The decomposition of SSSS during ageing is usually a complex process. The equilibrium precipitate, θ, normally does not form directly from the SSSS at commonly used ageing temperatures, because the nucleation barrier for its formation is too high (i.e., large energy is needed because of the surface energy required to create the surface of the critical-sized precipitate). The precipitation occurs in steps involving several transition (metastable) precipitates before the equilibrium-precipitate forms.
The microstructural changes which occur during the initial stages of ageing are very small sized and cannot be resolved by light microscope, yet it is in these early stages that the most profound changes occur in properties. Both, X-ray diffraction and electron microscopy are used to study the ageing process, and have shown that in virtually all age-hardening systems, the initial precipitate is not of the same structure as the equilibrium precipitate.
As an illustrative example, the following is the sequence of precipitates if an Al-4.5% Cu alloy is aged after obtaining SSSS by quenching from 550°C. This alloy system exhibits the greatest number of intermediate stages in its precipitation at low temperature of ageing.
GP Zones à θ” (CP zone 2) à θ’ à θ (CuAl2)
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Here, three distinct and identifiable precipitates form and dissolve prior to the formation of the equilibrium-precipitate (θ). This sequence occurs because the equilibrium-precipitate is incoherent with matrix, whereas the transition precipitates are either fully coherent (GP Zones, (θ”), or at least semi-coherent, i.e., the formation of a particular precipitate at a stage occurs because of the favourable nucleation conditions.
For example, the nuclei of an incoherent type of precipitate must exceed a certain minimum critical size before it can nucleate, but a coherent-precipitate instead can form faster for which the size effect is relatively unimportant (because surface energy requirements are negligible).
This is what causes the nucleation barrier to be too high for equilibrium-precipitate to form. Because of the importance of surface energy and strain energy, a system follows such a sequence so that it has lowest free-energy in all stages of precipitation.
Table 13.3 illustrates precipitation sequence in several other alloy systems:
GP Zones:
Guinier-Preston zones, earlier called GP1 zones-the early stage of ageing, are named thus, after the men who first independently studied their formation by X-ray diffraction. GP zones are plate-like clusters predominantly of copper atoms segregated on to (100) planes of the aluminium lattice (Fig. 13.3). The plates have a diameter of about 100°A and thickness of only 3-6 A°. These zones appear to form uniformly and homogeneously throughout the aluminium lattice with a density of around 1018 per cm3.
GP zones have same crystal lattice as of aluminium. As the atomic diameter of copper is less than that of aluminium, there occurs appreciable elastic-straining as the local changes in the interplant distance occurs as illustrated schematically in Fig. 13.3 and are called coherency-strains. GP zones, although also called pre-precipitates, are actually coherent-precipitates.
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They are seen to be formed quite rapidly upon quenching, probably part of the clusters form during quenching. GP zone formation occurs by diffusion of copper atoms aided by the quenched-in vacancies over relatively short distances.
GP zones are responsible for the first hardness peak shown in Fig. 13.6 (a) for alloys containing 4.0 and 4.5% copper aged at 130°C. As the elastic-fields, resulting from coherency, extend into the aluminium lattice, the effective size of the zone, in impeding the dislocation motion, is much greater than its actual physical size.
θ” Precipitate:
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It was earlier called as GP-2 zone, but since it has a definite but different crystal structure than matrix, it is more appropriate to call it a coherent-intermediate-precipitate with a symbol θ”. The overall composition is CuAl2. θ” precipitate is in plate form of maximum thickness 100 A° and up to a maximum diameter of 1500 A°. It has tetragonal (Fig. 13.7 b) crystal structure with a = 4.04° and C = 7.68 A°, i.e., it fits well with the aluminium unit cell in two directions but not along ‘C’ axis.
Thus, aluminium planes parallel to the plates are distorted by elastic coherency strains. Optimum hardness is attained in aluminium with 3 to 4.5% copper when aged at 130°C, for θ” precipitates are both more numerous and produce greater distortion than any other transition structure. θ” precipitate has ordered arrangements of copper and aluminium atoms (Fig. 13.7b).
θ‘ Precipitate:
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This transition precipitate is large enough to be resolved under the optical microscope. It has a tetragonal structure with a = 4.04 A° and C = 5.8 A°, (Fig. 13.7 c). The composition is slightly different than θ’. The disc-shaped θ’ precipitates are semi-coherent. The elastic strain around these precipitates is small as the long range strain fields of its dislocations and the precipitate largely cancel. Thus, the formation of θ’ structure leads to the softening of the alloy. θ’ precipitate form heterogeneously.
θ Precipitate:
The equilibrium precipitate, θ (CuAl2) has tetragonal (Fig. 13.7 d) structure (a = 6.07 A°, C = 4.87 A°). It is fully incoherent precipitate and thus, its formation always leads to softening as coherency strains disappear. It nucleates heterogeneously and is more easily formed while ageing at higher temperatures. θ precipitates are the ultimate result of over ageing.
Ageing Al-Cu alloys at 130°C as illustrated in Fig. 13.6 (a) shows that hardness increases with the formation of GP zones and the intermediate precipitates. The maximum hardness in an alloy is obtained when there is a critical-dispersion of GP zones, or an intermediate precipitate (θ” or θ’), or both. After the peak hardness, further ageing tends to decrease the hardness (Fig. 13.6). This phenomenon is called over- ageing. During overageing, the precipitate particles coarsen at the cost of neighbouring small precipitates.
Thus, during coarsening:
(i) The average precipitate size increases
(ii) The total number of precipitates decreases
(iii) The inter-precipitate distance increases
The presence of two peaks in Fig. 13.6 (a) for the alloys having 4.0% Cu and 4.5% Cu, indicates two different precipitates responsible for peaks respectively one after the other, i.e., precipitation occurs in stages. Two-stage hardening occurs at low ageing temperature (here 130°C) and results in higher peak hardness. Single stage-hardening occurs at higher ageing temperatures (here 190°C), or even at lower ageing temperatures for alloys having lesser solute contents (3% Cu, or 2% Cu in aluminium at 130°C) and results in lower peak hardness.
In alloys with 2% and 3% Cu, precipitates of θ’ phase are responsible for peak hardness when aged at 190°C. The increase in hardness up to the peak is due to an increase in the density as well as the size of the precipitates, but during overageing, only coarsening of the θ’ precipitates occurs, resulting in increasing the distance between particles.
Fig. 13.6 (b) also illustrates that the peak hardness in alloys with 4.0% Cu and with 4.5% Cu when aged at 190°C is due to θ” and θ’ precipitates. θ” induces more coherency strains than semi-coherent θ’ precipitates. θ’ precipitates are more difficult to be shared by dislocations. Roughly, at peak hardness, ratio is θ”: θ’:: 7: 3. Over ageing here is partly due to replacement of θ” precipitates by θ’ as well as by the coarsening of 0′ precipitates.
At low temperature of ageing, over ageing stage may not be attained, such as in naturally age-hardenable Duralumin. The hardening occurs due to the increased density of coherent precipitates (as well as some coarsening of them).
The peak-hardness in ageing-curve increases as the amount of solute contents increases in the alloy. The maximum amount of solute content could be a little less than the ultimate solubility limit point indicated by the solvus line in the phase diagram as it should be possible to heat the alloy to obtain a single-phase solid solution without causing burning of the alloy, i.e., the alloy should have the maximum super-saturation after quenching (with no burning).
Fig. 13.6 illustrates that by increasing the copper content from 2.0% to 4.5%, the peak hardness says by ageing at 190°C increases. Not only is the as-quenched hardness more, there occurs greater increase in hardness on ageing (due to larger volumes of precipitates formed).
Higher super-saturation also causes faster precipitation, and thus, the peak in the ageing curve is attained in shorter ageing time as illustrated by Fig. 13.6 (b). Also important is the rule that the degree of super-saturation decreases with the increase of ageing temperature resulting in lower peak hardness at higher temperatures.
Addition of trace elements (from thousandths to tenths of a percent), or controlling the impurities content can be effectively used as means to control the ageing process. For example, trace additions of Cd. In, Sn, Be retard the formation and growth of GP zones in Al-Cu alloys, and thus, retard hardening on ageing.
Actually, these trace elements form strong bonds with the quenched-in vacancies. This pinning of vacancies docs not leave enough vacancies for the diffusion of say copper in Al-Cu alloys and retards the formation of GP zones.
Addition of 0.1% Cd in AI-4% Cu alloy causes its segregation between precipitate and matrix, thereby reducing the interfacial energy of θ’ from 1530 to 250 ergs/cm2. This reduces the critical nuclei of precipitate leading to increased density of the precipitates. The reduction in interfacial energy also causes the decrease of rate of coarsening of θ’ precipitates by a factor of five in Al-Cu alloys. Such reduction in energy is used in high temperature age-hardenable alloys to resist softening.
Trace addition of cadmium retards formation of GP zones and thus retards hardening in natural ageing, but trace cadmium also increases the density of θ’ precipitate, i.e., enhances hardening in artificial ageing. Presence of iron in Al-Cu system removes a part of copper to form insoluble Al-Cu-Fe compound. Age- hardening effect gets reduced as lesser volume of precipitates form.
Effect of Composition in Binary System:
Fig. 13.8 illustrates how the increase in solute content B increases the hardness. Theoretically, the age- hardening effect is maximum in alloy C0 (maximum solid solubility limit at eutectic temperature), but it is impossible for such an alloy to be single phase at high temperature without melting (burning). As solution temperature is kept somewhat below the solidus point, the maximum super-saturation is obtained in an alloy having slightly lesser solute content than C0.
But if the solute content becomes more than C0 in the alloy, the amount of θ (equilibrium precipitate) remaining undissolved at the same solution temperature, increases, i.e., the amount of supersaturated α decreases. As the age-hardening occurs due to the precipitation from the α-solution, and as the amount of α-phase decreases, the hardness decreases as illustrated by line OP in Fig. 13.8.
GP Zones Solvus:
GP zones solvus is illustrated in Fig. 13.9. It defines the upper temperature limit of stability of the GP zones for different compositions, although the concentration of excess quenched-in vacancies can change it slightly (raise it with increase in concentration). It is the temperature for an alloy above which if ageing is started, then GP zones do not form.
Also if an alloy had been age-hardened at low temperature such as by natural ageing, if is reheated to above its solvus temperature (~ 250°C) for a short time, say 20-60 seconds, and then rapidly cooled, the alloy becomes soft. Its properties return to the values existed in the as-quenched state (before ageing).
This process of softening is called reversion or retrogression. Reversion is essentially the dissolution of GP zones which had formed during natural-ageing, by heating it above the GP zone solvus. Here, the time of holding at that temperature is quite important.
With a shorter holding time, not all the GP zones have sufficient time to be dissolved and softening may be incomplete. On the other hand, a prolonged holding time can cause hardening due to the precipitation of other transition precipitates characteristic of that temperature.
Reversion process is reversible if these new precipitates are not allowed to form at the temperature of reheating. Reversion process has been used to an advantage if the plasticity of the alloy must be restored before bending, flanging, riveting, etc.
Reversion does decrease the corrosion resistance of the alloys. Reversion of other transition precipitates also occurs depending on the temperature of heating. Solvus temperatures of other transition phases are illustrated in Fig. 13.9 for Al-Cu system.
The ageing temperature of an age hardened alloy has to be higher than the service temperature of a component to avoid the overageing and its effects during service. Supersonic-aircraft Concorde attains skin temperature of 100-110°C while flying and thus, the ageing temperature for the aluminium alloy of the body of the aircraft is 190°C.
Kinetics of Precipitation:
The entire process of precipitation is very complex as it depends on a large number of factors such as temperature and time of ageing, nature of alloy, the composition of alloy, amount of impurities, trace elements, the temperature of solutionising, rate of cooling alter solutionising, plastic deformation before, or after quenching, time and temperature of holding before artificial ageing, etc. and precipitation takes place through many stages with combinations of processes. It is difficult to make quantitative derivation.
However, Newkirk has made qualitative generalisations:
1. The rate of precipitation is faster at higher temperature of ageing.
2. The rate of precipitation is faster in alloys of widely dissimilar metals.
3. Impurities in soluble or in insoluble state invariably increase the rate of precipitation.
4. Plastic deformation just before ageing increases the rate of precipitation.
5. The rate of precipitation is faster at a given ageing-temperature in a low-melting alloy than in a high- melting alloy.