Precipitation in AI-Cu-Mg-Ag casting alloy

Taylor, Jason; Parker, B.A.; Polmear, I. J.

doi:10.1179/030634578790433341

Cited by 132 publications

(12 citation statements)

References 2 publications

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“…Suitable heat treatments of AI-Cu-Mg alloys with high Cu:Mg ratios and trace additions of silver can cause precipitates to nucleate and grow on {111} planes as well as {100} planes of the aluminium matrix (Taylor, Parker & Polmear, 1978;Chester & Polmear, 1983), and the occurrence of these thin disc-like precipitates significantly enhances the age-hardening characteristics of such alloys. The composition range within which these {111} precipitates have been found to occur suggests that they should be formed in the commercial AI-Cu-Mg-Ag alloys 201 and Avoir under suitable heat treatments, and indeed castings from such alloys are noted for their marked response to age hardening and for their good tensile properties (Iler, 0108-7681/88/030207-21503.00 1969; Taylor, Parker & Polmear, 1978).…”

Section: Introductionmentioning

confidence: 99%

“…The composition range within which these {111} precipitates have been found to occur suggests that they should be formed in the commercial AI-Cu-Mg-Ag alloys 201 and Avoir under suitable heat treatments, and indeed castings from such alloys are noted for their marked response to age hardening and for their good tensile properties (Iler, 0108-7681/88/030207-21503.00 1969; Taylor, Parker & Polmear, 1978). More recently, Brown Boveri have developed an experimental A1-Cu-Mg-Ag-Mn-Zr-Ti alloy for improved hightemperature strength and creep resistance, and the superior properties of this alloy over 2618 and 2219 arise from a fine dispersion of these {111} precipitates (Kubel, 1986;Polmear, 1986).…”

Section: Introductionmentioning

confidence: 99%

“…Taylor, Parker & Polmear (1978) suggested that the precipitates nucleate as Mg3Ag, before growing and attaining an overall composition close to A12Cu. However, in agreement with the earlier work of Williams (1972), a recent study of precipitation processes in Al-rich A1-Mg-Ag alloys by Cousland & Tate (1986) has failed to isolate Mg3Ag as a precipitating phase in ternary alloys, even in alloys with a high Mg to Ag ratio.…”

Section: Introductionmentioning

confidence: 99%

“…Instead, Cousland & Tate showed that the phases that they observed containing only Ag and Mg atoms were of composition MgAg (their type 1 and type 2 GP zones). This would suggest that the nucleation mechanism for the {111} precipitates suggested by Taylor, Parker & Polmear (1978)is unlikely.…”

Section: Introductionmentioning

confidence: 99%

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The structure of {111} age-hardening precipitates in Al–Cu–Mg–Ag alloys

Knowles¹,

Stobbs²

1988

Acta Crystallogr Sect B

196

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High-resolution transmission electron microscopy (HREM) of { 111 } precipitates in an A1-Cu-Mg-Ag alloy has been used to confirm by direct observation down (110) and (211) A1 matrix zone axes that the structure of these precipitates in peak-and over-aged material is consistent with the monoclinic structure proposed by Auld [Acta Cryst. (1972), A28, $98] of a =b= 4.96, c= 8.48 A, y= 120 ° , rather than the hexagonal structure with a=4.96, c= 7.01 A proposed by Kerry & Scott [Met. Sci. (1984), 18, 289-294]. Reexamination of the monoclinic structure suggested by Auld shows that the structure he proposes is in fact orthorhombic (a=4.96, b=8.59, c= 8.48 A), and is best regarded as a distortion of the structure of tetragonal 0-AI2Cu precipitates found in over-aged A1-Cu alloys. A detailed reanalysis of electron diffraction patterns from this alloy in the light of HREM observations confirms that this structure and the relative thinness of these preciptates perpendicular to the {111} planes can indeed together satisfactorily account for the extra spots and streaks in the patterns.

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Section: Introductionmentioning

confidence: 99%

Section: Introductionmentioning

confidence: 99%

Section: Introductionmentioning

confidence: 99%

Section: Introductionmentioning

confidence: 99%

See 2 more Smart Citations

The structure of {111} age-hardening precipitates in Al–Cu–Mg–Ag alloys

Knowles¹,

Stobbs²

1988

Acta Crystallogr Sect B

196

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“…Addition of trace elements Ag and Mg to Al-Cu/AlCu-Mg alloys may change the precipitation sequence from h/S to V. Taylor et al 114 proposed that Ag and Mg form Mg 3 Ag (possible hexagonal structure with a50 . 487 and c50 .…”

Section: Precursor To V Phasementioning

confidence: 99%

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

Wang¹,

Starink²

2005

International Materials Reviews

775

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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Precipitation in AI-Cu-Mg-Ag casting alloy

Cited by 132 publications

References 2 publications

The structure of {111} age-hardening precipitates in Al–Cu–Mg–Ag alloys

The structure of {111} age-hardening precipitates in Al–Cu–Mg–Ag alloys

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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