Nucleation kinetics of the α→γM massive transformation in a Ti-47.5 at.% Al alloy

Veeraraghavan, D.; Wang, P.; Vasudevan, Vijay K.

doi:10.1016/s1359-6454(02)00572-4

Cited by 56 publications

(26 citation statements)

References 51 publications

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“…Based on the reported observations [12,22], massive g grains are assumed to nucleate heterogeneously at a/a grain boundaries. The volumetric nucleation rate is calculated with the following expression coming from classical nucleation theory: Fig.…”

Section: Nucleation Of Massive Gmentioning

confidence: 99%

“…The nucleation rate is calculated with the following expression: (12) where J 0 g L is a prefactor expressing the lineal density of potential nucleation sites and s a is the surface fraction of grain boundaries uncovered by transformation products. The driving pressure of the…”

Section: Nucleation Of G Lamellaementioning

confidence: 99%

“…Analytical expressions for nucleation and growth rates of the massive grains [11][12][13] have been established. However, these expressions are not sufficient to calculate the overall kinetics of the phase transformation, which requires accounting for additional aspects such as grain impingement and the decrease of nucleation rates due to site coverage.…”

Section: Introductionmentioning

confidence: 99%

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A numerical model for the description of the lamellar and massive phase transformations in TiAl alloys

Rostamian

Jacot

2008

Intermetallics

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a b s t r a c tA phenomenological modelling approach has been developed to describe the massive transformation and the formation of lamellar microstructures during cooling in binary g titanium aluminides. The modelling approach is based on a combination of nucleation and growth laws which take into account the specific mechanisms of each phase transformation. Nucleation of massive and lamellar g is described with classical nucleation theory, accounting for the fact that nuclei are formed predominantly at a/ a grain boundaries. Growth of the massive g grains is based on theory for interface-controlled reactions. A modified Zener model is used to calculate the thickening rate of the g lamellar precipitates. The model incorporates the effect of particle impingement and coverage of the nucleation sites by the growing phase. The driving pressures of the phase transformations are obtained from Thermo-Calc based on the actual temperature and matrix composition. CCT diagrams and lamellar spacings calculated with the model are in good agreement with experimental data obtained from dedicated heat treatment experiments and from the literature. The model permitted investigating the influence of cooling rate, alloy chemistry and average a grain size upon the amount of massive g and the average thickness and spacing of the lamellae. In particular it indicates that the Al depletion of the a phase during lamellar precipitation seems to play an important role in the suppression of the massive transformation at moderate cooling rate and in the large lamellar spacings observed at low cooling rate.

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Section: Nucleation Of Massive Gmentioning

confidence: 99%

Section: Nucleation Of G Lamellaementioning

confidence: 99%

Section: Introductionmentioning

confidence: 99%

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A numerical model for the description of the lamellar and massive phase transformations in TiAl alloys

Rostamian

Jacot

2008

Intermetallics

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“…Usually at least one of these interfaces would have had a higher interfacial energy than plane-to-plane interfaces simply because a much larger proportion of the area of the latter interfaces is fully coherent. Hence, the sum of the interfacial energies of the two edge-to-edge interfaces might roughly approximate the sum of the energy of a coherent and an incoherent interface, as in the nucleus model of Figure 23(1)(c) [28] used by Veeraraghavan et al [148] to rationalize their nucleation rate data. The combination of the absence of independent information on the energies of ␣:␥ m (or even ␣:␥) boundaries in Ti-Al alloys and uncertainties as to the critical nucleus shape operative (particularly now that the option of E2E-LES boundaries at many boundary orientations [38] has become available) prevents any further discrimination from being made at present between these models on the basis of comparisons of nucleation theory with experiment.…”

Section: B E2e-les Vs Plane-to-plane Interfacial Structures and Tramentioning

confidence: 97%

“…Veeraraghavan et al [148] have measured the nucleation kinetics of ␥ m in a Ti-47.5 at. pct Al alloy at undercoolings of 231 °C, 284 °C, and 307 °C below the T o temperature, using stereological techniques (rather than the easy-to-use but far less reliable approach of back-calculation from the Johnson-Mehl-Avrami-Kolmogorov [149] [JMAK] equation [150] ).…”

Section: B E2e-les Vs Plane-to-plane Interfacial Structures and Tramentioning

confidence: 99%

Influence of crystallography and bonding on the structure and migration of irrational interphase boundaries

Aaronson

2006

Metall Mater Trans A

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Interphase boundary structure developed during precipitation from solid solution and during massive transformations is considered in diverse alloy systems in the presence of differences in stacking sequence across interphase boundaries. Linear misfit compensating defects, including misfit dislocations, structural disconnections, and misfit disconnections, are present over a wide range of crystallographies when both phases have metallic bonding. Misfit dislocations have also been observed when both phases have covalent bonding (e.g., US:␤ US 2 by Sole and van der Walt). These defects are also found when one phase is ionic and the other is metallic (Nb:Al 2 O 3 by Rühle et al.), albeit when the latter is formed by vapor deposition. However, when bonding is metallic in one phase but significantly covalent in the other, the structure of the interphase boundary appears to depend upon the strength of the covalent bonding relative to that in the metallically bonded phase. When this difference is large, growth can take place as if it were occurring at a free surface, resulting in orientation relationships that are irrational and conjugate habit planes that are ill matched (e.g., ZrN:␣ Zr-N by Li et al. and Xe(solid):Al-Xe by Kishida and Yamaguchi). At lower levels of bonding directionality and strength, crystallography is again irrational, but now edge-to-edge-based low-energy structures can replace linear misfit compensating defects (␥ m :TiAl:␣ Ti-Al by Reynolds et al.). In the perhaps still smaller difference case of Widmanstätten cementite precipitated from austenite, one orientation relationship yields plates with linear misfit compensating defects at their broad faces whereas another (presumably nucleated at different types of site) produces laths with poorly defined shapes and interfacial structures. Hence, Hume-Rothery-type bonding considerations can markedly affect interphase boundary structure and thus the mechanisms, kinetics, and morphology of growth.

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Phase Transformations and Microstructures

Appel

Paul

Oehring

2011

Gamma Titanium Aluminide Alloys

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Nucleation kinetics of the α→γM massive transformation in a Ti-47.5 at.% Al alloy

Cited by 56 publications

References 51 publications

A numerical model for the description of the lamellar and massive phase transformations in TiAl alloys

A numerical model for the description of the lamellar and massive phase transformations in TiAl alloys

Influence of crystallography and bonding on the structure and migration of irrational interphase boundaries

Phase Transformations and Microstructures

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