Role of Dopants

2.4.1. Role of Dopants

Due to the high amount of covalent bonding (~ 90%) in silicon carbide, it is not possible to obtain high densities with a ‘pure material’ by using a normal sintering process. It has been observed that the powder compact of very fine pure crystalline solids does not densify without any applied pressure. The reason for this lack of sintering characteristic involves kinetic as well as thermodynamic considerations. The ‘hot pressing’ technique with the addition of small amounts of boron, aluminium, nickel among others, is well known, but it is limited to the production of rather simple shapes, as found out by Alliegro et al [21] and Kriegesman [22]. The ‘hot isostatic’ pressing technique may be useful for the fabrication of more complicated geometric shapes.

The pressureless sintering of silicon carbide has been reported by Prochazka [23] who synthe- sized β-silicon carbide in the submicron range, which was sintered after the addition of 0.5 wt% boron and 1 wt% carbon at 2050°C to 2150°C in order to obtain densities of 85 to 96% of the theoretical density in the argon atmosphere. Three years later, Coppola and Mcmurtry [24] achieved the sintering of more readily accessible α-silicon carbide powder in a mixture containing boron and carbon additives.

However, in both boron doped sintered α- and β-silicon carbides, the exaggerated grain growth was observed by Johnson and Prochazka [25], if the sintering temperature exceeded a critical limit. The microstructure of boron doped SiC, sintered at higher than 2075°C, usually contains large tabular grains, which may act as ‘stress concentrators’, which seriously degrade the strength properties.

In its pure form, SiC powder will not sinter to a fully dense state. By heating at the temperature range of 1900—2300°C with pressures ranging from 100 - 400 MPa, Kriegsman et al [22] were able to make ‘pore-free’ dense particles of polycrystalline SiC, starting from green bodies of 50% green density without the use of any sintering aids. Prochazka [23] proposed that during the sintering of pure submicron powders of covalently bonded solids, the densification is prevented by a hypothetically high ratio of

‘grain-boundary’ to ‘solid-vapour’ surface energies, i.e. γ GB / γ SV (see the section 2.2.3). In order to get a pore ‘surrounded by three grains’ shrink to closure, the equilibrium dihedral angle ( θ) must be > 60° or γ GB / γ SV < 3. The equation relating these ‘important sintering’ parameters that is applicable at solid- vapour interface is written as :

The following argument was advanced by Prochaszka [23] in order to explain the role of boron and carbon :

1. As γ GB for SiC is expected to be high due to strong ‘directionality’ of Si-C bonds, γ GB / γ SV ratio is also too high, and the sintering is thus inhibited,

2. Boron segregates selectively along the grain-boundaries decreasing γ GB and consequently γ GB / γ SV ratio, and

3. Since γ SV for SiC is high, its surface has a tendency to adsorb impurities and eventually lower γ SV , thus increasing γ GB / γ SV ratio. Both silicon and silica which lower γ SV for SiC by this process should be removed. The carbon removes silica and silicon by forming SiC and CO.

NANO MATERIALS

On the other hand, Lange and Gupta [26] used mixtures containing β-SiC with different amounts of B and C or B 4 C. After sintering of these mixtures, their microstructural observations strongly sug- gested that a ‘boron-rich’ liquid was present during sintering. From these observations, they concluded that either ‘reaction’ sintering or ‘liquid-phase’ sintering is responsible for the densification of submicron silicon crbide with boron and carbon additions.

But, the solubility and lattice position of B in SiC are contradictory. Shaffer [27] reported on the solubility of 0.1 wt% B in SiC at 2500°C. By contrast, Murata and Smoak [28] found a solubility of 0.5 wt% at 2200°C. The lattice parameter measurements upon incorporation of B suggests that boron sub- stitutes for Si or C. A comparison of these data with the theoretical calculations by Tajima and Kingery [29] suggested that the aforementioned lattice contraction is less than that wold be expected if all the boron atoms replaced Si. The considerations of ‘strain energy’ and the ‘bonding stability’ favour the incorporation at C and Si sites simultaneously. Thus, it seems that boron may occupy both sites simul- taneously.

The Scanning Transmission Electron Microscope (STEM), Auger Electron Spectroscopy (AES), Wavelength-Dispersive X-ray Analysis (WDXRA) and Micro-Radiography (MRG) studies were con- ducted by Hamminger et al [30] on several α-SiC separately doped with Al and C or B and C, which showed a homogeneous region of ≈ 200 µm diameter containing B, O, N, Al, and/or C-enrichments.

Let us take a critical view on these findings. The pressureless sintering of silicon carbide has been reported by Prochazka [23] as explained above, and also by Prochazka and Scanlan [31], based on the grain-boundary and solid-vapour surface energies. This theory is inconsistent with the experimental measurements of ‘dihedral angles’ in partially sintered silicon carbide powder compacts, as reported by Greskovich et. al. [32]. Moreover, Prochazka and Scanlan [31] did not find the presence of any boron in the grain-boundaries through their experiment of Neutron Activation Auto-Radiography (NAAR), which contradicted Prochazka’s earlier theory that boron reduces the grain-boundary energies through absorp- tion at the grain-boundaries, and enhances the diffusion process in silicon carbide.

As also mentioned above that Lange and Gupta [26] predicted a liquid-phase sintering. The extensive study of Si-B-C phase equilibrium by Kieffer et. al. [33] shows that only the SiC - SiB 6 - Si compatibility triangle contains liquidus at temperatures < 2100°C. The lowest eutectic temperature in this compatibility triangle is 1380°C. Lange and Gupta [26] also found that Si + B 4 C powder composi- tions react to form SiC with traces of SiB 6 , as observed by the microscopic examination.

The ‘free silicon’, which is a necessary ingredient for both sintering mechanisms, is available at high temperature as a result of the decomposition of SiC. Coppola and Smoak [34] achieved the sintering of the more readily accessible α-silicon carbide powder in a mixture containing boron and carbon addi- tives. An explanation of the success of the densification, with the addition of boron nitride, boron phos- phate and boron carbide, was provided by Murata and Smoak [28].

Contrary to the view of Prochazka [23] as mentioned above, Suzuki and Hase [35] observed that both boron and carbon effectively inhibited the grain growth, but they did not cause sufficient densification, when used individually, indicating that the surface diffusion was not completed. By using high-resolu- tion microscopy, these workers found that at above 1900°C, a second phase with no more than 80 nm in size, that influences the material movement was formed. It was considered that boron dissolves very little into silicon carbide and that a large amount of silicon dissolves in this inter-granular phase. They did not observe a B-Si-C liquid phase formation. Above 1950°C, this inter-granular secondary phase dissolves in the matrix up to 0.2 to 0.3 %, and then distributed widely in the material.

Quite interestingly, these observations have been confirmed by Davis et al. [36]. Subsequently, Hamminger et al. [30] also investigated the fracture surfaces of B-C and Al-C doped silicon carbide to

85 show the formation of solid solution, which also showed a heterogeneous region of @ 200 mm diameter

SILICON CARBIDE

containing B, O, N, Al, and/or C enrichments. Tajima and Kingery [29] found by using lattice constant evaluations that aluminium entered into the ‘solid solution’ in β-silicon carbide, occupying the Si sites. Tajima and Kingery [37] also observed that there is some segregation of aluminium in the grain- boundaries.

The above reports show that the sintering of silicon carbide is not an easy affaire. Moreover, there is quite a difference of opinion among various important workers on the mechanism of sintering in the atomistic level. The sintering of SiC needs to be properly understood so that its various applications can be tailored at ease. It has already been mentioned above that the route to ‘nano particles sintering’ is

a better way to have an adequate knowledge on the sintering mechanism, after dealing properly with the microstructure of the sintered products with ‘nano-crystalline’ SiC. This amply justifies our intention of going to the “nano route” of material processing for an important material like SiC for various useful applications.

The studies on the sinterability of nano-crystalline α-silicon carbide with the addition of alu- minium nitride and boron carbide are presented here with their effect on the microstructure and on the mechanism of sintering (see later in the section 2.7.3). The next step is to understand the precise ‘role of carbon’ in the sintering process.