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able of contents1. Effect of Concurrent Precipitation on Recrystallization and Evolution of the P-Texture Component in a

Commercial Al-Mn Alloy.................................................................................................................................. 1

24 March 2014 ii ProQuest


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Document 1 of 1

Effect of Concurrent Precipitation on Recrystallization and Evolution of the P Texture Component in aCommercial Al Mn AlloyAuthor: Tangen, S; Sjlstad, K; Furu, T; Nes, EProQuest document link

Abstract: The recrystallization behavior of a cold-rolled Al-Mn alloy was investigated, focusing on the effect ofconcurrent precipitation on nucleation and growth of recrystallization and the formation of the P- ({011}[left angle

bracket]111[right angle bracket]) and ND-rotated cube ({001}[left angle bracket]310[right angle bracket]) texture

components. It was observed that if precipitation took place prior to or simultaneously with recovery and

recrystallization processes, i.e., by concurrent precipitation, this resulted in a delayed recrystallization, a coarse

and elongated grain structure, and an unusually sharp P-texture component. The P-texture component

sharpened with increasing initial cold rolling reduction, increasing initial supersaturation of Mn, and decreasing

annealing temperature. The P- and ND-rotated cube nucleation sites have an initial growth advantage

compared to the particle-stimulated nucleation (PSN) sites due to their 40 deg [left angle bracket]111[right angle

bracket]-rotation relationship to the Cu component of the deformation texture. The boundaries between such

sites and the surrounding matrix will be of the 7 type, and it is assumed that such highly perfect boundaries

will be less affected by solute segregation and precipitation, resulting in early growth advantage. It was further

observed that dispersoids present prior to cold rolling and annealing had a weaker effect on the recrystallized

grain size and texture compared to concurrent precipitation, even though the average dispersoid density was

higher in the pre-precipitation cases. The finer grain size was explained by the wider dispersoid free zones

surrounding the large constituent particles compared to the concurrent precipitation cases. Subsequent growth

of the nucleated grains, however, was more hindered due to the Zener drag, consistent with the higher

dispersoid densities. [PUBLICATION ABSTRACT]

Full text: HeadnoteThe recrystallization behavior of a cold-rolled Al-Mn alloy was investigated, focusing on the effect of concurrent

precipitation on nucleation and growth of recrystallization and the formation of the P- ({011}[left angle

bracket]111[right angle bracket]) and ND-rotated cube ({001}[left angle bracket]310[right angle bracket]) texture

components. It was observed that if precipitation took place prior to or simultaneously with recovery and

recrystallization processes, i.e., by concurrent precipitation, this resulted in a delayed recrystallization, a coarse

and elongated grain structure, and an unusually sharp P-texture component. The P-texture component

sharpened with increasing initial cold rolling reduction, increasing initial supersaturation of Mn, and decreasing

annealing temperature. The P- and ND-rotated cube nucleation sites have an initial growth advantage

compared to the particle-stimulated nucleation (PSN) sites due to their 40 deg [left angle bracket]111[right angle

bracket]-rotation relationship to the Cu component of the deformation texture. The boundaries between such

sites and the surrounding matrix will be of the 7 type, and it is assumed that such highly perfect boundaries

will be less affected by solute segregation and precipitation, resulting in early growth advantage. It was further

observed that dispersoids present prior to cold rolling and annealing had a weaker effect on the recrystallized

grain size and texture compared to concurrent precipitation, even though the average dispersoid density was

higher in the pre-precipitation cases. The finer grain size was explained by the wider dispersoid free zones

surrounding the large constituent particles compared to the concurrent precipitation cases. Subsequent growth

of the nucleated grains, however, was more hindered due to the Zener drag, consistent with the higherdispersoid densities.

DOI: 10.1007/s11661-010-0265-8

The Minerals, Metals &Materials Society and ASM International 2010

24 March 2014 Page 1 of 13 ProQuest
http://search.proquest.com/docview/863468633?accountid=46437http://search.proquest.com/docview/863468633?accountid=46437


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(ProQuest: ... denotes formulae omitted.)

I. INTRODUCTION

ALUMINUM-MANGANESE alloys (AA3xxx) dominate in the automobile heat exchanger industry. More than 95

pct of all heat exchangers, such as condensers, evaporators, radiators, and oil coolers, are based on Al-Mn

alloys due to their excellent combination of strength, ductility, corrosion resistance, brazeability, and

affordability. Grain-size control is essential for heat exchanger applications. A coarse grain size is desirable formost heat exchanger applications, as this gives a high corrosion resistance, high sagging resistance, and good

brazeability. In contrast, a fine grain size is preferred for products that require high formability as, for example,

thin wall tubes. To get the optimal grain size, it is crucial to control the alloy composition, the thermomechanical

processing, the size, the distribution of small and large particles, and the amount of elements in supersaturated

solid solution throughout the entire process route.

It is well established that second-phase particles have a strong effect on the recrystallization kinetics, final grain

size, and texture.[1-6] Deformation zones may develop around large particles (>1 m) during deformation,

which may activate particle-stimulated nucleation (PSN) of recrystallization,[1,2] whereas small, closely spaced

dispersoids have the ability to retard (Zener pinning) both low- and high-angle grain boundary motion.[7,8]

During thermomechanical processing or annealing of a deformed and supersaturated material, recovery as well

as recrystallization may be influenced by the precipitation reaction, a phenomenon which commonly is referred

to as concurrent precipitation. The different precipitation phenomena that may occur in the concurrent

precipitation regime were first reported by Hornbogen[9] and Kster,[10] while Nes and Embury[3] were the first

to report on the abnormal grain size that may result from this reaction. Later works by Nes and co-workers[5,6]

focused on the texture aspects associated concurrent precipitation. This phenomenon will here be subjected to

a new and focused study made possible by the recent developments in the electron backscattering diffraction

(EBSD) technique in high-resolution scanning electron microscopy (SEM).

Since the recrystallization texture determines the plastic anisotropy of the material, it has received increasing

interest over the last decades. For hot-rolled and annealed sheets, the most typical recrystallization texture

observed for materials with high stacking fault energy, such as aluminum, is the well-known cube texture. The

cube texture develops during hot deformation, as the cube orientation is metastable at high temperatures (e.g.,

References 11 and 12). After annealing of cold-rolled sheets, however, more random textures, which are

nucleated at grain boundaries and by PSN, are commonly observed. Over recent years, the P- and ND-rotated

cube textures have occasionally been observed in aluminum alloys after cold rolling and annealing.[6,12-18] In

general, these latter orientations have been observed after annealing of cold-rolled Al-Mn alloys with a highly

supersaturated solid solution, large cold rolling reductions, and low annealing temperatures.

The main objective of the present work was to investigate the texture formation during recrystallization of a

supersaturated commercial Al-Mn alloy. The focus was on establishing an understanding of the effect ofconcurrent precipitation on nucleation and growth of the P- and ND-rotated cube orientations, and the resulting

grain size. Furthermore, in order to contrast the effect due to concurrent precipitation with that caused by a fine

dispersion of particles present prior to annealing, this study also includes cases where a high number of Mn-

bearing dispersoids were precipitated prior to cold rolling and final annealing.

II. EXPERIMENTAL PROCEDURE

The examined material was a commercial DC-cast AA3103 extrusion ingot (0.57 wt pct Fe, 1.0 wt pct Mn, 0.12

wt pct Si, and other elements


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temperature-dependent term of Matthiessen's rule, the relationship between electrical conductivity and the solid

solution content (wt pct) for the alloy conditions studied here is given by (Althenpohl[19]):

... [1]

where is the electrical conductivity (m/ohm mm^sup 2^). It should be emphasized that the values of Mn in

solid solution given in Table I are only estimates, as the true effects of other elements in solid solution, particle

shape, number density, and type are not fully accounted for in the calculations.After the different heat treatments, the materials were water quenched to room temperature, and the four

material variants were deformed by cold rolling to true strains of 0.5, 1.5, and 3.0. Finally, the rolled sheets were

isothermally annealed (flash annealed) at different temperatures and times. Both hardness and electrical

resistivity measurements were carried out on the sheet surfaces in order to follow the softening and precipitation

reactions during annealing. Time-temperaturetransformation (TTT) diagrams were constructed on the basis of

these measurements, where 25 pct softening was defined as the onset of recrystallization and a 2.5 pct

increase in the electrical conductivity was defined as the start of precipitation. The recrystallized grain sizes

have been measured in the longitudinal cross section defined by the rolling and normal directions (RD-ND

plane) of the sheet midthickness, both by polarized light optical microscopy and the EBSD technique in a

scanning electron microscope. Grain-size diagrams have also been constructed. These sheet materials will in

the following be referred to as the "supersaturated sheets."

In addition, a set of samples from the as-cast (0) and A-homogenized (A) materials were pretreated to give

sheet variants with extremely high dispersoid number densities and nearly no Mn in supersaturated solid

solution. The pretreatment consisted of cold rolling to a true strain of 0.5 (39 pct reduction), annealing at either

573 K (300 C) or 623 K (350 C) for 10^sup 6^ s, and finally, cold rolling to an accumulated true strain of 3.0

(95 pct accumulated strain). An overview of these material conditions and their denotations is given in Table II.

These material variants will in the following be referred to as the "solute-free sheets" in order to distinguish them

from the supersaturated sheets mentioned previously.

The precipitation and softening behavior during annealing was investigated in greater detail by field emission

gun-scanning electron microscopy (FESEM), and the recrystallization textures were monitored by EBSD and X-

ray diffraction (XRD). Hence, both the local and bulk textures could be studied. In case of XRD, orientation

distribution functions (ODFs) were calculated from four incomplete pole figures, namely, the {200}, {111}, {220},

and {311} pole figures, according to the series expansion method.[20] All ODFs were further ghost corrected

using the method suggested by Lcke et al.[21]

III. EXPERIMENTAL RESULTS

A. Supersaturated Sheets

1. Characterization of the deformed state

The deformed state will have a strong influence on the subsequent annealing behavior and it is therefore ofspecial importance to characterize the microstructure and texture of this state. The average subgrain size (d)

and the misorientation across the subgrain boundaries ([varphi]) were calculated from EBSD orientation maps.

The critical misorientation to define a subgrain boundary was set to 1.5 deg. In the case of the C material (0.31

wt pct Mn in solid solution), increasing the true strain from 0.5 to 1.5, and 3.0, resulted in subgrain sizes of ...,

respectively. The corresponding misorientation between the subgrains varied from 2.4 to 4.1 deg and 4.5 deg.

Hence, the driving pressure for recrystallization, provided by the stored energy of the materials ..., generally

increases with strain and most of the increase in the stored energy is associated with the decreasing subgrain

size. The investigation showed that there was only a minor difference in the subgrain size and misorientation

between the different material conditions (0, A, B, and C) for a constant strain.

The particle size distributions of the deformed material conditions were measured by FESEM imaging and

automated image processing. These results showed that the as-cast sheet material contained considerably

fewer large particles and more small particles than the homogenized sheets. This difference will have an effect



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on the subsequent recrystallization annealing, as the Zener pinning pressure due to the small particles will

retard nucleation of recrystallization, while the largest particles will act as nucleation sites for recrystallization.

Further, it could be seen that there were no significant differences in the particle size distributions of the

homogenized and deformed conditions, indicating only minor particle breakup during rolling.

The development of the sheet midthickness deformation texture with increasing strain for the A-material

condition (0.72 wt pct Mn in solid solution), measured by XRD, is summarized as -fiber plots in Figure 1. Forthe investigated alloy, the deformation texture develops according to the theory of plane strain deformation of

aluminum alloys, starting with a random texture followed by a gradual buildup of the and fibers as the strain

increases. The deformation textures of the as-cast, B-homogenized, and C-homogenized material variants were

similar to those of the A condition. The only difference was a minor texture sharpening with increasing amount

of Mn in solid solution. These observations indicate that the Mn solid solution level has only a minor influence

on the texture formation during cold deformation of non-heat-treatable Al-Mn alloys. The strength of the cube

texture component after cold rolling was generally weak, around 1 to 3 times random, for all cold-rolled variants.

2. Softening and precipitation during annealing

The cold-rolled materials were isothermally annealed in salt baths, in the temperature range of 523 K to 773 K

(250 C to 500 C) and for times ranging from 5 to 10^sup 6^ s. The softening reactions were followed by

means of hardness testing, while the precipitation reaction was monitored by electrical conductivity

measurements. The interaction between precipitation and softening, in terms of the overall transformation

kinetics for the four Mn concentration levels, is summarized by the TTT diagrams in Figure 2. The background

experimental curves for generating these TTT diagrams can be found in Reference 22. These diagrams show

the effect of deformation on the softening and precipitation reactions for each of the four levels of Mn in solid

solution. The thick solid lines indicate the onset of precipitation, while the thin solid and broken lines represent

start and finished recrystallization, respectively.

A critical temperature, T^sub C^, was defined as the temperature where the precipitation curve of the TTT

diagram crosses the curve that indicates the start of recrystallization. Above this critical temperature,

recrystallization is little affected by concurrent precipitation of Mn-bearing dispersoids, while at temperatures

below T^sub C^, concurrent precipitation will retard recovery and recrystallization. Hence, the TTT diagrams

predict when recrystallization occurs prior to precipitation and when the softening reactions are slowed and

retarded by concurrent precipitation. The TTT diagrams confirm that an increased strain leads to a larger driving

pressure for recrystallization but simultaneously also to an increased driving pressure for precipitation of Mn-

bearing dispersoids. This is explained by a larger amount of stored deformation energy and a higher density of

microstructural heterogeneities available for nucleation of recrystallization and dispersoids. It can be seen that

the precipitation nose, i.e., the temperature where precipitation occurs most rapidly, is shifted toward shortertimes and higher temperatures with increasing initial solute Mn concentration. This results in a shift in the

finished recrystallization time toward longer times and higher temperatures. Consequently, recrystallization of

the as-cast strip (0.72 wt pct Mn in ss) becomes most retarded by concurrent precipitation, while the C strip

(0.31 wt pct Mn in ss) is more or less unaffected by concurrent precipitation at temperatures >573 K (300 C).

The information given by the TTT diagrams with respect to concurrent precipitation and the characteristic critical

temperature, T^sub C^, are summarized in Figure 3. T^sub C^ is plotted vs the nominal concentration of Mn in

solid solution prior to annealing for the three respective strain levels. The figure clearly demonstrates that the

critical temperature increases with initial supersaturation of Mn in solid solution. As a consequence, the

annealing temperature, which is necessary to avoid concurrent precipitation, increases with the initial solute Mn

concentration. No distinct variation in the critical temperature with strain is observed.

3. Effect of boundary misorientation on concurrent precipitation

When studying partially annealed sheet samples undergoing concurrent precipitation, it was observed that the



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dispersoids precipitated densely on some boundaries, while others were nearly dispersoid free (Figure 4 taken

from the A variant (0.47 wt pct Mn in ss) after annealing at 623 K (350 C) for 10^sup 3^ s). The combination of

backscattering electron (BSE) imaging and EBSD in an FESEM was used for this study. These investigations

demonstrated that the dispersoids precipitated preferentially on high-angle grain boundaries, typically with

misorientations >10 deg (cf. Figure 4). Microstructural observations of sheet samples annealed for longer times

showed, qualitatively, that the dispersoids were more homogeneously distributed on all boundaries with varyingboundary misorientations. These observations are is in accordance with recent work by Somerday and

Humphreys.[23]

The micrograph in Figure 4 also shows some quite thick, bright boundaries (arrows in the figure), which

correspond to high-angle grain boundaries. These boundaries are referred to as solute-rich zones. The reason

why these boundaries appear bright in the BSE micrograph is most probably due to a local clustering/

accumulation of elements with a higher atomic number than aluminum, i.e., clustering of Mn and Si. The grain

boundaries act as sinks for supersaturated elements due to a reduction of the total energy of the system. The

width of these boundaries/zones was typically ~1 to 10 nm, and they had a flakelike shape when viewed in

three dimensions, covering large parts of the boundaries. It is reasonable to believe that, with time at

temperature, the concentration of Mn eventually becomes so high at these zones or boundaries that Mn-bearing

dispersoids precipitate. Further, the dispersoids will grow as a result of continuous diffusion from the matrix

along the heterogeneities to the dispersoids. Hence, we can state that the solute zones are actually the

precursors of the grain boundary dispersoids. These solute zones and dispersoid effects will successfully retard

the softening processes by reducing the mobility of the grain boundaries. Since the elements diffuse along all

heterogeneities/ boundaries, there will be a retarding drag even on the low-angle grain boundaries. However,

the strongest drag will be on the most highly misoriented boundaries containing both a high concentration of

accumulated elements and a high number of dispersoids. As a result, the most potent high-angle grain

boundary nucleation sites experience the strongest Zener drag during annealing of a supersaturated material,

and the softening processes become efficiently retarded.

4. Recrystallization

a. Grain size and texture. The profound effect of concurrent precipitation on the recrystallized grain shape and

size is illustrated in Figure 5. When the sheets are annealed at temperatures above the critical temperature, TC,

the structure recrystallizes into a fine grain size (... cf. Figure 5(a)). In contrast, annealing below TC results in an

inhomogeneous and coarse grain structure, and the grains achieve a characteristic pancake shape (cf. Figure

5(b)). These observations are in accordance with earlier works (cf. References 3 through 6, 24, and 25). The

pancake-shaped grains resulting at low annealing temperatures are formed due to the lamellar alignment of the

dispersoids during concurrent precipitation. Because of the pancake shape of the deformed grains, the

dispersoids precipitate along the RD/TD plane, and hence, the growing grains experience the largest Zenerdrag in the direction normal to the rolling plane.[8] When annealing at temperatures below T^sub C^, the

recrystallized grain size generally increases with strain and with the initial solute Mn concentration, i.e.,

becoming coarsest for the as-cast variant (up to ...) and finest for the C material.

The recrystallized grain-size diagram in Figure 6 expresses in a condensed form the various grain structures

that may be expected for the current alloy system after cold deformation. The diagram indicates that in order to

achieve a fine-grained structure after isothermal annealing, regardless of the recrystallization temperature, only

~0.3 to 0.2 wt pct Mn can remain in solid solution after the thermal homogenization treatment.

The recrystallization textures were measured by means of XRD and in special cases by EBSD in an FESEM.

During isothermal annealing of a deformed and supersaturated condition, the selected annealing temperature is

crucial for the final recrystallization texture and grain structure in Al-Mn alloys. Figure 7 summarizes the

recrystallization textures achieved for the as-cast condition and the three homogenization variants (A, B, and C)

after cold rolling to strains of 3.0, as a function of the annealing temperature. The recrystallization textures are



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seen to follow the same tendency as the recrystallized grain size with respect to the initial solute content and

annealing temperature. At high annealing temperatures (T >T^sub C^), concurrent precipitation is avoided and

the recrystallization textures become more or less random with weak P ({011}[left angle bracket]111[right angle

bracket]) and cube ({001}[left angle bracket]100[right angle bracket]) components. The cube component is

generally seen to strengthen with decreasing initial concentration of Mn in solid solution and with increasing

annealing temperature, becoming as strong as 8 times random in the low solute B and C cases (0.38 wtpctMnand 0.31 wt pctMn in ss, respectively). It is interesting to note that the P component dominates with a

strength of 4 times random in the as-cast case (0.72 wt pctMn in ss) recrystallized at 773 K (500 C). At lower

annealing temperatures (T T^sub C^), the

recrystallized grain structures consisted of relatively small equiaxed grains with weak, predominantly cube

textures. Microstructural observations indicated that nucleation of recrystallization was controlled by PSN and

grain boundary nucleation under such circumstances (Figure 8). In contrast, the material conditions that

recrystallized under strong Zener drags from concurrent precipitation became coarse grained, with a sharp

texture controlled by the P- and ND-rotated cube components.

To study the nucleation efficiency of the particles under conditions of concurrent precipitation, the deformation

structure surrounding coarse constituent particles has to be considered. Figure 9(a) displays the deformation

structure around a constituent intermetallic particle after cold rolling to a strain of 3.0. The white broken lines in

the figure illustrate the extension of the characteristic deformation zone, which surrounds the particle after

deformation. Characteristic flow pattern can be seen to form around the hard particle. The deformation zone is,

according to Humphreys,[1] characterized by rotated zones to the left and right of the particle, and distorted

zones below and above the particle. A micrograph showing a closeup of a rotated zone close to a coarseparticle is presented in Figure 9(b). Such rotated zones correspond to a refined subgrain structure and high

lattice rotations, giving a high local stored energy, which is crucial for nucleation of recrystallization.

The micrograph in Figure 10, taken from the A-homogenized variant cold rolled to a strain of 3.0 and

subsequently annealed to a recovered state at 623 K (350 C) for 10^sup 3^ s, shows the formation of a

dispersoidfree zone surrounding the constituent eutectic particle after partial annealing. Note that the

surrounding matrix contains a very high number density of dispersoids, which to a large extent are situated on

the highangle grain boundaries. The dispersoid-free zone arises due to the depletion of Mn in supersaturated

solid solution around the constituent particles both during casting, homogenization, and annealing after cold

rolling (cf. References 26 through 28). Supersaturated Mn inside this zone will be drawn to the constituent

particle and contribute to its phase transformation and growth. The depleted and dispersoid-free zones will

further lead to a locally lower Zener drag during annealing of the deformed material. Consequently, nucleation

of recrystallization should be less affected by precipitation in these zones, as compared to the potential



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nucleation sites in the surrounding matrix that are hindered by a very high dispersoid number density.

B. Solute-Free Sheets

1. Characterization of the deformed state

The solute-free sheets were in contrast to the "supersaturated sheets" processed in such a way as to precipitate

a very high number density of small Mn-bearing dispersoids, which should in principle be randomly distributed,

prior to the final recrystallization annealing of the coldrolled strips. In Table II, the four different variants of thismaterial are denoted by their accumulated true strain, their initial state, and their intermediate annealing

temperature. Table II summarizes the calculated solute Mn contents in the initial state (after casting or

homogenization) and after thermomechanical processing to the deformed state (accumulated strain of 95 pct). It

can be seen that the two variants, which were intermediately annealed at 623 K (350 C), result in a very low

solute content in the final cold-rolled state (0.15 wt pct) and can hence be called solute-free. The two variants,

which were intermediately annealed at 573 K (300 C), do on the other hand contain a considerable soluteMn

content after processing (0.30 to 0.35 wt pct), which is relatively similar to the C variant (0.31 wt pct Mn in ss) of

the supersaturated sheets. This might result in some concurrent precipitation during final annealing and may

therefore enhance the total Zener drag.

The FESEM micrographs in Figure 11 represent the dispersoid structures resulting after completed processing

of the as-cast and A-homogenized variants, intermediately annealed at 623 K (350 C) for 106 s and cold rolled

to accumulated true strains of 3.0. It is clear that this process route results in a very high number density of

small (d ~ 50 to 100 nm) dispersoids, especially in the case of the as-cast variant, which initially had the highest

supersaturation of Mn prior to processing. The dispersoid number density per area was measured by counting

dispersoids on a number of micrographs for each of the four solute-free sheets (Table II). The dispersoid

number density of the two as-cast variants becomes around 21 to 23 m^sup -2^, which is higher than the

situation for the supersaturated sheet. In comparison the as-cast condition, after cold rolling and concurrent

precipitation annealing, typically gave a number density of around 10 to 16 m^sup -2^. The A-homogenized

variants of the solute-free sheets have dispersoid number densities of around 6 to 9 m^sup -2^ after annealing,

while annealing of the corresponding variant of the supersaturated sheet gave 1 to m^sup -2^.

2. Recrystallized structures and textures

The solute-free sheets were subsequently isothermally annealed in order to recrystallize the material for texture

and grain structure analyses. The effect of the high dispersoid number density is clearly seen in Figure 12,

where the softening kinetics of solute-free and supersaturated sheets at 673 K (400 C) are illustrated in a

hardness vs time plot. An annealing temperature above 723 K to 773 K (450 C to 500 C) was necessary to

fully soften the solute-free sheets within 106 s. The EBSD maps in Figure 13 represent the grain structure after

isothermal annealing at 773 K (500 C). The mean grain length (DRD) was measured to ~50 m in the as-cast

variants (Figure 13(a)) and ~12 m in the A-homogenized variant (Figure 13(b)).FESEM investigations of partially annealed sheet samples were performed in order to study the nucleation

behavior in the solute-free sheets containing a high number density of dispersoids prior to deformation and

annealing. Figure 14 shows a micrograph of a partially annealed specimen, and it is seen that nucleation of

recrystallization has taken place in bands along the constituent particles, while no nucleation is observed in the

areas without large particles. The image to the right shows an enlarged part of the area without any nucleation,

which contains a high fraction of dispersoids that effectively pin the structure and prevent nucleation of

recrystallization.

A more detailed study of the early stages of nucleation of recrystallization is shown in the FESEM micrograph in

Figure 15, illustrating the nucleation along bands of constituent particles in a partially annealed as-cast variant.

The dispersoid number density is clearly lower in the zones surrounding the coarse particles. These dispersoid-

free zones form during casting and further develop during intermediate annealing. Hence, similarly to the

supersaturated sheets, nucleation of recrystallization becomes less affected by Zener drag at the constituents



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than in the surrounding matrix.

The global recrystallization textures were measured by means of EBSD. In general, the recrystallization texture

becomes weaker for the solute-free sheets than for the supersaturated sheets. The two as-cast material

variants developed weak ND-rotated cube and P components, with intensities around 5 and 3 times random,

respectively. The two A-homogenized sheet variants, with a very fine recrystallized grain size, generally display

the same texture components but with a lower intensity, both around 2 times random. The appearance of theND-rotated cube and P components after recrystallization of the solute-free sheets confirm that there is indeed

an effect of the present dispersoids on the recrystallization texture evolution, although it is much weaker than in

the concurrent precipitation case described previously.

IV. DISCUSSION

The salient experimental observations regarding the annealing characteristics of the cold-rolled AA 3103 alloy

variants can be summarized as follows.

Annealing at temperatures so high that no concurrent precipitation took place (T >T^sub C^) resulted in a

finegrained recrystallized structure with a weak P texture and medium to strong cube texture. The strength of

the cube component increased with decreasing initial Mn in supersaturated solid solution.

Annealing at lower temperatures (T T^sub C^, become

so effectively restrained by concurrent precipitation at temperatures T


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due to a higher mobility of the reaction front than growing grains of random orientations. If that had been the

case, strong P and rotated cube textures observed in the present work should have resulted from

recrystallization also at temperatures T >T^sub C^. The idea proposed by Daaland and Nes was that the 40 deg

[left angle bracket]111[right angle bracket] grains had a nucleation and early growth advantage compared to

grains of more random orientations. This idea was based on the observation shown in Figure 16 that the

randomly oriented grains stopped growing after reaching an average size of about 5 to 7 m as a result ofconcurrent precipitation, while this precipitation reaction had less influence on the 40 deg [left angle

bracket]111[right angle bracket] grains. This idea will now be discussed in the context of the present, more

detailed metallographic observations.

EBSD examination by Daaland and Nes[6] demonstrated that in the very beginning of recrystallization the

density of randomly oriented grains, believed to be the result of PSN, was about an order of magnitude higher

than the density of grains with the rotated cube or P orientation. In the fully transformed condition, however, the

volume fraction of rotated cube grains was about 50 pct, while that of randomly oriented grains amounted to

only 9 pct. The fact that these randomly oriented grains indeed could be associated with PSN has been nicely

confirmed by the present investigation. During concurrent precipitation, the deformation zone recrystallizes

readily, while growth of recrystallization beyond and into the matrix is effectively stopped. This situation was

illustrated in Figure 10, taken from the A-homogenized condition. The solute depletion of the area just around

the particles is enhanced due to the homogenization treatment. In order to support this trend, an additional

experiment has been performed, selecting the as-cast variant cold rolled to a strain of 3.0 and subsequently

annealed at 623 K (350 C) for 10^sup 3^ s. The combination of FESEM BSE imaging and EBSD mapping was

used in order to display the recrystallized grains at an early stage of nucleation and growth and to measure their

corresponding orientations.

The EBSD maps in Figure 17 show a few examples of nucleated P- and ND-rotated cube grains. P- and ND-

rotated cube grains are circled with white and black dashed lines, respectively. The white nonindexed areas in

the EBSD maps correspond to constituent particles. Since the as-cast condition (0.72 wt pct Mn in ss) contains

a significantly higher initial amount of Mn in solid solution compared to the A-homogenized case (0.47 wt pct Mn

in ss) described previously, the effect of concurrent precipitation will be more severe. Consequently, an

annealing treatment of 103 s at 623 K (350 C) caused the as-cast variant to soften by only a 20 pct reduction in

hardness compared to 33 pct in the homogenized material, while the conductivity changes were 22 and 8 pct for

the as-cast and homogenized versions, respectively. Still, however, in this as-cast condition, cases of relatively

large recrystallized grains of the P and rotated cube orientations were found, as illustrated in Figure 17. This

confirms the hypothesis proposed by Daaland and Nes[6] that the nucleation and early growth of grains with this

40 deg h111i orientation relationship was much less affected by concurrent precipitation than grains of other

orientations. The nature of the nucleation sites for these orientations, however, has not yet been fully identified.It has been speculated that the P- and ND-rotated cube-oriented grains nucleate in the vicinity of the constituent

intermetallic particles by PSN (refer to the work of Engler and co-workers,[16,30] Ryu and Lee,[17] and

Sjlstad).[29] These investigators performed TEM and SEM studies and observed P-oriented grains to be

associated with coarse particles. However, one has to be aware that for any new grain of a few microns in

radius, the probability that such a grain will be in contact with a coarse particle is about 1, given that the density

of particles in commercial aluminum alloys with a size larger than 1 m in diameter is about 10^sup 16^ m^sup -

3^.

It is difficult to understand how 40 deg [left angle bracket]111[right angle bracket] grains can nucleate from the

turbulent zones of deformation zones. It is well established that 40 deg [left angle bracket]111[right angle

bracket] grains nucleate from bandlike features in commercial aluminum alloys.[5,6,31] However, still some

evidence has been provided that the grains of orientations close to the P have been detected in the deformation

zone[16,17] surrounding the large constituent particles. The explanation could very well be that the deformation



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zones recrystallized into a random orientation distribution due to the low Zener drag (depleted zones). However,

the growth of recrystallization beyond the deformation zones, in general, is stopped by a high solute and Zener

drag effect. Only the recrystallized deformations zones of the P and rotated cube orientations, which in addition

have a copper deformation component as their next neighbor, will be able to grow beyond the deformation

zones due to their 40 deg [left angle bracket]111[right angle bracket] orientation relationship. The reason why

these particular orientation relationships are less affected by concurrent precipitation is argued in the nextparagraph. One could also argue that P and rotated cube grains originate from matrix-type heterogeneities and

subsequently use a neighboring particle deformation zone as a growth environment.

The present observations clearly show that the P and rotated cube-oriented grains have an early growth

advantage. A plausible explanation for this can be found in the concurrent precipitation pattern shown

previously. First, solute elements segregate toward highangle boundaries, and then these boundaries become

decorated by precipitates. Somewhat later, precipitation is observed to take place also at low-angle boundaries.

On the assumption that the 40 deg [left angle bracket]111[right angle bracket] grains originate from matrix sites

(i.e., sites located with the Cu-deformation texture component as next neighbor), these sites will be separated

from the surrounding matrix by 7-type high-angle boundaries. It is reasonable to assume that such highly

perfect boundaries will be less affected by solute segregation and precipitation. Or, in other words, 40 deg h111i

grains will be less affected by concurrent precipitation, the result being a nucleation and early growth advantage

for these grains compared with grains of other orientations.

B. On the Effect of Dispersoids Present Prior to Annealing vs Concurrent Precipitation during Annealing

It was observed that annealing of the solute-free sheets gave a considerably finer grain size and a much weaker

recrystallization texture than the supersaturated sheets, even though the total number of dispersoids was

higher. On the other hand, the time to complete recrystallization was considerably longer in these "sheets" as

compared with the supersaturated sheets, which were affected by concurrent precipitation. Regarding the finer

grain size, nucleation of recrystallization was observed to take place during the early stages of annealing of the

solute-free sheets and thus mainly inside the dispersoid-free zones surrounding the constituent particles (cf.

Figure 15). It is believed that the high nucleation frequency at the constituent particles is caused by the

coarsening of the surrounding dispersoid-free zones during the intermediate annealing for 10^sup 6^ s, which

the solute-free sheets were subjected to prior to cold rolling. These zones are considerably larger and contain

fewer dispersoids than the corresponding zones in the supersaturated sheets. This might also explain why

these materials develop a rather weak texture after recrystallization. A higher number of PSN sites, with a large

variety of orientations, are able to grow to an overcritical size and become recrystallized grains. However, even

though the grains are easily nucleated and reach an overcritical size for growth in the dispersoid-free zones

close to the constituent particles, the continued expansion will be strongly retarded as the matrix contains a very

high dispersoid number density (cf. Figure 12). It should, however, be mentioned that the total driving pressurefor recrystallization is somewhat lower in the case of the solute-free sheets due to the "intermediate"

precipitation annealing at a true strain of 0.5. This can explain some of the difference in the recrystallization

kinetics between the two sheets seen in Figure 12.

Burger et al.[32] investigated the recrystallization texture after rolling of an Al-Mn alloy, pretreated differently to

achieve various dispersoid number densities prior to recrystallization annealing. A high temperature was used to

soften the material conditions. They found the P texture (termed RX texture in the actual article) to sharpen up

to a maximum intensity of 2 to 3 times random as the dispersoid density increased. The P texture was observed

together with the ND-rotated cube component, which was found to be less rotated as the dispersoid content

decreased. These observations further confirm that dispersoids present prior to recrystallization annealing have

a considerably less prominent effect on the development of the P-texture component compared to concurrent

precipitation.

V. CONCLUSIONS



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The details around the formation of the precipitation and recrystallization behavior of Al-Mn alloys can be

summarized as follows.

Concurrent precipitation in a supersaturated Al-Mn alloy strongly affects the recrystallization behavior, leading to

inhomogeneous and coarse-grained recrystallized structures. The degree of concurrent precipitation increases

with rising initial supersaturation of Mn and with raising prior cold reduction due to an enhanced diffusion rate

and higher number of heterogeneities available for nucleation of dispersoids.The recrystallization texture in the case of concurrent precipitation of Mn-rich dispersoids is dominated by the P-

and ND-rotated cube components. Unusually sharp P textures were observed in the present work. The intensity

of the P- and ND-rotated cube textures strengthens with increasing initial cold rolling strain, supersaturation of

Mn, and dispersoid density, but decreases with increasing annealing temperature.

The P- and ND-rotated cube nucleation sites have an initial growth advantage compared to the PSN sites due

to their 40 deg h111i-rotation relationship to the Cu component. The boundaries between such sites and the

surrounding matrix will be of the 7 type, and it is assumed that such highly perfect boundaries will be less

affected by solute segregation and precipitation, resulting in the early growth advantage.

Dispersoids present prior to annealing were seen to have a much weaker effect on the recrystallized grain size

and texture than concurrent precipitation, although the total dispersoid number density was much higher. The

growth of the recrystallized grains was thus more hindered than the nucleation process, giving a high nucleation

frequency but a very slow recrystallization rate. Thus, the sheets recrystallized into a rather fine grain size with a

weak texture. The high nucleation rate was attributed to the randomization of the dispersoids during rolling in

combination with the very large precipitate-free zones surrounding the coarse constituent particles after finalized

rolling, which made PSN relatively easy.

ACKNOWLEDGMENTS

This research was carried out as a part of the NFRproject Heat Treatment Fundamentals KMB project (Project

No. 143877/213) in the subproject Nucleation of Recrystallization. Funding by the industrial partners, Hydro

Aluminium, Raufoss ASA, and Elkem, is gratefully acknowledged.

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18. W.C. Liu and J.G. Morris: Metall. Mater. Trans. A, 2005, vol. 36A, pp. 2829-48.

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23. M. Somerday and F.J. Humphreys: Mater. Sci. Technol., 2003, vol. 19, pp. 20-29.24. E. Nes: Aluminium, 1976, vol. 52, pp. 560-63.

25. P. Furrer and H. Warlimont: Aluminium, 1978, vol. 54, pp. 135-42.

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AuthorAffiliationS. TANGEN, Project Manager, and T. FURU, Senior Advisor, are with Hydro Aluminium, Research and

Technology Development, N-6601 Sunndalsora, Norway. Contact e-mail: stian.tangen@hydro. com K.

SJLSTAD, Project Manager, and E. NES, Professor, are with the Department of Materials Technology, NTNU,

N-7491 Trondheim, Norway.

Manuscript submitted March 12, 2009.

Article published online July 15, 2010

Subject:Alloys; Metallurgy; Cold; Temperature; Studies; Scanning electron microscopy; Annealing; Corrosionresistance; Heat exchangers; Aluminum;

Publication title: Metallurgical and Materials Transactions

Volume: 41A

Issue: 11

Pages: 2970-2983

Number of pages: 14

Publication year: 2010

Publication date: Nov 2010

Year: 2010

Publisher: Springer Science & Business Media

Place of publication: Warrendale

Country of publication: Netherlands

Publication subject: Metallurgy, Engineering

ISSN: 10735623

CODEN: MMTAEB

Source type: Scholarly Journals



15/15

Language of publication: English

Document type: Feature

Document feature: Tables Graphs Photographs References

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