diffusion coefficient of the solut ormation c In alu. a decrease in the rate of formation of zones which minium-copper, for example, the nust mean that the dislocations introduced by cold is already apparent after only a few at room work are more effective as vacancy sinks than as temperature, and is complete after an hour or two, vacancy sources. Cold working or rapid quenching so that the copper atoms must therefore have moved therefore have opposing effects on the formation of responds to an apparent diffusion coefficient of copper of precipitation-hardening. For example, the excess in aluminium of about 10-20-10-22 m2s", which is vacancies, by condensing to form a high density of many orders of magnitude faster than the value of dislocation loops, can provide nucleation sites fc x 10-29 m2s-I obtained by extrapolation of high- intermediate precipitates. This leads to the interest- temperature data. Many workers have attributed this ing observation in aluminium-copper alloys that cold enhanced diffusion to the excess vacancies retained working or rapid quenching, by producing dislocations ion for the diffusion coefficient at a given temperature mation of the e phase but, as we have seen above, the ontains a factor proportional to the concentration of opposite effect on zone formation. It is also interesting acancies at that temperature. if the sample contains an to note that screw dislocations, which are not normally bnormally large vacancy concentration then the diffu- favourable sites for nucleation, can also become sites sion coefficient should be increased by the ratio co/co, for preferential precipitation when they have climbed where cq is the quenched-in vacancy concentration and into helical dislocations by absorbing vacancies, and co is the equilibrium concentration. the observed clus- have thus become mainly of edge character. The long tering rate can be accounted for if the concentration of arrays of 6 phase observed in aluminium-copper vacancies retained is about 10-3-10-4 alloys, shown in Figure 8.4c, have probably formed The observation of loops by transmission electron on helices in this way. In some of these alloys, defects microscopy allows an estimate of the number of containing stacking faults are observed, in addition to cies to be made and in all cases of the dislocation loops and helices, and examples have rapid quenching the vacancy concentration in these been found where such defects nucleate an interme alloys is somewhat greater than 10-4, in agreement diate precipitate having a hexagonal structure. In alu the excess vacancies are removed, the amount of dislocations introduced by quenching absorb silver and enhanced diffusion diminishes, which agrees with the degenerate into long narrow stacking faults on (111) observations that the isothermal rate of clustering planes; these stacking-fault defects then act as nuclei decreases continuously with increasing time. In fact, he hexagonal y precipitate it is observed that d decreases rapidly at first and Many commercial alloys depend critically on then remains at a value well above the equilibrium the interrelation between vacancies, dislocations and value for months at room temperature; the process is solute atoms and it is found that trace impurities therefore separated into what is called the fast and significantly modify the precipitation process. Thus slow reaction is that some of the vacancies quenched. inhibit zone formation, e.g. Cd, In, Sn prevent zone in are trapped temporarily and then released slowly. formation in slowly quenched Al-Cu alloys for up Measurements show that the activation energy in the to 200 days at 30'C. This delays the age-hardening fast reaction(0.5 eV) is smaller than in the slow process at room temperature which gives more time for reaction(l ev) by an amount which can be attributed mechanically fabricating the quenched alloy before it to the binding energy between vacancies and trapping gets too hard, thus avoiding the need for refrigeration sites. These tra e very likely small dislocation On the other hand, Cd increases the density of e loops or voids formed by the clustering of vacancies. precipitate by increasing the density of vacancy loops The equilibrium matrix vacancy concentration would and helices which act as nuclei for precipitation and by a factor exp lra/rkTl, where y is the surface energy, the interfacial ene k x-8 interfaces thereby reduci then be greater than that for a well-annealed crystal by segregating to the mat Q the atomic volume and r the radius of the defect Since grain boundaries absorb vacancies in many (see Chapter 4). The experimental diffusion rate can alloys there is a grain boundary zone relatively free be accounted for if r a 2 nm, which is much smaller from precipitation. The Al-Zn-Mg alloy is one com- than the loops and voids usually seen, but they do exist. mercial alloy which suffers grain boundary weakness The activation energy for the reaction would then but it is found that trace additions of Ag have a be Ep-(r/r)or approxi I ev for r≈2n eficial effect in refining the precipitate structure and Other factors known to ate free grain boundary zone early stages of ageing(e.g. altering the quenching rate, Here it appears that Ag atoms stabilize vacancy clus- interrupted quenching and cold work) may also be ters near the grain boundary and also increase the alized on the basis that thes tability of the gp zone thereby raising the gp zone
Strengthening and toughening 269 diffusion coefficient of the solute atoms. In aluminium-copper, for example, the formation of zones is already apparent after only a few minutes at room temperature, and is complete after an hour or two, so that the copper atoms must therefore have moved through several atomic spacings in that time. This corresponds to an apparent diffusion coefficient of copper in aluminium of about 10 -2~ -22 m 2 s -l, which is many orders of magnitude faster than the value of 5 • 10 -29 m 2 s -l obtained by extrapolation of hightemperature data. Many workers have attributed this enhanced diffusion to the excess vacancies retained during the quenching treatment. Thus, since the expression for the diffusion coefficient at a given temperature contains a factor proportional to the concentration of vacancies at that temperature, if the sample contains an abnormally large vacancy concentration then the diffusion coefficient should be increased by the ratio co/c o , where c o is the quenched-in vacancy concentration and Co is the equilibrium concentration. The observed clustering rate can be accounted for if the concentration of vacancies retained is about 10 -3-10 -4 . The observation of loops by transmission electron microscopy allows an estimate of the number of excess vacancies to be made, and in all cases of rapid quenching the vacancy concentration in these alloys is somewhat greater than 10 -4 , in agreement with the predictions outlined above. Clearly, as the excess vacancies are removed, the amount of enhanced diffusion diminishes, which agrees with the observations that the isothermal rate of clustering decreases continuously with increasing time. In fact, it is observed that D decreases rapidly at first and then remains at a value well above the equilibrium value for months at room temperature; the process is therefore separated into what is called the fast and slow reactions. A mechanism proposed to explain the slow reaction is that some of the vacancies quenchedin are trapped temporarily and then released slowly. Measurements show that the activation energy in the fast reaction (~0.5 eV) is smaller than in the slow reaction (~ 1 eV) by an amount which can be attributed to the binding energy between vacancies and trapping sites. These traps are very likely small dislocation loops or voids formed by the clustering of vacancies. The equilibrium matrix vacancy concentration would then be greater than that for a well-annealed crystal by a factor exp [yf2/rkT], where y is the surface energy, f2 the atomic volume and r the radius of the defect (see Chapter 4). The experimental diffusion rate can be accounted for if r ,~ 2 nm, which is much smaller than the loops and voids usually seen, but they do exist. The activation energy for the slow reaction would then be ED -- (Vf2/r) or approximately 1 eV for r ~ 2 nm. Other factors known to affect the kinetics of the early stages of ageing (e.g. altering the quenching rate, interrupted quenching and cold work) may also be rationalized on the basis that these processes lead to different concentrations of excess vacancies. In general, cold working the alloy prior to ageing causes a decrease in the rate of formation of zones, which must mean that the dislocations introduced by cold work are more effective as vacancy sinks than as vacancy sources. Cold working or rapid quenching therefore have opposing effects on the formation of zones. Vacancies are also important in other aspects of precipitation-hardening. For example, the excess vacancies, by condensing to form a high density of dislocation loops, can provide nucleation sites for intermediate precipitates. This leads to the interesting observation in aluminium-copper alloys that cold working or rapid quenching, by producing dislocations for nucleation sites, have the same effect on the formation of the 0' phase but, as we have seen above, the opposite effect on zone formation. It is also interesting to note that screw dislocations, which are not normally favourable sites for nucleation, can also become sites for preferential precipitation when they have climbed into helical dislocations by absorbing vacancies, and have thus become mainly of edge character. The long arrays of 0' phase observed in aluminium-copper alloys, shown in Figure 8.4c, have probably formed on helices in this way. In some of these alloys, defects containing stacking faults are observed, in addition to the dislocation loops and helices, and examples have been found where such defects nucleate an intermediate precipitate having a hexagonal structure. In aluminium-silver alloys it is also found that the helical dislocations introduced by quenching absorb silver and degenerate into long narrow stacking faults on {1 1 1 } planes; these stacking-fault defects then act as nuclei for the hexagonal ~,' precipitate. Many commercial alloys depend critically on the interrelation between vacancies, dislocations and solute atoms and it is found that trace impurities significantly modify the precipitation process. Thus trace elements which interact strongly with vacancies inhibit zone formation, e.g. Cd, In, Sn prevent zone formation in slowly quenched AI-Cu alloys for up to 200 days at 30~ This delays the age-hardening process at room temperature which gives more time for mechanically fabricating the quenched alloy before it gets too hard, thus avoiding the need for refrigeration. On the other hand, Cd increases the density of 0' precipitate by increasing the density of vacancy loops and helices which act as nuclei for precipitation and by segregating to the matrix-0' interfaces thereby reducing the interfacial energy. Since grain boundaries absorb vacancies in many alloys there is a grain boundary zone relatively free from precipitation. The A1-Zn-Mg alloy is one commercial alloy which suffers grain boundary weakness but it is found that trace additions of Ag have a beneficial effect in refining the precipitate structure and removing the precipitate free grain boundary zone. Here it appears that Ag atoms stabilize vacancy clusters near the grain boundary and also increase the stability of the GP zone thereby raising the GP zone solvus temperature. Similarly, in the 'Concorde' alloy, RR58 (basically AI-2.5Cu-1.2Mg with additions), Si
270 Modern Physical Metallurgy and Materials Engineerin addition(0. 25Si)modifies the as-quenched dislocation character. Thus, for example, the Babab (100) plane distribution inhibiting the nucleation and growth stacking sequence of the fcc structure can be changed dislocation loops and reducing the diameter of helices. to BAAba by the propagation of an a/ 2(100)shear ated in the presence of Si rather than heterogeneously gation of a pair of a/2(100)partials of opposite sign directly from zones, giving rise to improved and more from the double shear is precisely that required for the niform properties embryonic formation of the fluorite structure from the Apart from speeding up the kinetics of ageing, fcc lattice. and providing dislocations nucleation sites, vacan- In visualizing the role of lattice defects in the nucle cies may play a structural role when they precipi- ation and growth of plate-shaped precipitates, a tate cooperatively with solute atoms to facilitate the analogy with Frank and Shockley partial disle basic atomic arrangements required for transformin loops is useful. In the formation of a Frank loop, the parent crystal structure to that of the produ of hep material is created from the fcc lattice by the phase. In essence, the process involves the system-(non-conservative) condensation of a layer of vacan ic incorporation of excess vacancies, produced by the cies in (11 1 Exactly the same structure is formed by initial quench or during subsequent dislocation loop the(conservative)expansion of a Shockley partial loo nnealing, in a precipitate zone or plate to change the on a [l 1 ll plane. In the former case a semi-coherent omic stacking. A simple example of 6 formation in 'precipitate'is produced bounded by an a/3(1 11)dis Al-Cu is shown schematically in Figure 8. 12. Ideally, location, and in the latter a coherent one bounded by the structure of the g phase in Al-Cu consists of an a/ 6(11 2). Continued growth of precipitate plates layers of copper on (100) separated by three lay. occurs by either process or a combination of processes ers of aluminium atoms. If a next-nearest neighbour Of course, formation of the final precipitate structure layer of aluminium atoms from the copper layer is requires, in addition to these structural rearrangements. emoved by condensing cy loop, an embry the long-range diffusion of the correct solute atom con- 8 unit cell with Al in the correct AAA... stacking centration to the growing interface. sequence is formed(Figure 8. 12b). Formation of The growth of a second-phase particle with a dis- final CuAl2 8 fluorite structure requires only shuffling parate size or crystal structure relative to the matrix If of the copper atoms into the newly created next- is controlled by two overiding principles-the accom- earest neighbour space and concurrent relaxation of modation of the volume and shape change, and the the Al atoms to the correct 8 interplanar distances optimized use of the available deformation mecha- ure The structural incorporation of vacancies in a pre- are accommodated by vacancy or interstitial conden ipitate is a non-conservative process since atomic sation, or prismatic dislocation loop punching, while sites are eliminated There exist equivalent conserv deviatoric strains are relieved by shear loop pro tive processes in which the new precipitate structure is agation. An example is shown in Figure 8.13. The created from the old by the nucleation and expansion formation of semi-coherent Cu needles in Fe-1%C of partial dislocation loops with predominantly shear is accomplished by the generation of shear loops in 8-CUA 43 4.04A.c=7.68人 04A,c=5.80A Figure 8.12 Schematic diagram showing the transition of e" to 8 in Al-Cu by the vacancy mechanism. vacancies from annealing loops are condensed on a est Al plane from the copper laver in 8" to form the required AAA Al stacki Formation of the 0 fluorite structure then requires only slight redistribution of the copper atom layer and relaxation of the Al yer spacings(courtesy of K H. Westmacott
270 Modem Physical Metallurgy and Materials Engineering addition (0.25Si) modifies the as-quenched dislocation distribution inhibiting the nucleation and growth of dislocation loops and reducing the diameter of helices. The S-precipitate (A12CuMg) is homogeneously nucleated in the presence of Si rather than heterogeneously nucleated at dislocations, and the precipitate grows directly from zones, giving rise to improved and more uniform properties. Apart from speeding up the kinetics of ageing, and providing dislocations nucleation sites, vacancies may play a structural role when they precipitate cooperatively with solute atoms to facilitate the basic atomic arrangements required for transforming the parent crystal structure to that of the product phase. In essence, the process involves the systematic incorporation of excess vacancies, produced by the initial quench or during subsequent dislocation loop annealing, in a precipitate zone or plate to change the atomic stacking. A simple example of 0' formation in A1-Cu is shown schematically in Figure 8.12. Ideally, the structure of the 0" phase in AI-Cu consists of layers of copper on {1 00} separated by three layers of aluminium atoms. If a next-nearest neighbour layer of aluminium atoms from the copper layer is removed by condensing a vacancy loop, an embryonic 0' unit cell with AI in the correct AAA... stacking sequence is formed (Figure 8.12b). Formation of the final CuAI2 0' fluorite structure requires only shuffling half of the copper atoms into the newly created nextnearest neighbour space and concurrent relaxation of the AI atoms to the correct 0' interplanar distances (Figure 8.12c). The structural incorporation of vacancies in a precipitate is a non-conservative process since atomic sites are eliminated. There exist equivalent conservative processes in which the new precipitate structure is created from the old by the nucleation and expansion of partial dislocation loops with predominantly shear character. Thus, for example, the BABAB { 100} plane stacking sequence of the fcc structure can be changed to BAABA by the propagation of an a/2(100) shear loop in the {100} plane, or to BAAAB by the propagation of a pair of a/2(100) partials of opposite sign on adjacent planes. Again, the AAA stacking resulting from the double shear is precisely that required for the embryonic formation of the fluorite structure from the fcc lattice. In visualizing the role of lattice defects in the nucleation and growth of plate-shaped precipitates, a simple analogy with Frank and Shockley partial dislocation loops is useful. In the formation of a Frank loop, a layer of hcp material is created from the fcc lattice by the (non-conservative) condensation of a layer of vacancies in { 111 }. Exactly the same structure is formed by the (conservative) expansion of a Shockley partial loop on a { 111} plane. In the former case a semi-coherent 'precipitate' is produced bounded by an a/3(11 l) dislocation, and in the latter a coherent one bounded by an a/6(112). Continued growth of precipitate plates occurs by either process or a combination of processes. Of course, formation of the final precipitate structure requires, in addition to these structural rearrangements, the long-range diffusion of the correct solute atom concentration to the growing interface. The growth of a second-phase particle with a disparate size or crystal structure relative to the matrix is controlled by two overriding principles-the accommodation of the volume and shape change, and the optimized use of the available deformation mechanisms. In general, volumetric transformation strains are accommodated by vacancy or interstitial condensation, or prismatic dislocation loop punching, while deviatoric strains are relieved by shear loop propagation. An example is shown in Figure 8.13. The formation of semi-coherent Cu needles in Fe-1%Cu is accomplished by the generation of shear loops in e'-CuAI 3 e'-CuAI 2 a = 4.04~, c = 7.68A a = 4.04,&,, c = 5.80A f Figure 8.12 Schematic diagram showing the transition of 0" to O' in AI-Cu by the vacancy mechanism. Vacancies from annealing loops are condensed on a next-nearest A! plane from the copper laver in 0" to form the req.ired AAA AI stacking. Formation of the 0' fluorite structure then requires only slight redistribution of the copper atom layer and relaxation of the AI layer spacings (courtesy of K. H. Westmacott)
which have very good strength/weight ratio applica Ring and nickel alloys also develop better properties The basic idea of all heat-treatments is to * seed a uniform distribution of stable nuclei at the low temperature which can then be grown to optimum ze at the his ture. In most alloys there is a critical temperature Te above which homogeneous nucleation of precipitate does not take plac in some instances has been identified with zone solvus. On ageing above Te there is a critical zone size above which the zones are able act as nuclei for precipitates and below which the zones dissolve In general. the behaviour of Al-Zn-Mg alloys can be divide three classes which ca defined by the temp 1. Alloys quenched and aged above the gp zone solvus (i.e. the temperature above which the zones 155C in a typica Figure 8.13 The formation of semicoherent Ct ver formed during heat treatment, there are no Fe-1% Cu(courtesy of K. H. Westacot asy nuclei for subsequent precipitation and a very parse dispersion of precipitates results with nucle ion principally on dislocations. the precipitate/matrix interface. Expansion of the loops 2. Alloys quenched and aged below the GP zone solvus. GP zones form continuously and grow to tate interfaces leads to a complete network of disloc a size at which they are able to transform to pre- tions interconnecting the precipitates cipitates. The transformation will occur rather more slowly in the grain boundary regions due to the lower vacancy concentration there but since age In non-ferrous heat-treatment there is considerable ing will always be below the gP zone solvus, no interest in double (or duplex) ageing treatments to PFZ is formed other than a very small (30 nm) btain the best microstructure consistent with olute-denuded zone due to precipitation in the num properties. It is now realized that it is unlikely grain boundary that the optimum properties will be produced in alloys 3. Alloys quenched below the GP zone solvus and of the precipitation-hardening type by a single quench aged above it(e. g quenched to room temperature nd ageing treatment. For example, while the interior and aged at 180 C for Al-Zn -Mg). This is the most of grains may develop an acceptable precipitate size common practical situation. The final dispersion of nd density, in the neighbourhood of efficient vacancy precipitates and the Pfz width are controlled by the sinks, such as grain boundaries, a precipitate free zone nucleation treatment below 155C where GP zone (PFZ) is formed which is often associated with over size distribution is determined. A long nucleation ageing in the boundary itself. This heterogeneou reatment gives a fine dispersion of precipitates and structure s rise to poor properties, particularly under stress corrosion conditions Duplex ageing treatments have been used It is possible to stabilize GP zones by addition of come this difficulty. In Al-Zn-Mg, for example it trace elements. These have the same effect as raising was found that storage at room temperature before Te, so that alloys are effectively aged below T.One heating to the ageing temperature leads to xample is Ag to Al-Zn-Mg which raises Te from tion of finer precipitate structure and better 155C to 185C, another is Si to Al-Cu-Mg, another This is just one special example of two-step Cu to Al-Mg-Si and yet another Cd or Sn to Al-Cu geing treatments which have commercial alloys. It is then possible to get uniform distribution and have been found to be applicable to several alloys. and optimum properties by single ageing, and is an Duplex ageing gives better competitive mechanical oroperties in Al-alloys(e.g. Al-Zn-Mg alloys) with be done with physics during multiple ageing much enhanced corrosion resistance since the grain it is best to alter the chemistry or to change the physics boundary zone is removed. It is possible to obtain for a given alloy usually depends on other factors(e. g strengths of 267-308 MN/m2 in Mg-Zn-Mn alloys economics
Strengthening and toughening 271 Figure 8.13 The formation of semicoherent Cu needles in Fe- 1% Cu (courtesy of K. H. Westacott). the precipitate/matrix interface. Expansion of the loops into the matrix and incorporation into nearby precipitate interfaces leads to a complete network of dislocations interconnecting the precipitates. 8.2.5 Duplex ageing In non-ferrous heat-treatment there is considerable interest in double (or duplex) ageing treatments to obtain the best microstructure consistent with optimum properties. It is now realized that it is unlikely that the optimum properties will be produced in alloys of the precipitation-hardening type by a single quench and ageing treatment. For example, while the interior of grains may develop an acceptable precipitate size and density, in the neighbourhood of efficient vacancy sinks, such as grain boundaries, a precipitate-free zone (PFZ) is formed which is often associated with overageing in the boundary itself. This heterogeneous structure gives rise to poor properties, particularly under stress corrosion conditions. Duplex ageing treatments have been used to overcome this difficulty. In A1-Zn-Mg, for example, it was found that storage at room temperature before heating to the ageing temperature leads to the formation of finer precipitate structure and better properties. This is just one special example of two-step or multiple ageing treatments which have commercial advantages and have been found to be applicable to several alloys. Duplex ageing gives better competitive mechanical properties in Al-alloys (e.g. AI-Zn-Mg alloys) with much enhanced corrosion resistance since the grain boundary zone is removed. It is possible to obtain strengths of 267-308 MN/m 2 in Mg-Zn-Mn alloys which have very good strength/weight ratio applications, and nickel alloys also develop better properties with multiple ageing treatments. The basic idea of all heat-treatments is to 'seed' a uniform distribution of stable nuclei at the low temperature which can then be grown to optimum size at the higher temperature. In most alloys, there is a critical temperature Tc above which homogeneous nucleation of precipitate does not take place, and in some instances has been identified with the GP zone solvus. On ageing above Tc there is a certain critical zone size above which the zones are able to act as nuclei for precipitates and below which the zones dissolve. In general, the ageing behaviour of AI-Zn-Mg alloys can be divided into three classes which can be defined by the temperature ranges involved: 1. Alloys quenched and aged above the GP zone solvus (i.e. the temperature above which the zones dissolve, which is above ~155~ in a typical AI-Zn-Mg alloy). Then, since no GP zones are ever formed during heat treatment, there are no easy nuclei for subsequent precipitation and a very coarse dispersion of precipitates results with nucleation principally on dislocations. 2. Alloys quenched and aged below the GP zone solvus. GP zones form continuously and grow to a size at which they are able to transform to precipitates. The transformation will occur rather more slowly in the grain boundary regions due to the lower vacancy concentration there but since ageing will always be below the GP zone solvus, no PFZ is formed other than a very small ('-~30 nm) solute-denuded zone due to precipitation in the grain boundary. 3. Alloys quenched below the GP zone solvus and aged above it (e.g. quenched to room temperature and aged at 180~ for AI-Zn-Mg). This is the most common practical situation. The final dispersion of precipitates and the PFZ width are controlled by the nucleation treatment below 155~ where GP zone size distribution is determined. A long nucleation treatment gives a fine dispersion of precipitates and a narrow PFZ. It is possible to stabilize GP zones by addition of trace elements. These have the same effect as raising Tc, so that alloys are effectively aged below Tc. One example is Ag to AI-Zn-Mg which raises Tc from 155~ to 185~ another is Si to Al-Cu-Mg, another Cu to AI-Mg-Si and yet another Cd or Sn to Al-Cu alloys. It is then possible to get uniform distribution and optimum properties by single ageing, and is an example of achieving by chemistry what can similarly be done with physics during multiple ageing. Whether it is best to alter the chemistry or to change the physics for a given alloy usually depends on other factors (e.g. economics)