Precipitation in Cu-Ti and Cu-Ti -Al Alloys; Discontinuous and Localised Precipitation
T. K. V A I D Y A N A T H A N
Department of Biological Materials, Northwestern University, 311 E. Chicago Ave., Chicago, Illinois (U.S.A.)
K. M U K H E R J E E
Department of Physical and Engineering Metallurgy, Polytechnic Institute of New York, 333 Jay Street, Brooklyn, New York (U.S.A.)
(Received in revised fo rm F e b r u a r y 20, 1976)
Discontinuous precipitation in a Cu-4% Ti binary alloy is studied through transmission electron microscopy. Different types of contrast conditions are used and moire fringes and interfacial dislocations are revealed. The grain refinement that occurs during discontinuous precipitate growth is due to translational and rotational subgrain bound- aries formed on account of the different habit planes of precipitates from the different nucleation sites. It is also shown that discontinuous precipitation is absent in a Cu-2.1% Ti-2 .4%Al and a Cu-2.1% Ti-5% A1 alloys over a wide range of aging temperatures. However, a pronounced tendency for localised precipitation is observed in these alloys. It is found that the ternary Cu-Ti-A1 alloys have lower coherent solvus temperatures. Solute supersaturation is therefore relieved through precipitation of the equilibrium phase by heterogenous nucleation at grain boundaries, twin boundaries, dislocations and other structural imperfections where a lower activation energy barrier is present.
1. I N T R O D U C T I O N
It is known that discontinuous precipitation or recrystallisation is primarily responsible for overaging in Cu-Ti alloys [1, 2] . Recently, Korotayev e t al. [3] made an interesting proposit ion that A1 addition causes suppres- sion of discontinuous precipitation in Cu-Ti alloys. This conclusion is based on resistivity and optical metallographic studies; transmis- sion electron microscopy is probably the most elegant method of confirming this suggestion.
Since ternary additions are known to have a similar effect in other alloy systems such as Cu-Be [4] , the finer details of such effect of ternary additions on discontinuous precipita- tion should be of fundamental metallurgical interest.
2. E X P E R I M E N T A L
The alloys were prepared from 99.999% pure Cu and A1 and 99.97% pure iodide crystal grade Ti supplied by the United Mineral and Foote Mineral Corporations, respectively. Melting was carried out in an MRC arc melter using nonconsumable tungsten electrode after evacuating the fur- nace chamber and backfilling with argon to provide an argon atmosphere during melting. A titanium getter was used during the melting. To ensure sufficient homogenisation, melting was repeated at least three times. The bu t ton so obtained was again melted and cast into a rectangular fiat ingot, homogenised at 890°C for 40 hours in a vacuum furnace under a vacuum of 10 -~ Tort. Spectrographic analysis on typical samples of the nominal Cu-4% Ti, Cu-2.1% Ti-2.4% A1 and the Cu-2.1% Ti-5% A1 alloys indicated that the Ti content of the alloys were 3.99%, 2.03% and 2.14% respec- tively and the A1 contents of the ternary alloys were 2.35% and 4.90% respectively. The oxygen content of the alloys was less than 50 p.p.m. The ingots were cold rolled into 4 mil thick sheets. All aging was carried out with specimens encapsulated in evacuated quartz tubing (10-6 Tort). The sheet samples were solution treated at 890 °C for 40 minutes
144
(a)
ii
(c)
Fig. 1. BF transmission electron micrographs of aged Cu-4% Ti binary alloy. (a) Alloy aged 15 minutes at 500 °C. Incipient discontinuous precipitation at the
(b)
grain boundaries. (b) Alloy aged 30 hours at 450 °C. Discontinuous precipitation at the grain interior. (c) Alloy aged for 30 hours at 450 °C. Observe the incipient discontinuous precipitation at the grain boundaries and the absence of it at the coherent twin boundary.
and aged isothermally. The samples were electropolished using the window technique with an electrolyte of 2 parts methanol and one part nitric acid by volume at a temper- ature of less than --40 °C. The cathodes were michrome wires wound into a helix. The samples were examined in an electron micro- scope (Philips EM 200) operated at 100 kV.
3. EXPERIMENTAL RESULTS
Figure 1 shows three transmission electron micrographs illustrating the important features of incipient discontinuous precipita- tion in the Cu-4% Ti alloy on aging isother- mally at different temperatures. These micro- structures indicate that discontinuous preci- pitation starts predominantly at grain bound- aries and to a limited extent at sites within the grains. Coherent boundaries such as the twin boundary in Fig. 1(c) do not favour
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Fig. 2. BF transmission electron micrograph of the Cu-4% Ti alloy aged 41/2 minutes at 650 °C. Note the irregular initial precipitate development.
~i~iiiiii ̧ ii!i~iiiiiii,
Fig. 3. BF transmission electron micrograph of the Cu-4% Ti alloy aged at 500 °C for 25 hours. Comple- tion of discontinuous precipitation.
nucleation of these precipitates. These obser- vations are in agreement with those in other alloy systems [5]. Initial nucleation is often irregular as in Fig. 2 and propagation occurs by the motion of the precipitation front into the grain interior until the entire grain is consumed by the advancing front (Fig. 3). Precipitates form as plates on the { 111) habit planes in general. The spacing between the plates is characteristic of the reaction temper- atures and is typically 800 A at 450 °C, 1500 A a 500 °C and 3000 A at 700 °C. A closer lamellar spacing at a lower aging temper- ature is to be expected because of the increas- ed solute supersaturation at the lower temper-
A
I ~ / j \ \ ] / j /
Fig. 4. Model for grain refinement during the discontinuous precipitate growth
atures. Although the precipitates have a (111~ habit in general, the precipitates originating from a single site such as a particular grain boundary segment or an internal structural imperfection are uniquely associated with only one of the family of the (111} planes. An important consequence of the nucleation feature described above is the grain ref inement which occurs during precipitate propagation. Thus precipitates originating from different grain boundary segments with different mis- orientations and grain interior sites of a single grain may have different habits so that impingement of advancing fronts from dif- ferent nucleation sites will result in subgrain boundaries in a single grain of the original matrix. Since the precipitation is generally initiated at incoherent interface boundaries, it
146
Fig. 5. BF transmission electron micrograph of the Cu-4% Ti alloy aged for 14,500 minutes at 450 °C. Subgrain boundaries formed during discontinuous precipitate growth.
is to be expected that the selection of the precipitate habit may be determined by the misorientation between adjacent grains at dif- ferent grain boundary segments. The nature of the various subgrain boundaries formed may be different. In Fig. 4 several grain boundary segments surrounding a single grain are shown. Although the nuclei forming at segments A and A' may have the same precipitate habit, there is no continuity of the growing plates at XX' because the plates nucleated at A and A' have a translatory displacement relative to each other. Thus translational subgrain bound- aries may develop. On the other hand, the plates originating at B and B' may have dif- ferent habits altogether, say (111) and (111); a rotational subgrain boundary may develop at YY'. Figure 5 shows subgrain boundaries formed at XX' and YY' probably by such translational and rotational effects. Since the subgrain boundaries develop simply through change of precipitate habit, the above model can satisfactorily explain why solution treat- ment of the aged sample results in restoring the original grain size [1, 6].
(a)
(b)
Fig. 6. Different types of contrast of discontinuous precipitates in a Cu-4% Ti alloy. BF transmission electron micrographs. (a) Alloy aged for 30 hours at 450 °C. Precipitates revealed through moire fringe contrast. (b) Alloy aged for 25 hours at 500 °C. Dis- location contrast in one of the grains.
. ] u
147
(c) (d)
(e)
Fig. 6. (c) Alloy aged for 25 hours at 500 °C. Dis- placement fringe contrast. (d) Alloy aged for 25 hours at 500 °C. Composite contrast involving displacement fringes and moire fringes. (e) Precipitate contrast from strong precipitate diffraction and weak matrix diffraction.
4. CONTRAST OF PRECIPITATES
Incipient discontinuous precipitates in Fig. l (a) were probably revealed through selective thinning of recrystallised regions. However, depending on imaging conditions, other contrast effects were observed. Figure 6 shows a series of micrographs illustrat~ ing moire fringe contrast (Fig. 6(a)), inter- facial dislocations (Fig. 6(b)), displacement fringes (Fig. 6(c)), composi te contrast in which moire and displacement type fringes are superimposed (Fig. 6(d)) and precipitate contrast (Fig. 6(e)). The displacement fringes are reversed in contrast in a dark field using the operating matrix reflection (compare Fig. 6(c) with Fig. 7(a)). On the other hand, the contrast reversal occurs in the dark field using a precipitate reflection when imaging is done through precipitate contrast (compare Fig. 6(e) with Fig. 7(b)). In the case of imaging through precipitate contrast, no contrast reversal occurs in a dark field using a matrix reflection (compare Fig. 6(e) with Fig. 7(c)). In addition, precipitate contrast conditions often showed contrast reversals of precipitate and matrix regions across bend contours as in Fig. 7(d) owing to change of
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(a) (b)
(c) (d)
Fig. 7. Some features of contrast behavior of discontinuous precipitates under varying imaging conditions. Transmission electron micrographs. (a) DF from matrix reflection. Imaging through displacement fringes. Note the contrast reversal of the fringes in the DF compared with those in the BF of Fig. 6(c). (b) DF from precipitate reflection. Imaging through precipitate contrast. Observe the contrast reversal in the DF compared with the BF of Fig. 6(e). (c) DF from the matrix reflection. Imaging through precipitate contrast. Observe that there is no contrast reversal in the DF compared with BF in Fig. 6(e). (d) BF imaging through precipitate contrast. Observe the contrast reversal of precipitates and matrix in two parts of the micrograph due to a bend contour and consequent change in the diffraction conditions.
149
(a) toJ
Fig. 8(a) Selected area diffraction pattern of discontinuous precipitates from a region in Fig. 6(e). (b) Indexing of the SAD in Fig. 8(a).
1 . 0 u
(a) (b)
Fig. 9. BF transmission electron micrographs of Cu-2.1% Ti-2.4% A1 alloy. (a) Coherent precipitates at the grain interior and localised precipitates at the grain boundaries. Alloy aged for 1 hour at 500 °C. (b) Localised precipitat ion at the grain boundaries. Alloy aged at 600 °C for 20 hours.
150
(a) (b)
Fig. 10. BF transmission electron micrographs of the Cu-2.1% Ti-5% A1 alloy. (a) Coherent precipitat ion at the grain interior and localised precipitat ion at the grain boundaries. Alloy aged at 400 °C for 25 hours. (b) Localised precipitation at the grain boundaries. Alloy aged at 600 °C for 1 hour.
diffraction conditions from strong precipitate diffraction to strong matrix diffraction across the bend contour. In this way, the different contrast effects could be readily separated.
5. CRYSTAL STRUCTURE OF THE PRECIPITATE
PHASE
A selected area diffraction pattern of the recrystallised structure is shown in Fig. 8(a) and the indexing in Fig. 8(b) is based on an orthorhombic structure with a = 2.572 A, b = 4.503 A and c = 4.313 A [7]. On this basis the orientation relationship becomes
(111)M \\ ( 0 0 1 ) p
(211)M \\ (010)p.
6. INFLUENCE OF AI ADDITION
The Cu-2.1% Ti-2.4% A1 and the Cu-2.1% Ti-5% A1 ternary alloys do not decompose through discontinuous precipitation. Since the
discontinuous precipitation generally starts at the grain boundaries, the influence of A1 addition on discontinuous precipitation is most readily demonstrated at the grain boundary regions. Thus Figs. 9(a), (b) and 10(a), (b) show that equilibrium precipitates form at the grain boundaries at the early stages of aging in both the Cu-2.1% Ti-2.4% A1 and the Cu-2.1% Ti-5% A1 alloys over a wide range of aging temperatures. The metastable coherent precipitates may or may not form at the grain interior, depending on whether the aging temperature is below or above the coherent solvus temperature of the particular alloy. It was observed that the Cu-2.1% Ti-5% A1 alloy had the lowest coherent solvus temperature (~450 °C} followed by ~620 °C for the Cu-2.1% Ti- 2.4% A1 alloy and ~720 °C for the Cu-4% Ti alloy. Because of this low coherent solvus temperature, the Cu-2.1% Ti-5% A1 alloy shows ap ronounced tendency for localised precipitation of the equilibrium phase heterogeneously at the grain boundaries and other structural imperfections (see Fig. 10(b)
151
(a) (b)
and Fig. 11). Both d i scon t inuous prec ip i ta t ion and localised prec ip i ta t ion occur th rough he te rogeneous nuc lea t ion at grain boundar ies and o the r s t ructura l imperfec t ions . However , when localised prec ip i ta t ion of the equilib- r ium phase is energet ical ly favourable in the t e rna ry alloys, the available sites for heteroge- neous nuc lea t ion are t aken up by the localised t y p e o f prec ip i ta t ion , thus suppressing the possibil i ty of d i scon t inuous prec ip i ta t ion . The absence of d i scon t inuous prec ip i ta t ion and the marked t e n d e n c y for localised precipita- t ion in the Cu -T i -A 1 te rna ry al loys s tudied in this invest igat ion can be rat ional ised in this way.
(c)
Fig. 11. BF transmission electron microg~aphs of the Cu-2.1% Ti-5% Al alloy illustrating the localised precipitation and precipitate distribution. (a) Precipitation at the twin boundaries. Alloy aged for 18 hours at 600 °C. (b) Precipitation at dislocations. Alloy aged at 500 °C for 70 hours. (c) Precipitation at subgrain boundaries. Alloy aged at 500 °C for 4 hours.
7. CONCLUSIONS
Discon t inuous prec ip i ta t ion in the Cu -4 % Ti al loy can be revealed th rough d i f f e ren t con t ras t condi t ions . Th e prec ip i ta tes have an o r t h o r h o m b i c crystal s t ruc ture wi th a = 2 .572 A, b = 4 .503 A and c = 4 .313 A. Some grain r e f i n e m e n t occurs as a resul t o f the
152
difference in the habit planes of precipitates from different nucleation Sites within a single grain of the original matrix. There is no dis- continuous precipitation in the Cu-2.1% Ti-2.4% A1 alloy and the Cu-2.1% Ti-5% A1 alloy. These ternary alloys however show a marked tendency for localised precipitation through heterogeneous nucleation at grain boundaries and other structural imperfections.
ACKNOWLEDGEMENT
This research was supported by the NIH training grant DE-189.
REFERENCES
1 H. T. Michels, I. B. Cadoff and E. Levine, Met. Trans., 3 (1972) 667.
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3 A. D. Korotayev, A. T. Protasov. O. V. Tsinenko and M. V. Lyubchenko, Fiz. Metal. i Metalloved., 27 (1969) 127.
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5 J. W. Christian, The Theory of Transformations in Metals and Alloys, Pergamon Press, p. 608.
6 H. T. Michels, Ph.D. Thesis, New York University, 1971.
