Changes in properties accompanying temperatures the various property changes are not con- precipitation

8.2.1.2 Changes in properties accompanying temperatures the various property changes are not con- precipitation

sistent with such a simple picture of precipitation; the alloy may show two or more age-hardening peaks, and

The actual quenching treatment gives rise to small the electrical resistivity and lattice parameter may not changes in many of the mechanical and physical prop-

change in the anticipated manner. A hardening pro- erties of alloys because both solute atoms and point

cess which takes place in two stages is shown in alu- defects in excess of the equilibrium concentration are

minium–copper alloys (Figure 8.2a) where the initial retained during the process, and because the quench

hardening occurs without any attendant precipitation itself often produces lattice strains. Perhaps the prop-

being visible in the light microscope and, moreover, erty most markedly affected is the electrical resistance

is accompanied by a decrease in conductivity and no and this is usually considerably increased. In con-

change in lattice parameter. Such behaviour may be trast, the mechanical properties are affected relatively

accounted for if precipitation is a process involving much less.

more than one stage. The initial stage of precipitation, On ageing, the change in properties in a at the lower ageing temperatures, involves a cluster-

quenched material is more marked and, in particular, ing of solute atoms on the solvent lattice sites to form the mechanical properties often show striking

zones or clusters, coherent with the matrix; the zones modifications. For example, the tensile strength of

cannot be seen in the light microscope and for this rea- Duralumin (i.e. an aluminium–4% copper alloy

son this stage was at one time termed pre-precipitation. containing magnesium, silicon and manganese) may

At a later stage of the ageing process these clusters

be raised from 0.21 to 0.41 GN/m 2 while that of

break away from the matrix lattice to form distinct

a Cu–2Be alloy may be increased from 0.46 to 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 behaviour of the precipitation process is in agreement with that expected on thermodynamic grounds. From Figure 8.2 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 retrogres- sion. If an alloy hardened by ageing at low temperature is subsequently heated to a higher ageing temperature it softens temporarily, but becomes harder again on more prolonged heating. This temporary softening, or reversion of the hardening process, occurs because the very small nuclei or zones precipitated at the low tem-

Figure 8.1 Al-rich Al–Cu binary diagram showing GP [1],

00 0 solvus lines (dotted) perature are unstable when raised to the higher ageing . temperature, and consequently they redissolve and the

Strengthening and toughening 261 alloy developed streaks extending from an aluminium

lattice reflection along h1 0 0i Al directions. This was attributed to the formation of copper-rich regions of plate-like shape on f1 0 0g planes of the aluminium matrix (now called Guinier–Preston zones or GP zones). The net effect of the regrouping is to mod- ify the scattering power of, and spacing between, very small groups of f1 0 0g 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 interpreta- tion is made easier if the wavelength of the X-rays used is known. The second technique makes use of the phe- nomenon of scattering of X-rays at small angles (see Chapter 5). Intense small-angle scattering can often

be observed from age-hardening alloys (as shown in Figures 8.3 and 8.5) because there is usually a differ- ence in electron density between the precipitated zone and the surrounding matrix. However, in alloys such as aluminium–magnesium or aluminium–silicon the technique is of no value because in these alloys the small difference in scattering power between the alu- minium and silicon or magnesium atoms, respectively, is insufficient to give rise to appreciable scattering at small angles.

Figure 8.2 The ageing of aluminium–copper alloys at (a) 130 °

With the advent of the electron microscope the age- 1953–4) .

C and (b) at 190 °

C (after Silcock, Heal and Hardy,

ing of aluminium alloys was one of the first subjects to

be investigated with the thin-foil transmission method. Not only can the detailed structural changes which

alloy becomes softer; the temperature above which the occur during the ageing process be followed, but elec- nuclei or zones dissolve is known as the solvus tem-

tron diffraction pictures taken from selected areas of perature; Figure 8.1 shows the solvus temperatures for

the specimen while it is still in the microscope enable

00 0 further important information on the structure of the higher temperature larger nuclei, characteristic of that

precipitated phase to be obtained. Moreover, under temperature, are formed and the alloy again hardens.

some conditions the interaction of moving dislocations Clearly, the reversion process is reversible, provided

and precipitates can be observed. This naturally leads re-hardening at the higher ageing temperature is not

to a more complete understanding of the hardening allowed to occur.

mechanism.

Both the X-ray and electron-microscope techniques

8.2.1.3 Structural changes during precipitation show that in virtually all age-hardening systems the initial precipitate is not the same structure as the equi-

Early metallographic investigations showed that the librium phase. Instead, an ageing sequence: zones ! microstructural changes which occur during the initial

intermediate precipitates ! equilibrium precipitate is stages of ageing are on too fine a scale to be resolved

followed. This sequence occurs because the equilib- by the light microscope, yet it is in these early stages

rium precipitate is incoherent with the matrix, whereas that the most profound changes in properties are found.

the transition structures are either fully coherent, as in Accordingly, to study the process, it is necessary to

the case of zones, or at least partially coherent. Then, employ the more sensitive and refined techniques of

because of the importance of the surface energy and X-ray diffraction and electron microscopy.

strain energy of the precipitate to the precipitation pro- The two basic X-ray techniques, important in study-

cess, the system follows such a sequence in order to ing the regrouping of atoms during the early stages

have the lowest free energy in all stages of precipita- of ageing, depend on the detection of radiation scat-

tion. The surface energy of the precipitates dominates tered away from the main diffraction lines or spots

the process of nucleation when the interfacial energy is (see Chapter 5). In the first technique, developed

large (i.e. when there is a discontinuity in atomic struc- independently by Guinier and Preston in 1938, the

ture, somewhat like a grain boundary, at the interface Laue method is used. They found that the single-

between the nucleus and the matrix), so that for the crystal diffraction pattern of an aluminium–copper

incoherent type of precipitate the nuclei must exceed a

262 Modern Physical Metallurgy and Materials Engineering [010]

(b) Figure 8.3 (a) Small-angle X-ray pattern from aluminium–4% copper single crystal taken with molybdenum K˛ radiation at a

(a)

sample to film distance of 4 cm (after Guinier and Fournet, 1955; courtesy of John Wiley and Sons). (b) Electron micrograph of aluminium–4% copper aged 16 hours at 130 °

C, showing GP [1] zones (after Nicholson, Thomas and Nutting, 1958–9) .

certain minimum size before they can nucleate a new Precipitation reactions occur in a wide variety phase. To avoid such a slow mode of precipitation

of alloy systems as shown in Table 8.1. The

a coherent type of precipitate is formed instead, for aluminium–copper alloy system exhibits the greatest which the size effect is relatively unimportant. The

number of intermediate stages in its precipitation condition for coherence usually requires the precipi-

process, and consequently is probably the most tate to strain its equilibrium lattice to fit that of the

widely studied. When the copper content is high and matrix, or to adopt a metastable lattice. However, in

the ageing temperature low, the sequence of stages spite of both a higher volume free energy and a higher

0 2 ⊳ . On strain energy, the transition structure is more stable in

ageing at higher temperatures, however, one or more the early stages of precipitation because of its lower

of these intermediate stages may be omitted and, interfacial energy.

as shown in Figure 8.2, corresponding differences When the precipitate does become incoherent the

in the hardness curves can be detected. The early alloy will, nevertheless, tend to reduce its surface

stages of ageing are due to GP [1] zones, which energy as much as possible, by arranging the orienta-

are interpreted as plate-like clusters of copper atoms tion relationship between the matrix and the precipitate

segregated onto f1 0 0g planes of the aluminium matrix. so that the crystal planes which are parallel to, and sep-

A typical small-angle X-ray scattering pattern and arated by, the bounding surface have similar atomic

thin-foil transmission electron micrograph from GP [1] spacings. Clearly, for these habit planes, as they are

zones are shown in Figure 8.3. The plates are only called, the better the crystallographic match, the less

a few atomic planes thick (giving rise to the h1 0 0i will be the distortion at the interface and the lower

streaks in the X-ray pattern), but are about 10 nm long, the surface energy. This principle governs the precip-

and hence appear as bright or dark lines on the electron itation of many alloy phases, as shown by the fre-

micrograph.

quent occurrence of the Widmanst¨atten structure, i.e. GP [2] is best described as a coherent intermediate plate-shaped precipitates lying along prominent crys-

precipitate rather than a zone, since it has a defi- tallographic planes of the matrix. Most precipitates are

plate-shaped because the strain energy factor is least is often preferred. These precipitates, usually of max- for this form.

imum thickness 10 nm and up to 150 nm diameter, The existence of a precipitation sequence is reflected

have a tetragonal structure which fits perfectly with in the ageing curves and, as we have seen in

the aluminium unit cell in the a and b directions but Figure 8.2, often leads to two stages of hardening.

not in the c. The structure postulated has a central The zones, by definition, are coherent with the

plane which consists of 100% copper atoms, the next matrix, and as they form the alloy becomes harder.

two planes a mixture of copper and aluminium and The intermediate precipitate may be coherent with

the other two basal planes of pure aluminium, giv- the matrix, in which case a further increase of

ing an overall composition of CuAl 00 2 . Because of their hardness occurs, or only partially coherent, when either

precipitates are easily observed in the elec- hardening or softening may result. The equilibrium

tron microscope, and because of the ordered arrange- precipitate is incoherent and its formation always leads

ments of copper and aluminium atoms within the to softening. These features are best illustrated by a

structure, their presence gives rise to intensity max- consideration of some actual age-hardening systems.

ima on the diffraction streaks in an X-ray photograph.

Strengthening and toughening 263

Table 8.1 Some common precipitation-hardening systems

Base Solute Transition structure Equilibrium metal

precipitate Al

Cu (i) Plate-like solute rich GP [1] zones on

-CuAl 2

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

f 1 0 0g

0 -Ag 2 of hexagonal Al on f1 1 1g Al . Mg, Si

(i) GP zones rich in Mg and Si atoms on

ˇ -Mg 2 Si

f 1 0 0g Al planes; (ii) ordered zones of ˇ 0 .

(plates)

Mg, Cu (i) GP zones rich in Mg and Cu atoms on

S-Al 2 CuMg

f 1 0 0g planes; (ii) S 0

Al

platelets on

(laths)

f 0 2 1g Al planes.

Mg, Zn

(i) Spherical zones rich in Mg and Zn; (ii) platelets 0 -MgZn 2

phase on f1 1 1g Al .

(plates)

Cu

Be (i) Be-rich regions on f1 0 0g Cu planes; (ii) 0 .

-CuBe

Co Spherical GP zones. ˇ -Co plates

Fe 3 C plates (iii) ε-carbide.

Fe C (i) Martensite (˛ 0 ); (ii) martensite (˛ 00 );

cementite

(i) Nitrogen martensite (˛ 0 ); (ii) martensite

Fe 4 N

(˛ 00 Ni

) discs. cubes

Al, Ti

-Ni 3 (AlTi)

Since the c parameter 0.78 nm differs from that of hexagonal 0 ! equilibrium hexagonal . The hard- aluminium 0.404 nm the aluminium planes parallel to

ening is associated with the first two stages in which the plate are distorted by elastic coherency strains.

the precipitate is coherent and partially coherent with Moreover, the precipitate grows with the c direction

the matrix, respectively.

normal to the plane of the plate, so that the strain During the quench and in the early stages of ageing, fields become larger as it grows and at peak hard-

silver atoms cluster into small spherical aggregates and ness extend from one precipitate particle to the next

a typical small-angle X-ray picture of this stage, shown (see Figure 8.4a). The direct observation of coherency

in Figure 8.5a, has a diffuse ring surrounding the trace strains confirms the theories of hardening based on the

of the direct beam. The absence of intensity in the development of an elastically strained matrix (see next

centre of the ring (i.e. at ⊲0 0 0⊳) is attributed to the section).

0 is tetragonal; the true fact that clustering takes place so rapidly that there is left a shell-like region surrounding each cluster which unit cell dimensions are a D 0.404 and c D 0.58 nm

is low in silver content. On ageing, the clusters grow in

size and decrease in number, and this is characterized plates can be relieved, however,

and the axes are parallel to h1 0 0i 0 Al directions. The

by the X-ray pattern showing a gradual decrease in ring by the formation of a stable dislocation loop around

diameter. The concentration and size of clusters can be the precipitate and such a loop has been observed 0 followed very accurately by measuring the intensity

plates in the electron microscope as distribution across the ring as a function of ageing shown in Figure 8.4b. The long-range strain fields

time. This intensity may be represented (see Chapter 5) of the precipitate and its dislocation largely cancel.

by an equation of the form

Consequently, it is easier for glide dislocations to move through the lattice of the alloy containing an incoherent

l⊲ε⊳ D Mn 2 2 R 2 ε 2 / 2 ⊳

1 ε / 2 ⊳ ] 2 ⊲ 8.1⊳ and c D 0.487 nm. This equilibrium precipitate is

2 R 2 , and the hardness falls. 2

00 than a coherent precipitate such

and for values of ε greater than that corresponding incoherent with the matrix and its formation always

to the maximum intensity, the contribution of the leads to softening, since coherency strains disap-

second term, which represents the denuded region pear.

surrounding the cluster, can be neglected. Figure 8.5b shows the variation in the X-ray intensity, scattered at

C. An analysis of this Investigations using X-ray diffraction and electron

8.2.2 Precipitation-hardening of Al–Ag alloys

small angles (SAS) with cluster growth, on ageing an aluminium–silver alloy at 120 °

intensity distribution, using equation (8.1), indicates microscopy have shown the existence of three dis-

that the size of the zones increases from 2 to 5 nm in tinct stages in the age-hardening process, which may

just a few hours at 120 °

C. These zones may, of course,

be summarized: silver-rich clusters ! intermediate

be seen in the electron microscope and Figure 8.6a

264 Modern Physical Metallurgy and Materials Engineering

0.25 µ (a)

Figure 8.5 Small-angle scattering of Cu K˛ radiation by polycrystalline Al–Ag. (a) After quenching from 520 ° C (after Guinier and Walker, 1953). (b) The change in ring

1µ intensity and ring radius on ageing at 120 ° C (after Smallman and Westmacott, unpublished). (c) After ageing at (b)

C for 10 days (after Guinier and Walker, 1953) .

structure is hexagonal and, consequently, the precipi- tates are easily recognizable in the electron microscope by the stacking fault contrast within them, as shown in Figure 8.6b. Clearly, these precipitates are never fully coherent with the matrix, but, nevertheless, in this alloy system, where the zones are spherical and have little or no coherency strain associated with them, and where

no coherent intermediate precipitate is formed, the par- 0

precipitates do provide a greater resistance to dislocation movement than zones and a

tially coherent

(c)

Figure 8.4 Electron micrographs from Al–4Cu (a) aged second stage of hardening results.

The same principles apply to the constitution- 200 °

5 hours at 160 °

00 plates, (b) aged 12 hours at

ally more complex ternary and quaternary alloys 3 days at 160 °

00 plates, (c) aged

00 precipitated on helical as to the binary alloys. Spherical zones are found dislocations (after Nicholson, Thomas and Nutting, 1958–9) .

in aluminium–magnesium–zinc alloys as in alu- minium–zinc, although the magnesium atom is some 12% larger than the aluminium atom. The intermedi-

is an electron micrograph showing spherical zones ate precipitate forms on the f1 1 1g Al planes, and is in an aluminium–silver alloy aged 5 hours at 160 ° C;

partially coherent with the matrix with little or no the diameter of the zones is about 10 nm in good

strain field associated with it. Hence, the strength of agreement with that deduced by X-ray analysis. The

the alloy is due purely to dispersion hardening, and zone shape is dependent upon the relative diameters

the alloy softens as the precipitate becomes coarser. of solute and solvent atoms. Thus, solute atoms such

In nickel-based alloys the hardening phase is the as silver and zinc which have atomic sizes similar to

3 0 Al; this is an equilibrium phase in aluminium give rise to spherical zones, whereas solute

ordered 0 -Ni

the Ni –Al and Ni–Cr–Al systems and a metastable atoms such as copper which have a high misfit in the

phase in Ni–Ti and Ni –Cr–Ti. These systems form solvent lattice form plate-like zones.

the basis of the ‘superalloys’ (see Chapter 9) which With prolonged annealing, the formation and growth

owe their properties to the close matching of the 0

and the fcc matrix. The two phases have very simi- terized by the appearance in the X-ray pattern of short

of platelets of a new phase, 0 , occur. This is charac-

lar lattice parameters (⊲ 0.25%⊳, depending on com- streaks passing through the trace of the direct beam

position) and the coherency (interfacial energy 1 ³ (Figure 8.5c). The 0 platelet lies parallel to the f1 1 1g

10–20 mJ/m 2 ) confers a very low coarsening rate on planes of the matrix and its structure has lattice param-

the precipitate so that the alloy overages extremely eters very close to that of aluminium. However, the

slowly even at 0.7T m .

Strengthening and toughening 265 particles, when the dislocations bypass the particles,

the alloy strength is independent of the particle properties but is strongly dependent on particle size and dispersion strength decreasing as particle size or dispersion increases. The transition from deformable to non-deformable particle-controlled deformation is readily recognized by the change in microstructure, since the ‘laminar’ undisturbed dislocation flow for the former contrasts with the turbulent plastic flow for non-

deformable particles. The latter leads to the production

of a high density of dislocation loops, dipoles and other debris which results in a high rate of work-hardening. This high rate of work-hardening is a distinguishing feature of all dispersion-hardened systems.

(a)