Precipitation hardening of Al–Cu alloys

7.2.1 Precipitation hardening of Al–Cu alloys

7.2.1.1 Precipitation from supersaturated solid solution

The basic requirements of a precipitation-hardening alloy system is that the solid solubility limit should decrease with decreasing temperature, as shown in Figure 7.1 for the Al–Cu system. During

386 Physical Metallurgy and Advanced Materials

a⫹ liquid

0 1 2 3 4 5 Cu (wt%)

Figure 7.1 Al-rich Al–Cu binary diagram showing GP[1], θ ′′ and θ ′ solvus lines (dotted).

the precipitation-hardening heat treatment procedure the alloy is first solution heat-treated at the high temperature and then rapidly cooled by quenching into water or some other cooling medium. The rapid cooling suppresses the separation of the θ-phase so that the alloy exists at the low temperature in an unstable supersaturated state. If, however, after quenching, the alloy is allowed to ‘age’ for a sufficient length of time, the second phase precipitates out. This precipitation occurs by a nucleation and growth process, fluctuations in solute concentration providing small clusters of atoms in the lattice which act as nuclei for the precipitate. However, the size of the precipitate becomes finer as the temperature at which precipitation occurs is lowered, and extensive hardening of the alloy is associated with a critical dispersion of the precipitate. If, at any given temperature, ageing is allowed to proceed too far, coarsening of the particles occurs (i.e. the small ones tend to redissolve and the large ones to grow still larger, as discussed in Section 7.2.6) and the numerous finely dispersed, small particles are gradually replaced by a smaller number of more widely dispersed, coarser particles. In this state the alloy becomes softer, and it is then said to be in the overaged condition (see Figure 7.2).

7.2.1.2 Changes in properties accompanying precipitation

The actual quenching treatment gives rise to small changes in many of the mechanical and physi- cal properties of alloys, because both solute atoms and point defects in excess of the equilibrium concentration are retained during the process, and because the quench itself often produces lattice strains. Perhaps the property most markedly affected is the electrical resistance and this is usually considerably increased. In contrast, the mechanical properties are affected relatively much less.

On ageing, the change in properties in a quenched material is more marked and, in particular, the mechanical properties often show striking modifications. For example, the tensile strength of Duralumin (i.e. an aluminum–4% copper alloy containing magnesium, silicon and manganese) may

be raised from 0.21 to 0.41 GN m −2 , while that of a Cu–2Be alloy may be increased from 0.46 to

1.23 GN m −2 . The structure-sensitive properties such as hardness, yield stress, etc. are, of course, extremely dependent on the structural distribution of the phases and, consequently, such alloys usually exhibit softening as the finely dispersed precipitates coarsen.

A simple theory of precipitation, involving the nucleation and growth of particles of the expected new equilibrium phase, leads one to anticipate that the alloy would show a single hardening peak,

Mechanical properties II – Strengthening and toughening 387

GP.[1] GP.[2]

(VPN) 120

u⬘ 4.5% Cu

100 80 4.0% Cu 60 3.0% Cu

2.0% Cu Hardness 40

0.1 1 10 100 Ageing time (days) at 130⬚C (a)

3.0% Cu 40 2.0% Cu

Hardness

Ageing time (days) at 190⬚C (b)

Figure 7.2 The ageing of aluminum–copper alloys at 130 ◦

C (b) (after Silcock, Heal and Hardy, 1953–4).

C (a) and 190 ◦

the electrical resistivity a decrease and the lattice parameter an increase (assuming the solute atom is smaller than the solvent atom) as the solute is removed from solution. Such property changes are found in practice, but only at low supersaturations and high ageing temperatures. At higher supersaturations and lower ageing temperatures the various property changes are not consistent with such a simple picture of precipitation; the alloy may show two or more age-hardening peaks, and the electrical resistivity and lattice parameter may not change in the anticipated manner. A hardening process which takes place in two stages is shown in aluminum–copper alloys (Figure 7.2a), where the initial hardening occurs without any attendant precipitation being visible in the light microscope and, moreover, is accompanied by a decrease in conductivity and no change in lattice parameter. Such behavior may be accounted for if precipitation is a process involving more than one stage. The initial stage of precipitation, at the lower ageing temperatures, involves a clustering of solute atoms on the solvent lattice sites to form zones or clusters, coherent with the matrix; the zones cannot be seen in the light microscope and for this reason this stage was at one time termed pre-precipitation. At a later stage of the ageing process these clusters break away from the matrix lattice to form distinct particles with their own crystal structure and a definite interface. These hypotheses were confirmed originally by structural studies using X-ray diffraction techniques, but nowadays the so-called pre-precipitation effects can be observed directly in the electron microscope.

Even though clustering occurs, the general kinetic behavior of the precipitation process is in agreement with that expected on thermodynamic grounds. From Figure 7.2 it is evident that the rate of ageing increases markedly with increasing temperature, while the peak hardness decreases.

Two-stage hardening takes place at low ageing temperatures and is associated with high maximum hardness, while single-stage hardening occurs at higher ageing temperatures, or at lower ageing temperatures for lower solute contents.

Another phenomenon commonly observed in precipitation-hardening alloys is reversion or retro- gression. If an alloy hardened by ageing at low temperature is subsequently heated to a higher ageing

388 Physical Metallurgy and Advanced Materials temperature it softens temporarily, but becomes harder again on more prolonged heating. This tem-

porary softening, or reversion of the hardening process, occurs because the very small nuclei or zones precipitated at the low temperature are unstable when raised to the higher ageing temperature, and consequently they redissolve and the alloy becomes softer. The temperature above which the nuclei or zones dissolve is known as the solvus temperature; Figure 7.1 shows the solvus temperatures for GP zones, θ ′′ ,θ ′ and θ. On prolonged ageing at the higher temperature larger nuclei, characteristic of that temperature, are formed and the alloy again hardens. Clearly, the reversion process is reversible, provided re-hardening at the higher ageing temperature is not allowed to occur.

7.2.1.3 Structural changes during precipitation

Early metallographic investigations showed that the microstructural changes which occur during the initial stages of ageing are on too fine a scale to be resolved by the light microscope, yet it is in these early stages that the most profound changes in properties are found. Accordingly, to study the process, it is necessary to employ the more sensitive and refined techniques of X-ray diffraction and electron microscopy.

The two basic X-ray techniques, important in studying the regrouping of atoms during the early stages of ageing, depend on the detection of radiation scattered away from the main diffraction lines or spots (see Chapter 4). In the first technique, developed independently by Guinier and Pre- ston in 1938, the Laue method is used. They found that the single-crystal diffraction pattern of an aluminum–copper alloy developed streaks extending from an aluminum lattice reflection along

Al directions. This was attributed to the formation of copper-rich regions of plate-like shape on {1 0 0} planes of the aluminum matrix (now called Guinier–Preston zones or GP zones). The net effect of the regrouping is to modify the scattering power of, and spacing between, very small groups of {1 0 0} planes throughout the crystal. However, being only a few atomic planes thick, the zones produce the diffraction effect typical of a two-dimensional lattice, i.e. the diffraction spot becomes a diffraction streak. In recent years the Laue method has been replaced by a single-crystal oscillation technique employing monochromatic radiation, since interpretation is made easier if the wavelength of the X-rays used is known. The second technique makes use of the phenomenon of scattering of X-rays at small angles (see Chapter 4). Intense small-angle scattering can often be observed from age-hardening alloys (as shown in Figures 7.3 and 7.5) because there is usually a difference in elec- tron density between the precipitated zone and the surrounding matrix. However, in alloys such as aluminum–magnesium or aluminum–silicon the technique is of no value because in these alloys the small difference in scattering power between the aluminum and silicon or magnesium atoms, respectively, is insufficient to give rise to appreciable scattering at small angles.

With the advent of the electron microscope the ageing of aluminum alloys was one of the first subjects to be investigated with the thin-foil transmission method. Not only can the detailed structural changes which occur during the ageing process be followed, but electron diffraction pictures taken from selected areas of the specimen while it is still in the microscope enable further important information on the structure of the precipitated phase to be obtained. Moreover, under some conditions the interaction of moving dislocations and precipitates can be observed. This naturally leads to a more complete understanding of the hardening mechanism.

Both the X-ray and electron-microscope techniques show that in virtually all age-hardening systems the initial precipitate is not the same structure as the equilibrium phase. Instead, the ageing sequence zones → intermediate precipitates → equilibrium precipitate is followed. This sequence occurs because the equilibrium precipitate is incoherent with the matrix, whereas the transition structures are either fully coherent, as in the case of zones, or at least partially coherent. Then, because of the importance of the surface energy and strain energy of the precipitate to the precipitation process, the system follows such a sequence in order to have the lowest free energy in all stages of precipitation.

Mechanical properties II – Strengthening and toughening 389

Figure 7.3 (a) Small-angle X-ray pattern from aluminum–4% copper single crystal taken with molybdenum Kα radiation at a sample to film distance of 4 cm (after Guinier and Fournet, 1955; courtesy of John Wiley and Sons). (b) Electron micrograph of aluminum–4% copper aged 16 hours at 130 ◦

C, showing GP[1] zones (after Nicholson, Thomas and Nutting, 1958–9). The surface energy of the precipitates dominates the process of nucleation when the interfacial energy

is large (i.e. when there is a discontinuity in atomic structure, somewhat like a grain boundary, at the interface between the nucleus and the matrix), so that for the incoherent type of precipitate the nuclei must exceed a certain minimum size before they can nucleate a new phase. To avoid such a slow mode of precipitation a coherent type of precipitate is formed instead, for which the size effect is relatively unimportant. The condition for coherence usually requires the precipitate to strain its equilibrium lattice to fit that of the matrix, or to adopt a metastable lattice. However, in spite of both

a higher volume free energy and a higher strain energy, the transition structure is more stable in the early stages of precipitation because of its lower interfacial energy. When the precipitate does become incoherent the alloy will, nevertheless, tend to reduce its sur- face energy as much as possible, by arranging the orientation relationship between the matrix and the precipitate so that the crystal planes which are parallel to, and separated by, the bounding surface have similar atomic spacings. Clearly, for these habit planes, as they are called, the better the crystal- lographic match, the less will be the distortion at the interface and the lower the surface energy. This principle governs the precipitation of many alloy phases, as shown by the frequent occurrence of the

Widmanstätten structure, i.e. plate-shaped precipitates lying along prominent crystallographic planes of the matrix. Most precipitates are plate shaped because the strain energy factor is least for this form.

The existence of a precipitation sequence is reflected in the ageing curves and, as we have seen in Figure 7.2, often leads to two stages of hardening. The zones, by definition, are coherent with the matrix, and as they form the alloy becomes harder. The intermediate precipitate may be coherent with the matrix, in which case a further increase of hardness occurs, or only partially coherent, when either hardening or softening may result. The equilibrium precipitate is incoherent and its formation always leads to softening. These features are best illustrated by a consideration of some actual age-hardening systems.

Precipitation reactions occur in a wide variety of alloy systems, as shown in Table 7.1. The aluminum–copper alloy system exhibits the greatest number of intermediate stages in its precipitation

390 Physical Metallurgy and Advanced Materials Table 7.1 Some common precipitation-hardening systems.

Base Solute

Equilibrium metal

Transition structure

precipitate Al

Cu

(i) Plate-like solute-rich GP[1] zones on {1 0 0} Al ;

θ -CuAl 2

(ii) ordered zones of GP[2]; (iii) θ ′ -phase (plates). Ag (i) Spherical solute-rich zones; (ii) platelets of

γ -Ag 2 Al

hexagonal γ ′ on {1 1 1} Al .

Mg, Si

(i) GP zones rich in Mg and Si atoms on {1 0 0} Al

β -Mg 2 Si (plates)

planes; (ii) ordered zones of β ′ . Mg, Cu (i) GP zones rich in Mg and Cu atoms on {1 0 0} Al

S-Al 2 CuMg (laths)

planes; (ii) S ′ platelets on {0 2 1} Al planes.

Mg, Zn

(i) Spherical zones rich in Mg and Zn; (ii) platelets

η -MgZn 2 (plates)

of η ′ -phase on {1 1 1} Al .

Cu

γ -CuBe Co

Be (i) Be-rich regions on {1 0 0} Cu planes; (ii) γ ′ .

Spherical GP zones.

β -Co plates Fe C (i) Martensite (α ′ ); (ii) martensite (α ′′ ); (iii) ε-carbide.

Fe 3 C plates cementite N

(i) Nitrogen martensite (α ′ ); (ii) martensite (α ′′ ) disks. Fe 4 N Ni

Al, Ti

γ ′ cubes

γ -Ni 3 (AlTi)

process, and consequently is probably the most widely studied. When the copper content is high and the ageing temperature low, the sequence of stages followed is GP[1], GP[2], θ ′ and θ (CuAl 2 ). On ageing at higher temperatures, however, one or more of these intermediate stages may be omitted and, as shown in Figure 7.2, corresponding differences in the hardness curves can be detected. The early stages of ageing are due to GP[1] zones, which are interpreted as plate-like clusters of copper atoms segregated onto {1 0 0} planes of the aluminum matrix. A typical small-angle X-ray scattering pattern and thin-foil transmission electron micrograph from GP[1] zones are shown in Figure 7.3.

The plates are only a few atomic planes thick (giving rise to the but are about 10 nm long, and hence appear as bright or dark lines on the electron micrograph.

GP[2] is best described as a coherent intermediate precipitate rather than a zone, since it has a definite crystal structure; for this reason the symbol θ ′′ is often preferred. These precipitates, usually of maximum thickness 10 nm and up to 150 nm in diameter, have a tetragonal structure which fits perfectly with the aluminum unit cell in the a and b directions but not in the c. The structure postulated has a central plane which consists of 100% copper atoms, the next two planes a mixture of copper and

aluminum and the other two basal planes of pure aluminum, giving an overall composition of CuAl 2 . Because of their size, θ ′′ precipitates are easily observed in the electron microscope, and because of the ordered arrangements of copper and aluminum atoms within the structure, their presence gives rise to intensity maxima on the diffraction streaks in an X-ray photograph. Since the c parameter (0.78 nm) differs from that of aluminum (0.404 nm) the aluminum planes parallel to the plate are distorted by elastic coherency strains. Moreover, the precipitate grows with the c-direction normal to the plane of the plate, so that the strain fields become larger as it grows and at peak hardness extend from one precipitate particle to the next (see Figure 7.4a). The direct observation of coherency strains confirms the theories of hardening based on the development of an elastically strained matrix (see next section).

The transition structure θ ′ is tetragonal; the true unit cell dimensions are a = 0.404 and c = 0.58 nm and the axes are parallel to

Al directions. The strains around the θ ′ plates can be relieved, however, by the formation of a stable dislocation loop around the precipitate and such a loop has been observed around small θ ′ plates in the electron microscope, as shown in Figure 7.4b. The long-range

Mechanical properties II – Strengthening and toughening 391

Figure 7.4 Electron micrographs from Al–4Cu: (a) aged 5 hours at 160 ◦

C showing θ ′′ plates; (b) aged 12 hours at 200 ◦

C showing a dislocation ring round θ ′′ plates; (c) aged 3 days at 160 ◦ C

showing θ ′′ precipitated on helical dislocations (after Nicholson, Thomas and Nutting, 1958–9).

strain fields of the precipitate and its dislocation largely cancel. Consequently, it is easier for glide dislocations to move through the lattice of the alloy containing an incoherent precipitate such as θ ′ than a coherent precipitate such as θ ′′ , and the hardness falls.

The θ structure is also tetragonal, with a = 0.606 and c = 0.487 nm. This equilibrium precipitate is incoherent with the matrix and its formation always leads to softening, since coherency strains disappear.