Recently, a family of Al-Cu-Li alloys containing minor amounts of Ag, Mg, and Zr and having desirable combinations of strength and toughness were developed. The Weldalite (trademark) alloys exhibit a unique characteristic in that with or without a prior stretch, they obtain significant strength-ductility combinations upon natural and artificial aging. The ultra-high strength (approximately 690 MPa yield strength) in the peak-aged tempers (T6 and T8) were primarily attributed to the extremely fine T(sub 1) (Al2CuLi) or T(sub 1)-type precipitates that occur in these alloys during artificial aging, whereas the significant natural aging response observed is attributed to strengthening from delta prime (Al3Li) and GP zones. In recent work, the aging behavior of an Al-Cu-Li-Ag-Mg alloy without a prior stretch was followed microstructurally from the T4 to the T6 condition. Commercial extrusions, rolled plates, and sheets of Al-Cu-Li alloys are typically subjected to a stretching operation before artificial aging to straighten the extrusions and, more importantly, introduce dislocations to simulate precipitation of strengthening phases such as T(sub 1) by providing relatively low-energy nucleation sites. The goals of this study are to examine the microstructure that evolves during aging of an alloy that was stretch after solution treatment and to compare the observations with those for the unstretched alloy

Brown, S. A.

Kumar, K. S.

Pickens, Joseph R.

English

NASA Technical Reports Server

N91-24407
III. EFFECT OF A PRIOR STRETCH ON THE AGING RESPONSE
OF AN AI-Cu-Li-Ag-Mg-Zr ALLOY
K.S. Kumar, S.A. Brown, and J.R. Pickens
41
https://ntrs.nasa.gov/search.jsp?R=19910015094 2020-03-19T18:24:05+00:00Z
1. INTRODUCTION
Recently, a family of AI-Cu-Li alloys containing minor amounts of
Ag, Mg, and Zr and having desirable combinations of strength and
toughness has been developed (1). These "Weldalite m" alloys exhibit a
unique characteristic in that with or without a prior stretch, they
obtain significant strength-ductility combinations upon natural and
artificial aging. The ultra-high strength (~690 MPa yield strength) in
the peak-aged tempers (T6, T8) has been primarily attributed to the
extremely fine T1 (AI2CuLi) or Tl-type precipitates that occur in these
alloys during artificial aging (2), whereas the significant natural aging
response observed is attributed to strengthening from 5' (AI3Li) and GP
zones. (3) In recent work (4), the aging behavior of an AI-Cu-Li-Ag-Mg
alloy without a prior stretch was followed microstructurally from the
T4 to the T6 condition. Commercial extrusions, rolled plates, and
sheets of AI-Cu-Li alloys are typically subjected to a stretching
operation before artificial aging to straighten the extrusions and, more
importantly, introduce dislocations to stimulate precipitation of
strengthening phases such as T1 by providing relatively low-energy
nucleation sites (5,6).
The goals of this study are to examine the microstructure that
evolves during aging of an alloy that was stretched after solution
treatment and to compare the observations with those in sections I and
II of this report for the unstretched alloy.
2. EXPERIMENTALPROCEDURE
An extruded section of an aluminum alloy containing 5.3Cu-1.4Li-
0.4Ag-0.4Mg-0.17Zr (wt%) and incorporating a 3% stretch was
artificially aged at 160°C (433K) in a circulating air furnace. Prior to
stretching, the extrusions were solutionized at 504°C (777K) for 1 h
and water quenched to room temperature. The quenched and stretched
material was allowed to age naturally at room temperature for >1000
hours to obtain the T3 condition.
The aging response was monitored using Rockwell B hardness
measurements and select specimens along the aging curve [T3, T4
(without stretch), T4 + 15 min, T3 + 15 rain, T3 + 2 h, T3 + 4 h, T3 +
20 h, and T3 + 100 h] were examined using transmission electron
microscopy (TEM). Since the unstretched material in the previous
42
study (4) was aged at 180°C (453K), the stretched material in this
study was also examined after exposure at 180°C (453K) for 15 min to
allow a direct comparison with the unstretched material as well as the
stretched material exposed to 160°C (433K) for the same time.
Specimens to be analyzed by TEM were sliced from the extrusions
and mechanically ground to the desired thickness. These were then
twin jet electropolished to perforation at 243K and 12 15 volts in a
solution of 25% HNO3 in methanol. After polishing, the specimens were
dipped in a solution of 50% HNO3 in H20 to remove any Ag which may
have redeposited from the electrolyte. Selected area diffraction (SAD)
and centered dark field imaging (CDF) were used to follow the
microstructural evolution during aging. Due to space limitations, in
most cases only SAD patterns are provided, although CDF was used in
each instance to confirm SAD observations.
3. RESULTS
The change in hardness with aging time at temperature is shown
in Fig. 1 for both, the stretched and unstretched material. As
mentioned earlier, the stretched material was aged at 160°C (433K),
whereas the unstretched material was aged at 180°C (453K).* Several
interesting features are present in Fig. 1 including a higher hardness
for the T4 than the T3 condition and a reversion for the T3 and T4
conditions upon artificial aging for ...15 45 minutes, although the loss
in hardness is more dramatic for the T4 condition. Subsequent
artificial aging leads to equivalent peak hardness for both the
stretched (T3) and unstretched (T4) material, although the T4 material
ages faster than the T3 material at the temperatures investigated. In
both cases, a very high peak hardness (T6 or T8) of .-.95 - 97 RB is
reached. A sample of the unstretched material was also aged for 15
min at 160°C (433K) to provide a direct comparison of the reversion
behavior with that of the stretched material. In this case, the hardness
was found to be intermediate between that of T4 + 15 min at 180°C
(453K) and T3 + 15 min at 160°C (433K).
* [Strictly speaking, the T3 and T4 conditions, which are naturally
aged conditions, cannot be plotted in Fig. 1; however, for the purpose of
providing a reference, they have been plotted on the hardness axis (T3
and T4 should actually reflect zero time at temperature)].
43
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The SAD patterns using a [110] zone axis for the T4 condition (Fig.
2a), and those after 15-min exposure to 160°C (433K) (Fig. 2b) or
180°C (453K) (Fig. 2c) show that a) the naturally aged microstructure
consists of GP zones and e', b) upon 15 min artificial aging at 160°C
(433K), a significant amount of the 8' dissolves leaving a
microstructure of some 5' and GP zones (Fig. 2b) in the aluminum solid
solution, and c) 15 min artificial aging at the higher temperature of
180°C (453K) leaves a diffuse SAD pattern where <100> streaking due
to the GP zones (Fig. 2c) can still be discerned.
Similarly, SAD patterns for the stretched material in the T3
condition and after aging for 15 minutes at 160°C (433K) and 180°C
(453K) are compared in Figs. 3a-c. In the T3 condition (Fig. 3a), 5'
(AI3Li) spots are readily seen, although the GP zones, while present,
are not as prominent as in the T4 condition (compare <100> streaking in
Fig. 2a with Fig. 3a). Artificial aging at 160°C (433K) for 15 min
results in the dissolution of GP zones and some 5'. In addition, e' is
possibly present as can be seen from the [100] zone axis SAD pattern in
Fig. 3b. The strain and dislocation substructure introduced by the prior
stretching operation make it difficult to obtain and interpret the
diffraction patterns in this reverted condition. A [111] zone axis
pattern of the material exposed to 180°C (453K) for 15 minutes (Fig.
3c) reveals the presence of T1 precipitates, which suggests that this
time at the higher temperature is sufficient to place the stretched
material to the right of the reversion minimum on the aging curve (i.e.,
on the increasing hardness side). A hardness curve at this temperature
for the T3 material is not shown in Fig. 1.
Microstructural evolution in the stretched material aged at 160°C
(433K) for times longer than 15 min (i.e., to the right of the reversion
minimum in Fig. 1) was followed on the transmission electron
microscope using SAD (Figs. 4a-e). After 2 h at 160°C (433K), a [110]
zone axis SAD pattern (Fig. 4a) reveals the presence of all four variants
of the T1 phase and in addition, <100> streaking typical of the e' phase.
The maximum in intensity at the mid point of the <100> streaks derives
from e' but is also coincident with a 5' superlattice reflection. In this
case, the elliptical shape of this intensity maximum leads to the belief
that this is a consequence of the e' phase and not 5'.
A [100] zone axis SAD pattern (Fig. 4b) of the material exposed to
160°C (433K) for 4h reveals the presence of e' in the form of <100>
45
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streaks that intersect and give rise to a maximum in intensity that
coincides with the superlattice reflection that would arise from 6'.
However, the observed intensity is attributed to e', again because of
the shape of the bright spot, which appears as an intersection of two
ellipses rather than a circular spot which is characteristic of the 6'
phase. In addition to the e' phase, symmetrically distributed around the
<100> intersection maximum are the four spots corresponding to the T1
variants. A [112] zone axis SAD pattern (Fig. 4c) of the 4 h specimen
confirms the existence of the T1 phase and also reveals faint streaking
in the <210> directions. This family of streaks is attributed to the
AI2CuMg S' phase. Thus, after 4 h at 160°C (433K), the stretched
material contains e', T1, and the S' phase. It must be specified that
structure factor modifications due to compositional changes within
the precipitate phases can account for modified diffraction spot
intensities and work is currently in progress to measure the
compositions of the individual phases.
Moving further up the aging curve, in the slightly underaged
condition that corresponds to 20 h at 160°C (433K), an SAD pattern
based on a [110] zone axis (Fig. 4d) reveals the presence of T1 and e'.
The [112] zone axis SAD pattern confirms the presence of the S' phase.
In the overaged condition corresponding to a 102 h at 160°C
(433K) (Fig. 4e), the T1 phase is the only major strengthening phase
present; the presence of <100> streaks corresponding to the e' phase in
this condition could not be reliably confirmed and, if present, is in only
small amounts. Similarly, the <210> streaks due to the S' phase were
not reliably confirmed.
A bright-field and a corresponding T1 dark field micrograph in the
slightly underaged condition of 20 h at 160°C (433K) are shown in Figs.
5a and 5b, respectively. Even in this near peak-aged condition, the T1
precipitates are extremely fine and homogeneously distributed.
4. DISCUSSION
The observed higher hardness in the T4 condition compared to the
T3 condition is attributed to the greater amount of GP zones in the
former, although fine e' and GP zones are present in both the T3 and T4
conditions. Likewise, the large drop in hardness upon reverting the T4
material at 180°C (453K) is associated with the dissolution of e' and
5O
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51
some of the GP zones. The decrease in hardness for this material is not
as dramatic when reversion is performed at 160°C (433K), where more
GP zones are still present. A direct comparison of the reverted T4 with
the reverted T3 material (433K) reveals that the hardness of the latter
is higher. The exact reasons for this difference are unclear but must
arise from a complicated interplay of the work-hardening residual from
the stretch, (5' and what appears to be e' precipitates.
In the underaged condition of 2 h at 160°C (433K), the stretched
material contains the T1 and e' phases, which is in agreement with the
recent work of Gayle et al. (4) on the unstretched material. However,
they also found the S' phase (4) which, in this study, was observed
after 4 h of aging the T3 specimens. In addition, Gayle et al. (4) provide
evidence of a new phase, denoted v, but we made no effort to locate it
in this study. After 20 h at 160°C (433K), T1, e' and S' were present,
whereas Langan and Pickens (2) observed only the T1 phase in a similar
alloy after 24 h at 160°C (433K). However, over-aging the stretched
material in the present study at 160°C (433K) for 102 h causes the
dissolution of both e' and a significant amount of the S' phases leading
to a microstructure which is predominantly strengthened by the T1
phase dispersed in an a solid solution. Perhaps a significant amount of
the e' and S' dissolved after 24 h as observed earlier (2), even though
they are present after 20 h at 160°C (433K) in this study.
The observed high strengths in these alloys in the naturally aged
conditions (T3 and T4), are attributed to a combination of the fine 8'
and GP zones. In the artificially aged conditions (T6 and T8) high
strength is primarily a result of the extremely fine T1 precipitates
homogeneously distributed throughout the matrix.
Lithium, magnesium, and zirconium exhibit very high solute-
vacancy binding energies in aluminum (7-9). Consequently, vacancies
present at the solutionizing temperature are retained during quenching
rather than migrating to sinks such as grain boundaries. For the
specific case of binary AI-Li alloys, Suzuki et al. (10) suggest that
excess vacancies bound to lithium atoms are confined to the (5' phase
during aging. Excess vacancies may contribute to the precipitation
process in Weldalite 049-type alloys by two mechanisms. First, when
the (5' present during natural aging dissolves during reversion, the
released vacancies may act as nucleation sites for the plate-like T1
phase. Second, vacancies that were bound to Mg and Zr, as well as
52
released vacancies that are now likely bound to Li atoms, are
unavailable to assist lattice diffusion, thereby retarding growth of T1.
Consequently, the vacancies are effectively utilized for enhancing
nucleation and discouraging growth, which is consistent with the fine
distribution of T1 observed even in the overaged condition. The effects
of such trace elemental additions on precipitation have been
demonstrated previously (11, 12)in other aluminum alloys. In addition,
the role of these trace elements such as Ag and Mg in influencing
precipitate-matrix interfacial energies and/or the possibility of these
elements segregating to the interface and acting as diffusional
barriers cannot be ignored as both these factors can significantly
affect precipitation kinetics.
Prior to significant (5' precipitation after quenching, the
stretching operation introduces a significant number of dislocations
which may act as vacancy sinks by sweeping vacancies away, thereby
decreasing the vacancy concentration available for influencing the
natural aging response. However, dislocations are also effective
nucleation sites of the T1 precipitates. That is, in the stretched
material, heterogeneous nucleation on dislocations likely occurs in
addition to the somewhat reduced nucleation frequency at vacancies in
a way that approximately equals the nucleation at vacancies in the
unstretched material.
5. CONCLUSIONS
The microstructural evolution during artificial aging of an AI-Cu-
Li-Ag-Mg-Zr alloy has been characterized using TEM both with and
without prior cold work. In the stretched and near-peak aged condition,
a fine homogeneous distribution of T1, e' and S' phases is present in an
a solid solution matrix, but on over-aging, virtually all of the e' and
most of the S' dissolve leaving behind a microstructure of essentially
T1 precipitates in an a aluminum solid solution matrix. It is suggested
that trapped-in excess vacancies play a significant role in refining the
microstructure by enhancing precipitate nucleation and discouraging
their growth.
53
7. REFERENCES
[1] J.R. Pickens, F.H. Heubaum, T.J. Langan, and L.S. Kramer, in
"Aluminum-Lithium Alloys" (Proceedings of the Fifth
International Aluminum-Lithium Conference), T.H. Sanders and
E.A. Starke, eds., MCE Publications Ltd., Birmingham, U.K., 1989,
p. 1397.
[2] T.J. Langan and J.R. Pickens, ibid., p. 691.
[3] F.W. Gayle, F.H. Heubaum, and J.R. Pickens, op. cit., p. 701.
[4] F.W. Gayle, F.H. Heubaum, and J.R. Pickens, Scripta Metall., 24,
1990, p. 79.
[5] R.J. Rioja, P.E. Bretz, R.R. Sawtell, W.H. Hunt, and E.A. Ludwiczak,
in "Aluminum Alloys, Their Physical and Mechanical Properties,"
Vol. III, Proceedings of the International Conference on Aluminum
Alloys, Chameleon Press, London,U.K., 1986, p. 1781.
[6] W.A. Cassada, G.J. Shiflet, and E.A. Starke, Jr., J. de Physique,
colloq. C3, Vol. 48, 1987, p. 397.
[7] M. Ahmad and T. Ericsson, Scripta Metall., 19, 1985, p. 457.
[8] S. Ozbilen and H.M. Flower, Acta Metall., 37, 1989, p. 2993.
[9] S. Ceresara, A. Giarada, and A. Sanchez, Philos. Mag., 35, 1977,
p. 97.
[10] H. Suzuki, M. Kanno, and N. Hayashi, J. Jpn. Inst. Light Metals, 31,
1981, p. 122.
[11] B.C. Muddle and I.J. Polmear, Acta Metall., 37, 1989, p. 777.
[12] W.X. Feng, F.S. Lin, and E.A. Starke, Jr., Metall. Trans. A, 15, 1984,
p. 1209.
54


Effect of a prior stretch on the aging response of an Al-Cu-Li-Ag-Mg-Zr alloy

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