Tyranno SA3 fiber–ZrB2 composites. Part I: Microstructure and densification

Materials and Design 65 (2015) 1253–1263

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Tyranno SA3 fiber–ZrB2 composites. Part I: Microstructure and densification Laura Silvestroni ⇑, Daniele Dalle Fabbriche, Diletta Sciti CNR-ISTEC, Institute of Science and Technology for Ceramics, Via Granarolo 64, I-48018 Faenza, Italy

a r t i c l e

i n f o

Article history: Received 4 June 2014 Accepted 26 August 2014 Available online 6 September 2014 Keywords: Zirconium diboride Fibers Sintering additive Densification Transmission electron microscopy

a b s t r a c t ZrB2-ceramics containing 3rd generation SiC short fibers were hot pressed with different kinds of sintering agents in order to obtain an increase of fracture toughness. Several sintering temperatures were explored keeping fixed the amount of fiber to define the most suitable additive able to fully densify ZrB2 with minimal fiber deterioration. In Part I of this article, the microstructure of the composites was examined through scanning electron microscope, to study the distribution of the secondary phases, and transmission electron microscope, to analyse the microstructure at nanoscale level, with particular attention to the evolution of the fiber morphology and to the study of the interfaces with the matrix in high resolution mode. For the first time the chemical stability of Tyranno SA3 fibers towards different sintering additives and ZrB2 is presented. Comparison with Hi-Nicalon fiber is also discussed. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction Zirconium diboride (ZrB2), belonging to the ultra-high temperature ceramics (UHTCs), is considered among the most appealing materials for aerospace applications owing to its high melting temperature and better performances over metals in high-temperature environments [1]. One aspect that deserves major efforts regards the low fracture toughness of these materials, ranging from 2.5 to 4.5 MPa m1/2 [2], since higher defect tolerance would open a wider scenario of uses and expand the market for this class of compounds. In the last five years, we developed UHTC composites containing SiC chopped fibers [3–7]. Under appropriate processing conditions, the addition of short SiC fibers leads to significant increments of the fracture toughness, i.e. over 6 MPa m1/2. The main advantage of using short fiber instead of continuous fiber is the possibility to follow the same ceramic processing techniques to manufacture them, i.e. ball milling, mixing, and sintering. On the other hand, previous works on composites containing chopped Hi-Nicalon SiC fiber [6,7] pointed out that these fibers, being composed in large part by unstable amorphous Si–C–O phase and carbon, tend to react during sintering resulting in a highly modified structure strongly bonded to the matrix. A possible solution for this issue is the interposition of a weak interface between fiber and matrix, such as BN or pyrolytic carbon, to retard reactions between fiber and matrix during high thermal treatments and to trigger ⇑ Corresponding author. Tel.: +39 546 699723; fax: +39 546 46381. E-mail address: [email protected] (L. Silvestroni). http://dx.doi.org/10.1016/j.matdes.2014.08.068 0261-3069/Ó 2014 Elsevier Ltd. All rights reserved.

relevant toughening mechanisms, such as fiber pullout [8], in analogy with the CMC technology [9]. However, deposition of a BN coating is a very expensive procedure and not all the coatings are suitable for these composites. Indeed preliminary results in our labs revealed that amorphous BN coatings tend to detach from the fiber during milling and crystallize during sintering, forming phases that weaken the whole microstructure. As for pyrolytic carbon, again, preliminary results did not evidence any benefits over uncoated fibers. Therefore the use of refractory naked SiC fibers, able to maintain their shape and properties is at present the most feasible solution. In this context, the study of the interface between fiber, matrix and sintering additive is of paramount importance. Tyranno SA3 are classified among stoichiometric fibers with little intergranular residual carbon [10] and b-SiC crystallites with dimensions below 100 nm [11]. They have smaller diameter, 7.5 lm compared to 15 lm of the Hi-Nicalon, and are reported to possess higher stability at high temperature and thus are expected to display better performances in extreme environment. This type of refractory fibers is therefore considered more attractive owing to the excellent mechanical properties coupled with resistance to neutron irradiation, improved thermal conductivity and thermal stability [12,13]. In this paper, ZrB2-based composites were produced with various sintering agents, such as ZrSi2, Si3N4, TaSi2 and MoSi2. Different sintering temperatures were explored to identify the best conditions leading to minimal fiber degradation and maximal densification. The purpose of this study was hence to understand the chemical interaction between Tyranno SA3 fiber and ZrB2 matrices

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and compare it to the case of Hi-Nicalon fiber. A thorough study of the microstructural features is given in Part I, whereas Part II is devoted to the correlation of these microstructural features to the local and overall thermo-mechanical properties. 2. Experimental procedure The following commercial powders were used to produce ZrB2based composites, whose compositions and sintering parameters are reported in Table 1: – Hexagonal ZrB2 (H.C. Starck, Grade B, Germany), specific surface area 1.0 m2/g, particle size range 0.1–8 lm, impurities (wt%): 0.2 C, 1.5 O, 0.25 N, 0.1 Fe, 0.2 Hf. – Orthorhombic ZrSi2-F (Japan New Metals Co., LTD, Osaka, Japan), particle size 2–5 lm, impurities (wt%): 0.11 C, 1.00 O, 0.09 Fe. – a-Si3N4 (Baysinid Bayer, Germany), specific surface area 12.2 m2/g, mean particle size 0.15 lm, impurities (wt%): 1.5 O. – Hexagonal TaSi2 99.0%, (ABCR, Gmbh & Co. Karlsrhue, Germany), 5–10 lm FSSS, mean particle size 6.82 lm, impurities (wt%): 0.05 C, 0.21 O, 0.03 Fe. – Tetragonal MoSi2 (Aldrich, Germany), specific surface area 1.60 m2/g, impurities (wt%): 1 O. – Tyranno SA3 SiC fibers (UBE Europe GmbH, Dusseldorf, Germany) with composition (wt%) Si:C:O = 67:31:<0, Al < 2, diameter: 7.5 lm, chopped length: 100–200 lm. For scouring purpose, 3 vol% SiC chopped Tyranno fibers and 5 vol% of additive were added to ZrB2 to study the fiber-sintering additive-matrix interaction at various temperatures. For the optimized composites, 15 vol% of SiC fibers were added and the amount of sintering additive was kept as minimum as possible to enable full densification. The starting fiber was a continuous spool which was in house chopped, Fig. 1. The chopped fibers were mixed to the powders through ball milling for 24 h in absolute ethanol using ZrO2 milling media. Subsequently the slurries were dried in a rotary evaporator. After de-agglomeration of the powder mixtures, 30 mm-diameter pellets were shaped by uniaxial pressing with 20 MPa. The sintering was preceded by a desizing cycle in a graphite furnace (Onyx Furnace 22001C, LPA DVM, Seyssinet, France) at 500 °C for 1 h holding time, with a heating rate of 50 °C/h under flowing argon atmosphere (0.1 MPa). Subsequently, ZrB2 composites were hot pressed in low vacuum (100 Pa) using an inductionheated graphite die with an uniaxial pressure of 40 MPa during the heating and increased up to 50 MPa at the dwell temperature. The maximum sintering temperature was set on the basis of the shrinkage curve (see Table 1). Free cooling followed. The bulk den-

sities were measured by Archimedes’ method and confirmed by SEM inspection. The microstructure was analysed on fractured and polished surfaces by scanning electron microscopy (FE-SEM, Carl Zeiss Sigma NTS Gmbh, Oberkochen, DE) and energy dispersive X-ray spectroscopy (EDS, INCA Energy 300, Oxford instruments, UK). TEM samples were prepared by conventional procedure by cutting 3 mm discs from the bulk, mechanically grinding and ion beam thinning. Local phase analysis was performed using transmission electron microscopy (TEM, FEI Tecnai F20 ST) equipped with a Schottky emitter and operating at a nominal voltage of 200 keV, combined with a EDAX EDS X-ray liquid nitrogen cooled spectrometer PV9761/II with a super ultra-thin window. Scanning Transmission Electron Microscopy (STEM) measurements was performed on selected composites with the same TEM in nanoprobe mode providing a point resolution 0.19 nm. Quantitative calculations of the microstructural parameters, like residual porosity, mean grain size and secondary phase content, were carried out via image analysis with a commercial software package (Image-Pro PlusÒ version 7, Media Cybernetics, Silver Springs, MD, USA). 3. Results 3.1. Densification behavior Table 1 reports the final relative density of the ZrB2 composites as a function of sintering temperature and 3 vol% additive. It can be read that ZrSi2 is the compound allowing to achieve good density levels (>96%) at the lowest temperature, 1550 °C, followed by Si3N4 at 1700 °C, and TaSi2 at 1750 °C, whilst MoSi2 required 1800 °C and increased holding time to obtain density higher than 96%. Notable amount of porosity, around 5 vol%, was present in composites sintered with MoSi2 at 1750 and 1800 °C and with Si3N4 at 1650 °C. ZrB2 mean grain size as a function of the sintering temperature and the additive is reported in Table 1. Surprisingly, the composites sintered at the lowest temperature, ZZ-T, had the coarsest grains, 2.5–3.0 lm, indicating fast mass transfer and grain coarsening. The other systems had mean grain size from 1.5 to 2.2 lm with little difference from the two sintering temperatures. One interesting feature regards ZT-T system, where ZrB2 grains displayed elongated shape with maximum aspect ratio up to 4–5. Since MoSi2 induced notable fiber alteration (see Section 3.3), further materials optimization and microstructural analyses were focused only on the systems with ZrSi2, Si3N4 and ZrSi2. 3.2. Microstructural features of baseline composites In order to understand any microstructural change upon the introduction of fiber, it is useful to briefly summarize the

Table 1 Composition, hot pressing conditions, density and microstructural details of the reinforced composites estimated through image analysis. Label ZZ-T

Sint. add. (vol%)

SiC fiber (vol%)

Ton (°C)

Tmax (°C)

Dwell at Tmax (min)

q fin

q rel.

(g/cm3)

(%)

5 ZrSi2

3

1400

1550 1600 1600

10

5.71 5.77 5.53

1590

1650 1700 1720

10

1600

1750 1800 1750

1750

1750 1800

10 ZrSi2

15

ZS-T

5 Si3N4

3

ZT-T

5 TaSi2

3

10 TaSi2

15

5 MoSi2

3

15

ZM-T

ZrB2 m.g.s. (lm)

Core area (%)

SiC m.g.s. core (nm)

SiC m.g.s. rim (nm)

96 97 100

2.4 ± 0.8 2.7 ± 0.8

78.8 76.7

74 ± 17 82 ± 19

226 ± 42 244 ± 76

5.52 5.66 5.29

94 96 96

1.8 ± 0.5 2.1 ± 0.7

43.3 35.8

76 ± 14 54 ± 18

194 ± 47 223 ± 36

15

5.91 6.03 5.77

96 97 97

1.2 ± 0.8 2.4 ± 0.8

46.8 34.7

86 ± 17 64 ± 15

259 ± 72 268 ± 82

20

5.73 5.84

93 96

1.9 ± 0.7 1.8 ± 0.8

14.2 0.0

91 ± 39 –

293 ± 93 319 ± 76

TEM

X

X

X

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Fig. 1. Fiber section of Tyrannno SA3 fiber with EDS spectra of SiC crystallite and C-rich dark phase.

compositional structure of the unreinforced composites. A thorough characterization of the baseline composites can be found in [14–16] and the main features are summarized in Table 2. ZrB2–ZrSi2 system – This material was fully densified at 1600 °C. ZrB2 mean grain size was around 2.5 lm with less than 10 vol% of other secondary phases: ZrO2, ZrSi2, SiC and SiO2-based species, Fig. 2a. No wetted grain boundaries were noticed by SEM and TEM inspections, Fig. 2b and c. Often, SiO2 pockets surrounding elongated SiC crystals were found [15,16]. ZrB2–Si3N4 system – The baseline ZrB2 material doped with 5 vol% of Si3N4 was fully dense with a mean grain size around 2.7 lm. The secondary phases observed in the sintered microstructure were mainly concentrated at triple points and were identified as BN, ZrN, SiC, ZrO2, Zr–Si phases, discrete pockets of SiO2 and a borosilicatic glass containing Zr–Si–B–N–O, Fig. 2d–f. The formation of these secondary phases is a consequence of Si3N4 dissociation and interaction with oxides covering ZrB2 particles, B2O3 and ZrO2, as discussed in Section 4.1 [15,16]. ZrB2–TaSi2 system – This composite required 1850 °C to achieve the full density. ZrB2 grains had an average size of about 2 lm and were surrounded by a (Zr, Ta)B2 solid solution, Fig. 2g and h. According to quantitative EDS analysis and XRD, the composition of this solid solution was (Zr0.8Ta0.2)B2 [14]. Pure TaSi2 was not clearly identified in the composite, but bright TaxSiy phases were observed in amount around 5 vol%. About 1.6 vol% of ZrO2 particles were also found along with 2 vol% of silica-based phases containing various impurities. Spurious carbide phases, such as (Zr, Ta)C, SiC or Si–C–O, were also detected in limited amounts, below 2 vol%, Table 2. In most cases, wetted interfaces were found in the matrix, with variable glass composition in the Ta–Zr–Si–B–C– O system, Fig. 2i. 3.3. Fiber evolution Pristine fiber – After chopping, the average length of Tyranno SA3 fiber was 100–300 lm, with a stochastic nature of the diameter, from 5 to 8 lm. In the cross section of the SEM image of Fig. 1 it can be seen that the fibers are characterized by SiC crystallites with rounded shape and dimensions around 70 nm, surrounded by carbon-rich dark pockets. Image analyses confirmed that about

15 vol% of the fiber is composed by this dark phase. The fracture is mainly intergranular, owing to the presence of carbon itself acting as weak interphase. Excess of carbon is actually reported to inhibit grain growth and to maintain equiaxed microstructure [17]. Sintered fiber – The effect of the sintering temperature and the additive on the fiber modification is reported in Fig. 3. As a general trend it can be observed that the fiber section maintained a rounded shape with smooth interface, but with increasing temperature, the fiber shape tended to become more irregular and, at a temperature of 1750 °C or higher, coarsening and detachment of SiC crystallites and shrinkage of the fiber section were observed. By SEM imaging, it was ascertained that SiC crystallite coarsening was not homogeneous inside the fiber, but generally involved the outer part of the fiber, called rim, whereas in the central part, called core, was almost negligible. The rim was also characterized by lower amount of carbon pockets. The core and rim area as a function of the sintering additive were calculated by image analysis and Table 1 evidences that in ZZ-T composites the fibers remained closest to the original ones, whilst in ZM-T the fibers were highly modified. SiC dimensions in the core and rim were estimated through image analysis, Table 1. The crystallites mean grain size in the core was around 70 nm for all the systems, at all temperatures, for all sintering additives, whilst the grain size of SiC in the rim increased with the temperature, least for ZS-T, 195 nm, most for ZM-T, 350 nm. It can be noticed that ZZ-T composites showed similar mean grain size of the crystallites in the rim as ZT-T one, although sintered 200 °C lower. 3.4. Microstructural features of composites with 15 vol% SiC fiber ZrB2–ZrSi2–SiCf system – ZZ-T achieved the full density at 1600 °C, but despite the low the sintering temperature, discrete ZrB2 grain coarsening occurred achieving values up to 5.2 lm. A magnification of the fiber section is reported in Fig. 4a. Carbon pockets are homogeneously diffused all over the fiber up to about 200 nm from the fiber edge, where dense SiC grains compose a thin rim. The darker regions surrounding the coarsened SiC crystallites contain Zr traces. TEM images of the fiber in Fig. 4b evidence the core-rim regions and show the nature of the intergranular phase,

Table 2 Composition, sintering conditions and main microstructural features of the unreinforced ZrB2-based composites. m.g.s.: mean grain size. Label

Additive (vol%)

Sintering (°C, MPa, min)

Rel. density (%)

ZrB2 m.g.s. (lm)

Secondary phases (vol%)

ZZ ZS ZT

10 ZrSi2 5 Si3N4 10 TaSi2

1600,30,10 1700,30,15 1850,30,10

96.5 99.0 99.0

2.5 2.7 2.0

3.2 ZrO2, 3 ZrxSiy, 1.5 SiO2, 1.3 SiC 1.5 SiC, 1.5 BN, 1.3 ZrN, 1 ZrO2, 1 SiO2, 0.8 ZrxSy, Zr–Si–B–N–O glass (Zr, Ta)B2, 5.0 TaxSiy, 2.0 SiO2, 1.6 ZrO2, 1.5 SiC, 1 (Zr, Ta)C, Ta–Zr–Si–B–C–O glass

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Fig. 2. Microstructural features of baseline materials: (a–c) ZrB2–ZrSi2, (d–f) ZrB2–Si3N4, (g–i) ZrB2–TaSi2. (a), (d), (g) SEM images illustrating the overall microstructure. (b), (e), (h) TEM images showing clean triple point junction, secondary phases or core–shell grains. (c), (f), (i) HR-TEM evidencing non-wetted interfaces only in the case of ZrB2– ZrSi2. The inset in (e) is the FFT of the boxed area. The inset in (h) shows dislocations at the core-rim interface.

Fig. 3. SEM images of the fiber section as a function of the sintering additive and temperature. Circles indicate the core.

containing turbostratic carbon and amorphous SiC, Fig. 4c. At the fiber–matrix interface, secondary phases like ZrO2 or ZrSi2 were often found, Fig. 4a. The interfaces between fiber-ZrB2, fiber-ZrSi2, ZrB2–Zr5SiB2 and adjacent ZrB2 were generally clean, Fig. 5. ZrB2–Si3N4–SiCf system – ZS-T achieved the full density at 1720 °C, the matrix mean grain size was around 2 lm with the coarsest grains up to 3.7 lm, i.e. smaller than ZZ-T system

although sintered 90 °C above. An image of the fiber is shown in the SEM image of Fig. 6a. The compounds deriving from the sintering additive and from ZrB2 powder, identified as BN, ZrN, Si3N4 and ZrSi2, and ZrO2, displayed good adherence to the fiber with no visible cracks. Besides the above mentioned secondary phases, a thin dark amorphous phase surrounding the fiber could be seen and EDS revealed the presence Si–O–C–N–Al. Morphological analyses

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Fig. 4. (a) Fiber in ZZ-T by SEM, (b) TEM image evidencing morphological differences in the core and shell areas and (c) HR-TEM image of the intergranular phase in core with EDS spectrum.

Fig. 5. Examples of clean interfaces in ZZ-T between (a) SiC fiber and ZrB2, (b) SiC fiber and ZrSi2, (c) ZrB2 and Zr5SiB2 and (d) two adjacent ZrB2 grains.

confirmed that the outer rim area, about 200 nm thick, had the tendency to coalesce and this region was characterized indeed by tight SiC grains, coarser than in the core of the fiber and with lower content of dark phase. This behavior is quite reasonable since C removal can induce coarsening and coalescence of the SiC crystallites fiber [11,17]. The fiber morphology is displayed in TEM images in Fig. 6b and c. The fiber core maintained the pristine fiber characteristics, Fig. 6b. The rim is instead composed of coarsened grains which are a mixture of a- and b-SiC, around 250 nm, Fig. 6c, with amorphous phase confined to the triple junctions, Fig. 6d and e. Embedded in the rim, ZrO2 or ZrSi2 particles could be occasionally found showing non-wetted interface, Fig. 6c and f. TEM images of the fiber–matrix interface are reported in Fig. 7. Adjacent to the fiber, BN and ZrO2 were often observed as SEM analysis pointed out, Fig. 6a. In addition, Zr–Si–C particles were

trapped between the fiber and these secondary phases, as crystallization products of the liquid formed during sintering. The local chemistry of these particles varied from region to region, suggesting the formation of agglomerates of different phases. High resolution images of the interfaces between Zr–Si–C particle and ZrO2, fiber and BN revealed wetted grain boundaries, Fig. 7a–c. ZrB2–TaSi2–SiCf system – ZT-T resulted fully dense at 1750 °C. Fiber–matrix details are shown in the SEM images of Fig. 8a and b. For this composite, fiber debris detachment was observed as a consequence of excessive SiC coarsening, Fig. 8a; in particular it can be seen that the development of elongated SiC grains started to occur. The matrix of this composite was similar to the baseline material [14,16] with solid solution around ZrB2 grains and secondary phases like ZrO2, SiO2, (Ta, Zr)Si2 and other Ta–Si–C–B crystalline compounds concentrated at the triple junctions. Also in this

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Fig. 6. (a) SEM images of the fiber in ZS-T. TEM images showing (b) fiber core, (c) fiber rim, (d) HRTEM of a triple point in the rim with EDS spectra in (e), high resolution image of the boxed area in (c).

Fig. 7. HRTEM images of various interfaces at the fiber–matrix boundary in ZS-T.

case fiber and matrix were well adhered without evident microcracking. TEM images of the fiber section in Fig. 8c and d evidence the core, rich in carbon pockets, around 22%, and the outer rim composed by dense and coarsened SiC. SiC crystallites in the rim lie embedded in a darker phase that EDS revealed to be SiC plus Zr traces. STEM analyses were carried out across the fiber section to detect cations diffusion. Fig. 9a and b show that the element coming from the matrix, Zr and Ta, are confined to the first interphase. Carbon has brighter spots in the core region and silicon seems to have lower intensity in the outer region. The profile composition inside a fiber is shown in Fig. 9c and d and it can be noticed that in correspondence of the darker spots, silicon decrease and carbon increase were registered. Al was found in negligible amount. TEM analyses revealed further details of the fiber morphology and composition, Fig. 10. In this composite, line defects and a-SiC grains were present also in the core region, Fig. 10a, differently from the other composites. The fiber core, Fig. 10a and b, was composed of about 70 nm b-SiC crystallites surrounded by intergranular amorphous phase which formed pockets of dimensions around 150 nm. Inside the intergranular phase, amorphous phase and turbostratic carbon could be identified, Fig. 10a and b, and confirmed by the EDS analysis, which revealed important carbon enrichment and traces of oxygen and zirconium. In the rim region, SiC crystallites had dimensions around 250–300 nm, the intergranular phase was not present anymore, but only confined to the triple point junctions of the densified crystallites, Fig. 10c

and d. The crystallites showed typical stacking faults and lamellar grain growth morphology; diffraction patterns collected on these grown SiC crystallites seem to suggest the formation of a-SiC and always showed interfaces wetted by Si–O–C amorphous phase containing traces of Zr and Al. At the fiber–matrix interface, a crystalline Si–O–C–Zr phase was detected accordingly to what observed by SEM. The fiber–matrix interface was found to be mostly wetted, Fig. 10e and f.

4. Discussion 4.1. Effect of the sintering additive As previously outlined [2,6,14–16], the main mechanism inducing the densification of diborides, is the removal of oxide species from the surface particles by the sintering additives which form a liquid or a transient-liquid phase at characteristic temperatures. The nature of the liquid phase formed during sintering and its interaction with fiber and matrix is quite complex [14–16], but can simplified as follows. In the case of ZrSi2, low melting Zr–Si-based compounds can form already below 1400 °C, and at around 1600 °C ZrSi2 melts [18]. The presence of liquids very likely helped the removal of surface boron oxide from the diboride particles, according to reaction (1), and enhanced mass transfer mechanisms [15,16].

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Fig. 8. SEM images of fibers in ZT-T showing (a) debris detachment and (b) secondary phases adjacent to the fiber. TEM images evidencing microstructural details in the (c) core and (d) rim.

Fig. 9. STEM analyses on ZT-T: (a and b) elements mapping across a fiber and (c and d) profile composition in the fiber core.

ZrSi2 þ B2 O3 ! Zr—Si—B—OðlÞ

ð1Þ

The addition of SiC fibers introduced further sources of SiO2 which in turn increased the amount of the liquid phase. The presence of Zr–Si–C–O phases, Fig. 4a, suggests a strong interaction between the sintering agent and the fiber. When the melt came in contact

with the fiber, it reacted with the outgoing intergranular C phase, according to the general reaction (2) and left clean grain interfaces and solid crystalline phases at the triple point junctions upon cooling.

Zr—Si—B—OðlÞ þ C ! ZrO2 þ Zrx Siy Bz þ SiC

ð2Þ

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Fig. 10. (a and b) TEM images of (a) fiber core in ZT-T showing the amorphous nature and composition of the intergranular phase. (c) Overview of the fiber rim with diffraction pattern of a-SiC in the inset and (d) magnification of the triple point junction. (e and f) Examples of wetted interfaces between fiber and ZrB2.

As for Si3N4 addition to ZrB2-matrix, thermodynamic calculations revealed that these pure compounds cannot coexist under the process conditions [15,16]. During heating up, the main reaction occurring first, which is characterized by a negative Gibbs free energy (304 kJ/mol, at 1500 °C), could be the following:

2B2 O3 þ Si3 N4 ! 4BN þ 3SiO2

ð3Þ

Another competing reaction could be (4) leading to the formation of liquid phase:

ZrB2 þ SiO2 þ Si3 N4 ! Zr—Si—B—N—OðlÞ

ð4Þ

In fact, the occurring of reactions (3) and (4) led to the complete disappearance of Si3N4 and produced various compounds upon cooling (BN, ZrN, ZrO2, ZrSi2, borosilicate glass. . .), reaction (5).

Zr—Si—B—N—OðlÞ ! BN þ ZrN þ SiC þ ZrO2 þ ZrSi2

ð5Þ

The formation of liquid Si is also favored, but, in the final microstructure, only intergranular borosilicate glass and crystalline ZrxSiyCz phases were found at the triple junctions. In this case, the interaction with the fiber was limited despite the higher sintering temperature (1720 °C), owing to the presence of nitrogen in the glass which increased its viscosity. Finally, it has been demonstrated that in reducing environment (C/CO-rich) or in the presence of B2O3, TaSi2 tends to decompose into Ta and Si/SiO/SiO2, which favour the formation of liquid phases where ZrB2 is partially soluble [14]:

TaSi2 þ COðgÞ ! TaC þ 2SiðlÞ þ SiOðgÞ

ð6Þ

TaSi2 þ B2 O3 ! TaB2 þ SiOðgÞ þ SiO2

ð7Þ

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The liquid is then constituted by Ta–Zr–Si–B–C–O, which shows a wetting tendency towards the boride grains. The interaction of such liquid with the fiber was quite intimate and Ta could diffuse into the fiber (Fig. 9a and b). Solid products like TaC and TaB2 were not found as pure compounds, but rather formed solid solutions with the analogous ZrC and ZrB2, according to reactions (8) and (9) [14], as observed in the final core–shell type matrix morphology, Figs. 2g and h and 8a:

ð1  xÞZrC þ xTaC ¼ ðZr1x Tax ÞC

ð8Þ

ð1  xÞZrB2 þ xTaB2 ¼ ðZr1x Tax ÞB2

ð9Þ

Upon cooling, refractory compounds crystallized at the triple junctions leaving amorphous grain boundaries. Overall, the liquid phases formed at different temperatures, as indicated in Table 1 (Ton), flowed among ZrB2 powder and SiC fibers, but did not significantly interact with the fibers, i.e. they were not penetrated by the liquid. The only case where cation penetration was detected was for ZT-T, Fig. 9b, where Ta was found across the fiber. This can be due both to the nature of the cation, or to the higher sintering temperature, 1750 °C, enhancing cation mobility and fiber reactivity. Fiber evolution was clearly more accentuated with increasing the sintering temperature irrespective of the sintering aid used, Fig. 3. The extension of the rim varied almost linearly as the temperature increased, but generally, with silicides, coarsening and densification of the rim seemed to be faster than in the case of Si3N4, probably owing to a tighter glassy phase in this last case. The same sluggish interaction mechanism could be found also in the mean grain size of the crystallites in ZS-T, smaller than ZZ-T although sintered 50–100 °C higher, Table 1. Another interesting aspect is the morphology of crystallites in the fiber: in the core all the materials generally presented rounded crystallites, whilst in the rim, the crystallites had rounded shape in ZZ-T and ZS-T, but they tended to have elongated shape in ZT-T. Even if the reported temperature for SiC b ? a polytypic transformation is given for higher temperatures, above 1900 °C, the presence of impurities or cations, such as Zr or Ta, could boost the transformation at lower temperature [17,19]. As the composites sintered with Si3N4 and TaSi2 were sintered at close temperatures, 1720 and 1750 °C respectively, it is very probable that tantalum induced growth along the c-axis. This preferential growth was already outlined in the case of W-doping [20] and, given the proximity of the two cations in the periodic table, it is very likely that also Ta had the same effect.

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fiber, hence its crystallites are coarser, around 70 nm, and it possesses higher stability up to its pyrolysis temperature. It has been demonstrated that both kinds of fibers are temperature sensitive, i.e. the extent of fiber evolution increases almost linearly with the sintering temperature. This aspect was ascertained by specific studies conducted at increasing sintering temperatures on HfB2-Hi-Nicalon composites [6]. For Tyranno fibers, there is a progressive fiber transformation too, as pointed in Fig. 3, but always less pronounced than for Hi-Nicalon. This is very clear comparing fiber morphologies after sintering at the same temperatures. Fig. 11 shows the fiber section of ZrB2 composites sintered with ZrSi2 at 1600 °C in case of Tyranno SA3 fiber, Fig. 11a, or of Hi-Nicalon fiber, Fig. 11b. It is well evident that in the first case more than 80% of the fiber maintained the pristine fiber’s aspect and presumably properties too, whilst, in the second case, just 26% of fiber core was left, confirming the higher stability of Tyranno over Hi-Nicalon. Also in terms of chemistry, there are some similarities between the two classes of fibers, for example both kinds of fibers contain turbostratic carbon. The general fiber evolution observed in all the systems sees the outward migration of carbon-rich intergranular phase as the temperature increases. The removal of carbon enables SiC crystallites to densify and coarsen, resulting in two well distinct regions defined as core, carbon rich, and rim, denser and carbon-depleted. However, SEM–EDS mapping across the two fiber types clearly evidences that in Tyranno fiber, Fig. 12a, carbon and oxygen distribution is more homogeneous as compared

4.2. Fiber stability: comparison with Hi-Nicalon SiC fibers Previous works with Hi-Nicalon SiC fibers demonstrated strong interaction with UHTC matrices that partially reduced the final materials properties [3–7]. In order to weaken the interface between fiber and matrix, we explored the interactions between ZrB2 and these more refractory Tyranno fibers. The main differences between Hi-Nicalon and Tyranno SA3 fiber are related to the processing: both derive from a polycarbosilane precursor, but, for the second type of fiber, Al is present too as sintering agent [10–13,21]. Subsequently the polymer is spun and subjected to reticulation by electron irradiation in the first case and in air in the second one. Then the pyrolysis is conducted at 1500 °C for Hi-Nicalon and at 1800 °C for Tyranno [11,12]. These different processes generate major differences in the microstructural features of the fiber: in Hi-Nicalon, SiC mean grain size of the crystallites is around 5 nm and is much more reactive, owing to the still unstable system containing turbostratic carbon and amorphous Si–C–O. On the other hand, Tyranno SA3 fiber can be considered a sintered

Fig. 11. Fiber section in ZrB2–ZrSi2 composites containing (a) Tyranno SA3 fiber and (b) Hi-Nicalon fiber sintered at 1600 °C.

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Fig. 12. Carbon and oxygen SEM–EDS mapping of ZrB2–Si3N4 matrix containing (a) Tyranno SA3 fiber and (b) Hi-Nicalon fiber sintered at 1720 °C.

to Hi-Nicalon, Fig. 12b, where there is a clear carbon depletion in the rim, owing to diffusion towards the matrix, coupled to oxygen enrichment, owing to the trapping of oxides from the matrix and Si–C–O amorphous migration outwards. On the other side, in terms of interaction with liquid phases arising from the sintering aids, it can be stated that no reactions products formed at the interface between fiber core and rim, confirming higher thermal stability of these fibers over the Hi-Nicalon. Fig. 11 shows that the wavy morphology of SiC outer profile is due to exudation of liquid phase, particularly pronounced when HiNicalon were used. The rim region in Fig. 11b contains extra products which derive from reaction of C and Si–O based phases with oxides collected from the matrix. These particles are sometimes found also in the proximity of the Tyranno fiber/matrix interface, but very rarely in the fiber interior. Microstructural analyses have thus shown that Tyranno-SA3 fiber interacts with the boride matrix during sintering, but only in a limited extent. Besides the sluggish outward migration of the unstable fiber fraction, turbostratic carbon and amorphous Si–C, other concurring phenomena, like crystallization of amorphous phase, SiC coarsening, polytypic transformation and further fiber densification took place. As a matter of fact, the species coming from the matrix, boron in particular, could contribute to SiC shrinkage, as it is reported that B substitutes Si in the SiC lattice thereby inducing defects into the structure of SiC and enhancing the volume/lattice diffusion [17,19]. In addition, all the metallic cations plus boron and carbon could also modify the properties of the grain boundaries by lowering the grain energy and improve growth of a polytype [17,19]. One final remark concerns the limit sintering temperature that these fibers can bear. For Hi-Nicalon fibers, a temperature of 1600 °C was identified as limit to enable effective toughening mechanisms, above it the fibers strongly transform and can be considered a second phase [7]. On the contrary, although there is still dearth of studies on UHTCs-Tyranno composites, it seems that a temperature of 1750 °C is still acceptable not to compromise Tyranno SA3 fibers. 5. Conclusions ZrB2 composites containing chopped Tyranno SA3 fibers were hot pressed at various temperatures through addition of ZrSi2, Si3N4, TaSi2 and MoSi2. The best sintering temperature for each additive was defined, ranging from 1600 °C to 1750 °C, but MoSi2 induced too high fiber damage. Depending on the sintering additive, different liquid phases formed which interacted with the fiber in different specific ways,

but left the fiber with some common features. As general characteristics, the sintered fiber showed a well-rounded shape with sharp interface with the matrix. TEM investigations revealed that the core was typical of the as-received fiber, i.e. b-SiC crystallites of 70 nm surrounded by about 20 vol% of highly disordered Si–C phase and carbon. Moving from the fiber core outwards, a progressive removal of these intergranular pockets occurred accompanied by coarsening and densification of SiC crystallites and crystallization of amorphous Si–C. The fiber–matrix interfaces were clean only when ZrSi2 was employed, but in the other cases wetted interfaces were generally observed. The morphology of Tyranno-SA3 fiber after sintering resulted notably different from Hi-Nicalon fiber, for which multilayered morphology and jagged interface were observed in previous studies.

Acknowledgements We greatly acknowledge the financial support of the US Air Force Research Laboratory for TEM activity through grant N. FA8655-12-1-3004, with Dr. Ali Sayir as contract monitor.

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