Fibers made by chemical vapor deposition

Fibers made by chemical vapor deposition

24

Xian Luo, Na Jin Northwestern Polytechnical University, Xi’an, China

24.1

Introduction

High-performance ceramic fibers exhibit superior mechanical properties, thus can be used as reinforcements for metal, polymer, or ceramic matrix composites. Ceramic fibers can be prepared by pyrolysis of polymeric precursors, electrospinning, powder sintering, chemical vapour deposition (CVD), etc. CVD has the ability to produce dense, uniform deposits with well-controlled surface morphology and composition without relying on line-of-sight between the source material and the substrate. Furthermore, this technique offers the potential for rapid and continuous processing that would be desirable for largescale composite programmes. In general, the CVD method can produce continuous large-diameter fibers with high strength and tensile modulus (Choy, 2003). This chapter would focus on the CVD boron fibers and silicon carbide (SiC) fibers. CVD fibers of large diameter (i.e., 100e150 mm) have been used for many years for the reinforcement of both organic and metallic materials. The real impulse in CVD fiber, however, came in 1959, when Talley used the process of halide reduction to obtain amorphous boron fibers of high strength and stiff (Talley, 1959; Talley et al., 1960). The persistence of Talley and his coworkers at Texaco Experiment Incorporated (under sponsorship of the US Air Force Materials Laboratory) led to the development of a continuous boron fiber, which General Bernard Schriver in 1964 stated as “the greatest breakthrough in materials in the last 3000 years” (Schriver, 1964). Since then, interest in the use of strong and radiation shielding (Kowbel et al., 2005) but light boron fibers as a possible structural component in aerospace and other structures has been continuous, although boron fibers have some important drawbacks. In particular, when they are used within metal matrices for high-temperature applications, (1) the amorphous structure tends to recrystallize with a decrease in failure strength and (2) boron reacts with most metals to form brittle intermetallic phases. Then, a new class of CVD SiC fibers has been developed for the reinforcement of metal matrices. There are mainly two kinds of commercially available CVD-SiC filaments: the SCS family of fibers manufactured by Specialty Materials Inc. (formerly Textron) in Lowell (USA) and the Sigma family of fibers manufactured by the company Tisics (formally DERA) in Hampshire (UK). The carbon core is a monofilament melt-spun from coal-tar pitch (Ning et al., 1990). There is still another trade name of SiC fiber, which is named Trimarc fiber. Trimarc 1 SiC fiber is about 127-mm diameter fiber deposited by CVD onto a 12.5 mm tungsten core, and the manufacturer is Amercom Inc., Atlantic Research Corporation (Middleway, WV) (Liu and Bowen, 2003). Handbook of Properties of Textile and Technical Fibres. https://doi.org/10.1016/B978-0-08-101272-7.00024-9 Copyright © 2018 Elsevier Ltd. All rights reserved.

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This chapter will attempt to present the development of CVD boron and SiC fibers in the last decades from production, structure, properties, and applications. The effects of manufacturing processes and post-treatment of the fibers themselves on their microstructure and mechanical properties are described. Moreover, the effects of composite preparation processes, heat treatments, as well as creep/fatigue process of compositesdmainly titanium matrix compositesdon the tensile strength and fracture modes of the fibers are summarized. Finally, the developments of fiber coatings used for protecting fiber strength and improving mechanical properties of titanium-matrix composites are summarized.

Boron fibers

24.2

In Aerospace Conference of 2005, it was reported that future space systems developed under NASA Space Initiative needed to address the issue of radiation protection in conjunction with structural requirements (Kowbel et al., 2005). The boron fiber is the ideal form in a structural composite where it provides radiation shielding, mechanical strength, and lightweight (the density of boron fiber is only 2.5 g/cm3). Boron is an inherently brittle material and boron fibers can be made only by CVD of boron onto substrate fibers. The boron fibers are themselves a composite fiber, which contain a core material as a substrate and a deposited boron shell as the main material. In the last decades, the boron fiber has been used not only in structural requirements, but also in other fundamental applications. This section will present the development of boron fiber over the last few decades including production, structure, properties, and applications.

Boron fiber production

24.3 24.3.1

Undoped boron fibers

Commercially available boron fibers have tungsten (W) cores (w12 mm diameter) and a boron coating with a thickness of approximately 100 mm (Zhao, 2000; Suplinskas and Marzik, 1987). (The means of manufacturing W-core boron filament is the same as that for manufacturing doped boron fiber shown in Fig. 24.6, which will be introduced in Section 24.3.2.) This is accomplished by the hydrogen reduction of a boron trichloride gas onto a heated tungsten core by CVD method at w1200
C, where boron trichloride was the precursor gas because of its product consistency, handleability, and relative low cost compared with the other halides, and hydrogen was the carrier and reduction gas. A generalized reaction equation is 2BCl3(g) þ 3H2(g)/2B þ 6HCl

(24.1)

In an earlier investigation, Wawner (1967) emphasized the importance of the substrate surface for the nucleation of boron nodules. He even suggested that the ordered

Fibers made by chemical vapor deposition

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appearance of the surface of the final boron fibers could be attributed to die-marks on the substrate. Wawner showed that when depositing boron on a scratched tungsten sheet, preferential nucleation and initial growth could be observed “adjacent to the scratches.” Boggio et al. (1977) also observed a tight distribution of nodules along the die-mark ridges of tungsten wire. This is an essential feature for subsequent growth of the boron sheath. The phenomenon may be interpreted in terms of two competing mechanisms, one being based on boron diffusion and the other on surface energy. A very interesting fundamental phenomenon had been found during the CVD of boron in the “amorphous” temperature range, that of elongation or expansion of the (Talley, 1959; Eason and Johnson, 1980; Wawner et al., 1980, 1979; Eason, 1979; Mehalso, 1973). The researchers suggested that the elongation of boron during deposition was the result of diffusion of boron atoms from the surface of a fiber into its interior causing a density increase as well as a density gradient in the fiber. The boron atoms go into vacancies that create a “lattice” expansion leading to the elongation. Otherwise, Wawner et al. (Eason and Johnson, 1980; Wawner et al., 1980, 1979; Eason, 1979) also noted that substrate type and size and the quantity of boron deposition appeared to have an effect. The investigators developed a model to explain the elongation. The model is as follows: “Boron atoms are deposited as individual atoms in a random configuration. Each layer of atoms deposited is continually being covered by a fresh layer. The buried atoms then transform, at some rate, from a random distribution of individual atoms into a random distribution of 12-atom icosahedral units.” Commercial boron/tungsten fibers are valued for their high strength and modulus (stiffness), as shown in Table 24.1. It is worthy to note that the large diameters (100e140 mm) are a disadvantage in the fabrication of composites. However, the density of boron and tungsten is 2.34 and 19.35 g/cm3, respectively. The density of boron/ tungsten fiber will increase due to a diameter reduction, which does not follow the rule of mixtures since tungsten is converted into tungsten borides during deposition. X-ray diffraction analysis showed that core reaction phases progressed by the formation of

Table 24.1 Properties of large-diameter boron/tungsten fibers by chemical vapor deposition (CVD) (Nordine, 1995) Large-diameter boron/tungsten fibers Process and properties Fiber diameter (mm)

140

100

100

55

Tet data

Average

Average

Best

Average

Tensile strength (GPa)

3.5

3.5

4.8

(3.5)

Tensile modulus (GPa)

400.0

400.0

400.0

(400.0)

2.5

2.6

2.6

3.1

3 a

Density (g/cm ) a

CVD of boron on a hot, continuous tungsten wire

Tsirlin AM: In Watt W, Perov BV, editors: Handbook of composites, vol. 1. North Holland, 1985, Amsterdam, pp 115e199.

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delta-WB, WB4, and W2B5 in order. In the finished fiber (i.e., deposition times of 30 s or more) only WB4 and W2B5 coexist (Talley, 1959; Wawner, 1967; Witucki, 1967). From Table 24.1, it can be seen that a diameter reduction from 140e100 to 55 mm increases the fiber density from 2.48e2.59 g/cm3 to 3.10 g/cm3 and converts a diameter problem into a density problem. In addition, the unidirectional boron diffusion also leads to an increase in core diameter during boron deposition on tungsten core. This results in a compressive residual stress in the core and a dilational stress in the surrounding boron sheath (Boggio and Vingsbo, 1976a). The quenching, finally, implies a compressive residual stress in the outer part of boron sheath, as shown in Fig. 24.1 (Boggio and Vingsbo, 1977). Nordine (1995) reported a laser CVD (LCVD) method by which a continuous, substrate-free fiber consisting of amorphous boron could be obtained, having a tensile strength ranging from about 5100 MPa to over 6900 MPa and a diameter that was varied with the number or diffraction-limited laser focal spot size to provide a fiber diameter as small as 0.01 mm. Table 24.2 lists the properties of pure boron fiber by the LCVD method and other commercial ceramic fibers. The core of the boron/tungsten fiber is converted from tungsten into a mixed tungsten boride after the boron vapor σθ

(+) 0.7–0.8x103 MPa

r

(–) 0.3–0.5x103 MPa

(–) 1.0–1.4x103 MPa

σr σθ

σθ Core Sheath

Figure 24.1 Schematic residual stress pattern in the cross-section of a boron fiber. Reprinted from Boggio JV, Vingsbo O: Radial cracks in boron fibres, J Mater Sci 12: 2519e2524, 1977 with permission from Springer.

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Table 24.2 Properties of small-diameter pure boron and other commercial fibers (Nordine, 1995) Small-diameter boron and other commercial fibers Pure boron fibers

Intermediate modulus carbon fibers

Nicalon/ SiC fibers

S. Cryst. SiC whiskers

Average fiber diameter (mm)

<25

10

17

6

Fiber-forming process

LCVD continuous

PAN continuous

Carbosilane continuous

VLS discontinuous

Tensile strength (GPa)

5.17

3.52

2.62

7.44(@28 mm)

Max. strength (GPa)

7.58

4.14

3.10

15.9(@4 mm)

Avg. tensile modulus (GPa)

400

280

190

580

Density (g/cm3) (Diefendorf, 1974; Vrieze and Schob, 1974; Kuehl, 1976a,b)

2.29

1.80

2.38

3.18

Processes and properties

LCVD, laser chemical vapour deposition; VLS, vapor-liquid-solid.

deposition. The surface of the boron/tungsten fiber is nodular whilst that of the pure boron fiber is smooth, providing additional evidence for a more uniform structure. On the one hand, modulus (stiffness) is an intrinsic materials property that does not change much with the uniformity of a material (i.e., surface cracks, internal voids), so the small-diameter pure boron fibers by the LCVD method (Table 24.2) and commercial boron/tungsten fibers (Table 24.1) have the same modulus. On the other hand, strength is governed by the uniformity of a material. Pure boron fibers (Table 24.2) have up to twice the room temperature strength of boron/tungsten fibers (5.2e7.6 vs. 3.5e4.8 GPa) at &1/10th the diameter. Pure boron fibers tested at 1000
C had 95% of their room temperature strength, while boron/tungsten fibers has only 84%. Higher strength also reflects a greater uniformity. In contrast to the fibers in Table 24.1, the fibers in Table 24.2 have smaller diameters (6e17 mm), and it also can be found that the LCVD boron fibers exhibit higher strength and stiffness compared with commercial intermediate modulus carbon fibers made from PAN and Nicalon SiC fibers obtained by using polycarbosilane. The average tensile strength and modulus of LCVD boron fiber are <0.7 that of single crystal vapor-liquid-solid SiC whiskers, nevertheless no continuous process is available for the latter. As mentioned above, the core of boron/tungsten fiber is converted from tungsten into a mixed tungsten boride, and the fiber’s density increases with the reduction of

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Handbook of Properties of Textile and Technical Fibres

Figure 24.2 Amorphous and smooth boron coating on IM7 12k graphite fiber tow (Kowbel et al., 2005).

fiber diameter. On the other hand, because of the radiation phenomenon, tungsten is unacceptable as a substrate in addition to its weight penalty, whilst ideally carbon would be a good choice as a fiber-forming substrate because of both its weight and atomic number, relative to radiation resistance standpoint. However, monofilament carbon fiber is not generally available as a substrate to produce boron fibers. Consequently, Kowbel et al. (2005) tried to utilize commercially available spread graphite fiber tows as the substrate to deposit boron fiber by a CVD method. Fig. 24.2 shows the end view of the boron-coated spread fiber tow consisting of 5e7 mm individual fibers. As can be seen there were a few places where two or three individual fibers were bridged. It was indicated that the fiber surface morphology and strength were both related to the boron coating thickness. It has been suggested that the boron to carbon (B/C) ratio should be at least 4 to achieve the desired radiation shielding and high strength. Fig. 24.3 shows the tensile strength versus thickness of the boron coating for several graphite fiber substrates, where the boron coating thickness of 2.3 mm brings the B/C ratio to 4/1 for these size graphite fiber substrates, which is 4e5 mm in diameter. The results showed that although there might be some slight differences in strength of the composite boron fiber, at 2.3 mm boron coating, all substrates provided strength equal to an uncoated T300 graphite fiber, which was the typical fiber used in composites. Undoubtedly further refinement/optimization can produce composite fibers at a B/C ratio > 4 with strength >3.36 GPa. In Kowbel’s report of 2009, a novel CVD method was developed to coat large amounts of boron on carbon fibers, which involved a large batch reactor that heated

Fibers made by chemical vapor deposition

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7.0

B/C=4 M7

6.0

T1000 M8

Tensile strength (GPa)

5.0

4.0 T300 3.0 2.0

Effect of sizing removal on T1000 fiber

1.0 0.0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

Thickness (μm)

Figure 24.3 Tensile strength versus thickness of the boron coating for several graphite fiber substrate (Kowbel et al., 2005).

the carbon fibers (mounted on graphite) using inductive heating coils. The samples were Panex 33 48K carbon fiber tow (48,000 filament tow), which were coated with 2e4 microns of boron. To achieve surface structure and adhesion of the boron coating, a pyrolytic carbon coating was required on the carbon fibers before deposition of boron. The overall reaction to produce a boron coating on a fiber substrate was also followed as the formula (24.1). Table 24.3 shows CVD conditions during boron

Chemical vapor deposition conditions during boron deposition on Panex 33 48K carbon fiber tow and the thicknesses of boron coating obtained (Kowbel et al., 2009) Table 24.3

Run#

Deposition temperature (8C)

BCl3 flow rate (sccm)

Total pressure (torr)

Deposition time (min)

Coating thickness (mm)

1

950

50

50

30

0.1e0.2

2

950

50

50

60

0.1e0.3

3

1000

50

50

30

0.7e1.0

4

1000

50

50

60

1.3e2.1

5

1050

50

50

30

3e4

6

1050

50

50

60

5þ

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Handbook of Properties of Textile and Technical Fibres

(a)

(b)

Low magnification

High magnification

Figure 24.4 SEM micrograph of boron coating on Panex 33 48K carbon fiber tow (Kowbel et al., 2009).

deposition on 20 m long sections of Panex 33 48K and the thickness of boron coating obtained. Runs 1 and 2 were run at 950
C and the boron coating thickness was too thin. Runs 5 and 6 were run at 1050
C and the boron coating thickness was too thick. Run 3, 30 min at 1000
C and run 4, 30 min at 1000
C yielded a boron coating thickness that was less likely to cause bridging, as shown in Fig. 24.4. The coated fiber surfaces were smooth and a small amount of fiber bridging was seen. For boron coating thickness greater than 1.5 mm, the grain size increased. Very large grain sizes were seen with a 4.5 mm thickness. As a result of the large grain size growth and bridging problem the 1.5e2 mm boron thickness became a practical limit to fabricate boron coatings with the method explained in this work. Moreover, Fig. 24.5(a) shows that

Figure 24.5 Panex 33 48K carbon fiber tow coated with about 1.5 mm (a) and 3.5 mm (b) boron coating thickness (Kowbel et al., 2009).

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937

a boron coating with an average thickness of 1.5 mm, with good flexibility and with minimal fiber bridging (no breakage), while a boron coating with average 3.5 mm thick broke and was not flexible, as shown in Fig. 24.5(b).

24.3.2 Doped boron fibers Of more recent interest is the use of CVD boron fibers as precursors for the synthesis of MgB2 superconductors. Boron fibers can be converted to MgB2 via reaction with Mg vapor (Canfield et al., 2001), in principle eliminating a separate wire-forming manufacturing step. Marzik et al. (2005) had prepared doping boron fibers with different carbon dopant levels to investigate the effect of dopant additions on the microstructure of boron fibers before and after reaction to MgB2. Carbon-doped boron fibers were prepared similarly to undoped commercial CVD boron fibers described above, with the additions of methane to the BCl3 and H2 reaction mixture. Fig. 24.6 is the process diagram showing CVD synthesis of doped boron fiber. The values used for dopant levels were the gas phase atomic ratios of carbon to boron. However, initial analytical work indicated that the actual concentrations of carbon in the deposit might be approximately 50% higher than the gas phase concentration. Figs. 24.7 and 24.8 show the scanning electron microscopy (SEM) micrographs and X-ray diffraction patterns of undoped and carbon-doped boron fibers. The “corncob” surface texture was indicative of macroscopic growth morphology and not the size of individual grains. Below a carbon dopant level of 2%, the fiber surface morphology remained relatively

Payout spool for tungsten wire

BCI3 + H2 + dopant(s) Glass tube

Variable DC supply

Doped boron fiber 1200–1400°C

Exhaust gases

Mercury electrode (top & bottom)

Take-up spool for boron fiber on tungsten boride substrate

Figure 24.6 Process diagram showing CVD synthesis of doped boron fiber (Marzik et al., 2005).

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Handbook of Properties of Textile and Technical Fibres

(a)

(b)

50 μm

50 μm

(d)

(c)

50 μm

20 μm

Figure 24.7 Surface morphology of boron fibers at the following carbon dopant levels in atomic percent: (a) 0% (b) 1.4% (c) 2% and (d) 3%. Insets to 2(a) and 2(b) show a 25 mm wide area of the fiber surface (Marzik et al., 2005).

Intensity (Arb.units)

(a)

Deg 2θ

(b) Undoped CVD boron fiber

20

40

(c)

CVD boron fiber doped with 1.4% C

60

Deg 2θ 20

40

CVD boron fiber doped with 3% C

60

Deg 2θ

20

40

60

Figure 24.8 X-ray diffraction patterns of undoped and carbon-doped boron fibers. Undoped boron and boron doped with 1.4% carbon consist of amorphous boron plus crystalline tungsten boride (Marzik et al., 2005).

unchanged and the boron phase remained X-ray amorphous. At the 2% dopant level, however, there were isolated occurrences of much larger grains, 10e50 mm in size, that were most likely seeded by a threshold concentration of carbon dopant. At the 3% dopant level rapid crystal growth dominated and the fiber no longer retained a regular cylindrical shape, but rather was a string of larger 10e50 mm crystals irregularly arrayed around the tungsten boride substrate (Fig. 24.7(d)). X-ray diffraction of the 3% carbon-doped boron fiber indicated that it was well-crystallized tetragonal boron (Fig. 24.8(c)). Quantitative energy-dispensive detector (EDS) performed on spectra acquired from several thicknesses of crystallized boron determined that carbon was present within the crystal and was not merely a surface contaminant during the analysis. Quantitative EDS also established that the carbon concentration was significantly less than 20 atomic%, thereby indicating the crystalline phase was not B4C.

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24.3.3 Mechanical properties of boron fiber 24.3.3.1 As-produced boron fiber

Type B

5 10

50

Type A

IIIB

1

Fracture probability (%)

90 99

The original strength of the boron/tungsten fibers is determined by the defects formed during their preparation (The interface in metal composites, 1978), which are distributed nonuniformly over the length of a fiber. Such defects include surface cracks and also cracks and pores close to the tungsten core. Tanaka et al. (1991) had performed tensile tests on two kinds of boron/tungsten fibers fabricated by CVD method to discuss the relationship between the fracture strength and defect size. The Weibull plots of tensile strength for boron fibers were divided into three regions: the lowest strength one, medium strength one, and highest strength one, as shown in Fig. 24.9. Fig. 24.10 shows the microfractographs of boron fibers where crack origin was fractured from surface and voids situated near core tungsten fiber. It was revealed that in the lowest strength region the defect size was large, while in the highest strength region defect could not be observed. The defect size was linearly related to the fracture strength, showing the applicability of fracture mechanics. The critical stress intensity factor for the surface defects was about twice that for the inner defects observed near tungsten core, as shown in Fig. 24.11. This meant that the difference of residual stress would be about 1.5 GPa from near fiber surface to near tungsten core. The fracture surfaces of boron fibers essentially resemble those of glass rods and have been divided into three zones, according to appearance and crack propagation velocity (Andrews, 1959; Layden, 1973): (1) the mirror zone is very smooth, marks the initiation of the crack, and corresponds to low velocity; (2) the mist zone has a matt surface and corresponds to higher velocity; (3) the hackle zone, finally, displays surface steps and ridge, caused by the deviation from a single propagation plane for the highest crack velocities.

0.1

IIIA

1.5

2

IIB

IB IIA

3

IA

4

Tensile strength (GPa)

Figure 24.9 Weibull plot of tensile strength of boron fiber (Tanaka et al., 1991).

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Handbook of Properties of Textile and Technical Fibres

(a)

(b)

σ = 2.64 GPa.

σ = 2.24 GPa.

(c)

(d)

σ = 3.44 GPa.

σ = 3.14 GPa.

Stress intensity factor,KI(MPa ∙m1/2) 3 4 5 6 7 8 9 10 15

Figure 24.10 Microfractographs at crack origin of Type A boron fibers fractured from surface (a and c) and voids situated near core tungsten fiber (b and d) (Tanaka et al., 1991).

1

KImax = 9.8MPa ∙m1/2

KImax = 4.6MPa ∙m1/2 Surface

Internal

A B 2

3

4

5 6 7 8 910

20

30

Depth of mirror,b(μm)

Figure 24.11 Relationship between maximum stress intensity factor and depth of mirror zone (Tanaka et al., 1991).

Fibers made by chemical vapor deposition

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Since the beginning of the investigation of boron fibers, interest has been devoted to the initiation of fracture. With some generalization, the main observations of crack nucleation site can be summarized as follows (Layden, 1973; Wawner, 1965; Wawner and Satterfield, 1967; Line and Henderson, 1969): (1) surface flaws, such as accidental damage points or nodule boundaries; (2) crystalline or abnormally protruding amorphous nodules; (3) inclusions in the core/sheath interface or in the sheath, generally in the case of multistage fibers; (4) voids near the core/sheath interface: these voids are thought to result from the unidirectional flow of boron from the sheath into the core (Wawner, 1967); (5) Notches in the initial core surface: voids have been observed at the bottom of deep and narrow notches, possibly a result of diffusion-induced pinching-off of a notch neck (Wawner, 1967; Layden, 1973); (6) the interior of the core: crack nucleation has been observed without relation to any detectable imperfections. The strongest fibers (fracture stress, sf w 4.0 GPa) are related to core fracture according to point (6) above, while all types of defects successively weaken the fiber down to below 1.4 GPa (surface flaws) (Boggio and Vingsbo, 1976a). Voids in the boron sheath near core/sheath interface, so-called proximate voids (PV), have transverse irregularities, large enough to create stress concentrations (Boggio and Vingsbo, 1976b). The origin of failure in boron fibers can often be traced to a PV. Boggio and Vingsbo (1976b) found that the relation between the radial extensions, 2c, of fracture initiating proximate voids in mirror zones and sf was in good agreement with the Griffith criterion for brittle fracture.

24.3.3.2 Boron fibers with post-treatment In a previous report it was demonstrated that the fracture parameters for commercial boron/tungsten fibers could be significantly improved by simple secondary processing methods (Dicarlo, 1977a). Dicarlo (1979) concluded that probably the most simple and cost-effective secondary treatment was that in which the commercial fibers were heated in impure argon gas at temperatures near 900
C. It appeared that by reacting with an unidentified impurity in the argon, boron atoms were removed from within the boron sheath causing the fiber to contract axially and core flaws to be compressed. Using this heat treatment method and a subsequent surface etch treatment near 100
C, the room temperature strength of 203 mm diameter fibers increased proportionally with εz up to 0.3%, after which point the fiber strength began to drop off significantly due to the formation of new fracture-controlling flaws within the boron sheath. With this limitation, the average tensile strength of the 203 mm fibers was improved from 3.4 to 5.5 GPa. The strength of CVD boron sheath has the potential of being over 6.9 GPa, that is to say the possibility for still further improvement if the formation mechanism for the new sheath flaws could be understood and possibly avoided. The results from the boron fiber contraction study of Wawner et al. (1979) pointed to oxygen as the most likely impurity. Therefore, the approach taken by DiCarlo and Wagner (1981) was to broaden the impure argon studies by also measuring the effects of heat treating fibers in oxygen-argon gaseous mixtures containing much higher and better controlled oxygen contents. Boron fibers with a diameter of 203 mm commercially supplied by Avco Specialty Materials Division were used in this study, which were produced in a singlestage CVD reactor by the hydrogen reduction of BCl3 on a 13-mm diameter tungsten

Handbook of Properties of Textile and Technical Fibres

Probability of core-initiated fracture (%)

942

100

Contraction % 0.2

0.3

0.4

0.5

50 Extrapolated

0 700

800 900 1000 Processing temperature (°C)

1100

Figure 24.12 Approximate (solid) and extrapolated (dashed) curves for the effects of processing temperature and oxidation-induced contraction strain on the probability of observing high strength core-initiated fracture in slightly etched 203 mm boron fibers (DiCarlo and Wagner, 1981).

wire substrate at the temperature of w1300
C (Krukois, 1977). During deposition the substrate became completely borided to form a 17-mm diameter tungsten boride core. From previous studies (Smith, 1976; Dicarlo, 1977b) it was determined that the 203 mm boron produced in the above manner was excellent material for achieving strengthening by core compression because in the as-received condition the probability for observing core-initiated fracture after a slight surface etching was essentially 100%. Prior to contraction, the commercial fibers contain only two types of flaws. One type is located on the fiber surface and the other within the fiber core. Removal of the first type by a slight surface etching led to an increase of average tensile strength from near 3.4e4.2 GPa, and the coefficient-of-variability (COV) for the tensile strength decreased from w15% to less than 5%. Fig. 24.12 shows the probability of observing core-initiated fractures as a function of processing temperature and total contraction. By combining the high-temperature oxidation treatment with subsequent slight chemical surface etching near 100
C, the average tensile strengths of 203-mm diameter commercial fibers were increased to 5.5 GPa with a COV of less than 5%. This strength level was achieved by a compression of strength-limiting flaws within the fiber’s tungsten boride core at an axial contraction strain of 0.3%. However, if the oxidation treatment was allowed to continue too long (i.e., contractions greater than 0.3%) tensile strength would be obviously degraded due to interfacial void formation.

24.3.4

Performance characteristic of boron fiber in composites

24.3.4.1 Interaction with matrix In all metal-matrix/fiber systems undesirable interactions can occur at the matrix/fiber interface, which prevents their application at high temperatures. The interaction leads

Fibers made by chemical vapor deposition

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to dissolution of fibers in the matrix and/or to uncontrolled formation of brittle intermetallic compounds, both of which reduce mechanical properties of metal matrix composites. In the early 1960s, boron fiber had become available as a reinforcement of aluminum matrix composites, but because of its reactivity with aluminum during composite fabrication, a SiC-coated boron fiber sold under the trade name BORSIC by Hamilton Standard, Division of United Technologies, was developed to circumvent this problem. BORSIC was also needed for use in titanium matrices. In this case, one problem encountered was that BORSIC fiber reacted with titanium metal at temperatures above 900
C (Galasso and Pinto, 1969). Tan et al. (1991) had deposited a titanium coating on two kinds of SiC-coated boron fibers, BORSIC and SICABO (sold by Composites Inc.), by CVD method to investigate the interaction between boron fiber and titanium. The Auger electron spectroscopy (AES) results showed that the surface of as-received fibers was composed of a thin layer (150e200 Å thick) of silicon suboxides/oxide. Removal of this layer by argon ion sputtering caused a decrease in the relative intensity of the O peak to almost the noise level. The relative atomic concentration ratios for C/Si decreased from 7.4 to 1.0 and that of O/Si decreased from 1.9 to 0.0 when the SICABO fiber surface was argon ion sputter cleaned. A similar decrease in these two ratios was observed in the case of sputter-cleaned BORSIC fiber (C/Si from 12.8 to 1.0 and O/Si from 2.0 to 0.0). The changes in relative peak-to-peak heights were readily observed in survey scans (Fig. 24.13). It thus could be concluded that the

(b)

(i)

dN(E)/dE

Si

(i)

O

C

(ii)

dN(E)/dE

(a)

(ii)

Ar

130

330 530 730 930 Electron energy (eV)

130

330 530 730 930 Electron energy (eV)

Figure 24.13 Survey Auger electron spectra of (A) BORSIC and (B) SICABO fiber in the (i) asreceived state and (ii) after 15-min argon ion etch. Reprinted with permission from Tan BJ, Hwan L, Suib SL: Spectroscopic characterization of CVD Ti coating on SiC-coated boron fibers, Chem Mater 3:368e378, 1991. Copyright (2017) American Chemical Society.

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surfaces of the SICABO and BORSIC fibers were composed of an overlayer of adventitious carbon- and hydrocarbon-containing surface impurities beneath which was a thin layer (150e200 Å thick) of silicon oxides/suboxides. The AES sputter depth profile of titanium-coated SICABO and BORSIC fibers revealed that the Ti diffused deeper into the BORSIC fiber than the SICABO fiber (Fig. 24.14). This was evident from the high atomic concentration of Ti throughout the depth profile of Ti-BORSIC, whilst the Ti atomic concentration dropped dramatically on reaching the SiC coating of the Ti-SICABO fiber. This suggested that Ti had a more severe reaction with BORSIC than with SICABO, resulting in Ti diffusing deep into the SiC coating of the BORSIC fiber and reacting with the SiC to form Ti silicides and carbides. In an earlier study, Hwan and Suib (1989) reported that the surface composition of SICABO fiber was richer in carbon than that of the BORSIC fiber. Thus when Ti was deposited onto the SICABO fiber, there was an initial reaction of the Ti with the carbon-rich surface resulting in the formation of titanium carbide. The formation of TiC layer in the neighborhood of the SiC fiber resulted in a decrease of the reaction rate and the overall extent of reaction between Ti and Si. For the Ti-BORSIC fiber, the lack of a carbon-rich surface resulted in the reaction of Ti with SiC and concomitant formation of TiC and Ti silicides. Ti diffused into the SiC coating and continued to do so until the reaction zone grew thick enough and obstructed further reaction of Ti with SiC, because the supply of Ti to the SiC coating was decreased. Since TiC has the largest negative heat of formation compared to other phases, it is preferentially formed. The resulting free Si then reacted with fresh Ti, forming Ti silicides. According to the phase diagram, a solid solution exists between pure Ti and intermediate compounds in the system (Massalski et al., 1990). Additives to the matrix in small quantities can change the mixing energy of the solution. If an increase occurs, the mixing energy should act as an energy barrier against solid solution formation. Revzin et al. (1997) had applied this theory to evaluate the effect of several additives on the solubility of boron fibers in Ti-based composite materials. The value of mixing energy in a matrix-fiber system characterizes the solubility of fibers. In other words, the greater the mixing energy is the smaller the solubility of a fiber in a matrix at a given temperature. Therefore, if an additive increases the energy of mixing in the matrixfiber system, then fiber degradation will decrease. In this study, the excess energy of mixing, Eex, was calculated as follows, which was defined as the difference between the energy of a solid solution, Esol, and the energy of the heterogeneous mixture, Ehet. Eex ¼ EsolEhet

(24.2)

Table 24.4 shows the calculated mixing energy, E, in several Ti-based systems with different additives. The highest mixing energy value was that of the alloy containing Mo, whereas the lowest was that of the alloy with Si. The meaning of these values was that Mo addition was expected to be the most effective element to reduce B solubility, whereas Si was the least effective additive according to the calculations. The decrease of boron solubility in Ti by additions of Mo will result in the reduction of TiB2 formation at the interface. Thus, certain additions, such as Mo, Zr, Al, etc., which decrease B solubility in Ti, are expected to be beneficial in improving interface behavior, by reducing TiB2 formation also.

Fibers made by chemical vapor deposition

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(a) 70

Atomic concentration (%)

60 50 40 30 20 10 0 0

10

20

30

40

50

60

70

Etch time (min)

(b) 80

Atomic concentration (%)

70 60 50

Carbon

40

Oxygen Titanium

30

silicon

20 10 0 0

10

20

30 40 50 Etch time (min)

60

70

Figure 24.14 AES sputter depth profile for the titanium-coated (a) SICABO and (b) BORSIC fibers. Adapted with permission from Tan BJ, Hwan L, Suib SL: Spectroscopic characterization of CVD Ti coating on SiC-coated boron fibers, Chem Mater 3:368e378, 1991. Copyright (2017) American Chemical Society.

24.3.4.2 Mechanical properties One of the key questions in the theory of fiber composite materials is the influence of the structure of the fiber-matrix interface on the strength of the reinforcing fibers and on the process of their failure. As mentioned in Section 24.3.4.1, due to the carbon-rich surface of SICABO fiber, the TiC layer is initially formed and decreases the further diffusion of Ti atom into the fiber, while the lack of carbon-rich coating leads to Ti

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Handbook of Properties of Textile and Technical Fibres

Table 24.4 Calculated mixing energy, E, in several Ti-based systems System

DE (J/mol)

TieB

440.3

(Ti/Al)eB

3879

(Ti/Ge)eB

1468

(Ti/Mo)eB

6054

(Ti/Si)eB

678.3

(Ti/V)eB

4586

(Ti/Zr)eB

4914

Reproduced from Revzin B, Pelleg J, Fuks D: Non-empirical study of the solubility of boron fibers in Ti-based composite materials, Compos A 29A:627e630, 1997 with permission from Elsevier.

Room temperature tensile strength MPa

diffusing deep into the SiC coating of the BORSIC fiber. Therefore, after titanium coating and subsequent heat treatment for a certain period of time in a He atmosphere, the tensile strength of the SICABO fiber was about the same as that of the as-received SICABO fiber, 3024  540 MPa. However, the tensile strength of the BORSIC fiber decreased to 2884  423 MPa. After a 0.5 and 1.0h heat treatment in a He atmosphere, the tensile strength of the BORSIC decreased to 2335  864 MPa, and 2153  1026 MPa (Fig. 24.15; Hwan et al., 1990). However, the heat treatment was

4000

3000

2000

1000

SICABO BORSIC

As As 0.5 coated coated

0.5

1.0

1.0

2.0

(Hour)

Figure 24.15 Room temperature tensile strength of SICABO and BORSIC CVD coated with Ti and heat treated in He. Reprinted from Hwan L, Tan BJ, Suib SL, Galasso FS. Interaction of titanium with SiC coated boron fiber, MRS Proc 168:239e346, 1990 with permission from Ambridge University Press.

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even extended to 2.0 h and the tensile strength of Ti-coated SICABO fiber only decreased slightly to 2756  423 MPa. As described above, during preparation and subsequent operation of fiber composite materials, the strength of the fibers may decrease as the result of their interaction with the matrix, which leads to the formation of brittle reaction products on the surface of the fibers. Model concepts of the influence of interaction products on the properties of the fibers have been developed based on the assumption that the interphase interface is a continuous layer of reaction products, which fails in the early stages of deformation of a fiber composite material (The interface in metal composites, 1978; Ochiai and Murakami, 1981; Ustinov, 1979; Shorshorov et al., 1989). The formed cracks serve as a stress raiser. If the stress concentration caused by the cracks in the layer is greater than the stress concentration from the defects of the fibers themselves then the latter lose strength. The amount of loss of strength is a function of the layer thickness. In this case, the strength of fiber is only related to the thickness of the interaction layer and independent of the action of the fiber defects themselves. Sirenko et al. (1993) proposed a new model, which described the loss of strength of brittle fibers as the result of interaction of the reaction products with the surface cracks. In the model, the criterion of the start of failure was the reaching of some critical value KIc by the full stress concentration in the vicinity of a crack located in the fiber. Microcracks were formed in the fibers in the stage of their preparation and considered as a defect of the fiber itself. The model was based on the experimentally established fact that in a number of actual fiber composite materials the interface did not have a continuous structure but contained individual dispersed reaction products (Salibekov et al., 1978; Astanin et al., 1988). KI was used to evaluate the strength of the fiber in the proposed model by Sirenko et al. Then the condition of crack development was represented in the form KI ¼ KIc, where KIc was the critical stress intensity factor determined using the Irwin equation (Broek, 1980). In the case of plane strain KIc ¼

pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Eg=ð1  mÞ

(24.3)

where g was the intensity of the surface energy expended for failure, E was Young’s modulus, and m was Poisson’s ratio. KI for a single crack was determined from the known equation pffiffiffiffiffi KI ¼ s pl

(24.4)

where l is half the crack length. KI of a crack-inclusion paired defect is calculated by using the method of construction of a system of singular integral equations: pffiffiffiffiffi     ¼ s pl 4 þ l2 ð3  kÞð5  kÞ 16k  l4 13k2  148k þ 343 256k K11   þ 0 l6 4 (24.5)

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Handbook of Properties of Textile and Technical Fibres

pffiffiffiffiffi     K12 ¼ s pl 3  k  l2 ð5  kÞ 4 þ l4 k3  49k2 þ 379k  75 256k   þ 0 l6 4 (24.6) pffiffiffiffiffi     K21 ¼ s pl Hl3 ð3  kÞð5  kÞ 128k þ 0 l5

(24.7)

pffiffiffiffiffi    ¼ s pl Hl3 ð7 þ kÞ þ 0 l5 K22

(24.8)

where l ¼ 2 l/d. d is the distance between the lines of the crack and of the inclusion, k ¼ 3e4m for plane strain, the second subscript corresponds to the number of the defect, 1 for a crack and 2 for an inclusion, and the first to strain, 1 for normal strain and 2 for transverse shear strain. KI for a crack and an inclusion is calculated by the following equations: KI ¼

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi K 211 þ K 221 ;

KI ¼

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi K 212 þ K 222

(24.9)

Since the factor KI for an inclusion is significantly less than for a crack, this factor is not taken into consideration subsequently in the model. Similar equations may be obtained for crack-crack and inclusion-inclusion types paired defects. Whereas since the number of surface cracks in a fiber does not change and for an inclusion KI is small and may be neglected, this model is limited to the consideration of a crack-inclusion paired defect. It should be noted that the proposed model is quite universal and may be used without any changes of analysis of the strength loss of any brittle fibers with dispersed interaction products.

24.3.4.3 Fracture characteristic of boron fiber in composites For boron-fiber-reinforced metal (B/M) composites, the interfacial bonding status would affect the tensile fracture and the failure mode of boron fiber. Luo and Sun (2003) studied the effect of the interfacial bonding status on the tensile fracture characteristics of a boron-fiber-reinforced aluminum (B/Al) composite. In their work, there were four kinds of interfaces to be considered, i.e., interface with chemical reactants and three other types of interfaces without chemical reactants but with weak, strong, and appropriate B/Al interfacial bonding strengths, respectively. Since a direct contact between boron fibers with the aluminum matrix leads to chemical reactions during the composite preparation, these boron fibers were coated with about 3-mm thick B4C layer to prevent such interfacial reactions. For comparison, uncoated fibers were also used. Without coating on the boron fibers, the interaction of boron fiber with Al matrix occurred during the preparation of composite, leading to a flat fracture surface after tensile fracture. In this case, most of the cracks indeed propagated from one side of the fiber to the other side, without evident pullout of boron fibers. This type of fracture morphology of boron fibers was classified as type II fracture

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(Ochiai et al., 1983; Astanin and Imayeva, 1994). For B4C-coated boron fiber, the different interfacial bonding status in B/Al composites led to different fracture morphologies of boron fibers. When the B/Al interfacial bonding was weak, some of boron fibers separated from the matrix, leading to many pulled-out fibers, which was characterized as type I fracture morphology of boron fibers (Ochiai et al., 1983; Astanin and Imayeva, 1994). When the B/Al interfacial bonding was strong, most of the fibers were indeed broken into smaller pieces without pulled-out fibers. This type of fracture morphology of boron fiber was different from types I and II, thus, it was named as type III fracture of boron fibers (Luo and Sun, 2003). When the B/Al interfacial bonding was appropriate, the fracture morphology of boron fiber combined types I and II, and the broken and pulled-out fibers were both observed. Fig. 24.16 summarizes the relationship between the interfacial bonding status and the fracture features. The fracture characteristics of boron fiber in B/M composite are related to the interfacial bonding status, whilst the interfacial shear strength (IFSS) plays a significant role on the fracture characteristics of boron fiber in a boron fiber/epoxy (Bf/epoxy) resin composite (Wen et al., 2007a). According to the study of Wen et al., 2006, when the IFSS was small, the fracture surface of boron fiber was flat and showed the radiating pattern, which was the classical brittle fracture. Moreover, there were a few pulled-out fibers and the pullout lengths were very small (Fig. 24.17(a)). This was the nonaccumulating failure. With the increase of IFSS, the failure mode changed from the nonaccumulating failure into the accumulating failure. In this case, a large amount of fibers were pulled out,

Figure 24.16 Schematic illustration showing the interface-fracture relationship. (a) Interface with reactants, (b) interface with weak interfacial bonding strength, (c) interface with strong interfacial bonding strength, and (d) interface with appropriate interracial bonding strength. Reprinted from Luo ZP, Sun CY: Effect of the interfacial bonding status on the tensile fracture characteristics of a boronefiber-reinforced aluminum composite, Mater Charact 50(1):51e58, 2003 with permission from Elsevier.

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Handbook of Properties of Textile and Technical Fibres

(a)

20 KV

100 X

100 μm

KYKY–2800

20 KV

200 x

100 μm

KYKY–2800

0

200 x

100 μm

KYKY–2800

0

Bf/(E-51+IPDA)

(b)

20 KV

0

100 X

100 μm

KYKY–2800

0

20 KV

Bf/(E-51+IPDA+10%LNBR)

Figure 24.17 Fractographys of unidirectional Bf/epoxy composites under tension (Wen et al., 2006).

and the pullout lengths were large with many resin fragments left on the fiber surface (Fig. 24.17(b)). The IFSS can be controlled by using different curing agents and by modifying the resin matrix. Wen et al. (2007b) also investigated the effects of different curing agents and the modification of the resin on the IFSS. Table 24.5 lists the IFSS of Bf/epoxy resin composites with different curing agents and with matrix modification.

Interfacial shear strength (IFSS) of Bf/epoxy composites (Wen et al., 2007b)

Table 24.5

Material system

IFSS (MPa)

Standard deviation (MPa)

Dispersion coefficient (%)

Bf/(E-51 þ DETA)

10.4

1.26

12.1

Bf/(E-51 þ DETA þ 10%LNBR)

24.5

2.19

8.95

Bf/(E-51 þ IPDA)

15.1

1.53

10.2

Bf/(E-51 þ IPDA þ 10%LNBR)

29.8

1.79

6.00

Fibers made by chemical vapor deposition

951

Through using different curing agents and modifying the resin matrix, the interfacial shear strength was effectively increased, subsequently improving the fracture performance of boron fibers in Bf/epoxy resin composites.

24.4

Silicon carbide fiber

24.4.1 Production 24.4.1.1 For W-core silicon carbide fiber SiC filament fibers are made by CVD on a substrate heated between 1100 and 1350
C. The substrate can be a tungsten or carbon filament. Tungsten wire has been widely used as a substrate of CVD SiC fibers in the world, such as the USA, the UK, and China, but carbon filament as substrate of SiC fiber is just used in the USA. Maybe it is hard to obtain qualified carbon filament substrate in other countries. Some advantages of SiC fiber with tungsten core are the following: (1) tungsten wire of the appropriate diameter is available in relatively defect free, large spools, unlike carbon monofilament; (2) the cost of producing W-core SiC fiber is lower than C-core SiC; (3) the modulus of W-core SiC fiber is generally higher than that of C-core SiC fiber, and this may be desirable in stiffness-critical applications (Gambone and Gundel, 1997). Therefore, although carbon-core SiC fibers exhibit superior tensile properties, W-core SiC fibers are also well developed. Fig. 24.18 shows the schematic of a horizontal cold-wall CVD reactor for fabricating W-core SiC fibers, which is a typical device for Sigma SiC fibers. During the production, a W wire substrate is drawn continuously through the reactor via mercury seals at both ends of each stage at a certain speed through a relatively long reactor (2.5e4 m). The mercury seals act as gas seals, as well as the contact electrodes to heat up the filament resistively. Firstly, the W wire (about 13e18 mm diameter) needs to be cleaned via heating in a H2 atmosphere at the cleaning chamber, as the surface of as-received W wire has some oxides or oil contamination. Wang and coworkers’s study showed that the tensile strength of W-core SiC fibers is sensitive to the surface quality of W wire (Wang et al., 2016). After the cleaning is finished, a SiC sheath is

8

1

3

4

5

6 7

2

1-Supply reel, 2- take-up reel, 3-mercury electrode, 4-cleaning chamber, 5-SiC deposition chamber, 6-surface coating deposition chamber, 7-W core, 8-DC power

Figure 24.18 A horizontal cold-wall CVD reactor for the fabrication of W-core SiC fiber.

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Handbook of Properties of Textile and Technical Fibres

deposited onto the W substrate in the SiC deposition chamber where a mixture of H2, Ar, and silane is introduced. H2 is a kind of reducing gas, and it is also used as a carrier gas to carry the silane into the deposition chamber. Ar gas is a kind of dilute and protective gas, and it also has a special “heat preservation” effect due to its extremely low thermal conductivity and specific heat (The room temperature thermal conductivity and specific heat of Ar are just 0.016 W/m/K and 0.523 J/g/K, respectively), which is favorable for the deposition of SiC (Guo, 2015). Typical silanes are methyltrichlorosilane (CH3SiCl3) and methyldichlorosilane (CH3SiHCl2). CH3SiHCl2 is quite active and decomposes fast, so fibers produced from solely CH3SiHCl2 and H2 would have a fast production rate but the tensile strength would be low. The surface of these fibers would be bumpy and nodular and the SiC mantle would contain free silicon. CH3SiCl3 looks an ideal raw material, as it contains one silicon and one carbon atom, i.e., one would expect a stoichiometric SiC to be deposited. The chemical reaction is H2

CH3 SiCl3 ðgÞ ! SiCðsÞ þ 3HClðgÞ CH3SiCl3 is chemically more stable but its decomposing rate is too slow, so fibers produced from it grow at a very slow rate but has a very smooth surface. The too slow rate makes the process economically undesirable (Bunsell, 1988; Zheng and Zhang, 1996; Zhang, 2011). In addition, for W substrate, the prolonged production runs would aggravate the W/SiC interfacial reaction, which would decrease the fiber strength. So a mixture of CH3SiCl3 and CH3SiHCl2 is often used, which can increase the production rate by a factor of three over CH3SiCl3 alone, yet gives a fiber with a smooth surface. The optimized ratio of CH3SiCl3/CH3SiHCl2 is 3:1 (Zheng and Zhang, 1996). It should be noted that the temperature profile of the filament in the SiC deposition chamber has a distinct gradient over the reactor length, as shown in Fig. 24.19 (Zhang et al., 2010a), which is attributed to the insulating properties of the SiC and resistance changes with temperature in the core region. The large temperature gradient (160e300
C in terms of the length of deposition chamber and diameter variation rate of the filament) not only gives a low overall deposition rate for the fiber but also gave a material of low quality. In order to minimize the magnitude of the temperature gradient, very high frequency (VHF w 60 MHz) induction heating can be augmented to obtain an optimum temperature profile (Bunsell, 1988; Chawla, 2009). Du and coworkers in China just used radio-frequency (RF) heating to reduce the temperature gradient for preparing W-core SiC fiber, and a high strength of 3.6 GPa was obtained (Du et al., 1998). However, as RF heating area is short, and high-frequency electromagnetic radiation is serious and extremely unstable, it is difficult to achieve a large-scale production (Wang et al., 2016). Thus the RF heating technique alone has been discontinued. Another method of reducing the temperature gradient is multistage heating to deposit SiC, which has also proved to be effective (Zhang et al., 2010b; Cai et al., 2005). After the SiC mantle is obtained, the fiber goes into the third stage of production to obtain a carbon coating where a mixture of hydrocarbon and Ar is introduced. The carbon coating is used to protect the fiber from handling damage, to serve as a sacrificial layer in titanium matrix composites and to provide a weak interface for matrix crack

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Deposition temperature (K)

1600

1500

1400

1300

1200 0.0

0.2

0.4

0.6 0.8 1.0 Distance (m)

1.2

1.4

1.6

Figure 24.19 Deposition temperature profile along the filament’s travel direction. Adapted from Zhang RJ, Yang YQ, Shen WT, Wang C, Luo X: Microstructure of SiC fiber fabricated by two-stage chemical vapor deposition on tungsten filament, J Cryst Growth 313(1): 56e61, 2010a with permission from Elsevier.

deflection in some ceramic matrix composites (Bhatt and Hull, 1998). Most importantly, the carbon coating can increase the fiber strength nearly two-fold (Liu et al., 2007; Guo et al., 1998).

24.4.1.2 For C-core SiC fiber As the interaction between W and SiC forms brittle compounds, which is deleterious to tensile the strength of the fibers, carbon filament substrates were introduced, which has been proved not to react with the deposited SiC and have a much smaller temperature gradient in the deposition chamber. A vertical production device of carbon-core SiC fiber is schematically shown in Fig. 24.20, which is a typical device for SCS SiC fibers (Bunsell, 1988; Debolt and Henze, 1978) and resembles the production of W-core SiC fibers. A 33 mm carbon filament substrate is drawn continuously through the reactor, and a w140 mm fiber with an inner carbon-rich SiC layer, middle SiC layer, and outer carbon-rich SiC coating can be formed in the final stage. According to Guo and coworkers’ studies, the defects at the carbon core do not have much effect on the tensile strengths of the fibers (Guo et al., 1998). The inner carbon-rich SiC layer is constructed by combining with the silane and propane and outer carbon-rich SiC coating by a blend of argon, propane, and some residual silanes. The outer coating is produced for reducing sensitivity to surface abrasion (Debolt and Henze, 1978).

24.4.2 Microstructure and chemical composition The mechanical properties of CVD SiC fibers are strongly dependent on their microstructure and composition, whilst the microstructure and chemical composition are quite sensitive to a number of parameters, including reactant gas species, carrier and/or reducing gases, deposition pressure, temperature, gas phase concentrations,

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Handbook of Properties of Textile and Technical Fibres

Spooling device

C Silane, hydrogen Propane, argon Silane hydrogen

Glass tubular reactor Φ≈ 2 cm Exhaust gases to recovery

Propane argon

SiC

Mercury seal

Figure 24.20 The schematic of carbon-core SiC filament production reactor

flow rate, etc. However, the fiber fabrication process has been subjected to significant advanced development, although many details of which still remain proprietary. There have been many papers in the literature studying the microstructure and composition of CVD SiC fibers, and this section will mainly introduce or summarize the microstructures of SCS SiC fibers and Sigma SiC fibers.

24.4.2.1 SCS SiC fibers The SCS fibers have a very important family of CVD-derived SiC fibers manufactured by Speciality Materials Inc. in the USA. Due to the fact that the properties, especially the strength of the filaments, are extremely sensitive to surface defects and abrasion, after deposition of SiC, the fiber is coated with an outer layer, a 1.5e3.0 mm protective carbon-silicon coating (this is the reason for the SCS designation) (Flores et al., 2014). The company supplies mainly two types of SiC fibers, SCS-6 and SCS-Ultra. SCS-6 SiC fiber is the lowest cost fiber, with high strength properties, wettability for metals, low electrical conductivity, high heat resistance, and corrosion resistance/chemical stability. SCS-Ultra SiC fiber is a fiber for the most demanding applications. Produced at the same diameter as SCS-6, SCS-Ultra is 50% stronger and 10% higher in modulus (http://specmaterials.com/a).

SCS-6 SiC fiber

The microstructure of as-produced SCS-6 SiC fiber was firstly studied in detail by Ning and Pirouz using different kinds of techniques including transmission electron microscopy (TEM), high-resolution episcopic microscopy, scanning Auger microscopy

Fibers made by chemical vapor deposition

955

(a)

Outer coating Outer SiC Inner SiC

C-rich fine grain β-SiC

SiC4 35 μm

SiC4

Si-rich coarse grain β-SiC

Sublayer 3

(b)

Sublayer 1 1.7 μm Sublayer 2 0.1 μm

Carbon

SiC3 4.5 μm 142 μm

SiC2 4.5 μm SiC1 6 μm

Inner pyrolytic C 1.5 μm

C core 33 μm

Outer C coating with fine β-SiC grains 3 μm

Figure 24.21 Cross-sectional review of the SCS-6 SiC fiber. (a) SEM micrograph of a fractured surface; (b) schematic of drawing of the fiber microstructure. (a) is reprinted and (b) is redrawn from Ning XJ, Pirouz P, Lagerlof KPD: The structure of carbon in chemically vapor-deposited SiC monofilaments, J Mater Sci 5(12):2865e2876, 1990 with permission from Cambridge University Press.

etc. (Ning and Pirouz, 1991). Fig. 24.21(a) shows an SEM micrograph of a fractured surface of SCS-6 fiber (Ning et al., 1990). One can primarily see that the fiber is a composite that is composed of three main layers: the carbon core, SiC, and outer coating. The SiC sheath is mainly composed of two parts: the inner darker SiC and the outer bright SiC. Kim et al. (1997) have made an excellent review of the structure of SCS-6 SiC fibers. Ning et al. (1990) and Ning and Pirouz (1991) observed that the 33-mm diameter carbon core consists of 1e5 nm blocks of randomly oriented turbostratic carbon (TC). TC blocks have covalently bonded basal planes similar to graphite. However, these planes are not stacked regularly like graphite, but are rotated randomly about the caxis. The carbon core is then coated with a 1.5 mm inner pyrolytic carbon layer consisting of 30e50 nm TC blocks, as shown in Fig. 24.21(b). These blocks are

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Handbook of Properties of Textile and Technical Fibres

preferentially oriented with their c-axis parallel to the radial direction of the fiber. The purpose of this coating is to smooth the substrate surface and enhance the electrical conductivity of the fiber during deposition (Guo et al., 1998). In the next step, the coated substrate is exposed to silane (as a source of silicon and then SiC) and hydrogen to deposit a 50-mm thick SiC layer. This layer is the main structural element in the fiber, which consists of four layers identified by Ning and Pirouz (1991) from the innermost to the outermost layers as shown as SiC1, SiC2, SiC3, and SiC4 in Fig. 24.21(b). The SiC1 layer next to the inner carbon coating is 6 mm thick. The first 1 mm of this layer consists of randomly oriented rod-shaped b-SiC grains with 5e15 nm lengths and aspect ratios of 10. Further away from the carbon boundary, the SiC grains are 50e150 nm in length and become increasingly aligned with their major axes parallel to the radial direction. The SiC2 layer is 4.5 mm thick. At the boundary between this and the SiC1 layer there is a sharp increase in the size of the SiC2 elongated grains as well as an increase in the alignment with the radial direction. Another sharp increase in size was observed with the transition to the SiC3 layer. This layer is also 4.5 mm thick with elongated grains about 1 mm long. These SiC3 grains also showed substantial stacking faults on the <111> planes normal to the radial direction. The final SiC layer, SiC4, is 35 mm thick. The grains at the inner boundary of this layer are 250 nm in diameter (about twice the size of the SiC3 grains) and decrease to 90 nm diameter at the outer edge (Kim et al., 1997). Guo et al. (1998) also studied the morphology of SiC4 by TEM, and their results show that there are a great many stacking faults and/or micro-twin faults along the {111} plane of the grains (see Fig. 24.22). The planar disorder is common in b-SiC due Radial direction

0.2 μm

Figure 24.22 Bright-field TEM image of the SiC4 region of SCS-6 fiber adjacent to the outermost coating. Reprinted from Guo SQ, Kagawa Y, Tanaka Y, Masuda C: Microstructure and role of outermost coating for tensile strength of SiC fiber, Acta Mater 46:4941e4945, 1998 with permission from Elsevier.

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to the very low stacking defect energy of the {111} plane of b-SiC. SiC1eSiC3 layers together are usually called the inner SiC sheath, while SiC4 the outer SiC sheath, which is corresponding to Fig. 24.21(a). In the final deposition step, a 3.0-mm thick carbonaceous coating is added to the fiber. Basically, this coating consists of a carbon matrix with SiC particles embedded in it. The microstructure feature of the carbon matrix is that the basic structural unit blocks of the TC in the coating have a preferred orientation with the basal planes in each block being nearly parallel to the fiber axis. The coating can be divided into three sublayers, 1, 2, and 3, which differ in the fine microtexture of the carbon matrix and the distribution of SiC particles, as shown in Fig. 24.23 (Ning and Pirouz, 1991). Sublayer 1, with a thickness of 1.7 mm, is deposited on the rough saw-tootheshaped surface of SiC4 grains to heal the crystalline surface of SiC for improved surface strength. Within this sublayer, and in the region nearest to the SiC4 boundary, there is a zone of thickness w150 nm that consists of small crystallites of SiC embedded in a matrix of carbon. The size of the b-SiC crystallites is about 5e10 nm. Covering this is a w0.6 mm layerddenoted by 1A in Fig. 24.22dwhich contains larger SiC particles of diameters 20e50 nm. Next to region 1A is a 0.9 mm carbon layer denoted by 1B shown in Fig. 24.22. In this region the SiC particles are much finer, with average diameters of 3e5 nm. The thickness of sublayer 3 is w1.3 mm. It is subdivided into regions 3A and 3B, which is similar to those of 1A and 1B, respectively. Sublayers 1 and 3 are separated by a thin sublayer 2 (w100 nm) with no, or a lower density of, SiC

Figure 24.23 Cross-sectional view of the outermost fiber coating of SCS-6 fiber. Adapted from Ning XJ, Pirouz P: The microstructure of SCS-6 SiC fiber, J Mater Res 6(10): 2234e2248, 1991 with permission from Cambridge University Press.

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Handbook of Properties of Textile and Technical Fibres

SiC Subfiber layer I

Sublayer II

Sublayer III

Figure 24.24 A bright-field TEM image of the outermost coating of the as-received SCS-6 SiC fiber showing the multiple layer structure. Reprinted from Guo SQ, Kagawa Y, Tanaka Y, Masuda C: Microstructure and role of outermost coating for tensile strength of SiC fiber, Acta Mater 46:4941e4945, 1998 with permission from Elsevier.

particles. The formation of sublayer 2 is due to a change in the CVD conditions, possibly during the transition from deposition of sublayer 1 to that of sublayer 3 (Ning et al., 1990; Ning and Pirouz, 1991). Sublayer 2 has been found to be the weakest layer, with failure often found to occur here in ceramic matrix composites (Wither et al., 2010). Guo et al. (1998) also studied in detail the outermost coating of the SCS-6 fiber later by TEM. Their result of a bright-field TEM image of the outermost coating is shown in Fig. 24.24. The coating is also divided into three major sublayers: sublayer I, II, and III, and each sublayer can be further divided into A and B regions in terms of the size and distributions of SiC particles or needles. The main differences between Figs. 24.23 and 24.24 are the morphologies and distribution of SiC phases, which must be caused by different CVD conditions. So the different layering methods have no essential differences. Further study of the SCS-6 fiber by Ning et al. (1993) using a variety of microchemical analysis techniques which showed that the inner SiC region (SiC1eSiC3) was carbon rich. So the inner SiC region looks darker under SEM, as shown in Fig. 24.21(a). The carbon was amorphous or turbostratic in nature and is present only at the grain boundaries. In the SiC1 layer there was approximately 60 at.% carbon, which decreased in subsequent layers with the SiC4 layer consisting of stoichiometric SiC (50 at.% carbon) (Kim et al., 1997). Later in 2005, Chollon et al. studied the local chemical composition of the SCS-6 SiC fiber along its cross-section using electron probe microanalysis (EPMA), and the result is shown in Fig. 24.25 (Chollon et al., 2005). The result

Fibers made by chemical vapor deposition

Atomic concentration (%)

60

959

SCS-6

55 Si 50 C 45 PyC

C core

40 0

10

20

30

40 x (μm)

50

60

70

Figure 24.25 EPMA radial profile of the SCS-6 filament. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

further proved that the inner SiC sheath has excess of C while the outer SiC sheath has a little excess of Si. Raman spectroscopy has been applied successfully to characterize SiC crystal polytypes, SiC particles, and polymer-based fibers. The position, relative intensity, and bandwidth of Raman spectra are closely related to crystal polytypes, crystal size, and defects in the structure. The Raman technique also has the advantages of being very quick and efficient, as well as requiring little sample preparation. So Ramanspectroscopy has been widely used to analyze CVD SiC fibers (Ward et al., 2004). Ward and coworkers (Ward et al., 2001; Shatwell et al., 2001) systematically studied the microstructure of SiC fibers through the use of Raman microscopy. Fig. 24.26 shows the Raman mapping across an SCS-6 fiber starting from the fiber centre and moving toward the surface. All the spectra taken from the core region show only the carbon bands, indicating the fiber has a carbon core. It should be noted that the carbon bands in the core are intense and narrow. This indicates that the fiber core has wellordered graphite carbon crystallite structure. Moving away from the fiber core to the inner region, the Raman spectra reveal the presence of SiC bands as well as the carbon bands. This confirms that this region consists of a mixture of SiC and carbon. It should also be pointed out that, in addition to lower intensities and broader widths compared to the carbon core region, the carbon bands obtained moving away from the core show a decrease in intensity and an increase in bandwidth. On the other hand the SiC bands change very little in terms of intensity and bandwidth throughout the region. This indicates that the carbon structures in this region become more disordered. Furthermore, the Raman bands from SiC in this region are often weak in the presence of carbon. This can be explained by the fact that the Raman scattering efficiency of SiC is much lower than that of carbon. The carbon bands disappear abruptly at a distance about 45 mm away from the fiber centre and only SiC bands can be observed in this region. This

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Handbook of Properties of Textile and Technical Fibres

(a)

3500 3000 2500 2000 1500 1000 500 0

SiC optical phonons TO

LO

Intensity

Carbon bands

SCS-6 fiber

400

800

1200

1600

2000

the Dista fibe nce rc f en rom tre ( μm )

10 20 30 40 50 60 70 Raman wavenumber (cm–1) Carbon coating

(b) 2000

Raman wavenumber (cm–1)

1800

Carbon core

Core boundary

Mid-radius discontinuity

1600 Carbon

1400 1200

Sic optical phonons

Intensity 3160 -- 3500 2820 -- 3160 2480 -- 2820 2140 -- 2480 1800 -- 2140 1460 -- 1800 1120 -- 1460 780.0 -- 1120 440.0 -- 780.0 100.0 -- 440.0

1000 LO

800

TO

SiC optical phonons

600 SiC acoustic phonons

400 0 Centre

10

20 30 40 50 60 Distance from the fiber centre (μm)

70 Surface

Figure 24.26 (a) Raman line mapping and (b) Raman scattering intensity contour map from the centre to the surface of the fiber for an SCS-6 fiber. Reprinted from Ward Y, Young RJ, Shatwell RA: A microstructural study of silicon carbide fibres through the use of Raman microscopy, J Mater Sci 36:55e66, 2001 with permission from Springer.

indicates that the outer SiC region should be mainly SiC. However, as will be introduced in next paragraph, the outer SiC region has a slight excess of silicon, but the Raman technique cannot detect it. The reappearance of only carbon bands at the fiber surface indicates that the fiber has a carbon coating. The carbon bands from the outer coating are weak and extremely broad, indicating the carbon structure is highly

Fibers made by chemical vapor deposition

961

disordered (Ward et al., 2001). As has been mentioned before, the Raman scattering efficiency of carbon is so high that the Raman signal of very fine SiC particles in the coating cannot be detected in Fig. 24.26. However, Kim and coworkers’ studies showed that there existed a weak and broad SiC TO peak at about 790 cm1 when they analyzed the outer coating, which proved the presence of SiC (Kim et al., 1997). According to Fig. 24.25, the outer SiC sheath has a slight excess of Si. This result is different from earlier opinion, which thought that the outer SiC region was made up of only stoichiometric SiC (Ning et al., 1990; Ning and Pirouz, 1991). The presence of excess of Si was firstly reported by Lara-Curzio et al. (1993) according to their studies of high-temperature structural stability of SCS-6 filament. Thermal events observed by their differential scanning calorimetry analysis occurred at 1380
C on heating and at 1300
C on cooling were suggested as being related to the melting and solidification, respectively, of excess silicon. These results can explain the anomalous thermal expansion behavior of the filament between 1300 and 1450
C. Later Kim et al. (1996) also proved the presence of excess Si according to their Raman spectroscopy analysis to the SCS-6 filament processed at 1640
C for 40 min in a zircon matrix. In their work a strong peak in the outer SiC layer was identified as a Si signature. However, no Raman Si response was observed in the as-received fiber due to the small disordered and scattered Si. Bhatt and Hull also pointed out that the outer SiC region has an excess of Si according to their microprobe analysis of the cross-section of the as-received fibers (Bhatt and Hull, 1998). The existence of a slight excess of Si in the outer SiC region was also found by Chollon et al. (2007) who studied the annealed (1400
C for 2.2  104 s) and crept (at 1400
C for 2.2  104 s under a longitudinal tensile stress of 300 MPa) SCS-6 fiber by Raman microspectroscopy analyses. The presence of a small amount of carbon in SiC fibers can maintain the high strength in fibers by suppressing the growth of SiC crystals, whilst the presence of silicon may have an adverse effect (Kim et al., 1997; Sasaki et al., 1987).

SCS-Ultra SiC fiber

SCS-Ultra SiC fiber was developed to meet the most demanding applications for aerospace materials. In the mid-1990s, turbine engine manufacturers wanted a fiber that could withstand the fabrication conditions of titanium aluminde to make a composite for use in rotating engine parts. So Specialty Materials, Inc. developed SCS-Ultra, a fiber with very fine SiC crystallites (200 nm or less) that is 50% stronger and 10% stiffer than SCS-6 fiber and is stable at temperatures above 1371
C (http://specmaterials.com/b). However, as the SCS-Ultra fiber is the newest one and quite expensive, there is much less published literature on SCS-Ultra in comparison with that of SCS-6 and SM1140þ fibers. Fig. 24.27 shows cross-sectional SEM micrograph of the SCS-Ultra filament. Like the SCS-6 fiber, the deposition of SCS-Ultra takes place on a 33 mm carbon filament, which again is overcoated with a 1.5-mm thick pyrolytic carbon layer. However, the SUS-Ultra exhibits a relatively regular and fine-grained microstructure throughout the whole SiC layer. The first deposit apparently consists of very fine nanocrystallites. A radially oriented columnar structure gradually appears at an increasing distance from

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Handbook of Properties of Textile and Technical Fibres

Figure 24.27 SEM micrograph of the SCS-Ultra fiber. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

the core, with grains of up to 1 mm long at approximately 10 mm from the core. The columnar grains are still observed along the rest of the deposit but their size gradually decreases until the surface of the filament. This feature agrees well a previous TEM investigation, as shown in Fig. 24.28 (Shatwell et al., 2001), wherein the crystallites are columnar along the radial direction of the fiber and significantly smaller than those of SCS-6 and SM1140þ. Smith et al. (1998) also described the homogenous and fine microstructure as equiaxial SiC grains according to their SEM observation to the SCSUltra filament. Fig. 24.29 shows a summarized schematic diagram of microstructure of the Ultra SCS fiber.

Fibers made by chemical vapor deposition

963

0.2 μm

7449

Figure 24.28 TEM OF SCS-Ultra approximately 10 mm from the core/SiC interface, showing the small crystallite size (Shatwell et al., 2001).

Very fine β-SiC grains with a slight excess of carbon

140μm

C core Pyrolytic 33 μm C 1.5μm

Multi-layered coating 2-3μm

Figure 24.29 Schematic of microstructure of the Ultra SCS fiber.

An EPMA chemical composition analysis of the SCS-Ultra filament is shown in Fig. 24.30 (Chollon et al., 2005). It can be seen that the radial elemental composition of the SCS-Ultra differs from that of the SCS-6. The carbon concentration is significantly higher near the carbon core (about 20 at.% of free carbon). There is a gradual

964

Handbook of Properties of Textile and Technical Fibres

60 Atomic concentration (%)

SCS-ultra 55 C 50 Si

45 C core

40 0

10

PyC 20

30

40 x (μm)

50

60

70

Figure 24.30 EPMA radial profile of the Ultra SCS filament. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier

decrease in free carbon with radius, the bulk of the fiber being essentially stoichiometric SiC. The carbon concentration slightly increases again near the surface of the filament (w4 at.% of free carbon). The microstructure of the SCS-Ultra fiber can be also clearly illustrated by Raman mapping, as shown in Fig. 24.31. The spectra from the fiber core reveal only Raman 2000

2800 -- 3100 2500 -- 2800 2200 -- 2500 1900 -- 2200 1600 -- 1900 1300 -- 1600 1000 -- 1300 700.0 -- 1000 400.0 -- 700.0 100.0 -- 400.0

Carbon core

1800 Raman wavenumber (cm–1)

Intensity

Core boundary

Carbon

1600 1400

Sic optical phonons

1200 1000

LO

800

TO

600 400 0 Centre

10

20 30 40 50 60 Distance from the fiber centre (μm)

70 Surface

Figure 24.31 Raman scattering intensity contour map from the core to the surface of the fiber for an SCS-Ultra fiber. Reprinted from Ward Y, Young RJ, Shatwell RA: A microstructural study of silicon carbide fibres through the use of Raman microscopy, J Mater Sci 36:55e66, 2001 with permission from Springer.

Fibers made by chemical vapor deposition

965

bans characteristic of carbon, confirming the fiber has a carbon core. SiC bands can be observed throughout the rest of the fiber whilst the carbon bands disappear at about 35 mm from the centre and reemerge at around 15 mm from the surface. Thus the fiber has three different regions, for the inner and outer regions, a mixture of carbon and SiC, and for the middle region, almost stoichiometric SiC. This result agrees well with the composition analysis result of Fig. 24.30. It is interesting to see that the carbon bands of the carbon core are broadening and shifting toward higher frequencies toward the edge of the core. This means that carbon becomes more disordered near the core/SiC interface and is possibly in compression. The structures of carbon and SiC are distinctly different in different regions. In the inner region, carbon near the fiber core exhibits sharp and strong bands at 1330 and 1600 cm1 characteristic of well-ordered graphitic carbon structure. It becomes more disordered and possibly amorphous toward the fiber surface, as observed form the progressive broadening of the bands and appearance of a broad feature at 1400 cm1. On the other hand, SiC in the inner region remains poorly crystallized with small crystal sizes. The middle and outer regions contains almost no carbon or only a small amount of carbon going toward the fiber surface (Ward et al., 2001). SiC bands have relatively sharp TO and LO peaks whilst the acoustic modes (300e650 cm1) in comparison with those of SCS-6 indicate a better crystalline state of the CVD-SiC (i.e., closer to b-SiC) (Chollon et al., 2005).

24.4.2.2 Sigma SiC fibers For sigma CVD SiC fibers, instead on a carbon core, the SiC coating is deposited on the surface of a small-diameter (w15 mm) tungsten wire produced by hot drawing. The company Tisics in the United Kingdom manufactures four types of sigma CVD SiC fibers, named SM1140þ, SM1240, SM2156, and SM3256. Following a similar designation as the SCS family, the four types of sigma fibers differ in terms of the microstructure of the SiC deposited on the W core, which depends on the conditions of the CVD process (Flores et al., 2014; Rix et al., 2017). This section will mainly introduce the microstructure of typical SM1140þ and SM 2156 fibers.

SM1140þ SiC fiber

A typical cross-sectional fracture morphology of SM1140 þ SiC fiber is shown in Fig. 24.32 (Dyos and Shatwell, 1999). It can be seen that the fiber has a diameter of about 105 mm and is mainly composed of a W core, SiC mantle, and carbon coating. The W core has a diameter of about 15 mm, and the thickness of SiC sheath is 40 mm. The outer carbon coating is 5 mm thick, which is used as a sacrificial layer in titanium matrix composites to protect the inner SiC. The carbon coating has a strong bonding to adjust the coating conditions (Shatwell, 1994). Cheng et al. (1999) have detailedly investigated the microstructure of as-received sigma 1140þ SiC fibers by TEM. They revealed that the microstructure includes the tungsten core (W), a W/SiC interfacial layer, four distinct concentric annular rings of SiC (IeIV), and an outer turbostratic carbon coating. The W core consists of equiaxed grains z0.2 mm in size. Since the W wires were produced commercially by hot

966

Handbook of Properties of Textile and Technical Fibres

37.5μm

Figure 24.32 Scanning electron micrograph of Sigma SM1140þ filament (Dyos and Shatwell, 1999).

drawing, the grains elongated with their major axes parallel to the fiber axis. There is an interfacial reaction layer between the W and SiC, which is about 120 nm thick and consists of a W2C layer adjacent to W core and WSi2 layer close to SiC. The SiC sheath is divided into four regions in terms of the different morphologies of the SiC grains, which are coarse columnar, inner equiaxed, fine columnar, and outer equiaxed, respectively. Region I is the coarse columnar layer, which is the innermost and thickest (z20 mm) of the four and consists of coarse columnar grains with their long axes parallel to the radial directions. Fig. 24.33(a) and (b) are an image and diffraction pattern

Figure 24.33 (a) Bright-field TEM image showing the microstructure of the SiC deposit in the coarse columnar region (I); (b) diffraction pattern obtained from this region. Reprinted from Cheng TT, Jones IP, Shatwell RA, Doorbar P: The microstructure of sigma 1140þ fibres, Mater Sci Eng A 260:139e145, 1999 with permission from Elsevier.

Fibers made by chemical vapor deposition

967

obtained form this region, respectively, showing that the grains are heavily twinned and strongly textured with their <111> directions roughly parallel to the fiber radius. The diameter of the grains is z 200 nm at the SiC/W interface. This increases to z400 nm at a distance of z5 mm from the W surface and then decreases gradually to <50 nm at a distance of 15e20 mm from the W surface. A ring that corresponds to elemental Si can be seen in the diffraction pattern, which indicates the presence of excess Si in the form of small grains. However, no such Si grains have been observed directly in TEM images from this region. Inner equiaxed layer (region II) is about 10 mm thick and consists of equiaxed grains that are less than 20 nm in diameter. Fig. 24.34(a) and (b) are an image and diffraction pattern obtained from this region, respectively, showing no apparent texture. At the same time, free Si in the form of small grains can be detected in the diffraction pattern. The region III is about 5 mm across and consists of textured fine columnar grains. This is similar to layer I except that the grains are much smaller in size (<20 nm in diameter). However, the amount of free Si may be decreasing according to the diffraction pattern. The region IV is about 5 mm thick and consists of fine equiaxed grains (<10 nm) with very little, if any, free Si, as no obvious Si ring can be observed in the diffraction pattern. The outermost carbon coating has a high degree of graphitic order with basal planes aligned approximately parallel to the fiber surface (Cheng et al., 1999). According to above analyses, a schematic diagram of microstructure of the SM1140þ fiber is presented in Fig. 24.35. EPMA chemical composition analysis of the SM1140þ fiber studied by Shatwell et al. (2001) shows that the Si concentration gradually increases outwards. The result is quite similar to that of the SM1156 fiber. SM1156 SiC fiber is an early development version similar to the commercial SM1140þ but with an outer diameter of 147 mm instead of 105 mm. Fig. 24.36 (Chollon et al., 2005) shows the EMPA analysis results of SM1156. It can be seen that the SM1156 filament consists of stoichiometric SiC near the W core only. The Si concentration gradually increases outwards to reach a

(a)

(b)

111Si

Figure 24.34 Bright-field TEM image (a) and diffraction pattern (b) obtained from the inner equi-axed SiC region (II). Reprinted from Cheng TT, Jones IP, Shatwell RA, Doorbar P: The microstructure of sigma 1140þ fibres, Mater Sci Eng A 260:139e145, 1999 with permission from Elsevier.

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Handbook of Properties of Textile and Technical Fibres

Equiaxed SiC grains 10 μm

Coarse columnar SiC grains 20 μm

105 μm

Fine equiaxed SiC grains 5 μm

W core W/SiC reaction 15 μm layer 120 nm

Fine columnar SiC grains 5 μm

Carbon coating 5 μm

Figure 24.35 Schematic of microstructure of the Sigma 1140þ SiC fiber.

60 Atomic concentration (%)

SM1156 55

Si

50 C

45 W core

PyC

40 0

10

20

30

40 x (μm)

50

60

70

Figure 24.36 EPMA radial profile of the SM1156 filament, which is similar to that of SM1140þ. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

Fibers made by chemical vapor deposition

969

free Si excess of about 10 at.%. The result of SM1156 also reflects the chemical composition distribution feature of SM1140þ. This chemical composition analysis agrees well with Cheng et al.’s TEM analyses (Shatwell et al., 2001; Cheng et al., 1999), which confirms the existence of free Si in the SiC matrix. The microstructure of the SM1140þ fiber has also been clearly illustrated by Raman microscopy analyses, as shown in Fig. 24.37. Only the Raman bands assigned to SiC are observed near the W core, suggestive of stoichiometric SiC in the region. The intense features below 600 cm1 are attributable to free Si and SiC acoustic phonons. The presence of a sharp band at 520 cm1 and a broad band centred at around 480 cm1 indicate that there exist both crystalline and amorphous Si in the b SiC matrix. The only carbon bands are observed on the fiber surface from the outer coating, confirming that the fiber has a pure carbon coating. The SiC bands are very sharp and strong near the W core and become broad and weak moving out toward the fiber surface. This is due to the decrease in crystal size and perfection of SiC across the fiber section. This result agrees well with Cheng and coworkers’ TEM analyses (Cheng et al., 1999).

SM2156 SiC fiber

The SM2156 filament is a material developed by QinetiQ (now named as Tisics) in 2000s. In order to reduce the amount of free Si co-deposited with SiC, to subsequently improve the stiffness and thermal stability of the filament, propene (C3H6) was added to the dichloromethylsilane/hydrogen mixture used for the SM1140 or 1156 filament (Chollon et al., 2005; Chollon, 2007). SM2156 is 140 mm in diameter and has a tensile strength between 4 and 5 GPa. EPMA chemical composition analyses of the SM2156 are shown in Fig. 24.38. The results show a slight carbon excess (z8 at.% of free carbon) close to the W core. The free carbon amount fluctuates slightly along the first 20 mm of the coating. The composition remains almost stoichiometric throughout the rest of the SiC sheath, though with a very slight carbon excess (Chollon et al., 2005). The microstructure of the SM2156 filament is very similar to that of the Ultra SCS, but with a smaller SiC grain size (Fig. 24.39). The initial coating is nanocrystalline and radial columnar submicrometer grains appear at increasing distances from the core. A discontinuity in the CVD growth is observed at about 10 mm from the core interface. The columnar shape of the SiC crystallites is still observed along the rest of the deposit but their size progressively decreases until the surface. A schematic diagram of the structure of SM2156 filament is shown in Fig. 24.40. The Raman microspectroscopy analysis of the SM2156 filament is shown in Fig. 24.41. It can be seen that the Raman carbon features are visible near the W/SiC interface as a single broadband at 1100e1700 cm1, characteristic of a very disordered form of carbon. The evolution of the intensity of this band along the radius is consistent with the free carbon concentration obtained by EPMA. The SiC phase appears rather well crystallized, though in a more faulted form than in the SCS-Ultra (Chollon et al., 2005).

24.4.3 Tensile properties of as-produced fibers Fiber strength is one of the most important factors determining mechanical properties in continuous fiber reinforced composites. So tensile strengths of CVD SiC fiber will

970

Handbook of Properties of Textile and Technical Fibres

(a)

SiC

Optic

SM1140+

TO

3000

Acoustic LO

2000

0 μm

Intensity

10 μm 20 μm

1000 Silicon (c)

30 μm 40 μm

0

Silicon (am) Surface coating

Carbon

–1000

500

1000 1500 2000 2500 Raman wavenumber (cm–1)

3000 Intensity

(b)

995.0 -- 1100 890.0 -- 995.0 785.0 -- 890.0 680.0 -- 785.0 575.0 -- 680.0 470.0 -- 575.0 365.0 -- 470.0 260.0 -- 365.0 155.0 -- 260.0 50.00 -- 155.0

Wavenumber (cm–1)

1800 1600

Carbon from coating

1400 1200 1000 800

LO TO

SiC optical phonons crystalline Si

600

SiC acoustic phonons and Si

400 0 5 W/Sic interface

10

15

20

25

Distance/μm

30

35

40 Surface

Figure 24.37 (a) Raman spectra obtained at different positions from the W core to the surface of SM1140þ filament. (b) Raman scattering intensity contour map from the edge of W core to the surface of the SM1140þ fiber. (a) is reprinted from Ward Y, Young RJ, Shatwell RA: Application of Raman microscopy to the analysis of silicon carbide monofilaments, J Mater Sci 39:6781e6790, 2004 with permission from Springer. (b) is reprinted from Ward Y, Young RJ, Shatwell RA: A microstructural study of silicon carbide fibres through the use of Raman microscopy, J Mater Sci 36:55e66, 2001 with permission from Springer.

Fibers made by chemical vapor deposition

971

60 Atomic concentration (%)

SM2156 55 C 50 Si 45 W core

PyC

40 0

10

20

30

40 x (μm)

50

60

70

Figure 24.38 EPMA radial profile of the SM2156 filament. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

be summarized and discussed here. Ceramic fibers possess a wide spread of strengths because of inherent internal/surface flaws, which are usually assumed to be distributed randomly. The spread of fiber strengths about a mean value can be uniquely analyzed by the classic equationdthe two-parameter (s0 and m) Weibull distribution function (Baik and Grant, 2001; Gonzalez and Llorca, 2001) (see Chapter 1):  m  s PF ¼ 1  exp  s0

(24.10)

where PF is the cumulative failure probability up to a stress s, s0 is the reference strength or mean strength, and m is the Weibull modulus (shape parameter). The strength data were sorted in ascending order and the failure probability corresponding to each strength obtained by using the mean rank estimator: PF ¼

i Nþ1

(24.11)

where i is the ranking of the data and N is the total number of fibers tested. Eq. (24.1) can be transformed into the flowing equation by twice taking the natural logarithm: ln½ lnð1  PF Þ ¼ m ln s  m ln s0

(24.12)

Then the values of m and s0 can be determined from the slop and the y-intercept of the curve of ln[eln(1ePF)] versus L ¼ lns. The Weibull modulus m reflects the uniformity of the strength distribution, and a low value of m implies a large spread of measured strengths.

972

Handbook of Properties of Textile and Technical Fibres

Figure 24.39 SEM micrograph of the SM2156 filament. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

Table 24.6 summarizes properties of typical CVD SiC fibers. For reference, the properties of carbon core and tungsten wire are also given. It is clear that the W wire has higher density and tensile modulus than the carbon filament, so the W-core SiC fibers generally have a higher density and tensile modulus, too. The SCS-Ultra SiC fiber has the highest tensile strength due to its very fine SiC grains and nearly stoichiometric SiC composition, as have been described in the microstructure section.

Fibers made by chemical vapor deposition

973

SiC-251 μm SiC-110 μm 147 μm W/SiC interfacial reaction layer

W core 15 μm

Carbon coating 5 μm

Figure 24.40 Schematic of microstructure of the Sigma SM2156 SiC fiber. 1 +

W core

2 +

3 +

4 +

5 +

6 +

7 +

8 +

9 10 11 12 13 14 + + + + + +

PyC

W core

SM2156

1

SiC TO

2

LO 3

Free carbon

4 5

SiC-A

6 7 8 9 10 11 12 13 14

Surface 200 400 600 800 1000 1200 1400 1600 1800 Wavenumber (cm–1)

Figure 24.41 Raman microspectroscopy radial profile of the SM2156 filament. Reprinted from Chollon G, Naslain R, Prentice C, Shatwell R, May P: High temperature properties of SiC and diamond CVD-monofilaments, J Eur Ceram Soc 25:1929e1942, 2005 with permission from Elsevier.

Table 24.6

Properties of as-produced CVD SiC fibers Density (g/cm3)

Tensile strength (MPa)

Tensile modulus (GPa)

2350

390

3270

385

Weibull modulus

Outer coating

Coating thickness (mm)

Thermal expansion coefficient (10L6/8C)

None

/

4.1

Carbon rich

1.3

Carbon rich

3

Fiber type

Core material

Diameter (mm)

SCS-0

33 mm carbon

136

SCS-2

33 mm carbon

142

SCS-6

33 mm carbon

142

3.08

z4000

380

SCS-9A

33 mm carbon

78

2.8

z3500

z310

Carbon rich

3

SCS-Ultra

33 mm carbon

142

3.08

5900

415

Carbon rich

1.5e3

SM1040

14 mm W wire

96

3340

390

None

SM1140þ

14 mm W wire

105

3.4

3300

400  20

SM1240

14 mm W wire

100

3.4

3400

400

3.08

6.8

17.2

19.0

References Guo et al. (1998) Petitcorps et al. (1988)

4.1

Guo et al. (1998) and http:// specmaterials.com/a Flores et al. (2014) and Bhatti et al. (1999)

4.1

Guo et al. (1998) and http:// specmaterials.com/a Petitcorps et al. (1988)

Carbon

5

Flores et al. (2014) and García-Leiva et al. (2002)

C/TiBx

1/1.5

Guo and Derby (1998)

SM2156

14 mm W wire

147

3.4

4000e5000

SM3256

14 mm W wire

140

3.4

4000  200

Trimarc 1

12.5 mm W wire

127

3080

Chinese SiC fiber

18 mm W wire

110

z3500

W wire

/

14.5

/

26

/

34

Carbon filament

19.35

3300  200 3220  36

2.25

800

385  5 51

16.4

408  19

Carbon

4e5

Chollon et al. (2005)

Carbon

3e4

Rix et al. (2017)

Triple carbon rich

z4

Smith et al. (1998)

Carbon

2e3

Luo et al. (2015)

Faucon et al. (2001) 84 7.4

Petitcorps et al. (1988) Petitcorps et al. (1988)

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SCS-6 fiber is often used as reinforcements for titanium matrix composites, whilst SCS-Ultra fiber can be applied for titanium-aluminide matrix composites or ceramic matrix composites at higher temperatures. The SCS-9 fiber is similar to the SCS-6 filament except that it has a smaller diameter, and it can be used for resin-matrix or metal matrix composites. However, SCS-9 fiber is currently not commercially available. The SM2156 fiber has a similar structure to that of the SCS-ultra fiber and has a high tensile strength (4e5 GPa) due to its nearly stoichiometric SiC and with a little excess of free carbon. The chemical composition and microstructure of SiC in the SM2156 is comparable to that of the SCS-Ultra. However, as the W/SiC interfacial reaction is unavoidable, the tensile strength of the former is still distinctly lower than that of the latter. Recently it is reported that Tisics company has developed a new type of SiC fiber, named SM3256, which is the successor to SM1140þ and SM2156. The filament is fabricated in a single reactor using a high-speed CVD process at a rate of 8 m/ min. Stoichiometric SiC is deposited by using a mixture of CH3SHCl2, propene, argon, and hydrogen. The tensile strength of the fiber has a small variation of 4.0  0.2 GPa, and the Weibull modulus of the fiber is 51, which is exceptionally high compared to equivalent ceramic monofilaments (Rix et al., 2017). W-core CVD SiC fiber has also been successfully developed in China. Du et al. (1998) from the Institute of Metal Research, Chinese Academy of Sciences reported that a tensile strength of about 3600 MPa and Weibull modulus of 11.9 of CVD SiC fiber was obtained by using radio-frequency heating. Recently Wang et al. reported that they could fabricate W core SiC fibers with fine grains and with an outermost carbon coating in one reactor by DC Heating. The gas flow rates and temperature control are both digital and automatic, and the produced SiC fibers have high tensile strengths of >3800 MPa, which is between that of SM1140 and SCS-6 fibers. The Annual output of SiC fiber can be 70 kg (Wang et al., 2016). Our group has also developed W-core SiC fibers by three-stage CVD, and the average tensile strength and Weibull modulus are about 3500 MPa and 16.4 (Luo et al., 2015), respectively, which is comparable to that of SM1140 fiber.

24.4.4

Evolutions of strength and microstructure

It is well known that the properties of fiber-reinforced composites mainly depend on the tensile properties of fibers. During the high-temperature fabrication process and high-temperature working process, the microstructure and surface quality of fibers can vary, which directly influences the mechanical properties of fibers as well as composites. So increasing attention is being paid to the thermal and mechanical stabilities of CVD SiC fibers and there have been many papers in the literature on the subject. There have been primarily two ways to study the thermal and mechanical stabilities of SiC fibers. One is to study the fibers that are directly heat treated under a certain environment, such as in air, inert atmosphere, or vacuum, the other is to study the extracted fibers from the composites after high-temperature heat treatment. The former way is suitable for fibers used in ceramic-matrix composites as generally there is no deleterious interaction between fiber and matrix. The latter way is suitable for fibers used in metal-matrix composites as there usually exists distinct interfacial reactions

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between the fiber and metal-matrix. Severe interfacial reaction would damage the fiber surface, thus the fiber strength would be decreased. At the same time, the metal-matrix can be removed by chemical etching. Many studies have shown that carbon-core SiC fibers (mainly SCS) and W-core SiC fibers have different thermal stabilities due to their different substrates and different microstructures and chemical composition. Carbon-core fibers basically have better thermal stability than W-core SiC fibers as there is no interaction between carbon core and SiC, whilst a W/SiC interfacial reaction would form a brittle reaction layer (W2C and W5Si3) and micropores to reduce tensile strength of SiC fibers (Bunsell, 1988; Luo et al., 2015; Petitcorps et al., 1988). In 1988, Petitcorps et al. reported studies about the thermal and mechanical stability of W-core SiC fiber and carbon-core SiC fiber with or without outer carbon coating. Their results showed that the W-core SiC fibers without a surface coating annealed at 850
C in vacuum had a 25% strength decrease, whilst the carbon-core filaments with a carbon-rich coating annealed at 850
C under vacuum had a slight increase (7%e10%) in strength, even for durations as long as 100 h. Annealing at 750e950
C in the presence of titanium showed that the coated filaments (SCS-2, SCS-6) had superior behavior, as seen in Table 24.7. The strength of the coated filament remained almost unchanged even after annealing as long as 8 h at 850
C. In 1991, Bhatt and Hull reported the microstructural and strength stability of SCS-0, SCS-6, and SCS-9 SiC fibers after 1, 10, and 100 h heat treatment in an Ar environment at temperatures up to 2100
C. Their results show that all the fiber types showed tensile strength degradation beyond 1400
C. And at any temperature above 1200
C, strength decreased with increasing time of exposure. Recrystallization and growth of SiC grains in the outer region appear to be the dominant mechanism for the strength

Residual tensile strength (MPa) of chemical vapor deposition filaments annealed in presence of titanium Table 24.7

Annealing temperature (8C)

Annealing duration (min)

SM1040

SCS-2

SCS-6

3350

3270

3610

30

1900

e

e

60

1500

e

e

30

1360

3180

e

60

e

3210

e

30

1000

3200

3400

60

865

2660

3510

e

3400

Unannealed 750

800

850

950

60

e

Reproduced from Petitcorps YL, Lahaye M, Pailler R, Naslain R: Modern boron and SiC CVD filaments: a comparative study, Compos Sci Technol 32:31e55, 1988 with permission from Elsevier.

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degradation. Nonetheless, the crystal structure SiC remained cubic to 2000
C. At 2200
C, partial transformation of b to a SiC was noticed (Bhatt and Hull, 1991). In 1992, Ning and coworkers also reported similar results of SCS-6 fibers after 2000
C/1 h heat treatment in an argon atmosphere. Their results show that, after annealing, the tensile strength of the fiber decreased to just 0.7 GPa, and the grains in the SiC1 region transformed to equiaxed grains approximately 50e150 nm in size. Unlike the SiC1 region, the microstructures of the SiC2 and SiC3 regions remained essentially the same as that in the untreated fiber due to the segregation of the excess carbon atoms at grain boundaries. Grains in the SiC4 region distinctly coarsened by a factor of 5e20 into equiaxed grains (e.g., coarsening in a range from 250 nm to 3 mm), as shown in Fig. 24.42 (Petitcorps et al., 1988). In 1998, Bhatt and Hull further reported the strength degradation of SCS-6 SiC fibers after 1e400 h heat treatment in 0.1 MPa of argon at temperatures up to 2100
C. Their results show that the fibers heat treated for 1 h above 1400
C and those heat treated for 400 h above 1300
C showed strength degradation. SEM and TEM analyses of the degraded fibers showed formation of a recrystallization region followed by grain growth within SiC4 region and the growth of SiC particles in the outer carbon-rich coating. The tensile strength of the fibers varied as an inverse function of the recrystallized zone thickness. However, the inner SiC grains with carbon-rich boundaries seemed to resist grain growth even at temperatures up to 2100
C (Bhatt and Hull, 1998). Also, in 1998, Smith et al. from the Wright Laboratory Directorate at WrightPatterson Air Force Base reported their detailed research results concerning heat

3 μm

Figure 24.42 Cross-sectional optical micrograph of the 2000
C/1 h annealed SCS-6 fiber. Reprinted from Ning XJ, Pirouz P, Bhatt RT: Effect of high temperature annealing on the microstructure of SCS-6 SiC fibers. Mater Res Soc Symp Proc 250:187e192, 1992 with permission from Cambridge University Press.

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Table 24.8 The tensile strength silicon carbide fibers after heat treatment or extracted from heat treated composites (MPa) Extracted D10858C heat treatment

11608C heat treatment

Extracted D 11608C heat treatment

Fiber type

Asreceived

10858C heat treatment

Trimarc 1

3080

2766

2272

2175

1693

SCS-6

4403

No decrease

No decrease

No decrease

No decrease

SCS-Ultra

5610

A slight decrease

A slight decrease

A slight decrease

A slight decrease

Largediameter SCS-Ultra

6313

5450

<2000

5071

<2000

treatment effects on SiC fibers. The fibers examined included Trimarc 1, SCS-6, SCSUltra, and an experimental large-diameter version of Ultra SCS. The fibers were heat treated under two conditions: 1085
C/2 h þ [email protected]
C/min to 815
C/8h/FC (shortened to 1085
C hereafter) and 1160
C/2 h þ [email protected]
C/min to 815
C/8h/FC (shortened to 1160
C hereafter). The fibers were evaluated for room-temperature tensile strength in the following conditions: (1) as-received; (2) heat treated in vacuum; and (3) consolidated into Tie22Ale23Nb, heat treated in vacuum, and chemically extracted. The tested results are shown in Table 24.8. It can be seen that the Trimarc 1 exhibited significant fiber strength degradation for all heat treatments primarily due to W/SiC reaction and possibly fiber/matrix chemical reaction. The b SiC microstructure in Trimarc 1 appeared to be thermally unstable, as shown in Fig. 24.43. The SCS-6 showed no discernible effect of heat treatment on the tensile strength. Microstructure

(a)

(b)

(c)

10 μm

Figure 24.43 BSE SEM photos of Trimarc 1 b SiC region in: (a) as-consolidated, (b) 1085o C solution HT, and (c) 1160o C solution HT conditions. Reprinted from Smith PR, Gambone ML, Williams DS, Garner DI: Heat treatment effects on SiC fiber, J Mater Sci 33:5855e5872, 1998 with permission from Springer.

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(a)

Handbook of Properties of Textile and Technical Fibres

10 μm

(b)

(c)

Figure 24.44 Secondary SEM images of b SiC portion of SCS-6 for: (a) as-received, (b) 1085o C HT, and (c) 1160o C HT. Reprinted from Smith PR, Gambone ML, Williams DS, Garner DI: Heat treatment effects on SiC fiber, J Mater Sci 33:5855e5872, 1998 with permission from Springer.

observations in Fig. 24.44 shows that it has a very slight thermal instability of the b SiC structure. The SCS Ultra showed only a very modest effect of heat treatment on the tensile strength, too. The microstructure of the b SiC is very fine equiaxed grains with no distinct variation after heat treatment, as seen in Fig. 24.45. The large-diameter (184 mm) SCS-Ultra exhibited the highest as-produced tensile strength. However, this fiber exhibited significant reduction in tensile strength when heat treated, particularly within the matrix alloy. The microstructure of the b SiC in the fiber was the same as that of the SCS Ultra and also had no distinct variations after the heat treatment. The strength debits was suspected to be the result of an unoptimized fiber surface coating, which may be too thin to protect the fiber from chemical reaction with the matrix alloy. In 1999, Lewinsohn et al. studied high-temperature creep and microstructural evolution of SCS-6 fibers. When the fibers were heat treated at 1400
C for 500 h, there was little evidence of actual grain growth throughout the fiber. However, when the SCS-6 fibers were crept at 1400
C for 500 h, the entire inner portion (SiC1, SiC2 and SiC3) transformed to a distinct two-phase region that consists of graphitic carbon segregated between equiaxed b-SiC grains, and the SiC grain size increased by at least five times. Moreover, the grain size of SiC4 region increased to 4e10 mm in width,

(a)

(b)

(c)

10 μm

Figure 24.45 Secondary SEM images of b SiC portion of SCS-Ultra fiber in: (a) as-received, (b) 1085o C solution HT and, (c) 1160o C solution HT conditions. Reprinted from Smith PR, Gambone ML, Williams DS, Garner DI: Heat treatment effects on SiC fiber, J Mater Sci 33:5855e5872, 1998 with permission from Springer.

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Table 24.9 Tensile properties of batches of SCS-6 and SM1140D fibers under different conditions (Vassel and Pautonnier, 2001) SCS-6

SM1140D

Condition

Mean UTS (MPa)

St. dev. (MPa)

Mean UTS (MPa)

St. dev. (MPa)

As-received

4330

350

3170

110

After triode sputtering coating

4280

520

3110

120

After MPa/1 h HIP consolidation

4100

430

2970

320

1000
C/2 h heat treatment

4360

410

2970

80

1000
C/80

HIP, hot isostatic pressing.

which is almost a two-order-of magnitude increase, and the b-SiC phase in region SiC4 transformed to a-SiC polytype. Transformation of the 3C polytype to the 6H polytype is possible by motion of associated Shockley partial dislocation defects (Lewinsohn et al., 1999). Nevertheless, the grain growth and phase transformation is caused by dislocation motion and atomic diffusion induced by the creep stress. In 2001, Vassel and Pautonnier evaluated the property variations of SCS-6 and SM1140þ fibers after each step of the SiC/Ti-22Al-27Nb composite processing. The results are shown in Table 24.9. It can be seen that the tensile strength of SCS-6 fibers basically shows no decrease. However, there is a slight decrease in the SM1140þ fiber strength that can be attributed to the beginning of the W-core/SiC reaction. This study also confirms that the triode sputtering and hot isostatic pressing consolidation have little effect on the tensile strengths of the fibers. Also in 2001, Baik and Grant reported the strength degradation of SM1140þ fibers during manufacture of titanium matrix composites by low-pressure plasma spraying (LPPS) and hot pressing. The fibers were extracted from the titanium matrix composites (TMCs) at various processing stages. The tested tensile strength of the fibers is shown in Table 24.10. The results indicate that the LPPS processing caused a decrease in mean fiber strength and Weibull modulus in comparison with as-produced fibers. A number of fiber surface flaws, primarily in the outer C layer of the fiber, formed as a result of mechanical impact of poorly melted particles from the plasma spray. Coarse feedstock powders promoted an increase in the population of fiber surface flaws, leading to significant reduction in fiber strength. The vacuum hot pressing (VHP) consolidation promoted further development of fiber surface flaws by fiber bending and stress localization because of nonuniform matrix shrinkage, resulting in further degradation in fiber strength. Table 24.11 also shows the tensile strengths of the SM1140 þ and SCS-6 fibers extracted from Ti-6Al-4V composites manufactured by foil-fiber-foil (FFF) and matrix-coated fiber (MCF) methods. It can be seen that, compared with the FFF route, LPPS/VHP led to less degradation of fiber strength. This was mainly attributed to a better fiber distribution and lower porosity level in the TMC performs.

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Table 24.10

Handbook of Properties of Textile and Technical Fibres

Summary of tensile strength of SM1140 D SiC fibers

Fiber

Mean strength (MPa)

Weibull modulus

As-received fiber

3171  135

26.3

Exposure to plasma flame only

3155  146

23.2

Fine powder spraying

3062  183

18.5

Coarse powder spraying

2841  156

19.8

Fine powder spraying

2955  193

16.8

Coarse powder spraying

2809  214

14.4

Fine powder TMC

2882  275

11.5

Coarse powder TMC

2774  354

8.4

LPPS monotapes

LPPS multiplayer preforms

900
C/50 MPa/1h VHP consolidation

LPPS, low-pressure plasma spraying; TMC, titanium matrix composite; VHP, vacuum hot pressing. Reproduce from Baik KH, Grant PS: Strength degradation of SiC fiber during manufacture of titanium matrix composites by plasma spraying and hot pressing, Metall Mater Trans A 32:3133e3142, 2001 with permission from Springer.

Tensile strength of SiC fibers for titanium matrix composite manufacture routes

Table 24.11

Fiber

Manufacturing route

Mean strength (MPa)

Weibull modulus

SM1140þ

FFF As-received

3217  163

23.2

After processing

2604  484

5.5

As-received

3185  125

27.0

After processing

3236  151

22.9

As-received

3157  112

30.1

After processing

2882  275

11.5

As-received

4635  411

11.9

LPPS

4437  393

12.5

VHP

4499  344

14.5

MCF

LPPS/wind

SCS-6

LPPS/wind

FFF, foil-fiber-foil; LPPS, low-pressure plasma spraying; MCF, matrix-coated fiber; VHP, vacuum hot pressing. Reproduce from Baik KH, Grant PS: Strength degradation of SiC fiber during manufacture of titanium matrix composites by plasma spraying and hot pressing, Metall Mater Trans A 32:3133e3142, 2001 with permission from Springer.

Fibers made by chemical vapor deposition

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The MCF route basically had no fiber damage and retained initial fiber strength. This study indicates that MCF is an ideal route for the fabrication of TMCs. In addition, SCS-6 fiber is more resilient to fiber damage than SM1140þ fiber during the LPPS/ VHP, because of differences in fiber diameter and damage tolerance of the fiber coatings (Baik and Grant, 2001). Our group also studied the thermal and mechanical stability of W-core SiC fibers after 900 and 1000
C thermal exposure in a vacuum. The strength variations are shown in Table 24.12. The results show that the 900
C heat treatment had little effect on the fiber strength whilst notable strength degradation occurred after 1000
C thermal exposure. The degradation is due to the growth of W/SiC interfacial reaction as well as the coarsening of SiC grains in the middle region (Luo et al., 2015). The results agree well with Baker’s work who also found out that the 1140þ fiber could retain tensile strength after thermal exposure up to 900
C (Baker, 1999). The increase of the W/ SiC interphase thickness decreases the strength of the filaments that can be attributed to a notch effect of the brittle interfacial zone on the SiC sheath (Faucon et al., 2001). Kinetic analysis of the W-core/SiC reaction permitted a calculation of the maximum exposure time as a function of temperature prior to fiber strength degradation. Results indicate that at 750
C, the expected exposure time before significant strength degradation is less than 10,000 h (Sirenko et al., 1993). However, the as-received strength of SCS-6 fibers can be retained up to about 1100
C after 10,000 h exposure according to the recrystallization kinetics analysis (Bhatt and Hull, 1998).

Effects of various heat treatments on fiber strength of W-core SiC fiber

Table 24.12

T (8C)

Time (h)

Diameter (with C/no C) (mm)

s (with C/no C) (GPa)

Weibull modulus (with C/no C)

Conditions

As-produced

_

110/102

3.5/3.2

16.4/11.9

_

900

1

110/_

3.5/_

16.4/_

Vacuum

900

2

110/_

3.5/_

16.4/_

Vacuum

900

10

110/_

3.2/_

15.2/_

Vacuum

900

50

110/_

3.1/_

13.6/_

Vacuum

900

200

110/_

3.0/_

11.9/_

Vacuum

1000

1

110/102

2.8/2.5

8.7/4.6

Vacuum

1000

2

110/102

2.3/2.0

6.6/3.4

Vacuum

1000

10

110/102

1.6/1.3

7.3/5.9

Vacuum

1000

20

110/_

1.4/_

9.0/_

Vacuum

1000

50

110/_

1.2/_

18.9/_

Vacuum

Reprinted from Luo X, Guo P, Yang Y, Jin N, Liu S, Kou Z, Wu S: Microstructure, tensile strength and thermostability of Wcore SiC fibers with or without carbon coating, Mater Sci Eng A 647:265e276, 2015 with permission from Elsevier

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Handbook of Properties of Textile and Technical Fibres

Protective coating

CVD SiC filament has great strength and stiffness; however, it is brittle and its strength is extremely sensitive to surface conditions, as have been mentioned above. To overcome this, the manufacturers or researchers have applied different kinds of coatings, of which the prevailing one is carbonaceous material, such as the SCS-6 fiber and SM1140 þ fiber, as mentioned above. The carbon coating can heal the rough sawtootheshaped surface of SiC mantle, thus significantly improve the fiber tensile strength. In order to further improve the high-temperature service life of SiC fibers in metalmatrix composites (mainly TMC), many different kinds of coating systems have been tried to add onto the carbon-coated fibers. For example, Sigma SM1240 SiC filament has a 1 mm carbon inner coating and a 1.5 mm TiBx outer coating. The TiBx coating is not a stoichiometric TiB2 compound layer, but a mixture of chemically co-deposited borides, mostly TiB2, and free boron (Guo and Derby, 1998). Boron-rich TiB2 coating led to the formation of needle-like TiB with the matrix (B þ Ti / TiB), but such needle-like products were not observed when stoichiometric TiB2 coating was produced, which can inhibit the fiber-titanium matrix interfacial reaction up to 1100
C (Choy and Derby, 1994). However, the stoichiometric TiB2-coated fiber can only retain 85% of the strength of the virgin SiC fibers (Choy, 2003). TiB2 has also been added onto SCS-6 SiC fibers by magnetron sputtering. The adding of TiB2 also reacted with the matrix to form needle-like TiB (Yang et al., 1999, 2001). This is also due to the existence of free B. Jeng et al. (1993) studied the effect of Ag/Ta bimetal coating on the mechanical behavior of SCS-6/Ti-25-10 composites. The results show that the Ag/Ta coating is an effective diffusion barrier in preventing fiber/matrix interfacial reaction during composite consolidation and can be used as a transition layer to effectively adjust the thermal residual stress, because Ag can prevent the reaction between SCS-6 and Ta, and Ta is a beta phase stabilizer (Chiu et al., 1993). However, the strong interfacial shear strength decreases the toughness of the composite (Jeng et al., 1993; Chiu and Yang, 1995). Upadhyaya studied the effects of rare earth elements and boride Gd/GdBx coating on the interface and properties of SCS-6/Tie6Ale4V and SCS-6/TiAl composite. The results indicated that Gd/GdBx can effectively protect fibers at above 1000
C by acting as a reaction barrier, and the coating can adjust the thermal residual stress, so there were no cracks in the matrix of the two kinds of composites. However, there existed growth cracks in the GdBx coating. In addition, Gd/GdBx is easily oxidized, thus it is difficult to prepare (Upadhyaya, 1995). Choy has prepared a C/TiC/Ti compositionally graded coating system on SiC fibers, which can (1) solve the interfacial problems encountered in SiC/Ti, as well as (2) conserve the strength of the as-received SiC fiber and (3) provide effective protection for SiC fibers in Ti matrices against deleterious interfacial reactions at high temperatures up to 1100
C (Choy, 1996). Majumdar has proposed five considerations for the design of composite interface/ coating: (1) The coating must allow fiber strength retention, i.e., it should not react

Fibers made by chemical vapor deposition

985

with the fiber and thereby reduce fiber strength; (2) the coating must provide adequate bonding strength between the fiber and matrix; (3) the coating should possess some degree of elongation, in order to compensate for higher interface toughness that would accompany higher interfacial bonding strength; (4) the coating should allow for some degree of debonding, so that matrix and coating cracks can be deflected over some distance along the fiber-coating interface; (5) it is desirable that the coating be under residual axial compressive stress, so that stress concentrations from incoming cracks or imperfections have difficulty in penetrating the coating and proceeding on to the fiber (Majumdar, 1999). However, due to the contradiction between the above requirements, it is difficult to meet all the design requirements with single coating at the same time. So double coating or composite coating can be used to optimize the performance of composites. Majumdar designed two coating systems: Ti/Y/Ti and C/SiC. The results of the Ti/Y/Ti coating system showed that the in-situ strength of the fiber decreased sharply and the transverse strength of the composite was lower, which may be due to the oxygen layer of Y layer and the sputtered columnar structure. The C/SiC coating (C coating close to the fiber side, SiC close to the matrix side) significantly improved transverse tension performance but decreased longitudinal tension performance compared with conventional carbon-coated SiC fibers. Although the outer sheath cracked at relatively low stresses, the inner SiC fiber retained the strength of a virgin carbon-coated fiber. Thus, the sheathed system may perform well in situations that require high stiffness at elevated temperatures, but where applied loads are lower (Majumdar, 1999). Guo et al. studied the effects of Cu/Mo and Cu/W double metal coating on the interface and mechanical properties of SCS-6/Ti-15-3 composite. The results showed that the two kinds of double metal coating can delay the interfacial reaction rate (Guo et al., 1999; Guo and Kagawa, 2001), but the Cu/Mo coating is more conducive to improve the axial tensile properties of composites and fatigue properties, as shown in Table 24.13 (Guo and Kagawa, 2001). Our group added a Mo coating of about 1.3 mm in thickness by magnetron sputtering on to carbon-coated SiC fibers, which was used to study the effect of C/Mo duplex coating on the interface and mechanical properties of SiC/Ti

Mechanical properties of duplex metal-coated SCS-6 fiberereinforced Ti-15-3 composites

Table 24.13

Composite

Young’s modulus (GPa)

Tensile strength (MPa)

Strain to failure (%)

Fatigue crack growth rate (m/cycle)

SCS-6/Ti-15-3

128  8

1326  66

1.19  0.05

4.2  108

SCS-6/Cu/Mo/Ti-15-3

131  2

1538  5

1.41  0.03

1.5  108

SCS-6/Cu/W/Ti-15-3

124  2

1321  78

1.23  0.09

2.2  108

Reproduced from Guo S, Kagawa Y: Fatigue behaviour of duplex metal coated SiC fibre reinforced Ti-15-3 matrix composites, Mater Sci Technol 17:1107e1113, 2001 with permission from Taylor & Francis Ltd.

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composites. Our results showed that C/Mo duplex coating can significantly reduce the interfacial reaction and protect the C coating of SiCf/g-TiAl and SiCf/Ti2AlNb composite (Luo et al., 2016, 2017).

24.5

Conclusions

Boron and SiC fibers produced by CVD exhibit unique properties for the reinforcement of metal, polymer, and ceramic matrix composite systems. Their combination of extremely high and consistent mechanical properties with light density has generated composites with outstanding properties. Especially, boron is known as a good shielding material for neutron and cosmic radiation. Therefore, boron fiber is a potential material to address the issue of radiation protection in conjunction with structural requirements in future space systems developed under NASA Space Initiative. Both W-core SiC fiber and carbon-core SiC fiber (such as Sigma SM2156 and SCS-Ultra) can possess extremely high tensile strength once the microstructure is composed of very fine SiC grains and a slight excess of free carbon in the grain boundaries to prevent grain growth. However, due to the unavoidable W/SiC interfacial reaction, W-core SiC fiber is more suitable as reinforcement for low and intermediate temperature composites systems (below 900
C), whilst carbon-core SiC fiber can be used at temperatures up to 1400
C. On the other hand, the development and research of protective coatings would further help fiber strength retention and in increasing the service life of composites.

References Andrews EH: Stress waves and fracture surfaces, J Appl Phys 30(5):740e743, 1959. Astanin VV, Imayeva LA: Two stages of interfacial reaction in B-Al composite, J Mater Sci 29(12):3351e3357, 1994. Astanin VV, Sirenko AA, Zaitsev KB: Use of replicas for studying interaction in composite materials reinforced with boron fibers, Zavod Lab 11:72e73, 1988. Baik KH, Grant PS: Strength degradation of SiC fiber during manufacture of titanium matrix composites by plasma spraying and hot pressing, Metall Mater Trans A 32:3133e3142, 2001. Baker AM (Ph.D. thesis), Oxford, United Kingdom, 1999, University of Oxford. Bhatt RT, Hull DR: Microstructural and strength stability of CVD SiC fibers in argon environment, 1991. NASA Technical Memorandum 103772. Bhatt RT, Hull DR: Strength-degrading mechanisms for chemically-vapor-deposited SCS-6 silicon carbide fibers in an argon environment, J Am Ceram Soc 81(4):957e964, 1998. Bhatti AR, Moss RN, Piller R, Shatwell RA, Slightam R: Processing and characterization of silicon carbide fibre reinforced MoSi2-based composites, Key Eng Mater 164e165: 117e120, 1999. Boggio JV, Vingsbo O: Tensile strength and crack nucleation in boron fibres, J Mater Sci 11: 273e282, 1976a. Boggio JV, Vingsbo O: Application of the Griffith criterion to fracture of boron fibres, J Mater Sci 11:2242e2246, 1976b.

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