In this deformation state it can be observed that only localised necking occurs (*AB*¯) and there is no reduction in width over the whole specimen length (*BB*¯) which would be denoted as

diffuse necking [6]. This is **in contrast** to most metallic materials, where localised throughthickness necking arises e.g. on aluminium alloy EN AW-5xxx in an annealed condition, but not over the whole length. As a consequence, a pattern of localised through-thickness striations can be observed oriented at around ± 55° to the direction of the principal normal stress (Figure 2 a)) [2] [3]: these are also known as PLC bands. Figure 2 b) illustrates a formation of this localised affect which is co-occurring with shear deformations in several planes. The condition

a **55°**

**A**

**Figure 1.** Fracture elongation for aluminium wrought alloys of the EN AW-5xxx, EN AW-6xxx and EN AW-7xxx series at T=298 K and deep-drawing steels at T=298 K. Yield strength in an optimised hardened condition for all mentioned groups at T=298 K.

the EN AW-6xxx group. So, ways of enhancing the fracture elongation of wrought aluminium alloys to levels that are at least comparable with those of deep-drawing steels and additionally possess corresponding strength values are sought. Forming aluminium alloys below room temperature would be an approach, where strain hardening losses and precipitation processes are not an issue. In this regard there is no danger of a reduction in strength through the mechanism of recovery, recrystallization and aging which is in most cases induced by forming at higher temperatures. Furthermore, in previous investigations during tensile testing it was observed that in certain aluminium alloys, which were present in various heat treatment conditions, the Lüders and the Portevin-LeChatelier (PLC) effect occurs at room temperature [2] [3]. This PLC effect not only causes a bad surface quality on sheet metals, but also creates a reduction of the uniform elongation and fracture elongation [2]. It is expected that in tensile testing at low temperatures the negative impact caused by the PLC effect can be reduced and therefore the values of uniform elongation and fracture elongation can be increased.

#### **2. Portevin-LeChatelier (PLC) Effect**

#### **2.1. Phenomena of the PLC effect**

Considering a tensile testing specimen, consisting of an isotropic ductile material, it can be observed that straining in the direction of the principal normal stress is accompanied by a cross-sectional reduction in width and thickness. In the region of uniform elongation the prismatic geometry of the specimen remains, since a *random localisation* stabilises itself according to the Considère-criteria. Leaving this region of the stress-strain curve and moving into the region which signifies the region of necking behaviour a localised and random reduction in width and thickness results in an overall reduction in cross-section. It is of note that for a tensile specimen, where the width is much greater than the thickness, a reduction in the cross-section only appears at a volume of material inclined at an angle of approximately 55° measured in direction of the principal normal stress. Localised necking predominantly in thickness will finally fracture the specimen when constricted [4] [5]. Testing an annealed aluminium alloy with the designation of EN AW-5xxx shows that a cross-sectional reduction occurs in fact before reaching the necking area. Figure 2 a) shows localised necking over the whole specimen length.

the EN AW-6xxx group. So, ways of enhancing the fracture elongation of wrought aluminium alloys to levels that are at least comparable with those of deep-drawing steels and additionally possess corresponding strength values are sought. Forming aluminium alloys below room temperature would be an approach, where strain hardening losses and precipitation processes are not an issue. In this regard there is no danger of a reduction in strength through the mechanism of recovery, recrystallization and aging which is in most cases induced by forming at higher temperatures. Furthermore, in previous investigations during tensile testing it was observed that in certain aluminium alloys, which were present in various heat treatment conditions, the Lüders and the Portevin-LeChatelier (PLC) effect occurs at room temperature [2] [3]. This PLC effect not only causes a bad surface quality on sheet metals, but also creates a reduction of the uniform elongation and fracture elongation [2]. It is expected that in tensile testing at low temperatures the negative impact caused by the PLC effect can be reduced and

**Figure 1.** Fracture elongation for aluminium wrought alloys of the EN AW-5xxx, EN AW-6xxx and EN AW-7xxx series at T=298 K and deep-drawing steels at T=298 K. Yield strength in an optimised hardened condition for all mentioned

therefore the values of uniform elongation and fracture elongation can be increased.

Considering a tensile testing specimen, consisting of an isotropic ductile material, it can be observed that straining in the direction of the principal normal stress is accompanied by a

**2. Portevin-LeChatelier (PLC) Effect**

**2.1. Phenomena of the PLC effect**

groups at T=298 K.

62 Light Metal Alloys Applications

**Figure 2.** a) Cross-sectional reduction of EN AW-5xxx specimen in annealed condition which shows localised necking at an angle of around 55° measured in direction of principal normal stress before reaching the necking area. b) Sche‐ matic cross-sectional reduction in thickness co-occurring with shear deformations in several planes.

In this deformation state it can be observed that only localised necking occurs (*AB*¯) and there is no reduction in width over the whole specimen length (*BB*¯) which would be denoted as diffuse necking [6]. This is **in contrast** to most metallic materials, where localised throughthickness necking arises e.g. on aluminium alloy EN AW-5xxx in an annealed condition, but not over the whole length. As a consequence, a pattern of localised through-thickness striations can be observed oriented at around ± 55° to the direction of the principal normal stress (Figure 2 a)) [2] [3]: these are also known as PLC bands. Figure 2 b) illustrates a formation of this localised affect which is co-occurring with shear deformations in several planes. The condition of the volume of material where localised necking arises is that no strain be imposed on the non-deforming material adjacent to the flowing section which leads to dεw´=0 as illustrated in Figure 3 [6]. The localising volume only deforms in thickness (dεt ) and lengthways (dεl ´), which characterises a biaxial strain state. Provided that there is an ideal plastic material behaviour and a negligible elastic fraction compared with the plastic fraction, Lévy-von Mises equations can be applied.

$$\frac{d\,\varepsilon\_{l}}{d\sigma\_{1}\cdot\frac{1}{2}\,\bullet\,(\sigma\_{2}\cdot\sigma\_{3})} = \frac{d\,\varepsilon\_{w}}{\sigma\_{2}\cdot\frac{1}{2}\,\bullet\,(\sigma\_{3}\cdot\sigma\_{1})} = \frac{d\,\varepsilon\_{l}}{\sigma\_{3}\cdot\frac{1}{2}\,\bullet\,(\sigma\_{1}\cdot\sigma\_{2})}\tag{1}$$

let σ1≠ 0 and σ2=σ3=0 then

$$\frac{d\,\,\varepsilon\_{l}}{\,\,\,\sigma\_{1}} = \frac{d\,\,\varepsilon\_{w}}{\,\,\,\frac{1}{2}\,\bullet\,\,\sigma\_{1}} = \frac{d\,\,\varepsilon\_{l}}{\,\,\,\,\frac{1}{2}\,\bullet\,\,\sigma\_{1}}\tag{2}$$

Thus, the angle between the direction of the localised neck and the direction of the principal normal stress is ideally 54°44´ which is in accordance with the hypothesis of Hill [7]. Eventual

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The microscopic background of the PLC effect is still a subject of discussion. This is because the PLC effect appears at temperatures where interstitial or substitutional diffusion of solute alloying elements do not offer sufficient mobility by bulk diffusion. A brief introduction of the development of ideas concerning the PLC effect is illustrated in [8] [9]. The first model developed by Cottrell and Bilby [10] assumed that the mobility of solute atoms is actually high enough to follow mobile dislocations. Subsequently, McCormick recognised that the mobility of solute atoms is barely sufficient to follow mobile dislocations, and he postulated that clustering mobile dislocations only occurs if they are arrested by obstacles such as by forest dislocations, for example [11]. Van den Beukel presupposed in his model, which was based on McCormick's work, that arresting mobile dislocations at forest dislocations is only caused by volume diffusion of solute atoms [12]. Following a proposal of Sleeswyk [13], the model created by Mulford [14] not only postulates the arresting of the stopped dislocation by volume

diffusion, but additionally by a simultaneously running pipe diffusion (see Figure 4).

**Figure 4.** Substitutional solute atoms at stationary forest dislocations generate Cottrell-atmospheres, whose atoms

A requirement for such a condition is a prior formation of a Cottrell-atmosphere at stationary forest dislocations. The pipe diffusion occurs without the acceleration caused by the vacancy mechanism. Modelling the pipe diffusion of the substitutionally dissolved magnesium in the

cluster at a mobile dislocation which meets a forest dislocation via pipe diffusion.

variations can be attributed to textures of the material.

**2.2. Microscopic background of the PLC-effect**

If only dεl and dεw are considered (because of a small sheet thickness), the expression can be rewritten as

$$d\,\varepsilon\_l = -2\bullet d\,\varepsilon\_w \tag{3}$$

Since dεw< 0 a Mohr´s circle of strain can be developed as illustrated in Figure 3. Applying a coordinate transformation it can be shown that α=70°32´if dεw´=0.

**Figure 3.** Cross-sectional reduction of a flat specimen according Mohr´s circle of strain

Thus, the angle between the direction of the localised neck and the direction of the principal normal stress is ideally 54°44´ which is in accordance with the hypothesis of Hill [7]. Eventual variations can be attributed to textures of the material.

#### **2.2. Microscopic background of the PLC-effect**

of the volume of material where localised necking arises is that no strain be imposed on the non-deforming material adjacent to the flowing section which leads to dεw´=0 as illustrated in

characterises a biaxial strain state. Provided that there is an ideal plastic material behaviour and a negligible elastic fraction compared with the plastic fraction, Lévy-von Mises equations

> = *d εt* - 1 2 ∙ *σ*1

Since dεw< 0 a Mohr´s circle of strain can be developed as illustrated in Figure 3. Applying a

= *d εt σ*3 - 1

and dεw are considered (because of a small sheet thickness), the expression can be

2 ∙ (*σ*1 + *σ*2)

*dεl* =-2∙*dεw* (3)

) and lengthways (dεl

(1)

´), which

(2)

Figure 3 [6]. The localising volume only deforms in thickness (dεt

= *d εw σ*2 - 1

*d εl σ*1 = *d εw* - 1 2 ∙ *σ*1

coordinate transformation it can be shown that α=70°32´if dεw´=0.

**Figure 3.** Cross-sectional reduction of a flat specimen according Mohr´s circle of strain

2 ∙ (*σ*3 + *σ*1)

*d εl*

2 ∙ (*σ*2 + *σ*3)

*σ*1 - 1

can be applied.

64 Light Metal Alloys Applications

If only dεl

rewritten as

let σ1≠ 0 and σ2=σ3=0 then

The microscopic background of the PLC effect is still a subject of discussion. This is because the PLC effect appears at temperatures where interstitial or substitutional diffusion of solute alloying elements do not offer sufficient mobility by bulk diffusion. A brief introduction of the development of ideas concerning the PLC effect is illustrated in [8] [9]. The first model developed by Cottrell and Bilby [10] assumed that the mobility of solute atoms is actually high enough to follow mobile dislocations. Subsequently, McCormick recognised that the mobility of solute atoms is barely sufficient to follow mobile dislocations, and he postulated that clustering mobile dislocations only occurs if they are arrested by obstacles such as by forest dislocations, for example [11]. Van den Beukel presupposed in his model, which was based on McCormick's work, that arresting mobile dislocations at forest dislocations is only caused by volume diffusion of solute atoms [12]. Following a proposal of Sleeswyk [13], the model created by Mulford [14] not only postulates the arresting of the stopped dislocation by volume diffusion, but additionally by a simultaneously running pipe diffusion (see Figure 4).

**Figure 4.** Substitutional solute atoms at stationary forest dislocations generate Cottrell-atmospheres, whose atoms cluster at a mobile dislocation which meets a forest dislocation via pipe diffusion.

A requirement for such a condition is a prior formation of a Cottrell-atmosphere at stationary forest dislocations. The pipe diffusion occurs without the acceleration caused by the vacancy mechanism. Modelling the pipe diffusion of the substitutionally dissolved magnesium in the aluminium bulk showed that the effectiveness of this diffusion process is not high enough to explain the experimental results [9]. As a consequence of this model the diffusion velocity of magnesium atoms is too slow to form Cottrell-atmospheres around the core of the arrested dislocation during representative times. Considering these results Picu [8] developed a new model which is based on the estimation that a forest dislocation embedded in a Cottrellatmosphere reacts with a mobile dislocation by forming *dislocation jogs*. The higher the intensity of Cottrell-atmosphere around a forest dislocation the greater the energy necessary for forming jogs.

For the purpose of evaluating potential strain rate dependencies the uniaxial tensile tests for the aluminium alloys EN AW-6016 and EN AW-7021 in the solid solution heat treated and naturally aged condition T4, and the EN AW-5182 aluminium alloy in the strain hardened condition H111 (as currently implemented in industrial applications), were carried out with the following three forming velocities; v1 =200*mm* / *min* (maximum velocity of the tensile

Mechanical Behaviour of Commercial Aluminium Wrought Alloys at Low Temperatures

( 1 10 v2

*ε*1(˙ v1) =6.60×10-2 s-1,*ε*2(˙ v2) =1.70×10-2 s-1 and *ε*3(˙ v3) =0.17×10-2 s-1 respectively (see Table 2). By way of comparison, the highest strain rate illustrated, v1 , approaches the forming velocity used in a typical industrial process. Determining the effect of precipitations for the 6000 and

30 minutes after a solid solution heat treatment (W30 condition). As a reference the aluminium

a potential influence of the strain rates, the reference aluminium alloy was also tested with the

EN AW-1050A H14 X X EN AW-5182 H111 X X X

Flat sheet specimens were CNC machined showing a test section width of 12.5mm in accord‐ ance with DIN EN ISO 6892-1 (A50) [16]. Attention was given towards the rolling direction of the sheet metal which was perpendicular to the applied load. The length of the 50mm wide clamp sections of the specimens were designed to fit with the requirements of the cryogenic

A specially designed cryogenic dewar (Figure 5) was used to obtain and maintain the sub-zero temperatures required in these tests. A variation of testing temperatures was reached by way of a *medium-pipe*, which was used as a vessel carrying isothermal media, and was cooled by liquid nitrogen. The isothermal media used for the experiment shown in Table 3 were first precooled in an external chamber. The final temperature was reached in the cryogenic dewar, where the specimen remained for around 5 minutes, before the test started. Furthermore,

W30 X

W30 X

) =5*mm* / *min* resulting in strain rates

**strain rates** 6.60 . 10-2 s-1 1.70 . 10-2 s-1 0.17 . 10-2 s-1

T4 X X X

T4 X X X

˙ were carried out within

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67

˙ . To be thorough regarding

) =50*mm* / *min* and v3

7000 series a number of subsequent tests with the highest strain rate *ε*1

alloy EN AW-1050A was tested in half-hard condition (H14) with *ε*1

testing machine), v2

lowest strain rate (*ε*3

EN AW-6016

EN AW-7021

set-up.

( 1 4 v1

˙ )

**Aluminium alloy condition**

**Table 2.** Different strain rates used in the experiment

#### **3. Experimental methods**

A set of experimental tests were conducted based around a tensile test machine. Four types of wrought aluminium alloys were examined which are detailed in Table 1. In this table the sheet thickness and detailed composition are shown.

**Table 1.** Heat treatment, thickness and chemical composition of tested aluminium alloys

For the purpose of evaluating potential strain rate dependencies the uniaxial tensile tests for the aluminium alloys EN AW-6016 and EN AW-7021 in the solid solution heat treated and naturally aged condition T4, and the EN AW-5182 aluminium alloy in the strain hardened condition H111 (as currently implemented in industrial applications), were carried out with the following three forming velocities; v1 =200*mm* / *min* (maximum velocity of the tensile testing machine), v2 ( 1 4 v1 ) =50*mm* / *min* and v3 ( 1 10 v2 ) =5*mm* / *min* resulting in strain rates *ε*1(˙ v1) =6.60×10-2 s-1,*ε*2(˙ v2) =1.70×10-2 s-1 and *ε*3(˙ v3) =0.17×10-2 s-1 respectively (see Table 2). By way of comparison, the highest strain rate illustrated, v1 , approaches the forming velocity used in a typical industrial process. Determining the effect of precipitations for the 6000 and 7000 series a number of subsequent tests with the highest strain rate *ε*1 ˙ were carried out within 30 minutes after a solid solution heat treatment (W30 condition). As a reference the aluminium alloy EN AW-1050A was tested in half-hard condition (H14) with *ε*1 ˙ . To be thorough regarding a potential influence of the strain rates, the reference aluminium alloy was also tested with the lowest strain rate (*ε*3 ˙ )

aluminium bulk showed that the effectiveness of this diffusion process is not high enough to explain the experimental results [9]. As a consequence of this model the diffusion velocity of magnesium atoms is too slow to form Cottrell-atmospheres around the core of the arrested dislocation during representative times. Considering these results Picu [8] developed a new model which is based on the estimation that a forest dislocation embedded in a Cottrellatmosphere reacts with a mobile dislocation by forming *dislocation jogs*. The higher the intensity of Cottrell-atmosphere around a forest dislocation the greater the energy necessary for forming

A set of experimental tests were conducted based around a tensile test machine. Four types of wrought aluminium alloys were examined which are detailed in Table 1. In this table the sheet

**Condition** H14 H111 W30– T4

**Specimen thickness [mm]** 0,95 1,15 1,15 0,95

**Si [wt.-%]** 0,25 0,20 1,20 0,07 **Fe [wt.-%]** 0,40 0,35 0,40 0,16 **Cu [wt.-%]** 0,05 0,15 0,15 0,12 **Mn [wt.-%]** 0,05 0,35 0,15 0,02 **Mg [wt.-%]** 0,05 4,40 0,40 - **Cr [wt.-%]** - 0,10 0,10 - **Zn [wt.-%]** 0,07 0,25 0,15 6,24 **Ti [wt.-%]** 0,05 0,10 0,15 0,02 **Zr [wt.-%]** - - - 0,14

H111 slightly strainhardened (less than H11) – no thermal treatment

H14 strain-hardened to a half hard condition - no thermal treatment

**Table 1.** Heat treatment, thickness and chemical composition of tested aluminium alloys

**EN AW-1050A EN AW-5182 EN AW-6016 EN AW-7021**

T4 solution heat-treated at 465°C for 20 minutes and naturally aged to a substantially stable condition W30 solution heat-treated at 465°C for 20 minutes and tested within 30 minutes

jogs.

66 Light Metal Alloys Applications

**3. Experimental methods**

**Temper** [15]

thickness and detailed composition are shown.

Flat sheet specimens were CNC machined showing a test section width of 12.5mm in accord‐ ance with DIN EN ISO 6892-1 (A50) [16]. Attention was given towards the rolling direction of the sheet metal which was perpendicular to the applied load. The length of the 50mm wide clamp sections of the specimens were designed to fit with the requirements of the cryogenic set-up.

A specially designed cryogenic dewar (Figure 5) was used to obtain and maintain the sub-zero temperatures required in these tests. A variation of testing temperatures was reached by way of a *medium-pipe*, which was used as a vessel carrying isothermal media, and was cooled by liquid nitrogen. The isothermal media used for the experiment shown in Table 3 were first precooled in an external chamber. The final temperature was reached in the cryogenic dewar, where the specimen remained for around 5 minutes, before the test started. Furthermore, homogeneity of the temperature within the medium-pipe was accomplished by injecting inert gas. In case of a sudden temperature drop the chamber included an inlet for compressed air, which ensured a replacement of the cold environment between medium-pipe and cryogenic dewar. The investigations were conducted using a type Z100 tensile testing machine from Zwick.

6

Medium

**Table 3.** Application temperature for different media **[17]**

**4. Results**

elongation *ef*

approaches 50% at 77K.

**Results**

dry ice paraffin oil acetone isobutyl alcohol n‐pentane methylcyclopentane isopentane liquid nitrogen liquid hydrogen liquid helium

8

Temperature range of fluid consistence ( temperatures in K) 243 193 178 165 143 131 113 77 20 4

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69

Figure 6 a) through f) shows for the aluminium alloys tested the temperature dependency of the yield strength *σy*, ultimate tensile strength (UTS),

Figure 6 a) through f) shows for the aluminium alloys tested the temperature dependency of the yield strength *σy*, ultimate tensile strength (UTS), uniform elongation *eu* and fracture

It can be seen that with one exception and for all materials the yield strength, ultimate tensile strength and elongation rises consistently when temperature is decreased. The strength values of the aluminium alloys in W30 condition are lower than in T4 condition, which is as expected. For the aluminium alloys EN AW-5182 and EN AW-7021 in W30 condition a uniform elongation of over 40% and a fracture elongation of considerably more than 50% can be reached at 77K. A similar trend can be observed for the aluminium alloy EN AW-6016 in both heat treatment conditions, where the elongation in T4 condition is less than in W30 condition. For a decrease in temperature the aluminium alloy EN AW-1050A shows a rather shallow rise of the ultimate and yield strengths over the whole temperature range; whereas curves of uniform and fracture percentage elongation of this alloy indicate significant increases in the magnitudes of their gradients. The strength and elongation of the aluminium alloys EN AW-7021 also show a growing trend if temperature is reduced. For both heat treatment conditions, T4 and W30, a gradual elongation rise of 25% at 152 K can be seen. If the temperature is lower than 152 K the elongation of the peak aged aluminium (T4) alloy decreases again. In turn, when present in W30 condition, the uniform elongation rises to more than 40% and the fracture elongation

which was determined at a strain rate of 6.60×10-2 s-1.

uniform elongation *eu* and fracture elongation *ef* which was determined at a strain rate ofǤͲ ൈ ͳͲିଶ s-1.

**Table 3.** Application temperature for different media [17]

**Figure 5.** Cryogenic set-up

**Figure 5.** Cryogenic set-up

Mechanical Behaviour of Commercial Aluminium Wrought Alloys at Low Temperatures http://dx.doi.org/10.5772/58362 69

**Table 3.** Application temperature for different media **[17] Table 3.** Application temperature for different media [17]

#### **4. Results**

**Results**

6

homogeneity of the temperature within the medium-pipe was accomplished by injecting inert gas. In case of a sudden temperature drop the chamber included an inlet for compressed air, which ensured a replacement of the cold environment between medium-pipe and cryogenic dewar. The investigations were conducted using a type Z100 tensile testing machine from

1) Specimen

6) Gas inlet 7) Air inlet

2) Cryogenic dewar 3) Medium-pipe 4) Nitrogen filling 5) Medium filling

Zwick.

68 Light Metal Alloys Applications

**Figure 5.** Cryogenic set-up

**Figure 5.** Cryogenic set-up

Figure 6 a) through f) shows for the aluminium alloys tested the temperature dependency of the yield strength *σy*, ultimate tensile strength (UTS), uniform elongation *eu* and fracture elongation *ef* which was determined at a strain rate ofǤͲ ൈ ͳͲିଶ s-1. Figure 6 a) through f) shows for the aluminium alloys tested the temperature dependency of the yield strength *σy*, ultimate tensile strength (UTS), uniform elongation *eu* and fracture elongation *ef* which was determined at a strain rate of 6.60×10-2 s-1.

8

It can be seen that with one exception and for all materials the yield strength, ultimate tensile strength and elongation rises consistently when temperature is decreased. The strength values of the aluminium alloys in W30 condition are lower than in T4 condition, which is as expected. For the aluminium alloys EN AW-5182 and EN AW-7021 in W30 condition a uniform elongation of over 40% and a fracture elongation of considerably more than 50% can be reached at 77K. A similar trend can be observed for the aluminium alloy EN AW-6016 in both heat treatment conditions, where the elongation in T4 condition is less than in W30 condition. For a decrease in temperature the aluminium alloy EN AW-1050A shows a rather shallow rise of the ultimate and yield strengths over the whole temperature range; whereas curves of uniform and fracture percentage elongation of this alloy indicate significant increases in the magnitudes of their gradients. The strength and elongation of the aluminium alloys EN AW-7021 also show a growing trend if temperature is reduced. For both heat treatment conditions, T4 and W30, a gradual elongation rise of 25% at 152 K can be seen. If the temperature is lower than 152 K the elongation of the peak aged aluminium (T4) alloy decreases again. In turn, when present in W30 condition, the uniform elongation rises to more than 40% and the fracture elongation approaches 50% at 77K.

**Figure 6.** Temperature dependence of the yield strength σ*y*, ultimate tensile strength UTS, uniform elongation *eu* and fracture elongation *ef* which was determined at a strain rate of 6.60. 10-2 s-1.

The influence of strain rates regarding strength values can be considered as small (see Figure 7 a), c), e) & g)). A slightly more distinct effect can be observed for the elongation values, especially for EN AW-5182 and EN AW-6016 T4. Nevertheless, those variations converge as the temperature approaches 77K (see Figure 7 b), d), f)&g). 7

11

**Figure 7.** Temperature dependency of yield strength, ultimate strength and fracture elongation at different strain

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71

rates (black icon 6.60 × 10-2 s-1, grey icon 1.70 × 10-2 s-1, white icon 0.17 × 10-2 s-1)

Mechanical Behaviour of Commercial Aluminium Wrought Alloys at Low Temperatures http://dx.doi.org/10.5772/58362 71

**Figure 7.** Temperature dependency of yield strength, ultimate strength and fracture elongation at different strain rates (black icon 6.60 × 10-2 s-1, grey icon 1.70 × 10-2 s-1, white icon 0.17 × 10-2 s-1)

11

The influence of strain rates regarding strength values can be considered as small (see Figure 7 a), c), e) & g)). A slightly more distinct effect can be observed for the elongation values, especially for EN AW-5182 and EN AW-6016 T4. Nevertheless, those variations converge as

**Figure 6.** Temperature dependence of the yield strength σ*y*, ultimate tensile strength UTS, uniform elongation *eu* and

which was determined at a strain rate of 6.60.

fracture elongation *ef*

70 Light Metal Alloys Applications

the temperature approaches 77K (see Figure 7 b), d), f)&g). 7

10-2 s-1.

Some overall observations where made as a result of the tests from stress-strain data. Firstly aluminium alloy EN AW-1050A-H14 showed no pronounced yielding for all strain rates and temperatures. The same effect was observed for EN AW-6016 in both heat treatment condi‐ tions, T4 and W30. In contrast the aluminium alloy EN AW-5182-H111 showed pronounced yielding at all strain rates and temperatures. Aluminium alloy EN AW-7021 in W30 condition indicated pronounced yielding only at 298K and 248K. When the temperature was reduced this effect was not evident. On the other hand, EN AW-7021 in T4 condition showed similar yielding effects for all strain rates as was also observed for EN AW-6016. If one now considers Figure 8, the results of increasing the uniaxial load on an EN AW-5182-H111 specimen can be seen over the elastic and plastic regions. It is clear that there is a pronounced yielding behaviour accompanied by a region of stress oscillations. For the results, shown at T=298K, the amplitude is as high as ± 2 MPa; whereas for T ≤ 248K the magnitude reduces to ± 0.2 MPa.

and uniform elongation *eu*. It seems that the dissolved atoms had not the opportunity to form Cottrell-atmospheres which confirms the assertion of [8]. If the temperature is reduced, the diffusivity of the dissolved atoms is generally decelerated. Therefore, the formation of Cottrellatmospheres, of the same intensity during the same period of time, is unlikely: this would correspond to a reduced PLC-effect. Figure 7 b), d) &f) illustrates this effect in terms of a growth

*n* rises simultaneously as illustrated in Figure 9 a), b), c), d) & f). The overall increase in strength throughout the tested aluminium alloys could be explained *inter alia* by the activation energy which reduces as temperature is lowered. Consequently, the force needed to move a disloca‐ tion through the crystal lattice (Peierls Nabarro force), reaches a maximum as temperature

Ludwik, Hollomon and Backofen identified that the flow stress *Kf* can be elevated by raising the strain *ε* as well as lifting the strain rate *ε*˙ [19]. The general relationship between the physical

A relationship which was empirically developed by Ludwik modified by Hollomon and

where *c* is a material dependent constant (strength coefficient), *n* being the strain-hardening coefficient and *m* represents the strain rate sensitivity parameter. At temperatures where a metallurgical softening takes place to a limited extent (*T*≤ 0.4*Tm*; *Tm*=melting point [K]), an increase of the strain rate will not affect the true stress. Therefore, the *m* value can be assumed

Using this relationship most stress-strain curves can be described accurately in the region of uniform elongation a*t T*≤ 0.4*Tm.* As exemplified in Heiser [5] the strain hardening coefficient *n*

The strain hardening coefficient *n* describes to what extent a material can be deformed before localised necking will start, i.e. the flow curve. Figure 9 shows correlations between experi‐ mental evaluated strain hardening coefficients using Eq. (7) and uniform elongation *eu* in relation to temperature and strain rate of EN AW-1050A, EN AW-5182, EN AW-6016 and EN

to be much smaller than 1. The results in Figure 7 confirm this statement.

can be written as true uniform strain *εu* or related to the uniform elongation *eu*.

and uniform elongation *eu*; thereby, the strain hardening coefficient

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73

*K f* = *K f* (*ε*, *ε*˙) (4)

*K f* =*c* ∙*ε n* ∙*ε*˙ *m* (5)

*K f* =*c* ∙*ε n* (6)

*n* =*εu* =*ln* (1 + *eu*) (7)

in the fracture elongation *ef*

converges to 0 K [4] [18].

extended by Backofen is

If *m* is set to be 0 it follows:

values is:

**Figure 8.** Pronounced yielding and stress amplitudes in the plastic region of EN AW-5182

#### **5. Discussion**

The PLC-effect, caused by Cottrell-atmospheres, offers a plausible mechanism to explain the resulting reduction of fracture elongation *ef* and uniform elongation *eu*, which leads to a decrease in the strain hardening coefficient *n*. The intensity of Cottrell-atmospheres is pro‐ portional to the quantity of dissolved atoms with a sufficiently high diffusivity, and the time to form Cottrell-atmospheres. Thus, in the case of the precipitation hardenable alloy EN AW-6016 and EN AW-7021, it was intended to increase the amount of dissolved atoms via solution heat-treatment, quenching and immediate testing of the mechanical properties at this condition. Comparing the two heat-treatment conditions T4 and W30, it can be seen in Figure 6 c) through f) that this procedure had the expected positive effect on fracture elongation *ef* and uniform elongation *eu*. It seems that the dissolved atoms had not the opportunity to form Cottrell-atmospheres which confirms the assertion of [8]. If the temperature is reduced, the diffusivity of the dissolved atoms is generally decelerated. Therefore, the formation of Cottrellatmospheres, of the same intensity during the same period of time, is unlikely: this would correspond to a reduced PLC-effect. Figure 7 b), d) &f) illustrates this effect in terms of a growth in the fracture elongation *ef* and uniform elongation *eu*; thereby, the strain hardening coefficient *n* rises simultaneously as illustrated in Figure 9 a), b), c), d) & f). The overall increase in strength throughout the tested aluminium alloys could be explained *inter alia* by the activation energy which reduces as temperature is lowered. Consequently, the force needed to move a disloca‐ tion through the crystal lattice (Peierls Nabarro force), reaches a maximum as temperature converges to 0 K [4] [18].

Ludwik, Hollomon and Backofen identified that the flow stress *Kf* can be elevated by raising the strain *ε* as well as lifting the strain rate *ε*˙ [19]. The general relationship between the physical values is:

$$\mathbf{K}\_f = \mathbf{K}\_f \begin{pmatrix} \varepsilon & \dot{\varepsilon} \end{pmatrix} \tag{4}$$

A relationship which was empirically developed by Ludwik modified by Hollomon and extended by Backofen is

$$K\_f = \mathfrak{c} \bullet \mathfrak{e}^n \bullet \dot{\mathfrak{e}}^m \tag{5}$$

where *c* is a material dependent constant (strength coefficient), *n* being the strain-hardening coefficient and *m* represents the strain rate sensitivity parameter. At temperatures where a metallurgical softening takes place to a limited extent (*T*≤ 0.4*Tm*; *Tm*=melting point [K]), an increase of the strain rate will not affect the true stress. Therefore, the *m* value can be assumed to be much smaller than 1. The results in Figure 7 confirm this statement.

If *m* is set to be 0 it follows:

Some overall observations where made as a result of the tests from stress-strain data. Firstly aluminium alloy EN AW-1050A-H14 showed no pronounced yielding for all strain rates and temperatures. The same effect was observed for EN AW-6016 in both heat treatment condi‐ tions, T4 and W30. In contrast the aluminium alloy EN AW-5182-H111 showed pronounced yielding at all strain rates and temperatures. Aluminium alloy EN AW-7021 in W30 condition indicated pronounced yielding only at 298K and 248K. When the temperature was reduced this effect was not evident. On the other hand, EN AW-7021 in T4 condition showed similar yielding effects for all strain rates as was also observed for EN AW-6016. If one now considers Figure 8, the results of increasing the uniaxial load on an EN AW-5182-H111 specimen can be seen over the elastic and plastic regions. It is clear that there is a pronounced yielding behaviour accompanied by a region of stress oscillations. For the results, shown at T=298K, the amplitude

is as high as ± 2 MPa; whereas for T ≤ 248K the magnitude reduces to ± 0.2 MPa.

**Figure 8.** Pronounced yielding and stress amplitudes in the plastic region of EN AW-5182

The PLC-effect, caused by Cottrell-atmospheres, offers a plausible mechanism to explain the

decrease in the strain hardening coefficient *n*. The intensity of Cottrell-atmospheres is pro‐ portional to the quantity of dissolved atoms with a sufficiently high diffusivity, and the time to form Cottrell-atmospheres. Thus, in the case of the precipitation hardenable alloy EN AW-6016 and EN AW-7021, it was intended to increase the amount of dissolved atoms via solution heat-treatment, quenching and immediate testing of the mechanical properties at this condition. Comparing the two heat-treatment conditions T4 and W30, it can be seen in Figure 6 c) through f) that this procedure had the expected positive effect on fracture elongation *ef*

and uniform elongation *eu*, which leads to a

**5. Discussion**

72 Light Metal Alloys Applications

resulting reduction of fracture elongation *ef*

$$K\_f = \mathfrak{c} \bullet \mathfrak{s}^n \tag{6}$$

Using this relationship most stress-strain curves can be described accurately in the region of uniform elongation a*t T*≤ 0.4*Tm.* As exemplified in Heiser [5] the strain hardening coefficient *n* can be written as true uniform strain *εu* or related to the uniform elongation *eu*.

$$
\Delta \eta = \varepsilon\_u = \ln\left(1 + e\_u\right) \tag{7}
$$

The strain hardening coefficient *n* describes to what extent a material can be deformed before localised necking will start, i.e. the flow curve. Figure 9 shows correlations between experi‐ mental evaluated strain hardening coefficients using Eq. (7) and uniform elongation *eu* in relation to temperature and strain rate of EN AW-1050A, EN AW-5182, EN AW-6016 and EN AW-7021 specimens. In addition, the Ludwik curve is included which represents *n* calculated using Eq. (7).

**Figure 9.** Temperature and strain rate dependence of strain hardening exponent n and uniform strain ε*u*. Additionally included is the dependence of strain hardening coefficient and uniform elongation after Ludwik if only strain harden‐ ing is considered

23

Figure 9 a) through f) shows only a slight strain rate dependency. Apart from the temperature range between 298K and 248K and from EN AW-7021 in the precipitation hardened condition (T4), an increase of the strain hardening coefficient and uniform elongation as temperature is reduced can be seen. Except for the aluminium alloy 1050A, (Figure 9 a)) which is consistent with the Ludwik curve, tested aluminium alloys indicate a parallel offset. It is assumed that this shift is due to the PLC-effect. Therefore Eq. 6, which takes only the strain hardening process into account, seems not to be applicable for all aluminium alloys. A modification to this equation would seem appropriate to allow for the description of materials which exhibit a PLC-effect. Eq. (8) illustrates a *simple* modification of Eq. (7). The addition of a constant, *p*,

Mechanical Behaviour of Commercial Aluminium Wrought Alloys at Low Temperatures

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75

The modified Ludwik curve is illustrated on Figure 9 b) through f) which incorporates a value of *p*=0.1. Whilst it is clear that the new fit corresponds more appropriately (recalling that the 7021 Figure 9 e) is in a naturally aged condition (T4) and is not in an ideal state for cryogenic forming processes), further analyses of flow curves presented in this and additional studies will provide the basis of a more refined relationship. This will be presented in a subsequent

**Figure 10.** Fracture elongation for aluminium wrought alloys of the EN AW-5xxx, EN AW-6xxx and EN AW-7xxx series at T=77 K and deep-drawing steels at T=298 K. Yield strength in optimised hardened condition for all mentioned

*n* =*ln* (1 + *eu*) + *p* (8)

provides a shift to the Ludwik curve and can now be written

publication.

groups at T=298 K.

Figure 9 a) through f) shows only a slight strain rate dependency. Apart from the temperature range between 298K and 248K and from EN AW-7021 in the precipitation hardened condition (T4), an increase of the strain hardening coefficient and uniform elongation as temperature is reduced can be seen. Except for the aluminium alloy 1050A, (Figure 9 a)) which is consistent with the Ludwik curve, tested aluminium alloys indicate a parallel offset. It is assumed that this shift is due to the PLC-effect. Therefore Eq. 6, which takes only the strain hardening process into account, seems not to be applicable for all aluminium alloys. A modification to this equation would seem appropriate to allow for the description of materials which exhibit a PLC-effect. Eq. (8) illustrates a *simple* modification of Eq. (7). The addition of a constant, *p*, provides a shift to the Ludwik curve and can now be written

AW-7021 specimens. In addition, the Ludwik curve is included which represents *n* calculated

**Figure 9.** Temperature and strain rate dependence of strain hardening exponent n and uniform strain ε*u*. Additionally included is the dependence of strain hardening coefficient and uniform elongation after Ludwik if only strain harden‐

23

using Eq. (7).

74 Light Metal Alloys Applications

ing is considered

$$m = ln \; \begin{pmatrix} 1 + e\_u \end{pmatrix} + p \tag{8}$$

The modified Ludwik curve is illustrated on Figure 9 b) through f) which incorporates a value of *p*=0.1. Whilst it is clear that the new fit corresponds more appropriately (recalling that the 7021 Figure 9 e) is in a naturally aged condition (T4) and is not in an ideal state for cryogenic forming processes), further analyses of flow curves presented in this and additional studies will provide the basis of a more refined relationship. This will be presented in a subsequent publication.

**Figure 10.** Fracture elongation for aluminium wrought alloys of the EN AW-5xxx, EN AW-6xxx and EN AW-7xxx series at T=77 K and deep-drawing steels at T=298 K. Yield strength in optimised hardened condition for all mentioned groups at T=298 K.

#### **6. Conclusion**

The plastic behaviour of commercial aluminium wrought alloys in annealed or solution annealed condition can be positively enhanced if the forming temperature is below 123K. This advantage can be reached without any interference of higher strain rates.

[8] R. C. Picu, "A mechanism for the negative strain-rate sensitivity of dilute solid solu‐

Mechanical Behaviour of Commercial Aluminium Wrought Alloys at Low Temperatures

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77

[9] R. C. Picu, G. Vincze, F. Ozturk, J. J. Gracio, F. Barlat and A. M. Maniatty, "Strain rate sensitivity of the commercial aluminum alloy AA5182-O," Mat. Sci. Eng. A390 (2005)

[10] A. H. Cottrell and B. A. Bilby, "Dislocation theory of yielding and strain ageing of

[11] P. G. Mc Cormick, "A model for the Portevin-Le Chatelier effect in substitutional al‐

[12] A. van den Beukel, "Theory of the effect of dynamic strain ageing on mechanical

[13] A. W. Sleeswyk, "Slow strain hardening of ingot iron," Acta Metall 6 (1958) 598-603. [14] R. A. Mulford and U. F. Kocks, "New observations on the mechanisms of dynamic

[15] DIN EN 515: Aluminium und Aluminiumlegierungen-Halbzeug-Bezeichnung der

[16] DIN EN ISO 6892-1: Zugversuch-Metallische Werkstoffe-Prüfverfahren bei Raum‐

[17] "IFA Institut für Arbeitsschutz der Deutschen Gesetzlichen Unfallversicherung," GESTIS Substance Database, [Online]. Available: http://gestis-en.itrust.de [Accessed

[18] B. Heine, Werkstoffprüfung Ermittlung von Werkstoffeigenschaften, München: Carl

[19] J. K. Gregory, Superplastic deformation in oxide dispersion strengthened nickel base

strain aging and of jerky flow," Acta Metall 27 (1979) 1125-1134.

tions," Acta Mat. 52 (2004) 3447-3458.

iron," Proc. Phys. Soc. A62 (1949) 49-62.

properties," Status Solidi A30 (1975) 197-206.

Werkstoffzustände, Berlin: Beuth Verlag, 1993.

superalloys, Stanford University: Ph.D. Thesis, 1983.

temperatur, Berlin: Beuth Verlag, 2009.

Hanser Verlag, 2011 pp.170-172.

loys," Acta Metall 20 (1972) 351-354.

334-343.

10 02 2014].

Figure 10 demonstrates the fracture elongation values which are currently attainable from aluminium alloys compared with values achievable from deep-drawing steels. A competitive advantage of cryogenic forming of annealed or solution annealed aluminium alloys is clearly shown. Flow curves showing PLC-effects can be described more precisely with addition of a constant p which represents an extension of the initial Ludwik relationship.

#### **Author details**

R. Schneider1,2, B. Heine2 and R.J. Grant1

1 Glyndŵr University, Department of Engineering and Applied Physics, United Kingdom

2 Aalen University of Applied Sciences, Institut für Oberflächen-und Werkstofftechnik, Germany

#### **References**

[8] R. C. Picu, "A mechanism for the negative strain-rate sensitivity of dilute solid solu‐ tions," Acta Mat. 52 (2004) 3447-3458.

**6. Conclusion**

76 Light Metal Alloys Applications

**Author details**

Germany

**References**

R. Schneider1,2, B. Heine2

The plastic behaviour of commercial aluminium wrought alloys in annealed or solution annealed condition can be positively enhanced if the forming temperature is below 123K. This

Figure 10 demonstrates the fracture elongation values which are currently attainable from aluminium alloys compared with values achievable from deep-drawing steels. A competitive advantage of cryogenic forming of annealed or solution annealed aluminium alloys is clearly shown. Flow curves showing PLC-effects can be described more precisely with addition of a

1 Glyndŵr University, Department of Engineering and Applied Physics, United Kingdom

[1] W. S. Miller, L. Zhuang, J. Bottema, A. J. Wittebrood, P. De Smet, A. Haszler and A. Vieregge, "Recent development in aluminium alloys for the automotive industry,"

[2] E. Pink and A. Grinberg, "Praktische Aspekte des Portevin-LeChatelier-Effektes (1),"

[3] E. Pink and A. Grinberg, "Praktische Aspekte des Portevin-LeChatelier-Effektes (2),"

[4] G. E. Dieter, Mechanical Metallurgy, London: McGraw-Hill BookCompany (UK)

[5] M. Heiser and G. Lange, "Scherbruch in Aluminium-Legierungen infolge lokaler

[6] S. P. Keeler and W. A. Backofen, "Plastic instability and fracture in sheets stretched

[7] R. Hill, "On discontinuous plastic states, with special reference to localized necking

2 Aalen University of Applied Sciences, Institut für Oberflächen-und Werkstofftechnik,

advantage can be reached without any interference of higher strain rates.

constant p which represents an extension of the initial Ludwik relationship.

and R.J. Grant1

Mat. Sci. Eng. A280 (2000) 37-49.

Aluminium 60 (1984) 687-691.

Aluminium 60 (1984) 764-768.

plastischer Instabilität," Z. Metallkd. 83 (1992) 115.

over rigid punches," Trans. ASM 56 (1963) 25-48.

in thin sheets," J. Mech. Phys. Sol. 1 (1952) 19-30.

Limited, 1988 pp. 289-291.

**Chapter 4**

**Recrystallization Behavior of Commercial Purity**

Commercial purity aluminium alloys are largely used in the forms of foil for food packaging industries. Aluminium ingots are processed through the multi-operational steps in the consequence of foil preparation. The ingot of 7mm thickness, obtained from either direct chill (DC) casting or continuous casting process, is first cold rolled to reduce the thickness approx‐ imately 0.6 mm followed by an intermediate annealing treatment and then final roll pass is given to produce the foils of desired thickness. Therefore, the intermediate annealing treatment and consequent recrystallization behavior are prime factors for controlling microstructures as well as properties of foils. In general, the recrystallization behavior is the resultant of three steps, viz., deformation, recrystallization and grain growth. It is affected by various impurities

Iron and silicon are the inevitable impurities in commercial purity aluminium alloys. The precipitation reaction due to these impurities easily occurs during heat treatment at the

of the alloys. The type of intermetallic phases formed during solidification and volume fraction, amount and size of the individual particles can be controlled by changing the silicon (Si) and iron (Fe) content, i.e., by adjusting the Fe/Si ratio [1]. The silicon influences the nature of eutectic constituent, solid solubility of other elements, formation of precipitates and dispersoids, and the transformation characteristics of the precipitates [2]. It is necessary to control the Si content of high quality aluminium foils where material characteristics are determined by Fe/Si ratio. If the material is properly processed, then it can only be assumed that the material properties are not actually deteriorated with increased amount of Fe, Si and/ or Mn. In general, composition ratio of the various alloying elements is maintained with respect

C, affecting cold working, softening, and corrosion resistance

© 2014 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

**Aluminium Alloys**

http://dx.doi.org/10.5772/58385

Additional information is available at the end of the chapter

and the precipitates forming at the operational steps.

temperature range of 200 – 600o

Rajat K. Roy

**1. Introduction**

## **Recrystallization Behavior of Commercial Purity Aluminium Alloys**

Rajat K. Roy

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/58385

### **1. Introduction**

Commercial purity aluminium alloys are largely used in the forms of foil for food packaging industries. Aluminium ingots are processed through the multi-operational steps in the consequence of foil preparation. The ingot of 7mm thickness, obtained from either direct chill (DC) casting or continuous casting process, is first cold rolled to reduce the thickness approx‐ imately 0.6 mm followed by an intermediate annealing treatment and then final roll pass is given to produce the foils of desired thickness. Therefore, the intermediate annealing treatment and consequent recrystallization behavior are prime factors for controlling microstructures as well as properties of foils. In general, the recrystallization behavior is the resultant of three steps, viz., deformation, recrystallization and grain growth. It is affected by various impurities and the precipitates forming at the operational steps.

Iron and silicon are the inevitable impurities in commercial purity aluminium alloys. The precipitation reaction due to these impurities easily occurs during heat treatment at the temperature range of 200 – 600o C, affecting cold working, softening, and corrosion resistance of the alloys. The type of intermetallic phases formed during solidification and volume fraction, amount and size of the individual particles can be controlled by changing the silicon (Si) and iron (Fe) content, i.e., by adjusting the Fe/Si ratio [1]. The silicon influences the nature of eutectic constituent, solid solubility of other elements, formation of precipitates and dispersoids, and the transformation characteristics of the precipitates [2]. It is necessary to control the Si content of high quality aluminium foils where material characteristics are determined by Fe/Si ratio. If the material is properly processed, then it can only be assumed that the material properties are not actually deteriorated with increased amount of Fe, Si and/ or Mn. In general, composition ratio of the various alloying elements is maintained with respect

© 2014 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

to each other, but the absolute value of the single alloying element is not critical as long as primary crystallization of coarse intermetallic phases is avoided.

#### **2. Plastic deformation**

#### **2.1. Microstructure of deformed metals**

The cold-rolled alloys are consisting high density of dislocations clearly examined through transmission electron microscopy (TEM). During deformation the dislocations are arranged as cell boundaries whereas some cell interiors are free of dislocations (Fig. 1). Due to heavy deformation, the cell boundaries are diffuse in nature instead of simple dislocation.

**Figure 2.** Schematic diagram of deformation bands, transition bands and kink bands [4]

the cell might be called subgrain.

formed.

**2.3. Deformation texture**

**2.2. The effect of particles on deformed microstructure**

After plastic deformation by slip mechanism, the walls of high dislocation density separate the region of low dislocation density. Such types of microstructure have been referred to as cell or subgrain. In the case of cell structures, the boundary consists of a tangled array of dislocations and appears to be a diffuse in nature. For subgrain, the boundary is sharp and consists of a well ordered dislocation array. The cell structure develops in the bulk material during or after the dislocation movement induced by the applied stress. With increasing strain, the walls become quite sharp and interiors become reasonably free of dislocations. At this stage

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Due to rapid cooling rate in the continuous casting and direct chill casting, the alloying elements remain as supersaturated solid solution. Upon the processing operation of alumini‐ um foils, the increased cold rolling reduction gives rise to a more pronounced dislocation entanglement. It results in the acceleration of both stored energy and precipitation. Therefore, during plastic deformation of two-phase alloys, particles affect the overall dislocation density, inhomogeneity of deformation in the matrix and deformation structure. Subsequently, the recrystallization behaviour of the alloy is also affected due to the influence of driving force as well as nucleation sites for recrystallization. If the particle is strong enough to withstand the applied stress on it, the dislocation then proceeds to encircle the particle and leaves an Orowan loop, otherwise it deforms. When the particle does not deform then extra dislocations are generated at the particle-matrix interface. On the other hand, if the particle deforms either before or after the formation of Orowan loop, no extra dislocations are generated. Generally, smaller particles are weaker than larger particles, which results in weaker slip plane. Move‐ ment of subsequent dislocations occurs on the same plane. As a consequence, slip bands are

The cold rolling texture of aluminium alloy sheets has been characterized as β fiber, which is associated with plane strain deformation [5]. The texture is varied from layer to layer of the

**Figure 1.** Dislocation structures of 92% cold rolled commercial purity aluminium alloy (AA1235 :0.67 Fe, 0.16Si, 0.01 Ti and rest Al) [3]

However, the orientation of plastically deformed alloy changes from grain to grain or in different regions in a same grain. It happens owing to different rotations during the deforma‐ tion by use of a different combination of slip systems to achieve the imposed strain. Therefore, the deformation band is nothing but a volume of constant orientation that is significantly different than the orientation of other parts of the grain, which is explained by a schematic diagram (Fig. 2). It shows the variation of orientation at different regions of a same grain. The orientation of one part (X) of a deformed grain changes rapidly to that of a differently oriented part (Y) of the same grain across a thin boundary of a finite width called the transition bands (T) or microbands. Sometimes the orientation of deformation bands is changed twice, such as from X to Y and then Y to X. This special type of deformation band is called kink band.

**Figure 2.** Schematic diagram of deformation bands, transition bands and kink bands [4]

After plastic deformation by slip mechanism, the walls of high dislocation density separate the region of low dislocation density. Such types of microstructure have been referred to as cell or subgrain. In the case of cell structures, the boundary consists of a tangled array of dislocations and appears to be a diffuse in nature. For subgrain, the boundary is sharp and consists of a well ordered dislocation array. The cell structure develops in the bulk material during or after the dislocation movement induced by the applied stress. With increasing strain, the walls become quite sharp and interiors become reasonably free of dislocations. At this stage the cell might be called subgrain.

#### **2.2. The effect of particles on deformed microstructure**

Due to rapid cooling rate in the continuous casting and direct chill casting, the alloying elements remain as supersaturated solid solution. Upon the processing operation of alumini‐ um foils, the increased cold rolling reduction gives rise to a more pronounced dislocation entanglement. It results in the acceleration of both stored energy and precipitation. Therefore, during plastic deformation of two-phase alloys, particles affect the overall dislocation density, inhomogeneity of deformation in the matrix and deformation structure. Subsequently, the recrystallization behaviour of the alloy is also affected due to the influence of driving force as well as nucleation sites for recrystallization. If the particle is strong enough to withstand the applied stress on it, the dislocation then proceeds to encircle the particle and leaves an Orowan loop, otherwise it deforms. When the particle does not deform then extra dislocations are generated at the particle-matrix interface. On the other hand, if the particle deforms either before or after the formation of Orowan loop, no extra dislocations are generated. Generally, smaller particles are weaker than larger particles, which results in weaker slip plane. Move‐ ment of subsequent dislocations occurs on the same plane. As a consequence, slip bands are formed.

#### **2.3. Deformation texture**

to each other, but the absolute value of the single alloying element is not critical as long as

The cold-rolled alloys are consisting high density of dislocations clearly examined through transmission electron microscopy (TEM). During deformation the dislocations are arranged as cell boundaries whereas some cell interiors are free of dislocations (Fig. 1). Due to heavy

**Figure 1.** Dislocation structures of 92% cold rolled commercial purity aluminium alloy (AA1235 :0.67 Fe, 0.16Si, 0.01 Ti

However, the orientation of plastically deformed alloy changes from grain to grain or in different regions in a same grain. It happens owing to different rotations during the deforma‐ tion by use of a different combination of slip systems to achieve the imposed strain. Therefore, the deformation band is nothing but a volume of constant orientation that is significantly different than the orientation of other parts of the grain, which is explained by a schematic diagram (Fig. 2). It shows the variation of orientation at different regions of a same grain. The orientation of one part (X) of a deformed grain changes rapidly to that of a differently oriented part (Y) of the same grain across a thin boundary of a finite width called the transition bands (T) or microbands. Sometimes the orientation of deformation bands is changed twice, such as from X to Y and then Y to X. This special type of deformation band is called kink band.

deformation, the cell boundaries are diffuse in nature instead of simple dislocation.

primary crystallization of coarse intermetallic phases is avoided.

**2. Plastic deformation**

80 Light Metal Alloys Applications

and rest Al) [3]

**2.1. Microstructure of deformed metals**

The cold rolling texture of aluminium alloy sheets has been characterized as β fiber, which is associated with plane strain deformation [5]. The texture is varied from layer to layer of the sheet. The surface layer of the rolled sheet is known to have the shear texture, which is different from that of the plane strain deformed center layer [6-10]. According to Grewen and Huber, the ideal orientations in 95% cold rolled aluminium are characterized by {112}<111>, {110}<112> and {123}<412> [11]. The penetration depth of shear texture increased with increasing friction coefficients and rolling temperature [7]. Above all, material parameters also influence the shear texture formation. The most important factors among material parameters are the yield strength and strain hardening exponent [8,9]. The texture inhomogeneity has often been of considerable importance in rolling of pure aluminium, which has a low yield strength and work hardening exponent and tends to develop the shear texture in the surface layer [12-14]. Aluminium and copper single crystals of the S-orientations develop cube texture after coldrolling by more than 97.5% and subsequent annealing [15, 16]. This occurs owing to the close correlation between the formation of {001}<100> deformation structure and the development of cube texture. The cube-oriented material formed during rolling is considered to be the preferential nucleation sites of cubically aligned recrystallized grains. The factors which affect the generation of the {001}<100> deformation structure and the equivalent components during rolling are process parameters like friction between roll and sheet, rolling temperature, lubricants and the L/d value (where L and d represent the contact length and the thickness of specimens, respectively). Low viscosity lubricant such as kerosene suppresses the formation of {001}<100> deformation structure, whereas using machine oil as lubricant a cube component is formed. Thereafter, subsequent annealing treatment develops retained rolling and random texture in kerosene treated alloy and sharper cube texture in machine oil treated alloy [17]. Truszkowski et al. [9, 10] and Asbeck and Mecking [8] have reported that the shear texture is developed when L/d ≥ 5 or L/d ≤ 0.5.

ature a well-defined substructure is formed by recovery [19]. Coarsening of substructure takes

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83

The cell structure of a heavily cold rolled alloy is modified as a function of annealing time, examining the subgrain growth. In a commercial purity aluminium alloy (AA1235), after annealing at 250ºC for 1 h, the dislocations are examined to rearrange themselves to form a subgrain structure with an average diameter of 0.60 μm. Dislocation networks form the low angle boundary also known as subboundary. The dislocation rearrangement in the regions where the cell structure is not well-developed is described as a "disentanglement of disloca‐ tions" (Fig. 3a). Consequently, the dislocation density decreases inside the subgrain and particles are precipitated at subboundary as well as inside of subgrains (Fig. 3b). After 4 h of annealing at same temperature homogeneous subgrain growth takes place with an average diameter of 2 μm. Generally, the sharp delineation of many subboundaries decreases upon annealing and ultimately fades away, in agreement with the proposed coalescence mechanism (Figs. 3c). At this stage, subgrain growth is hindered by the pinning action of Al3Fe precipitates.

**Figure 3.** TEM micrographs (BF) showing subgrain formation of 92% cold rolled alloy AA1235 annealed at 250ºC for

(a) & (b) 1h and (c) 4h.

place by two methods-subgrain boundary migration and subgrain rotation.

#### **3. Recovery**

#### **3.1. Subgrain formation and its growth**

During plastic deformation of polycrystalline material (specially the alloy of medium and high stacking fault energy), unequal numbers of dislocations of two signs are generated, and the dislocations are rearranged as three-dimensional cell structure with complex dislocation tangles as cell walls. At recovery stage, the dislocations of opposite sign annihilate each other by combination of gliding and climbing mechanisms. The excess dislocations are left in the material at the end of the first stage of recovery. Upon progressing stages, these excess dislocations are arranged in a low energy configuration in the form of regular arrays or low angle grain boundaries (LAGBs). This mechanism is called polygonization [18], and the newly formed cells are called subgrains. In other words, the dislocation movement from the cell interiors to the cell boundary causes low angle grain boundaries, resulting in the subgrain formation. Sometimes, the dynamic recovery also helps to form a well-developed subgrain structure during deformation. In this case, recovery is the only involvement for the coarsening of subgrain structure. Although low stacking fault energy material shows poor subgrain structure, but when fine second-phase particles inhibit recrystallization then at high temper‐ ature a well-defined substructure is formed by recovery [19]. Coarsening of substructure takes place by two methods-subgrain boundary migration and subgrain rotation.

sheet. The surface layer of the rolled sheet is known to have the shear texture, which is different from that of the plane strain deformed center layer [6-10]. According to Grewen and Huber, the ideal orientations in 95% cold rolled aluminium are characterized by {112}<111>, {110}<112> and {123}<412> [11]. The penetration depth of shear texture increased with increasing friction coefficients and rolling temperature [7]. Above all, material parameters also influence the shear texture formation. The most important factors among material parameters are the yield strength and strain hardening exponent [8,9]. The texture inhomogeneity has often been of considerable importance in rolling of pure aluminium, which has a low yield strength and work hardening exponent and tends to develop the shear texture in the surface layer [12-14]. Aluminium and copper single crystals of the S-orientations develop cube texture after coldrolling by more than 97.5% and subsequent annealing [15, 16]. This occurs owing to the close correlation between the formation of {001}<100> deformation structure and the development of cube texture. The cube-oriented material formed during rolling is considered to be the preferential nucleation sites of cubically aligned recrystallized grains. The factors which affect the generation of the {001}<100> deformation structure and the equivalent components during rolling are process parameters like friction between roll and sheet, rolling temperature, lubricants and the L/d value (where L and d represent the contact length and the thickness of specimens, respectively). Low viscosity lubricant such as kerosene suppresses the formation of {001}<100> deformation structure, whereas using machine oil as lubricant a cube component is formed. Thereafter, subsequent annealing treatment develops retained rolling and random texture in kerosene treated alloy and sharper cube texture in machine oil treated alloy [17]. Truszkowski et al. [9, 10] and Asbeck and Mecking [8] have reported that the shear texture is

During plastic deformation of polycrystalline material (specially the alloy of medium and high stacking fault energy), unequal numbers of dislocations of two signs are generated, and the dislocations are rearranged as three-dimensional cell structure with complex dislocation tangles as cell walls. At recovery stage, the dislocations of opposite sign annihilate each other by combination of gliding and climbing mechanisms. The excess dislocations are left in the material at the end of the first stage of recovery. Upon progressing stages, these excess dislocations are arranged in a low energy configuration in the form of regular arrays or low angle grain boundaries (LAGBs). This mechanism is called polygonization [18], and the newly formed cells are called subgrains. In other words, the dislocation movement from the cell interiors to the cell boundary causes low angle grain boundaries, resulting in the subgrain formation. Sometimes, the dynamic recovery also helps to form a well-developed subgrain structure during deformation. In this case, recovery is the only involvement for the coarsening of subgrain structure. Although low stacking fault energy material shows poor subgrain structure, but when fine second-phase particles inhibit recrystallization then at high temper‐

developed when L/d ≥ 5 or L/d ≤ 0.5.

**3.1. Subgrain formation and its growth**

**3. Recovery**

82 Light Metal Alloys Applications

The cell structure of a heavily cold rolled alloy is modified as a function of annealing time, examining the subgrain growth. In a commercial purity aluminium alloy (AA1235), after annealing at 250ºC for 1 h, the dislocations are examined to rearrange themselves to form a subgrain structure with an average diameter of 0.60 μm. Dislocation networks form the low angle boundary also known as subboundary. The dislocation rearrangement in the regions where the cell structure is not well-developed is described as a "disentanglement of disloca‐ tions" (Fig. 3a). Consequently, the dislocation density decreases inside the subgrain and particles are precipitated at subboundary as well as inside of subgrains (Fig. 3b). After 4 h of annealing at same temperature homogeneous subgrain growth takes place with an average diameter of 2 μm. Generally, the sharp delineation of many subboundaries decreases upon annealing and ultimately fades away, in agreement with the proposed coalescence mechanism (Figs. 3c). At this stage, subgrain growth is hindered by the pinning action of Al3Fe precipitates.

**Figure 3.** TEM micrographs (BF) showing subgrain formation of 92% cold rolled alloy AA1235 annealed at 250ºC for (a) & (b) 1h and (c) 4h.

#### **3.2. The effect of second phase particles**

In two phase alloys, second phase particles inhibit recovery by the pinning of individual dislocation or dislocations of the low angle boundaries. Therefore, the dispersion of fine particles exerts a strong pinning effect on the subgrains [20-22]. Moreover, the particles stabled at high temperature improve the strength and creep properties of the alloy by pinning and stabilizing the recovered substructure. If the particles prevent the subgrain to reach the critical size for the formation of recrystallization nucleus, then recrystallization is also hindered. The evidence of dislocation pinning and restriction of subgrain growth is examined for a commer‐ cial purity aluminium alloy (AA8011) after annealing below 400ºC [23]. Similarly, the welldefined subgrains are examined in another commercial purity aluminium alloy (AA1235) after 4h annealing at 250ºC [24], as shown in Fig.4. The dislocation density change is clearly examined at some portions (X) of boundaries. The distribution of dislocation is also another observation at some points (Y). At this stage, dislocations accumulate at the grain boundary region and sometime they result in the formation of grain boundary lines at Z [25]. It can be said that the effect of particles on subgrain growth is similar to the effect of particles on grain growth.

zation decreases with increasing the minimum strain level. It is known that at low strain (<5%) high activation energy for nucleation results in a low nucleation rate and a relatively high growth rate [26]. With varying cold rolling percentage for different commercial purity aluminium alloys (AA1145, AA1200, AA8011), the degree of softening is identical for low reduction levels (<95%) and less steep for reduction levels above 95% (Fig.5) [27]. It corresponds to the gradual coarsening of microstructure with increasing temperature for high reduction

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**Figure 5.** Yield stress variation with annealing temperature for three commercial purity aluminium alloys at different

The recrystallization nuclei originate at preferred sites such as prior grain boundaries, transition bands, and shear bands of deformed microstructure. The nucleation process at the

**1. Strain induced boundary migration,** where a nucleus forms when an original grain

**2. Subgrain coarsening by coalescence at the original grain boundary,** where a nucleus

levels (>95%), relating to the continuous recrystallization.

percentage of cold rolling [27]

**4.2. Nucleation of recrystallization**

boundary bulges out [28].

grain boundary region by two different mechanisms:

forms in the consequence of subgrain growth [29-31].

**Figure 4.** TEM micrographs (BF) showing the effect of annealing at 250°C for 4 h on subgrain structure of 92% cold rolled bulk specimen (AA1235).

#### **4. Recrystallization**

#### **4.1. The effect of deformation percentage**

Nucleation and growth of new recrystallized grains are closely dependent on the distribution of dislocations within the deformed grains, relating to the deformation percentages. Recrys‐ tallization does not happen below a minimum strain and the start temperature of recrystalli‐ zation decreases with increasing the minimum strain level. It is known that at low strain (<5%) high activation energy for nucleation results in a low nucleation rate and a relatively high growth rate [26]. With varying cold rolling percentage for different commercial purity aluminium alloys (AA1145, AA1200, AA8011), the degree of softening is identical for low reduction levels (<95%) and less steep for reduction levels above 95% (Fig.5) [27]. It corresponds to the gradual coarsening of microstructure with increasing temperature for high reduction levels (>95%), relating to the continuous recrystallization.

**Figure 5.** Yield stress variation with annealing temperature for three commercial purity aluminium alloys at different percentage of cold rolling [27]

#### **4.2. Nucleation of recrystallization**

**3.2. The effect of second phase particles**

84 Light Metal Alloys Applications

growth.

rolled bulk specimen (AA1235).

**4. Recrystallization**

**4.1. The effect of deformation percentage**

In two phase alloys, second phase particles inhibit recovery by the pinning of individual dislocation or dislocations of the low angle boundaries. Therefore, the dispersion of fine particles exerts a strong pinning effect on the subgrains [20-22]. Moreover, the particles stabled at high temperature improve the strength and creep properties of the alloy by pinning and stabilizing the recovered substructure. If the particles prevent the subgrain to reach the critical size for the formation of recrystallization nucleus, then recrystallization is also hindered. The evidence of dislocation pinning and restriction of subgrain growth is examined for a commer‐ cial purity aluminium alloy (AA8011) after annealing below 400ºC [23]. Similarly, the welldefined subgrains are examined in another commercial purity aluminium alloy (AA1235) after 4h annealing at 250ºC [24], as shown in Fig.4. The dislocation density change is clearly examined at some portions (X) of boundaries. The distribution of dislocation is also another observation at some points (Y). At this stage, dislocations accumulate at the grain boundary region and sometime they result in the formation of grain boundary lines at Z [25]. It can be said that the effect of particles on subgrain growth is similar to the effect of particles on grain

**Figure 4.** TEM micrographs (BF) showing the effect of annealing at 250°C for 4 h on subgrain structure of 92% cold

Nucleation and growth of new recrystallized grains are closely dependent on the distribution of dislocations within the deformed grains, relating to the deformation percentages. Recrys‐ tallization does not happen below a minimum strain and the start temperature of recrystalli‐

The recrystallization nuclei originate at preferred sites such as prior grain boundaries, transition bands, and shear bands of deformed microstructure. The nucleation process at the grain boundary region by two different mechanisms:

With an increasing degree of deformation, the number of potential nucleation sites such as deformation bands and grain boundary bands increases significantly and the number of such sites is larger in fine grain than in coarse grain specimens [32]. The effectiveness of nucleation sites is enhanced by the presence of intermetallic particles (Al3Fe) with increasing degree of deformation. Potential nucleation sites are more important factor than number of nuclei. If the same numbers of nuclei form at the grain boundaries of both fine grain material and coarse grain material, then the fine grain material shows more homogeneous recrystallization. In case of heterogeneous recrystallization, all grains are not recrystallized at the same rate. It is known that crystallographic orientation affects the slip systems and strain path during deformation. Therefore, distribution and density of dislocations, large scale microstructural inhomogenet‐ ies, availability of nucleation sites, and growth rate of recrystallized grains are also dependent on crystallographic orientation.

which viable nuclei formation is inhibited owing to the proximity of particles. In such circumstances nucleation is likely to occur predominantly at the original matrix grain boun‐ daries. As a consequence, the nucleation as well as recrystallization rate drastically get reduced

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At the time of deformation, large particles (> 1 μm) lead to a heterogeneous distribution of dislocations, whereas fine particles (< 0.1 μm) give rise to a homogeneous distribution of dislocations. As a consequence, the number of possible recrystallization nucleation sites increases for large particles and decreases for fine particles. This can be related to the degree of deformation (Fig. 6), explaining the nucleation of recrystallized grains at the particles greater than 2 μm for a highly deformed metal [37]. The fine particles also inhibit sub-boundary migration and thus the nucleation process is retarded. As the particles become more finely spaced, nearly all cell boundaries will be pinned by the particles at the end of the deformation stage. In this case, when recrystallization occurs, the mechanism of nucleation is not well understood, but the kinetics of the process is certainly very slow [38]. Closely spaced, thermally stable particles preserve the deformed/ recovered microstructure up to the melting point of the matrix. Kim et al. have revealed that the retention of dislocation substructure at high temperature provides an additional strengthening mechanism to the dispersion hardening of the alloys used at high temperature structural applications [39]. With the decrease in the particle size and spacing between particles, recrystallization kinetics is retarded, but final grain size becomes quite large [37, 40]. Several researchers have studied the relationship between

dispersion characteristics, deformation substructure, and recrystallization [41-43].

**Figure 6.** The conditions of deformation and particle size for which nucleation of recrystallized grains is observed to

if the inter particle spacing falls below C2 [36].

**4.5. The effect of particle size**

occur at particles [37].

#### **4.3. Sequence of precipitation and recrystallization**

It is well established that the recrystallization behavior of a deformed and supersaturated alloy is largely dependent on whether or not precipitation of the second phase particles take place simultaneously with recrystallization [33]. Hornbogen and Köster have suggested that recrystallization occurs prior to precipitation at high temperatures, whereas precipitation takes place prior to recrystallization at low temperatures [34]. If particles are precipitated during recrystallization they may hinder both the formation and migration of recrystallization fronts. Alternatively, if the second phase particles are precipitated in the matrix prior to cold rolling, the recrystallization behaviour will depend on the size and dispersion of the second phase particles.

#### **4.4. The effect of inter particle spacing**

At wide inter particle spacing, when only a few particles are present, there is basically no difference between the recrystallization behaviour of two-phase and single phase alloys. In this case, nucleation generally occurs at the original matrix boundaries and the ultimate recrystallized grain size may vary owing to inhomogeneous distribution of nucleation sites. With an increase in particle content of the alloy (i.e., decrease in interparticle spacing) nucle‐ ation occurs more rapidly at the lattice curvature of particle-matrix (for the particles >1 μm) interface than at the grain boundary region [35]. In addition, due to the increase in quantity and uniform distribution of particle-matrix nucleation sites, the final recrystallized grains become more uniform and finer. However, the trend towards increased nucleation due to decrease in interparticle spacing occurs until particles maintain a critical spacing (C1) to allow the nucleation to occur simultaneously and independently at each particle. After reaching the critical spacing (C1), a nucleus attached to one particle would be viable itself, if it left to develop by itself. It would be non-viable if nuclei start to form simultaneously at neighbouring particles. They would then interfere with each other before reaching a viable size. Thus, due to the formation of fewer amount of viable nuclei at particle-matrix interfaces during increase in particle content (i.e., decrease in interparticle spacing) the overall nucleation rate of recrystal‐ lization decreases. Further increase in particle content leads to a second critical spacing (C2) at which viable nuclei formation is inhibited owing to the proximity of particles. In such circumstances nucleation is likely to occur predominantly at the original matrix grain boun‐ daries. As a consequence, the nucleation as well as recrystallization rate drastically get reduced if the inter particle spacing falls below C2 [36].

#### **4.5. The effect of particle size**

With an increasing degree of deformation, the number of potential nucleation sites such as deformation bands and grain boundary bands increases significantly and the number of such sites is larger in fine grain than in coarse grain specimens [32]. The effectiveness of nucleation sites is enhanced by the presence of intermetallic particles (Al3Fe) with increasing degree of deformation. Potential nucleation sites are more important factor than number of nuclei. If the same numbers of nuclei form at the grain boundaries of both fine grain material and coarse grain material, then the fine grain material shows more homogeneous recrystallization. In case of heterogeneous recrystallization, all grains are not recrystallized at the same rate. It is known that crystallographic orientation affects the slip systems and strain path during deformation. Therefore, distribution and density of dislocations, large scale microstructural inhomogenet‐ ies, availability of nucleation sites, and growth rate of recrystallized grains are also dependent

It is well established that the recrystallization behavior of a deformed and supersaturated alloy is largely dependent on whether or not precipitation of the second phase particles take place simultaneously with recrystallization [33]. Hornbogen and Köster have suggested that recrystallization occurs prior to precipitation at high temperatures, whereas precipitation takes place prior to recrystallization at low temperatures [34]. If particles are precipitated during recrystallization they may hinder both the formation and migration of recrystallization fronts. Alternatively, if the second phase particles are precipitated in the matrix prior to cold rolling, the recrystallization behaviour will depend on the size and dispersion of the second phase

At wide inter particle spacing, when only a few particles are present, there is basically no difference between the recrystallization behaviour of two-phase and single phase alloys. In this case, nucleation generally occurs at the original matrix boundaries and the ultimate recrystallized grain size may vary owing to inhomogeneous distribution of nucleation sites. With an increase in particle content of the alloy (i.e., decrease in interparticle spacing) nucle‐ ation occurs more rapidly at the lattice curvature of particle-matrix (for the particles >1 μm) interface than at the grain boundary region [35]. In addition, due to the increase in quantity and uniform distribution of particle-matrix nucleation sites, the final recrystallized grains become more uniform and finer. However, the trend towards increased nucleation due to decrease in interparticle spacing occurs until particles maintain a critical spacing (C1) to allow the nucleation to occur simultaneously and independently at each particle. After reaching the critical spacing (C1), a nucleus attached to one particle would be viable itself, if it left to develop by itself. It would be non-viable if nuclei start to form simultaneously at neighbouring particles. They would then interfere with each other before reaching a viable size. Thus, due to the formation of fewer amount of viable nuclei at particle-matrix interfaces during increase in particle content (i.e., decrease in interparticle spacing) the overall nucleation rate of recrystal‐ lization decreases. Further increase in particle content leads to a second critical spacing (C2) at

on crystallographic orientation.

86 Light Metal Alloys Applications

**4.4. The effect of inter particle spacing**

particles.

**4.3. Sequence of precipitation and recrystallization**

At the time of deformation, large particles (> 1 μm) lead to a heterogeneous distribution of dislocations, whereas fine particles (< 0.1 μm) give rise to a homogeneous distribution of dislocations. As a consequence, the number of possible recrystallization nucleation sites increases for large particles and decreases for fine particles. This can be related to the degree of deformation (Fig. 6), explaining the nucleation of recrystallized grains at the particles greater than 2 μm for a highly deformed metal [37]. The fine particles also inhibit sub-boundary migration and thus the nucleation process is retarded. As the particles become more finely spaced, nearly all cell boundaries will be pinned by the particles at the end of the deformation stage. In this case, when recrystallization occurs, the mechanism of nucleation is not well understood, but the kinetics of the process is certainly very slow [38]. Closely spaced, thermally stable particles preserve the deformed/ recovered microstructure up to the melting point of the matrix. Kim et al. have revealed that the retention of dislocation substructure at high temperature provides an additional strengthening mechanism to the dispersion hardening of the alloys used at high temperature structural applications [39]. With the decrease in the particle size and spacing between particles, recrystallization kinetics is retarded, but final grain size becomes quite large [37, 40]. Several researchers have studied the relationship between dispersion characteristics, deformation substructure, and recrystallization [41-43].

**Figure 6.** The conditions of deformation and particle size for which nucleation of recrystallized grains is observed to occur at particles [37].

#### **4.6. Type of particles in commercial purity aluminium**

In commercial purity aluminium alloys, the main impurity Fe combines with both Al and Si to form a large variety of phases during solidification or during subsequent thermomechanical processing [44, 45]. Shoji and Fujikura have identified three types of precipitates Si, α-AlFeSi, and Al3Fe in cold rolled commercial alloy AlFe0.6Si0.16 (all in wt%) after annealing in salt bath at different temperatures 300, 400 and 500ºC, respectively [3, 46]. Some researchers have identified only Al3Fe precipitates during recrystallization of commercial Al-Fe-Si alloy [47]. In AA8011 alloy, Al3Fe and Si precipitates are observed in cold rolled and annealed conditions [48]. After heavy plastic deformation (~92%), very small particles of 0.17 μm size are distributed in the dislocations [3]. Upon progress of annealing, the particle size increases, and both smaller and larger, spherical and plate-shaped Al3Fe particles are inhomogeneously distributed and situated at subboundaries and subgrain interiors. Since the dislocations cannot cross the grain boundary, during annealing a large variation of dislocation density occurs from grain to grain owing to the pinning effect of precipitate particles. After completion of recrystallization, particle also inhibits the grain boundary migration along with pinning of the dislocations. Even after a long time of high temperature annealing for heavily deformed alloy, the Al3Fe particles may pin down the dislocations, resulting in the presence of a large amount dislocations inside the grains (Fig. 7) [3, 24]. During the intersection of grain boundary by particles, a Zener drag force is generated to restrain the boundary migration [49]. The coherent particle generally loses the coherency when a high angle grain boundary moves past a coherent particle. The coherent particles are twice as effective in pinning a grain boundary as incoherent particles of the same size [4]. There are many alternative situations during particle-boundary interaction. The particle may dissolve during passage of the boundary and re-precipitate in a coherent orientation, it may reorient itself to a coherent orientation, or the boundary may cut through the particle [50].

**4.7. Particle distributions in bimodal alloys**

particle size for the growth of a nucleus is

*ρ* is dislocation density,

Pz is Zener pinning force, is grain boundary energy,

*r* is radius of small particles

**4.8. Recrystallization texture**

*f* is volume fraction of small particles,

*G* is shear modulus, *b* is burgers vector,

Many commercial alloys are bimodal alloys which contain both large (>1 μm) and small particles. Large particles act as nucleation sites for recrystallization and small particles hinder the grain boundary migration, i.e., retard the recrystallization. Therefore, recrystallization behaviour of bimodal alloys is affected by particle distribution. During deformation, the fine dispersion of small particles does not affect significantly the degree of lattice curvature generated at the large particle-matrix interface. Later, by altering the dispersion parameters of particles, recrystallization kinetics and microstructures are controlled. Chan and Humphreys have reported that in the bimodal alloy, nucleation of recrystallization takes place at the large particles, but fine particles determine the time for completion of recrystallization [51]. They have observed a large number of small island grains and coarse irregular grains in the microstructure of bimodal alloys. With increasing coarsening of the fine precipitates, the grain shapes become more regular and the mean size decreases. Nes has taken the help of the parameter *f/r* (where *f* and *r* are volume fraction and radius of small particles, respectively) to describe a model to account for the grain size of bimodal alloys [52]. In bimodal alloys, critical

2

As *f/r* of fine dispersion increases, the large particles act as the nucleation sites, but growth of the grains is slowed down at the early stage of annealing. However, the earlier developed

Annealing of heavily deformed material leads to a wide range of recrystallization textures. Different types of recrystallization textures can be produced in similar alloys each having almost identical and very strong deformation textures [11]. In f.c.c. metals, sometimes recrys‐ tallization texture contains cube texture along with retained rolling texture and in some cases

 g

3

*r*

g

= = - - (1)

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4 4

g

*d*

where *Pd* is driving force (stored energy in the subgrain boundaries),

grains either consume the smaller grains, or form island grains.

2 2 *z d*

r

*P P Gb f*

**Figure 7.** TEM micrographs (BF) of 92% cold rolled alloy AA1235 after annealing at 480ºC for 8 h showing (a) pinning of dislocations by a particle, and (b) presence of dislocations even after completion of recrystallization

#### **4.7. Particle distributions in bimodal alloys**

**4.6. Type of particles in commercial purity aluminium**

88 Light Metal Alloys Applications

the particle [50].

In commercial purity aluminium alloys, the main impurity Fe combines with both Al and Si to form a large variety of phases during solidification or during subsequent thermomechanical processing [44, 45]. Shoji and Fujikura have identified three types of precipitates Si, α-AlFeSi, and Al3Fe in cold rolled commercial alloy AlFe0.6Si0.16 (all in wt%) after annealing in salt bath at different temperatures 300, 400 and 500ºC, respectively [3, 46]. Some researchers have identified only Al3Fe precipitates during recrystallization of commercial Al-Fe-Si alloy [47]. In AA8011 alloy, Al3Fe and Si precipitates are observed in cold rolled and annealed conditions [48]. After heavy plastic deformation (~92%), very small particles of 0.17 μm size are distributed in the dislocations [3]. Upon progress of annealing, the particle size increases, and both smaller and larger, spherical and plate-shaped Al3Fe particles are inhomogeneously distributed and situated at subboundaries and subgrain interiors. Since the dislocations cannot cross the grain boundary, during annealing a large variation of dislocation density occurs from grain to grain owing to the pinning effect of precipitate particles. After completion of recrystallization, particle also inhibits the grain boundary migration along with pinning of the dislocations. Even after a long time of high temperature annealing for heavily deformed alloy, the Al3Fe particles may pin down the dislocations, resulting in the presence of a large amount dislocations inside the grains (Fig. 7) [3, 24]. During the intersection of grain boundary by particles, a Zener drag force is generated to restrain the boundary migration [49]. The coherent particle generally loses the coherency when a high angle grain boundary moves past a coherent particle. The coherent particles are twice as effective in pinning a grain boundary as incoherent particles of the same size [4]. There are many alternative situations during particle-boundary interaction. The particle may dissolve during passage of the boundary and re-precipitate in a coherent orientation, it may reorient itself to a coherent orientation, or the boundary may cut through

**Figure 7.** TEM micrographs (BF) of 92% cold rolled alloy AA1235 after annealing at 480ºC for 8 h showing (a) pinning

of dislocations by a particle, and (b) presence of dislocations even after completion of recrystallization

Many commercial alloys are bimodal alloys which contain both large (>1 μm) and small particles. Large particles act as nucleation sites for recrystallization and small particles hinder the grain boundary migration, i.e., retard the recrystallization. Therefore, recrystallization behaviour of bimodal alloys is affected by particle distribution. During deformation, the fine dispersion of small particles does not affect significantly the degree of lattice curvature generated at the large particle-matrix interface. Later, by altering the dispersion parameters of particles, recrystallization kinetics and microstructures are controlled. Chan and Humphreys have reported that in the bimodal alloy, nucleation of recrystallization takes place at the large particles, but fine particles determine the time for completion of recrystallization [51]. They have observed a large number of small island grains and coarse irregular grains in the microstructure of bimodal alloys. With increasing coarsening of the fine precipitates, the grain shapes become more regular and the mean size decreases. Nes has taken the help of the parameter *f/r* (where *f* and *r* are volume fraction and radius of small particles, respectively) to describe a model to account for the grain size of bimodal alloys [52]. In bimodal alloys, critical particle size for the growth of a nucleus is

$$d = \frac{4\gamma}{P\_d - P\_\perp} = \frac{4\gamma}{\frac{\rho G b^2}{2} - \frac{3f\gamma}{2r}}$$

where *Pd* is driving force (stored energy in the subgrain boundaries),

*ρ* is dislocation density,

*G* is shear modulus,

*b* is burgers vector,

Pz is Zener pinning force,

is grain boundary energy,

*f* is volume fraction of small particles,

*r* is radius of small particles

As *f/r* of fine dispersion increases, the large particles act as the nucleation sites, but growth of the grains is slowed down at the early stage of annealing. However, the earlier developed grains either consume the smaller grains, or form island grains.

#### **4.8. Recrystallization texture**

Annealing of heavily deformed material leads to a wide range of recrystallization textures. Different types of recrystallization textures can be produced in similar alloys each having almost identical and very strong deformation textures [11]. In f.c.c. metals, sometimes recrys‐ tallization texture contains cube texture along with retained rolling texture and in some cases nearly random textures [11]. It is reported that with increasing deformation, the volume fraction of the cube texture decreases preferably for low annealing temperatures and that of the rolling texture increases for all temperatures, whereas the random part of the texture strongly decreases [53]. The possible origin of recrystallization texture has been reviewed several times [54, 55]. It is notable that "oriented nucleation" and "oriented growth" are two possible theories for the development of recrystallization texture [56, 57]. The recrystallization texture develops from the competition between cube oriented grains which nucleate at the cube bands, R-oriented grains (i.e., retained rolling texture) which stem from the grain boundaries between the former deformed grains, and randomly oriented grains due to particle stimulated nucleation (PSN).

a single particle of a lightly deformed polycrystals, a spread of orientation is observed

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**2.** In the case of heavily deformed material, the presence of different deformation zones of different grains or different deformation regions of the same grain results in a single grain with a wide range of orientations. As a consequence, either a weak texture or randomly

The recrystallization texture is dependent on the particle size, strength and spacing. When the particles precipitate during annealing, the recrystallization texture developed is similar to the deformation texture. There is no clear reason for the retention of the rolling texture in the recrystallized alloys. It appears that a strong cube texture is rarely formed and particles may be responsible for suppressing the formation or viability of cube sites. This enables other

Although PSN nuclei are oriented randomly in the heavily rolled polycrystal, the final texture of this alloy is not random. Therefore, it is concluded that grains from other sites with different orientation affect the final texture. From Table 2.2 it is observed that for the alloys (e.g., Al-Fe-Si or Al-Si) containing low volume fraction of particles, cube and rolling components are developed, whereas rolling and random components develop for high volume fraction of particles containing alloys (e.g., Al-Ni). The cube grains are larger than randomly oriented grains owing to the faster growth rate of former [59]. It results in formation of small island grains inside the large cube grains. Some researchers have shown that the strength of the cube component decreases with an increase in the number of supercritical sized particles in the case of hot-rolled aluminium alloy AA3004 [60]. On the other hand, volume fraction of randomly oriented grains increases with an increase in the number of supercritical sized particles. Hornbogen et al. have distinguished between 'discontinuous' (conventional recrystallization) and 'continuous' recrystallization [61]. The former cases give rise to normal recrystallization textures, such as the cube texture, while the latter cause retention of the rolling texture.

A typical recrystallization texture of highest purity aluminium is the cube-orientation {001}<100> [5, 11]. In commercial purity aluminium the iron and silicon contents have been found to be important factors in controlling the recrystallization texture [62, 63]. It is reported that a small amount of Fe may cause a change almost from pure cube to retained rolling texture [4]. During recrystallization, precipitation of Al3Fe particles inhibits the growth of the early formed cube nuclei, thereby forcing nucleation in the abundant rolling texture components. Thus, the recrystallization textures of most commercial aluminium alloys are composed of three texture types, namely the cube texture, the retained rolling texture and the random texture [64, 65]. Cube-oriented subgrains are known to rapidly recover either dynamically during deformation or statically during the early stage of recrystallization, which gives rise to a size advantage of cube nuclei [66]. Growth of the dominant cube grains is lowered by solute drag and/or precipitation, which leads to the development of rolling component and reduces the strength of the cube component. In a sheet, a mixture of cube and rolling components is

components including the retained rolling components to dominate the texture.

around the particles in the deformed matrix.

*4.8.2. Effect of Fe on the commercial purity aluminium*

oriented grains form.

#### *4.8.1. The effect of particles*

It is already discussed in previous section (§ 4.4) that widely spaced coarse particles (>1 μm) enhance nucleation and the rate of recrystallization if it is present before deformation. The finely dispersed particles normally reduce recrystallization kinetics having a greater retarding influence on nucleation than on growth. Therefore, the texture changes can be rationalized on the basis of the orientation dependent nucleation. If nucleation of all components is retarded in equal proportions, then the time available for growth of the first formed grains is increased. The recrystallization texture will therefore show increased selectivity and will become even more strongly biased towards the most favoured orientation nuclei. The effect of particle size on the formation of major components of recrystallization texture is shown in Table 1 [58].

**Table 1.** Major components of the recrystallization texture for some particle containing alloys [58]

The effect of PSN on recrystallization texture is as follows [59]:

**1.** PSN results in a sharp recrystallization texture for low strained material or in case of particle containing single crystals. When PSN originates in several deformation zones at a single particle of a lightly deformed polycrystals, a spread of orientation is observed around the particles in the deformed matrix.

**2.** In the case of heavily deformed material, the presence of different deformation zones of different grains or different deformation regions of the same grain results in a single grain with a wide range of orientations. As a consequence, either a weak texture or randomly oriented grains form.

The recrystallization texture is dependent on the particle size, strength and spacing. When the particles precipitate during annealing, the recrystallization texture developed is similar to the deformation texture. There is no clear reason for the retention of the rolling texture in the recrystallized alloys. It appears that a strong cube texture is rarely formed and particles may be responsible for suppressing the formation or viability of cube sites. This enables other components including the retained rolling components to dominate the texture.

Although PSN nuclei are oriented randomly in the heavily rolled polycrystal, the final texture of this alloy is not random. Therefore, it is concluded that grains from other sites with different orientation affect the final texture. From Table 2.2 it is observed that for the alloys (e.g., Al-Fe-Si or Al-Si) containing low volume fraction of particles, cube and rolling components are developed, whereas rolling and random components develop for high volume fraction of particles containing alloys (e.g., Al-Ni). The cube grains are larger than randomly oriented grains owing to the faster growth rate of former [59]. It results in formation of small island grains inside the large cube grains. Some researchers have shown that the strength of the cube component decreases with an increase in the number of supercritical sized particles in the case of hot-rolled aluminium alloy AA3004 [60]. On the other hand, volume fraction of randomly oriented grains increases with an increase in the number of supercritical sized particles. Hornbogen et al. have distinguished between 'discontinuous' (conventional recrystallization) and 'continuous' recrystallization [61]. The former cases give rise to normal recrystallization textures, such as the cube texture, while the latter cause retention of the rolling texture.

#### *4.8.2. Effect of Fe on the commercial purity aluminium*

nearly random textures [11]. It is reported that with increasing deformation, the volume fraction of the cube texture decreases preferably for low annealing temperatures and that of the rolling texture increases for all temperatures, whereas the random part of the texture strongly decreases [53]. The possible origin of recrystallization texture has been reviewed several times [54, 55]. It is notable that "oriented nucleation" and "oriented growth" are two possible theories for the development of recrystallization texture [56, 57]. The recrystallization texture develops from the competition between cube oriented grains which nucleate at the cube bands, R-oriented grains (i.e., retained rolling texture) which stem from the grain boundaries between the former deformed grains, and randomly oriented grains due to particle

It is already discussed in previous section (§ 4.4) that widely spaced coarse particles (>1 μm) enhance nucleation and the rate of recrystallization if it is present before deformation. The finely dispersed particles normally reduce recrystallization kinetics having a greater retarding influence on nucleation than on growth. Therefore, the texture changes can be rationalized on the basis of the orientation dependent nucleation. If nucleation of all components is retarded in equal proportions, then the time available for growth of the first formed grains is increased. The recrystallization texture will therefore show increased selectivity and will become even more strongly biased towards the most favoured orientation nuclei. The effect of particle size on the formation of major components of recrystallization texture is shown in Table 1 [58].

**Volume fraction**

**Table 1.** Major components of the recrystallization texture for some particle containing alloys [58]

**1.** PSN results in a sharp recrystallization texture for low strained material or in case of particle containing single crystals. When PSN originates in several deformation zones at

The effect of PSN on recrystallization texture is as follows [59]:

**(f)**

Al (99.9965) (90%) 0 - Cube Fe-AlN (70%) 0.06 0.017 Rolling Al-Al2O3 (30-90%) 0.4 0.1 Deformation Cu-SiO2 (70%) 0.5 0.23 Rolling + Twins Al-FeSi (90%) 0.5 0.2- 7 Cube, Rolling, Random Al-Si (90%) 0.8 2 Cube, Rolling, Random Al- Ni (80%) 10.0 1 Rolling, Random Al- SiC (70% 20.0 10 Random

**Particles Recrystallization Texture**

**Diameter (μm)**

stimulated nucleation (PSN).

**Material (% of reduction)**

*4.8.1. The effect of particles*

90 Light Metal Alloys Applications

A typical recrystallization texture of highest purity aluminium is the cube-orientation {001}<100> [5, 11]. In commercial purity aluminium the iron and silicon contents have been found to be important factors in controlling the recrystallization texture [62, 63]. It is reported that a small amount of Fe may cause a change almost from pure cube to retained rolling texture [4]. During recrystallization, precipitation of Al3Fe particles inhibits the growth of the early formed cube nuclei, thereby forcing nucleation in the abundant rolling texture components. Thus, the recrystallization textures of most commercial aluminium alloys are composed of three texture types, namely the cube texture, the retained rolling texture and the random texture [64, 65]. Cube-oriented subgrains are known to rapidly recover either dynamically during deformation or statically during the early stage of recrystallization, which gives rise to a size advantage of cube nuclei [66]. Growth of the dominant cube grains is lowered by solute drag and/or precipitation, which leads to the development of rolling component and reduces the strength of the cube component. In a sheet, a mixture of cube and rolling components is desired to prohibit earing during deep drawing of that sheet [54]. It is interesting to note from the research work of Hirsch and Lücke that 95% cold-rolled alloy Al-0.007% Fe shows the retained rolling texture with a small amount of cube texture at 360ºC compared to strong cube texture at 280 and 520ºC [67]. It is due to the formation of iron-rich phase at 360ºC, which restricts cube grain growth. When precipitation occurs before or after completion of recrys‐ tallization, the effect on boundary migration is less drastic and a strong cube texture develops. It has been observed that the growing grains of cube component can consume the recrystallized grains of non-cube component along with the surrounding deformed material [68]. It is reported that coarse iron rich particles inhibit the formation of cube texture in cold rolled aluminium alloys [11]. It is evident that increasing iron content decreases the strength of the cube texture, especially when it is present in solid solution form before rolling [11]. Even the addition of 0.1% iron to aluminium may change the recrystallization texture [69].

4 3 *Zener r D*

Equation 2 is only applicable to those materials where the second-phase particles are stable during grain growth. Therefore, stability of the particles with temperature is one of the important factors for growth mechanism. The instability of the second-phase particles, i.e., precipitation after the formation of grain and subgrain structure or coarsening of particles during grain growth, affects the grain growth by a different way and it has been discussed

Orientation gradient has a greater effect on the grain boundary mobility. High stacking fault energy alloys (Al, Ni and α-iron etc.) readily form a cellular or subgrain structure during deformation. This structure is often not uniform and orientation gradients are usually developed here. In addition, many alloys produced by thermomechanical processing develop a preferred orientation or texture by recrystallization and grain growth. A model developed by Ferry illustrates that the grain coarsening is rapid for large orientation gradient and particlefree systems, and it is markedly reduced in a system containing a large volume fraction of fine particles, despite the presence of the orientation gradient [72]. The latter case occurs due to

Grain growth may be divided into two types, normal grain growth and abnormal grain growth. When the grains grow uniformly with a narrow range of size and shapes and the grain size distribution is independent of time, then, this type of grain growth is called "normal" or "continuous" grain growth. In "abnormal" or "discontinuous" grain growth, few grains are large compared to the rest and grain size distribution is bimodal. Since this discontinuous growth of selected grains has similar kinetics to primary recrystallization and has some microstructural similarities, abnormal grain growth is sometime called secondary recrystalli‐ zation. The normal grain growth theory is based on the grain boundary interfacial free energy as the driving force. If the initial grain size distribution is too wide, a fraction of large grains will grow in an abnormal manner until all the other grains have been consumed. When completed, this process has resulted in a more narrow size distribution and at longer times the steady state may be approached asymptotically. It appears convenient to define this state as a normal grain growth. Abnormal grain growth may sometime be a necessary initial stage

In spite of considerable research efforts, the origin of abnormal grain growth is yet to be fully understood. However, by Monte Carlo simulation it has been observed that two different conditions may initiate abnormal grain growth [74]. Firstly, anisotropy in the grain boundary energy may lead to rapid growth of grains having boundary energies much lower than the average [75]. This is often the case when the material exhibits a strong primary recrystallization

elsewhere [4].

retardation of the onset of recrystallization.

toward find normal grain growth [73].

**5.2. The competition between normal and abnormal grain growth**

*f* = (2)

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93

#### *4.8.3. Effect of recrystallization texture on microstructure and mechanical properties*

Grain size and mechanical properties can be controlled by maintaining proper texture in recrystallized microstructure. It is already discussed in §4.8.2 that Al3Fe particles suppress the cube texture formation. Cube grain growth is also inhibited by "orientation pinning" [59]. Orientation pinning occurs when a recrystallized grain of a given orientation grows into deformed material of its own orientation. It results in a low angle boundary and the growth of the recrystallized grain will practically stop due to the low mobility of low angle boundaries. The case where orientation pinning is expected to be important is for the growth of grains with the same orientation as a main component in the deformation texture, like retained rolling components. Vatne et al. have also reported a case where the growth of cube grains has been most likely reduced due to orientation pinning. This may be due to an extremely high asdeformed cube fraction of 35% [70]. Therefore, fine grains can be generated in the foil by controlling recrystallization texture during intermediate annealing treatment, which improves the mechanical properties of the foil. Blade has demonstrated that earing tendency of alumi‐ nium alloy increases with increase in cube texture component [71]. It is shown that earing tendency becomes zero when 25% cube texture is present.

#### **5. Grain growth**

#### **5.1.  As the driving pressure of grain growth is extremely low, particles may have a very large influence on both the kinetics of grain growth and the resultant microstructure. Humphreys et al. have elaborately discussed the growth related formulation, i.e., driving and pinning pressures for growth, grain growth rate, limiting grain size due to particles, etc. [4]. In the case of planar grain boundary, Zener limiting grain size (*DZener*) which equals to the critical radius of island grains, arises from the balance between the driving pressure for grain growth (*P*) and Zener pinning pressure (*Pz*).

Recrystallization Behavior of Commercial Purity Aluminium Alloys http://dx.doi.org/10.5772/58385 93

$$D\_{Zener} = \frac{4r}{3f} \tag{2}$$

Equation 2 is only applicable to those materials where the second-phase particles are stable during grain growth. Therefore, stability of the particles with temperature is one of the important factors for growth mechanism. The instability of the second-phase particles, i.e., precipitation after the formation of grain and subgrain structure or coarsening of particles during grain growth, affects the grain growth by a different way and it has been discussed elsewhere [4].

Orientation gradient has a greater effect on the grain boundary mobility. High stacking fault energy alloys (Al, Ni and α-iron etc.) readily form a cellular or subgrain structure during deformation. This structure is often not uniform and orientation gradients are usually developed here. In addition, many alloys produced by thermomechanical processing develop a preferred orientation or texture by recrystallization and grain growth. A model developed by Ferry illustrates that the grain coarsening is rapid for large orientation gradient and particlefree systems, and it is markedly reduced in a system containing a large volume fraction of fine particles, despite the presence of the orientation gradient [72]. The latter case occurs due to retardation of the onset of recrystallization.

#### **5.2. The competition between normal and abnormal grain growth**

desired to prohibit earing during deep drawing of that sheet [54]. It is interesting to note from the research work of Hirsch and Lücke that 95% cold-rolled alloy Al-0.007% Fe shows the retained rolling texture with a small amount of cube texture at 360ºC compared to strong cube texture at 280 and 520ºC [67]. It is due to the formation of iron-rich phase at 360ºC, which restricts cube grain growth. When precipitation occurs before or after completion of recrys‐ tallization, the effect on boundary migration is less drastic and a strong cube texture develops. It has been observed that the growing grains of cube component can consume the recrystallized grains of non-cube component along with the surrounding deformed material [68]. It is reported that coarse iron rich particles inhibit the formation of cube texture in cold rolled aluminium alloys [11]. It is evident that increasing iron content decreases the strength of the cube texture, especially when it is present in solid solution form before rolling [11]. Even the

addition of 0.1% iron to aluminium may change the recrystallization texture [69].

Grain size and mechanical properties can be controlled by maintaining proper texture in recrystallized microstructure. It is already discussed in §4.8.2 that Al3Fe particles suppress the cube texture formation. Cube grain growth is also inhibited by "orientation pinning" [59]. Orientation pinning occurs when a recrystallized grain of a given orientation grows into deformed material of its own orientation. It results in a low angle boundary and the growth of the recrystallized grain will practically stop due to the low mobility of low angle boundaries. The case where orientation pinning is expected to be important is for the growth of grains with the same orientation as a main component in the deformation texture, like retained rolling components. Vatne et al. have also reported a case where the growth of cube grains has been most likely reduced due to orientation pinning. This may be due to an extremely high asdeformed cube fraction of 35% [70]. Therefore, fine grains can be generated in the foil by controlling recrystallization texture during intermediate annealing treatment, which improves the mechanical properties of the foil. Blade has demonstrated that earing tendency of alumi‐ nium alloy increases with increase in cube texture component [71].  As the driving pressure of grain growth is extremely low, particles may have a very large influence on both the kinetics of grain growth and the resultant microstructure. Humphreys et al. have elaborately discussed the growth related formulation, i.e., driving and pinning pressures for growth, grain growth rate, limiting grain size due to particles, etc. [4]. In the case of planar grain boundary, Zener limiting grain size (*DZener*) which equals to the critical radius of island grains, arises from the balance between the driving pressure for grain

*4.8.3. Effect of recrystallization texture on microstructure and mechanical properties*

tendency becomes zero when 25% cube texture is present.

growth (*P*) and Zener pinning pressure (*Pz*).

**5.1. The effect of second-phase particles and orientation gradient**

**5. Grain growth**

92 Light Metal Alloys Applications

Grain growth may be divided into two types, normal grain growth and abnormal grain growth. When the grains grow uniformly with a narrow range of size and shapes and the grain size distribution is independent of time, then, this type of grain growth is called "normal" or "continuous" grain growth. In "abnormal" or "discontinuous" grain growth, few grains are large compared to the rest and grain size distribution is bimodal. Since this discontinuous growth of selected grains has similar kinetics to primary recrystallization and has some microstructural similarities, abnormal grain growth is sometime called secondary recrystalli‐ zation. The normal grain growth theory is based on the grain boundary interfacial free energy as the driving force. If the initial grain size distribution is too wide, a fraction of large grains will grow in an abnormal manner until all the other grains have been consumed. When completed, this process has resulted in a more narrow size distribution and at longer times the steady state may be approached asymptotically. It appears convenient to define this state as a normal grain growth. Abnormal grain growth may sometime be a necessary initial stage toward find normal grain growth [73].

In spite of considerable research efforts, the origin of abnormal grain growth is yet to be fully understood. However, by Monte Carlo simulation it has been observed that two different conditions may initiate abnormal grain growth [74]. Firstly, anisotropy in the grain boundary energy may lead to rapid growth of grains having boundary energies much lower than the average [75]. This is often the case when the material exhibits a strong primary recrystallization texture [76]. Secondly, abnormal grain growth may arise from anisotropy in the grain boun‐ dary mobility [74, 77].

**Acknowledgements**

INDIA.

**Author details**

Rajat K. Roy

**References**

Society), (1993), p. 3.

p. 333.

The author acknowledges helpful discussions with Prof. S. Das, Prof. K. Das, Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kaharagpur 721302,

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95

[1] P. Furrrer, G. L. Schenks, and R.E. Upton, *Aluminium Alloys for Packaging*, Eds. By J. G. Morris, H. D. Merchant, E. J. Westerman and P. L. Morris (Warrendale, PA: The

[2] E. J. Westerman, *Aluminium Alloys for Packaging*, Eds. By J. G. Morris, H. D. Merchant, E. J. Westerman and P. L. Morris (Warrendale, PA: The Minerals, Metals & Materials

[4] F. J. Humphreys and M. Hatherley, *Recrystallization and Related Annealing Phenomen‐*

[3] R. K. Roy, S. Kar, K. Das and S. Das, *Mater. Let.,* Vol. 59 (2005), p. 2418.

[6] G. Von Vargha and G. Wassermann, *Metallwirtschaft,* Vol. 12 (1933), p. 511.

[9] W. Truszkowski, J. Krol and B. Major, *Metall. Trans. A,* Vol. 11A (1980), p. 749.

[10] W. Truszkowski, J. Krol and B. Major, *Metall. Trans. A,* Vol. 13A (1982), p. 665.

[11] J. Grewen and J. Huber, in : F. Hassener (Ed.), Recrystallization of metallic materials,

[13] H. W. F. Heller, J. H. van Drop, G. Wolff and C. A. Verbraak, *Met. Sci*., Vol. 15 (1981),

[7] I. L. Dillamore and W. T. Roberts, *J. Inst. Met*., Vol. 92 (1963-64), p. 193.

[8] H. O. Asbeck and H. Mecking, *Mater. Sc. Engg*., Vol. 34 (1978), p. 111.

MST Division, CSIR-National Metallurgical Laboratory, Jamshedpur, India

Minerals, Metals & Materials Society), (1993), p. 144.

*on,* Published by Elsevier Science Ltd., (2002).

[5] J. Hirsch and K. Lücke, *Acta Metall*., Vol. 36 (1988), p. 2863.

Dr. Riederer Verlag GmbH, Stuttgart, 1978, p. 111.

[12] H. Hu and R. S. Cline, *Trans. TMS-AIME*, Vol. 224 (1962), p. 784.

Hillert has deduced three conditions for the development of abnormal grain growth in a material as follows [73]:

Whether these conditions are automatically fulfilled in a material where the normal grain growth has stopped growing due to the presence of second phase particles, is a question of considerable practical importance. Again, the growth of a very large grain in a material where the normal grain growth has stopped depends upon the value of the final grain size reached by the normal grain growth. The mathematical analysis, on the other hand, predicts that normal grain growth should proceed up to the limit 1/2*z*. This causes abnormal grain growth to be impossible for a limited time period when the normal grain growth has started to slow down but has not yet reached the limit. Therefore, initiation of abnormal grain growth is possible by a continuous decrease in the *z* value. This effect can be accomplished increasing the particle size *r* through coalescence or decreasing the volume fraction *f* by dissolving the second phase. Any process that leads to a slow increase in the grain size limit may initiate the development of abnormal grain growth. Most cases of abnormal grain growth, met within practice, seem to be connected with the dissolution of the second phase rather than the coalescence [73].