In this article we will discuss about:- 1. Introduction to Age-Hardening Treatment 2. Requirements of Age-Hardening Treatment 3. Steps 4. Classification of Precipitates 5. Duplex Ageing 6. Barriers.
Contents:
- Introduction to Age-Hardening Treatment of Metals
- Requirements of Age-Hardening Treatment of Metals
- Steps in Age-Hardening Treatment of Metals
- Classification of Precipitates in Age-Hardening Treatment of Metals
- Duplex Ageing Treatment of Metals
- Barriers of Age-Hardening Treatment of Metals
1. Introduction to Age-Hardening Treatment:
In 1906, Alfred Wilm, a German engineer, discovered by accident, the phenomenon of natural-ageing. He found that a quenched alloy of aluminium with copper and magnesium-called Duralumin, increased its hardness with ageing (time). This is called age-hardening (or precipitation hardening), because the hardness of the quenched alloy increases as a function of ageing time. The effectiveness of age-hardening in an aluminium alloy with 4.5% copper, as the yield strength of the alloy becomes more than four times after age-hardening.
It is this mechanism that has made possible the use of light-weight aluminium alloys, called Duralumin as air-craft-frame-materials since the Second World War. Copper- beryllium alloys may be made as strong as some tool steels by similar heat treatment to be used as non- sparking tools in coal-mines. The effects of various treatments on the properties of copper-2% beryllium alloy. Wilm could not observe (optical) microstructural changes during the strengthening period. It remained a mystery until the development of electron microscope in 1950s, when the nature of precipitation became physically clear. Age-hardening is now the base of heat-treatment of a large number of non-ferrous as well as ferrous-alloys.
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2. Requirements of Age-Hardening Treatment:
Two basic requirements are:
1. The main basic requirement of a precipitation-hardening alloy system is that the solid solubility limit should decrease with the decrease in temperature, i.e., the phase diagram as illustrated in Fig. 13.1. The alloys susceptible to precipitation-hardening are those which can form supersaturated-solid-solutions, and then reject finely-dispersed-precipitates when aged at room, or intermediate temperatures.
The aluminium-copper system, the aluminium-rich portion of which is illustrated in Fig. 13.1, is a typical example of a precipitation-hardening system. Here, such as an alloy X (4.5% Cu) exists as a homogeneous α-solid solution at high temperatures (say at T2), hut on cooling becomes saturated (at T3) with respect to the second phase, θ (CuAl2).
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If this alloy is cooled slowly further to room temperature, coarse precipitates of θ are formed, and mostly at the grain boundaries of a as illustrated by the schematic illustration in Fig. 13.1 b(iii), and this does not cause any improvement in mechanical properties as shown is table 13.1.
If this alloy is quenched from temperature T2, then the high temperature phase, α-solid solution is retained at room temperature because the rapid cooling suppresses the separation of θ, and as no time is available for the diffusion to occur to bring about the compositional changes.
At room temperature, thus, this alloy is present as supersaturated (metastable) solid solution (SSSS) as illustrated in Fig. 13.1 b (II). If this alloy is aged for a sufficient length of time at room temperature, or at slightly higher temperature, line precipitation occurs inside the grains to cause age-hardening as illustrated in Fig. 13.1 b (IV).
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The super-saturation is the driving force for the rejection of the excess solute in the form of the precipitates. Due to limited diffusion rates at these low temperatures, the solute atoms move through only a few tens of the interatomic distances (helped by the quenched-in vacancies), giving rise to extremely fine precipitation.
This is called ageing. The precipitation occurs by the nucleation and growth process. The fluctuations in the solute concentration provide small clusters of solute atoms in the crystal lattice of solvent (here aluminum), which act like nuclei for the precipitation. The growth rate of these nuclei is controlled by the rate of atomic migration, so that precipitation increases with the increasing ageing temperature.
However, the size of the precipitates becomes finer as the ageing temperature at which precipitation occurs is lowered. Extensive hardening of the alloy (as illustrated in table 13.1) is associated with the critical-dispersion of the precipitates.
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2. The precipitates of the second phase should be coherent in nature. As the precipitate size increases, partial and later complete loss of coherency at the interface occurs, when finally the equilibrium-precipitate θ forms, but by then over ageing had already begun. Alloy systems Al-Mn, Al-Mg, Mg-Pb show decrease of solid solubility with decrease of temperature, but coherent-precipitates are not formed. Alloys of such systems cannot be age-hardened.
3. Steps in Age-Hardening Treatment:
Controlled precipitation from a supersaturated solid solution hardens the alloys, while its critical dispersion causes extensive hardening.
The following are the steps in age-hardening treatment after choosing a proper composition of the alloy:
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1. Solutionizing:
It is the process of heating the alloy just above the solvus temperature (Ti) to obtain a single phase solid solution, say a. The alloy should not be heated above solidus temperature as melting and oxidation shall occur to cause adverse effect on ductility. Heating the alloy much above the solvus temperature such as Ti in Fig. 13.1 (a) is also not advised as it causes grain-growth in the alloy, and the refinement in grain size later is a difficult process as no phase-change occurs during heating in these alloys.
Many problems are faced while effectively solutionizing the Al-Cu-Mg alloys within a few degrees of the solidus temperature. If coring is present in the alloy, then the solutionizing temperature close to the solidus causes ‘burning’, i.e., melting and oxidation caused at the grain boundaries seriously decrease the ductility of the alloy.
2. Quenching:
The solutionised alloy is cooled fast to retain the high temperature single-phase solid solution at room temperature as metastable supersaturated solid solution (SSSS). As this super- saturation is normally required even in the centre of the section of a part, cold, hot or boiling water or even air cooling may be used. As cold water gives maximum super-saturation, it may be used for thick sections.
3. Ageing:
It is the process of controlled decomposition of SSSS to form finely-dispersed-precipitates usually at one and sometimes at two intermediate temperatures for a suitable time period. A. Wilm’s natural-ageing is the process of age-hardening by holding the quenched alloy at room temperature.
Artificial-ageing is the process of ageing by holding the alloy at slightly higher temperature than room temperature. The effect of ageing on hardness of an Fe-0.02% N alloy at three different temperatures.
If at any given temperature, ageing is allowed to take place for much longer time, then coarsening of the precipitates occurs, i.e., precipitates grow in size at the cost of dissolution of smaller precipitates, so that the numerous fine precipitates are gradually replaced by a few coarse precipitates which are quite widely placed. The alloy, in this state becomes softer by drooping of the curves after attaining peak hardness. This is called over ageing.
4. Classification of Precipitates in Age-Hardening Treatment:
Depending on the structure of the boundary between the precipitate and the matrix, precipitates could be classified as:
(i) Coherent precipitate
(ii) Semi-or partially coherent precipitate
(iii) Incoherent precipitate.
With a coherent precipitate, its whole interface with the matrix is coherent, i.e., there is a one-to-one matching of the lattice planes across the interface. This generally produces elastic lattice-strains called coherency-strains around the boundary where the lattice planes must be ‘bent’ to give this one-to-one matching.
In an incoherent boundary, there is no regularity of lattice-plane matching across the boundary, i.e., there is no coherent boundary, as illustrated in Fig. 13.5 (b). This is a large angle boundary, between the precipitate and the matrix, i.e., normal interface boundary.
The formation of semi-coherent boundary is illustrated in Fig. 13.4. When two crystal lattices (here, α and β) touch each other to form an interface, the lattice planes like to have one-to-one matching across this interface. As the lattice parameters of two lattices are different (αα > αβ), elastic lattice-strains are produced as shown by dashed positions in Fig. 13.4 (c), but this creates an edge dislocation at the point where the β plane is located symmetrically between two α-planes. It clearly illustrates that on both sides of this edge dislocation, there are coherency-strains.
On such an interface, the edge dislocation is obtained almost periodically with a spacing D between the two neighbouring dislocations (D in above case of Fig. 13.4 = 6). Thus, a semi-coherent boundary consists alternately of region of coherency and region of disregistry (region around dislocation). The term disregistry, δ, is defined by
Problem:
Prove that the spacing between the dislocations.
D is related to the disregistry. δ on a semi-coherent interface as
Solution:
(Just like a vernier caliper) looking at Fig. 13.4,
An important conclusion can be drawn from the above equation, that when D becomes sufficiently small, (i.e., there is large difference between αα, and αβ), then complete disregistry is obtained across the boundary. i.e., it becomes an incoherent boundary.
Normally, when the precipitation starts, the precipitate formed may be initially coherent, which might become semi-coherent as the size of the precipitate increases to ultimately form incoherent precipitate, unless there is a large difference between αα and αβ. There normally is a change of the crystal lattice of precipitate simultaneously with respect to matrix.
5. Duplex Ageing Treatment of Metals:
Duplex-ageing has developed a lot of interest in age-hardening of non-ferrous alloys, as it is able to overcome the inferior properties such as stress-corrosion cracking in Al-Zn-Mg alloys, which developed by a single quench and age-hardening at a single temperature.
In normal age-hardening treatment, the interior of the grains developed the required fine distribution and density of the precipitates, but the places near grain boundaries (a place for the sink of vacancies) had precipitate free zones.
The boundary itself was seen to be overaged. This leads to inferior properties. Duplex-ageing in Al-Zn-Mg alloys, produced finer precipitates and good corrosion resistance as the grain-boundary precipitate-free-zone is absent to attain a strength of 267-308 MNm-2, enhancing their applications.
Duplex-ageing consists of holding the alloy first at one and then, at another temperature, the first temperature being lower than the second-stage ageing. The process aims at forming a large number of precipitate- nuclei at the low-temperature stage, when the solid solution is appreciably supersaturated and diffusion process is low.
Thus, a large number but very fine-sized nuclei form. These are then allowed to grow to optimum size at a fast rate by ageing at a second but higher temperature. Thus, duplex ageing provides a higher density and grater uniformity of distribution of precipitates than is possible in one step ageing at an elevated temperature.
If an alloy has to be artificially-aged by precipitates of intermediate phase, then the density of precipitates shall be coarse, because this intermediate phase can nucleate heterogeneously at dislocations, and grain- and sub-grain-boundaries at the optimum temperature of ageing. If somehow, this phase can be made to additionally nucleate at G.P. zones obtained due to earlier low temperature ageing, or even natural ageing, then there is large increase in the density of intermediate phase precipitates.
It has been seen that there is a critical temperature, Tc, of first ageing, above which such uniform homogeneous nucleation cannot occur (normally Tc ≈ G.P. zone solvus in some alloys). There is also a certain critical size of the first formed nuclei (normally the G.P. zones) above which, these are able to act as nuclei for precipitates while being aged at second as well as higher temperature. But below this size, the zones dissolve, when the ageing is done above Tc temperature.
Al-Zn-Mg alloys, after solution treatment, are quenched to room temperature. It is first aged at 100°C for 10-20 hours or naturally-aged for a month, or more. It results in uniform distribution of critical-sized zones, or even larger sized. A longer nucleation treatment gives a finer-dispersion of precipitates and very narrow precipitate-free-zones.
The final ageing may be done at 175°C to develop optimum properties with high resistance to corrosion. It is possible to stabilise GP zones by the addition of trace elements, such as by adding Si to Al-Cu-Mg alloy, so that ageing may be done even below Tc temperature, i.e., even a single ageing may be able to produce uniform properties.
Some Al-Mg-Si alloys show inferior properties being developed, and thus, duplex-ageing can produce useful or bad effects, depending on the composition and the temperature of ageing stage. Thus, before adopting the duplex-ageing, it should be carefully thought of.
6. Barriers to Age-Hardening Treatment of Metals:
According to dislocation theory, the strength of a material is controlled by the generation and mobility of the dislocations. The increased strength of an age-hardened alloy is due to the interaction of the moving dislocations with dispersed precipitates.
In an age-hardened alloy, the barriers to the motion of dislocations could be either, or both of the following:
(i) Coherency-strains around the GP zones.
(ii) GP zones or precipitates.
The mechanism of interaction is a complex process and thus, some very important basic concepts are explained here.
There could be at least three reasons of hardening by ageing:
1. Internal strain-hardening by elastic coherency-strains around zones,
2. Dispersion-hardening due to formation of loops of dislocations around precipitates.
3. Chemical-hardening due to precipitates being sheared (cut) by moving dislocation,
1. Internal Strain Hardening due to Coherency Strains around Zones:
The coherency strains (around coherent or even semi-coherent precipitates) act as barriers to the movement of moving dislocations. For a dislocation to pass through such regions of internal stress, the applied stress must be at least equal to average internal stress.
The internal stress increases as the misfit between the precipitate and matrix increases, as the elastic modulus of the matrix increases, or the surface area of the coherent boundary increases.
2. Dispersion Hardening:
When the stress required, let the dislocation cut through the precipitate particles is too high and thus cannot cut them (non-deformable particles), it then bows to bypass them as illustrated in Fig. 13.10. As the applied shear stress is increased, the dislocation bows sufficiently so that it begins to meet at points such as X and Y (Fig. 13.10 iv).
The nature of dislocation at X and Y are of opposite type, and consequently when these dislocation-segments meet, they annihilate each-other resulting in the main dislocation to separate from the looped regions A loop of dislocation is left behind around the precipitate particles. This is called Orowan mechanism.
The stress required to let the dislocation bypass the precipitate particles is given by:
where,
G is the shear modulus of the matrix,
b is the Burgers vector of the dislocation,
l is the distance between the precipitate particles.
As I decreases, the shear stresses to bypass increases. Normally, this process is of significance when the precipitates become incoherent after coherency strains are lost. Every time a dislocation bypasses, it leaves behind a loop of dislocation, thereby decreasing the distance between particles, and the next dislocation can be forced to bypass only at a higher shear stress. As the distance between the coarsening particles during over ageing increases, the strength of the alloy decreases.
3. Chemical-Hardening:
It is due to the short-range interaction of dislocation with a zone, or fine precipitate as illustrated in Fig. 13.11. As the dislocation cuts through the zone, the zone gets offset across the slip plane by one 
vector, and a change in the number of solute-solvent near-neighbours takes place across the slip plane.
This tries to reverse the process of clustering that means, additional work has to be done by the applied stress to bring about this change. Excess energy is also required as additional surface between the precipitate and the matrix has been created. Surface imperfections like stacking-faults are created too, which also resist the motion of dislocations.
The lattice of the zone or precipitate is not exactly the same as of the matrix, even in a fully coherent precipitate. The shearing disturbs the arrangement of atoms along the slip plane. Greater the difference in the lattices, greater is the disturbance in the arrangement of atoms, higher is the stress required to shear the precipitate. The shear modulus of the zone is usually higher than that of matrix. The more hard is a precipitated particle, the more difficult it is for a dislocation to cut it.
Taking help of equation 13.4, and assuming G/100 as the yield strength of the hardest age-hardened alloy, the radius of curvature of bending of a dislocation calculates to 50 atomic spacings, or 10 nm, i.e., it is the spacing between particles which gives maximum impedance. It has been seen that at this stage of fineness of precipitates (i.e., when distance between particles is ≈ 20 nm), normally the GP zones are present, but these are invariably sheared by the dislocations. Thus, to obtain high yield strength at this stage, the density of GP zones should be high.
Once the GP zones are sheared, the dislocations continue to pass through the particles on the active slip-planes and work-hardening is small, i.e., deformation tends to become localised on only a few active-planes. On the other hand if precipitates are widely spaced, these can be readily by-passed by moving dislocations by Orowan mechanism (Fig. 13.10).
The yield strength of the alloy is low in such a state, but the rate of work hardening is high, and the plastic deformation takes place more uniformly throughout the grains. This occurs when over ageing occurs, associated with the transition from shearing to by-passing the precipitates.
The cutting-through the precipitate-particles is possible, only when the slip plane is continuous from the matrix through the precipitate particle, and when the stress to move a dislocation in the crystal structure of the precipitate is comparable to that in the matrix. It occurs in small precipitate-particles like zones (less than 50 A° in size).
Cutting-through is not possible if there is an interface between the precipitate and the matrix, and if the orientation changes abruptly at the interface. Under such conditions, the dislocation can bend and bypass them. It occurs when particles are 100-500 A° in size.
The most interesting case is, when the precipitates resist the cutting through by dislocations and yet be too closely spaced not to allow bypassing by dislocations. Effective strengthening is obtained in bending process when the precipitate particles are submicroscopic in size.
The spacing between the particles should be typically a few hundred angstroms. Optimum ageing results when the antiparticle distance is proper. In the overaged alloy, particles are of larger sizes (> 1000 A°), and bypassing is easier. Thus, the strength and hardness decrease.
The degree of strengthening by age-hardening depends on the nature of the strengthening phases, and the size, quantity (volume) and distribution of these particles. In aluminium alloys, the maximum strengthening effect is due to MgZn2, Mg2Si, and Al2CuMg particles, which have a complex structure, and composition much different than a-solid solution.
Normally, the recrystallisation temperature is higher than the ageing-temperature of the alloy, and thus, the alloys retain polygonised structure too which has high dislocation density, if the alloy had been cold- worked earlier. The precipitation is more fine and uniform due to their precipitation on dislocation substructure. This causes additional strengthening of the alloys.