Accomodation of the misfit strain surrounding {III} precipitates (Ω) in AlCuMg(Ag)

Fonda, R. W.; Cassada, W. A.; Shiflet, G. J.

doi:10.1016/0956-7151(92)90323-7

Cited by 85 publications

(62 citation statements)

References 13 publications

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“…zones and '' and ' phases. A HRTEM observation of the broad faces of  disks revealed the presence of small coherent ledges [9]. Figures 5 (a) and (b) show the coherent ledges observed on the face of a  precipitate and its schematic illustration.…”

Section: Sadp Corresponding To (A)mentioning

confidence: 99%

“…According to Fonda et al [9], to reduce the overall misfit strain produced by the coherent ledges, the coalescence of coherent ledges takes place, resulting in a taller ledge containing misfit dislocations, which is called the misfit-compensating ledge. As  precipitates grow, such misfit compensating-ledges were observed on the  precipitates, as exemplified in Fig.…”

Section: Sadp Corresponding To (A)mentioning

confidence: 99%

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Precipitation Process in a Cu-Ni-Be Alloy

Watanabe

Monzen

2011

SSP

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Abstract. The precipitation process in an aged Cu-1.9wt%Ni-0.3wt%Be alloy has been examined by high-resolution transmission electron microscopy. The precipitation sequence found is: GunierPreston (G.P.) zones → γ'' → γ' → stable γ. The disk-shaped G.P. zones and the disk-shaped γ'', γ' and γ precipitated phases are composed of monolayers of Be atoms on {100} α of the Cu matrix and alternative Be and Ni matrix layers parallel to {100} α . The γ'' phases consisting of two to eight Be-layers has a body-centered tetragonal (bct) lattice with a=b=0.24 nm and c=0.28 nm. The γ' or γ phase is bct with a=b=0.24 nm and c=0.26 nm or a=b=0.26 nm and c=0.27 nm. The γ'', γ' or γ phase aligns with the matrix according to the Bain orientation relationship. The growth kinetics of disk-shaped γ precipitates on aging at 500 o C has been also investigated. The {001} α habit planes of the γ precipitates migrate by a ledge mechanism. The average thickness of the γ disks increases with aging time t as t

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Section: Sadp Corresponding To (A)mentioning

confidence: 99%

Section: Sadp Corresponding To (A)mentioning

confidence: 99%

Precipitation Process in a Cu-Ni-Be Alloy

Watanabe

Monzen

2011

SSP

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“…28. Figure 29a and b shows the diffraction patterns of V phase in [112] Al by Fonda et al 106 and Kerry and Scott. 103 In fact, Fig.…”

Section: Phasementioning

confidence: 99%

“…the c value is 0 . 8576 nm for V after aging at 375uC for 1 h. 106 Second, in recent HREM findings, Yoshimura…”

Section: Phasementioning

confidence: 99%

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Precipitates and intermetallic phases in precipitation hardening Al–Cu–Mg–(Li) based alloys

Wang¹,

Starink²

2005

International Materials Reviews

797

276

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. Although many suggestions for the existence of pre-precipitates in Al-Cu-Mg alloys with a Cu/Mg atomic ratio close to 1 have been made, a critical review reveals that evidence exists for only two truly distinct ones. The precipitation sequence is best represented as: supersaturated solid solutionRco-clustersRGPB2/S"RS where clusters are predominantly Cu-Mg co-clusters (also termed GPB or GPB I zones), GPB2/S" is an orthorhombic phase that is coherent with the matrix (probable composition Al 10 Cu 3 Mg 3 ) for which both the term GPB2 and S" have been used, and S phase is the equilibrium Al 2 CuMg phase. GPB2/S" can co-exist with S phase before the completion of the formation of S phase. It is further mostly accepted that the crystal structure of S' (Al 2 CuMg) is identical to the equilibrium S phase (Al 2 CuMg). The Perlitz and Westgren model for S phase is viewed to be the most accepted structure. 3DAP analysis showed that Cu-Mg clusters form within a short time of natural and artificial aging. Cu-Mg clusters and S phase contribute to the first and second stage hardening during aging. In Al-Cu alloys, the h phase (Al 2 Cu) has I4/ mcm structure with a50 . 607 nm and c50 . 487 nm, and h' phase with tetragonal structure and a50 . 404 nm, c50 . 58 nm, the space group is I4 m2. Gerold's model for h" (or GPII) appears to be favourable in terms of free energy, and is consistent with most experimental data. The transformation from GPI to GPII (or h") seems continuous, and as Cu atoms will not tend to cluster together or cluster with vacancies, the precipitation sequence can thus be captured as: supersaturated solid solutionRh" (Al 3 Cu)Rh' (Al 2 Cu)Rh (Al 2 Cu). The V phase (Al 2 Cu) has been variously proposed as monoclinic, orthorhombic, hexagonal and tetragonal distorted h phase structures. It has been shown that V phase forms initially on {111} Al with c50 . 935 nm and on further aging, the c lattice parameter changes continuously to 0 . 848 nm, to become identical to the orthorhombic structure proposed by Knowles and Stobbs (a50 . 496 nm, b50 . 858 nm and c50 . 848 nm). Other models are either wrong (for example, monoclinic and hexagonal) or refer to a transition phase (for example, the Garg and Howe model with c50 . 858 in a converted orthorhombic structure). For Al-Li-Cu-Mg alloys, the L1 2 ordered metastable d' (Al 3 Li) phase has been observed by many researchers. The Huang and Ardell model for T 1 phase (space group P6/ mmm, a50 . 496 nm and c50 . 935 nm), appears more likely than other proposed structures. Other proposed structures are perhaps due to the T 1 phase forming by the dissociation of Kn110m dislocations into 1/6n211m Shockley partials bounding a region of intrinsic stacking fault, in which copper and lithium enrichment of the fault produces a thin layer of the T 1 phase.

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The Role of Ledge Nucleation/Migration in Ω Plate Thickening Behaviour in Al‐Cu‐Mg‐Ag Alloys

Hutchinson

Fan

Pennycook

et al. 2001

Lightweight Alloys for Aerospace Application

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The thickening kinetics of Q plates in an Al-4Cu-0.3Mg-0.2Ag (wt. %) alloy have been measured at 200 OC, 250 "C and 300 "C using conventional transmission electron microscopy techniques. At all temperatures examined the thickening showed a linear dependence on time. At 200 "C the plates remained less than 6nm in thickness after lOOOh exposure. At temperatures above 200 "C the thickening kinetics are greatly increased. Atomic resolution Z-contrast microscopy has been used to examine the structure and chemistry of the (OOl), 11 ( 11 l)c( interphase boundary in samples treated at each temperature. In all cases, two atomic layers of Ag and Mg segregation were found at the broad face of the plate. The risers of the growth ledges and the ends of the plates were free of segregation. No significant levels of Ag or Mg were detected inside the plate at any time. The necessary redistribution of Ag and Mg accompanying a migrating thickening ledge occurs at all temperatures and is not considered to play a decisive role in the excellent coarsening resistance exhibited by the S2 plates at temperatures up to 200 "C. Plates transformed at 200 'C rarely contained ledges and usually exhibited a strong vacancy misfit normal to the plate. A large increase in ledge density was observed on plates transformed at 300 'C, concomitant with accelerated plate thickening kinetics. The high resistance to plate coarsening exhibited by $2 plates at temperatures up to 200 'C, is due to limited ledge nucleation under these conditions. The prohibitively high barrier to coherent ledge nucleation on the broad faces of plates aged at 200 'C arises from the contribution to the total free energy change attending nucleation from elastic interactions between the misfitting coherent ledges and the significant strain field that can exist normal to the broad face of the B plate.

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Accomodation of the misfit strain surrounding {III} precipitates (Ω) in AlCuMg(Ag)

Cited by 85 publications

References 13 publications

Precipitation Process in a Cu-Ni-Be Alloy

Precipitation Process in a Cu-Ni-Be Alloy

Precipitates and intermetallic phases in precipitation hardening Al–Cu–Mg–(Li) based alloys

The Role of Ledge Nucleation/Migration in Ω Plate Thickening Behaviour in Al‐Cu‐Mg‐Ag Alloys

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