Within the zone is ample evidence of a subsequent P' crystal dissolution into the glass matrix. The P' volume fraction is notably reduced, producing a large continuous glass matrix. This occurs by the general g' morphology alteration, with the loss of sharp faceted edges, which must have dissolved due to their associated high chemical
potential. More complex P' dissolution also occurs, which is
characterised by liquid intrusion formation on crystal facets. This is probably initiated by local liquid composition variations (as 8'
solubility in liquid is enhanced by Y ^ + concentrations). Once within the P', the dissolution front appears to preferentially proceed along the prism's c axis, dissolve the core, and leave a skeletal shell similar to the oxide g - Y2Si207 grains (Fig 6.14c). This consistently
observed 'coring' has several explanations. (1) A radial compositional gradient or distribution of fine inclusions exist in the p' grain.
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developed during its initial growth. However, S.T.E.M. micro-analysis shows no conclusive variation in Si or Al nor the presence of impurity elements, although subtle variations may be sufficient. (2) More likely, preferred dissolution is due to chemical potential differences associated with various crystallographic planes. The critical
observation of the dissolution front, within the P' grain, progressing parallel to the facet planes to produce a hexagonal shaped glass volume
(Fig 6.14d), suggests preferred dissolution along (1010) type planes. These planes have the highest linear density and largest interplanar spacing (~ 6.6 A Jin the 0' hexagonal structure and therefore the weakest interplanar bonding.
Closer to the oxide, the degree of 0' dissolution is observed to increase. The driving force for the reaction and the variability are readily explained by the extent of oxygen indiffusion. In the 3-D prism representation (Fig 5.4b) increasing only eq. % 0 takes the overall 0' + glass composition (c) closer to the eutectic glass composition. Applying the phase rule to the 0' - glass tie-line dictates an increase in glass volume. However, 0* dissolution at these oxidation temperatures is only possible because of the high Al^ and Y^ + levels in the liquid conferred by the YAG reversion. Only limited dissolution is observed in oxidised as-sintered glass matrix materials due to their rapid Y, Al outdiffusion and matrix crystallisation.
6.4.3 Kinetics of YAG Reversion
The kinetics of the reverted-YAG zone growth show a dependence on material type at 1300°C, which diminishes at 1350°C (Fig 6.15). An accurate rate-limiting mechanism cannot be determined from zone dimensions due to the transient nature of the oxide. Non-protective oxide shell formation both rapidly increases the zone, indicating the dependence of surface oxygen potential, and decreases it via the rapid
7 0
S u b - o x i d e l a y e r t h i c k n e s s
(fim)
Figure 6.15 Kinetics of sub-oxide (reverted YAG) zone at 1300°C and 1350 °C.
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conversion of the reverted-YAG zone to an oxide scale. The complete YAG reversion and linear front requires the maintenance of an SiC^-rich phase at the reaction interface and is facilitated by 3 mechanisms.
(1) The 0' dissolution, with consequential glass enrichment of Si, Al, 0, N. (2) The subsequent Y, Al outdiffusion leaving 0 2” and Si *** ions. (3) The 0 2- indiffusion to react with the S i ^ - r i c h reaction zone glass. The driving force for indiffusion is simply chemical equilibration between the oxide and g.b. glass. Enhanced 0
indiffusion (and therefore zone kinetics) would be expected if rapid diffusion channels of residual g.b. glass were present. There is direct correlation between degree of matrix crystallisation and zone kinetics, which explains the 1300°C variability and suggest 0 2_ indiffusion as the rate-limiting process. At 1350°C, the increased reaction rate lessens the importance of residual glass content, with the YAG volume and its continuity becoming critical (Chapter 9).
6.4.4 Microetxuctxiral Stability in the Absence of Oxidation Reactions
In the centre of a large specimen, where any oxidation effect is limited, the microstructure is still observed to be unstable. From supplementary experiments in Argon atmosphere, for varying temperature and alloys,the YAG is found to revert to liquid above 1525°C and this temperature defines the solidus for the 8' (z “ 0.75) and YAG binary. However, at sub-solidus temperatures above 1475°C, an important YAG morphology change is observed where the semi-continuous matrix progresses to a more isolated granular form (Fig 6.16a) with a corresponding reduction in surface area. This YAG "morphological stabilisation" is characterised by 3 interrelated features. (1) Primarily, the development of symmetric~120° 0' - 8* - YAG grain junctions, which dictates (2) curved p' YAG interfaces which results ultimately in (3) the loss of the p' faceted morphology and narrow
Figure 6.16. a) T.E.M. section of 'bulk' P + YAG alloy annealed at 1475°C (with no oxidation interaction) showing 'isolated' YAG morphology. b) Schematic of route to isolated YAG via
crystallisation of glass matrix and a reduction in surface energy anisotropy.
/
Figure 6.16. a) T.E.M. section of 'bulk' 0 + YAG alloy annealed at 1475°C (with no oxidation interaction) showing 'isolated' YAG morphology. b) Schematic of route to isolated YAG via
crystallisation of glass matrix and a reduction in surface energy anisotropy.
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interconnected "YAG" channels. The driving force for the
transformation is the loss of surface energy anisotropy associated with the initial solld/liquid interface (of the as-sintered microstructure), and the replacement with approximately isotropic solid/solid interfaces between p'/B' and B' /YAG (Fig 6.16b). The B'/YAG g.b. movement will not be by atomic jumping across the interface (due to the lack of B'/YAG solid solution), tut by transport along the g.b. with atomic removal and redeposition, producing the characteristic 'S' shaped boundaries. Morphological stabilisation is an important
microstructural modification. In combatting the oxidation degradation mechanisms above 1300°C this loss of an interconnected YAG path through the microstructure will be shown (Chapter 9) to significantly reduce oxidation and reverted-YAG zone kinetics.
6.5 SUB-OXIDE MICROSTKPCTURAL STABILITY IN ff t GLASS MATERIALS 6.5.1 As-sintered Materials
In materials SO and S13, the sub-oxide zone contrast in b.s.e. mode is explained by out-diffusion of heavy elements through the matrix to the oxide driven by the SiOj g.b. glass diffusion couple. The kinetics of the zone's development appear controlled by the parabolic rate law (Fig 6.17a,b) and have an associated activation energy. Zone growth rate is therefore diffusion controlled, with the same mechanism operative over the temperature range studied. The rate limiting diffusive species is therefore concluded to be Y 3*". The initial high, non-parabolic, zone growth is probably due to the rapid oxide crystallisation to 8 - Y2S which tends to "draw” Y ^ + ions to the oxide.