1200 °C) is corroborated. These findings demonstrate that the interplay of reinforcement affects the generation of residual stresses at the interface (analytically quantified), which, then, governs the strength of ZrB2-SiC36

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CNT composites, making them a suitable candidate for application in high-damage tolerant structures.

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Acknowledgements: Authors at Indian Institute of Technology Kanpur (IITK) acknowledge the financial support received from IITK-Space Technology Cell and Indian Space Research Organization (ISRO),

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Trivandrum, India. Mr. Vincent Xavier (Technical Staff) and Mr. Fazil Mohammad are acknowledged for helping with plasma-arc jet exposure testing and thermal conductivity

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measurements, respectively, at ISRO. KB acknowledges P.K. Kelkar fellowship, IITK. Authors also acknowledge TEM Facility, MSE Department, IITK and the support from the Advanced Centre for Materials Science (ACMS) at IITK for extending characterization facilities (SEM, mechanical testing and instrumented indentation testing).

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Supporting Information:

Appendix 1: elemental mapping along the cross-section of Z20S10C sample after plasma jet exposure.

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Appendix 2: Load versus displacement curve showing the displacement and maximum load

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before fracture for SPSed ZrB2-based composites during diametrical compression test.

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