1 Introduction

Continuous cast (CC) technology for production of Al alloys has attracted extensive attention from Al alloy makers, because it offers a significant energy and economic savings in Al alloy sheet production over the direct chill (DC) ingot casting method, which is conventionally used to produce Al alloy sheet products.[1,2] Compared with the Al sheet products produced by the traditional DC route, the CC Al alloys retain a high level of supersaturation due to the rapid solidification rate and omission of homogenization treatment in the CC method. Fine precipitates could then be generated in these supersaturated CC Al alloys during their downstream thermomechanical processing, which could have a significant effect on the development of microstructure and texture in the final O-temper Al alloys. The mechanical properties of the O-temper CC Al alloy sheet could also be influenced by these precipitates, depending on their sizes, densities, and distributions in the alloys. Fine second-phase particles (<1 μm) may exert a strong pinning effect on dislocations and grain boundaries, thus retarding the recrystallization process, while large particles (≥1 μm) may act as potential grain-nucleation sites by the so-called particle stimulation nucleation (PSN), thereby promoting recrystallization in Al alloys.[3–10]

As compared to the commonly occurring cube texture in the recrystallized DC Al alloys, annealing the deformed and supersaturated CC Al alloys could result in the formation of uncommon recrystallization textures due to precipitation taking place upon heating in recrystallization, which is called the concurrent precipitation effect. These uncommon textures include P {011}\( \left\langle {566} \right\rangle \), normal direction (ND)–rotated cube {001}\( \left\langle {310} \right\rangle \), and {113}\( \left\langle {110} \right\rangle \) orientations.[1,2,10–12] It has been reported that these “uncommon” textures occur in Al alloys due to the effect of solid solution. Magnesium in solid solution reduces recovery, which leads to formation of shear bands during cold rolling, and during subsequent annealing, uncommon texture components occur at the shear bands at the expense of cube texture components.[13,14] However, it is still desirable to further study the formation of recrystallization textures and microstructure without the influence of concurrent precipitation in CC Al alloys, with a view to optimizing the texture and microstructure in these alloys.

The occurrence of P orientation has previously been observed in DC Al-Cu and Al-Mn alloys,[6–9] though the P orientation is relatively weak, as compared with that in the CC Al alloys.[5,12,16] It has been proposed that the formation of P orientation in CC Al alloys is attributed to concurrent precipitation taking place upon slow heating during recrystallization.[10,12–16] The precipitates so produced have a strong pinning effect on dislocations and grain boundaries, thereby retarding the recrystallization process in the CC Al alloys. Engler[17] suggested that P orientation was originated from the extensive plastic deformation zones in the vicinity of large particles due to PSN during annealing. However, it is still unclear how the pinned dislocations in these deformation zones lead to nucleation of P-oriented grains.

In the present study, a heat treatment prior to cold rolling and final annealing was carried out to minimize the concurrent precipitation effect taking place during the final annealing in a CC AA3004 Al alloy, in an attempt to further understand the formation mechanism of P orientation in the alloy. The results obtained from this study revealed that a strong P orientation was also formed in the alloy with reduced concurrent precipitation, and that the precipitates of Al6(Mn,Fe) formed either before or during annealing could lead to formation of a strong P texture in this alloy.

2 Experimental

The material used in this work was a hot band of a commercial CC 3004 Al alloy (2.54-mm thick) produced by industrial practice. In the CC production process of Al sheet, the molten Al alloy was poured into two rotating steel belts to produce a cast slab, which was then immediately fed into three consecutive hot rolling mills to form hot band Al products. The chemical composition of this alloy is given in Table I. Since the slab was formed under water cooling and no homogenization treatment was conducted on the CC Al alloy hot band, the as-received hot band showed a high level of solid solution of alloying elements such as Mn, Si, and Fe. Annealing treatment therefore could lead to the formation of second-phase particles in the alloy.[18] In order to investigate the effect of precipitates on the microstructures and textures of the alloy during downstream thermomechanical processing, some of the as-received hot band was given a prior annealing treatment at 420 °C for 3 hours with a heating rate of 1 °C/min from room temperature before cold rolling, while the rest of the hot band was directly cold rolled. This annealing temperature was chosen because it led to the maximum precipitation in the alloy, i.e., the minimum resistivity of the alloy.[19] Figure 1 illustrates the different thermomechanical processing routes of these specimens. The specimens with or without prior heat treatment were cold rolled at reductions of 25, 50, 70, 80, 85, and 90 pct, respectively, and subsequently annealed at 510 °C for 4 hours to recrystallize these samples, followed by air cooling.

Table I Chemical Composition of the As-Received CC AA 3004 Al Alloy
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Fig. 1
figure 1

Schematic illustration of two thermomechanical processing routes of the as-received CC AA 3004 Al alloy hot band

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The cold-rolled and annealed samples for microstructural examination were cut along the plane of the rolling direction (RD) and ND. These samples were electropolished and anodized to reveal their microstructures using an optical microscope with a polarized light.

The measurement of macrotextures was performed using the Schulz reflection method on a Rigaku diffractometer (Rigaku, Akishima City, Tokyo, Japan) using Cu K α radiation.[20] Four incomplete pole figures, (111), (200), (220), and (311), were measured up to a maximum tilt angle of 75 deg. A two-step method based on the series expansion method was employed to calculate orientation distribution functions (ODFs), from which the volume fractions of texture components were determined.[21] The ODFs were presented as plots of constant φ 2 sections with isointensity contours in Euler space defined by the Euler angles φ 1, Φ, and φ 2. The volume fraction of each texture component was calculated by integrating the intensities of the ODF f(g) within 15.5 deg of its theoretical orientation in the subset of Euler space.

3 Results

3.1 Microstructures and Textures of CC AA 3004 Al Alloys before Cold Rolling

Figure 2 shows the microstructures of CC AA 3004 Al alloy hot band in the as-received and annealed (at 420 °C) conditions, respectively. The as-received hot band (Figure 2(a)) exhibited a grain structure severely elongated in the RD, which was typical for a hot-rolled Al alloy. Some fine newly formed grains were observed in the annealed hot band (Figure 2(b)), indicating that the annealing temperature used was the onset of recrystallization in the hot band.

Fig. 2
figure 2

Microstructures of the as-received CC AA 3004 Al alloy hot band observed using optical microscopy with polarized light: (a) hot band and (b) hot band with prior heat treatment at 420 °C

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Figure 3 presents the ODFs of the hot band samples as received and annealed at 420 °C, respectively. The as-received hot band (Figure 3(a)) showed a typical β fiber texture (running from copper orientation {112}\( \left\langle {111} \right\rangle \) through S orientation {123}\( \left\langle {634} \right\rangle \) to brass orientation {011}\( \left\langle {211} \right\rangle \)) and a weak cube component {001}\( \left\langle {100} \right\rangle , \) which was consistent with the deformed grain structure observed in the sample (Figure 2(a)). The ODF of the sample annealed at 420 °C (Figure 3(b)) also showed a strong β fiber texture, though there were some fine newly formed grains in this sample (Figure 2 (b)).

Fig. 3
figure 3

ODFs of CC 3004 Al alloy hot band: (a) as-received and (b) heat treated at 420 °C

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Since the as-received CC Al alloy hot band had a high level of supersaturation due to the high cooling rate in slab casting, annealing the as-received hot band at an appropriate temperature could reduce its solid solution level by precipitating second-phase particles, which was revealed by measurement of the electrical resistivity of the alloy before and after annealing.[18] The electrical resistivity was measured to be 5.26 μΩ cm for the as-received hot band and 3.60 μΩ cm (i.e., the minimum resistivity obtained) for the alloy annealed at 420 °C. The difference in electrical resistivity between these two samples was 32 pct, indicating that a significant amount of second-phase particles were precipitated in the sample after annealing at 420 °C. Transmission electron microscopy (TEM) observations verified that densely distributed second-phase precipitates ranging from 20 nm to 0.2 μm in size existed in the preannealed sample, as shown in Figure 4(b). In contrast, few fine precipitates were observed in the as-received hot band (Figure 4(a)). These precipitates have previously been identified as Al6(Mn,Fe) in CC Al alloys.[22–25] Figure 5 demonstrates profoundly that these precipitates interact with dislocations in the preannealed sample. As discussed later in this article, these fine precipitates could have a significant effect on the formation of textures and grain structure in the alloy after subsequent cold rolling and recrystallization, since they had a strong pinning effect on dislocations as well as grain boundaries.

Fig. 4
figure 4

TEM micrographs showing precipitates of CC AA3004 Al alloy hot band: (a) the as-received and (b) preannealed at 420 °C. Precipitates are pinning dislocations in (b)

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Fig. 5
figure 5

TEM micrograph showing precipitates interacting with dislocations in the alloy preannealed, 50 pct cold rolled and annealed at 371 °C

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3.2 Recrystallization Microstructures after Cold Rolling

Both the as-received and preannealed hot band were cold rolled at different reductions between 25 and 90 pct, followed by annealing at 510 °C for 4 hours. Figure 6 shows the final recrystallized grain structures in these samples. It can be seen in Figure 6 that grains were coarse and severely elongated in the preannealed samples after cold rolling at all the rolling reductions and recrystallization, while those in the samples without the prior heat treatment appeared to be relatively less elongated and finer. This demonstrated that the precipitates produced during the preannealing, as shown in Figure 5, were more effective in interfering with the recrystallization process than those generated during concurrent precipitation in this alloy.

Fig. 6
figure 6

Recrystallization grain structures of the CC AA 3004 Al alloys without (I) and with (II) the heat treatment at 420 °C prior to cold rolling at reductions of (a) 25 pct, (b) 70 pct, and (c) 90 pct followed by annealing at 510 °C for 4 h

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As shown in Figure 6, the grain size decreased with increased rolling reduction in the final recrystallized samples. The coarsest grain structure, typically 400 × 320 × 100 μm3, appeared in the preannealed sample at 25 pct rolling reduction, after recrystallization. It was possible that rolling reduction of 25 pct was close to the critical strain for secondary grain growth in the preannealed alloy. Although the grain size decreased with increased rolling reduction, the average aspect ratio of grains still increased with rolling reduction in the preannealed samples, while that in the samples without the prior heat treatment did not change significantly with the rolling reduction, as shown in Figure 7. At 90 pct rolling reduction, the average aspect ratio of grains in the preannealed sample was as large as 7 (Figures 6(c) and 7). The elongated grain structures are commonly formed as a result of the concurrent precipitation effect taking place upon heating during recrystallization in CC Al alloys.[26–28] Fine precipitates are generated at large-angle grain boundaries, which are predominantly parallel to the RD during slow heating, thereby making it hard for recrystallized grains to grow in the direction perpendicular to the rolling plane. However, the severely elongated grains observed in the preannealed sample can not simply be accounted for by the concurrent precipitation effect, as the preannealed sample was not supersaturated before cold rolling and recrystallization, i.e., the concurrent precipitation was minimized in this alloy during recrystallization. The formation of the elongated grains was related to the existence of a large amount of precipitates produced during the preannealing in the alloy (Figures 4(b) and 5). These precipitates interacted with large-angle grain boundaries, which were predominantly aligned in the RD, especially after cold rolling, which reduced the spacing between precipitates in the direction perpendicular to the rolling plane in the preannealed sample. Figure 4(b) demonstrates that these precipitates had a strong pinning effect on dislocations in the preannealed alloy before cold rolling. It was evident that these precipitates strongly pinned dislocations and large-angle grain boundaries after cold rolling and during subsequent recrystallization (Figure 5). During recrystallization, the boundaries between the newly formed grains and the deformed grains were hardly advanced across these pinned large-angle grain boundaries in the direction perpendicular to the RD, hence giving rise to the formation of an elongated grain structure in this alloy. Consequently, the aspect ratio of grains increased with the increase in rolling reduction, since the spacing between precipitates was decreased in the direction normal to the rolling plane with the increase in rolling reduction. In the samples without the prior heat treatment, fine precipitates were also formed at dislocations and grain boundaries in the deformed samples upon heating. Therefore, their distribution in this sample was less anisotropic than that in the preannealed sample in the RD, since there was no further rolling deformation after the concurrent precipitation taking place, unlike, in the preannealed sample where the rolling deformation was conducted after the precipitation process. The postdeformation significantly reduced the spacing in the direction normal to the rolling plane. Therefore, the recrystallized grain structure was finer and less elongated in the RD, since resistance to grain growth was more isotropic in this alloy than in the preannealed sample.

Fig. 7
figure 7

Aspect ratio of grains as a function of rolling reduction in samples with and without prior heat treatment

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3.3 Recrystallization Textures after Cold Rolling

Figure 8 shows the ODFs, at sections of φ 2 = 0 and 45 deg, respectively, of the recrystallized alloys without and with the prior heat treatment. It can be seen in Figure 8 that P orientation is the dominant texture in both these samples. Although the strong P orientation could be anticipated to appear in the sample without the prior heat treatment due to the concurrent precipitation effect,[5,11,29] the occurrence of a strong P orientation has never been observed in CC Al-Mn-Mg alloys without concurrent precipitation effect. The specimen without the prior heat treatment at 25 pct rolling reduction also showed two fiber textures, the axes of which were parallel to \( \left\langle {100} \right\rangle \) and \( \left\langle {110} \right\rangle \) in the ND, respectively. Minor Goss and P orientations were parts of the \( \left\langle {110} \right\rangle \)//ND fiber. With increasing rolling reduction, the intensity of the P orientation was strengthened at the expense of the remaining components, as evident in Figure 8(a), where the P orientation became stronger and sharper with an increase in rolling reduction. A very strong P orientation accompanied by minor ND-rotated cube, Goss, and cube orientations was observed in the sample without prior annealing at a large rolling reduction (over 80 pct).

Fig. 8
figure 8

ODFs at sections of φ 2 = 0 and φ 2 = 45 deg of the CC AA 3004 Al alloys: (a) without the prior heat treatment and (b) with the prior heat treatment at 420 °C, after cold rolling at different reductions followed by annealing at 510 °C

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Figure 8(b) shows the evolution of recrystallization texture in the preannealed samples as a function of rolling reduction. Analogous to the sample without the prior heat treatment, P orientation was also strengthened by increasing rolling reduction in the sample. However, the cube component was stronger (about 5.15 pct in volume fraction) in this sample than that (about 2.33 pct) in the corresponding sample without prior heat treatment (Figure 10). This was likely to be caused by the fact that the particles around 0.2 μm or larger in the preannealed sample were much more than those in the sample without the prior annealing, as shown in Figure 4. It has been recognized that coarse particles, as discussed earlier in this article, could promote nucleation of cube grains in Al alloys.

Figure 9 depicts a plot of the volume fraction of P orientation as a function of cold rolling reduction in the alloys with or without the prior heat treatment. It can be seen that the volume fraction of P texture increased with increasing rolling reduction in both these samples with or without prior heat treatment, which indicated that rolling reduction was an important factor controlling the formation of P texture. Since the rolling reduction was directly associated with the plastic strain imparted by cold rolling in the alloy, the formation of P orientation was related to the dislocation density in the alloy. Namely, the formation of P orientation should also be a function of dislocation density generated by cold rolling in the alloys. It can also be seen in Figure 9 that the volume fraction of P orientation in the preannealed sample is higher than that in the sample without prior heat treatment when the cold rolling reduction exceeds 75 pct. This might be attributed to the precipitates produced by the prior annealing being more effective than those generated by concurrent precipitation in pinning dislocations and grain boundaries in the subsequent recrystallization treatment, when the rolling reduction was over 75 pct. The precipitates generated by concurrent precipitation were less in quantity than those by prior annealing; thereby, they could not effectively pin as many dislocations and grain boundaries as those by the precipitates formed in prior annealing. More P orientation therefore was generated in the preannealed sample when the rolling reduction exceeded 75 pct.

Fig. 9
figure 9

Volume fractions of major texture components as a function of rolling reduction in the samples with or without prior heat treatment after annealing at 510 °C

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4 Discussion

It has been recognized that P orientation can be formed as a result of concurrent precipitation in the CC Al alloys containing Mn, especially CC AA3000 series Al alloys.[6–10,12–15,26] The hot band of CC 3004 Al alloy without prior heat treatment retained a high level of solid solution, especially supersaturated Mn, due to rapid solidification in slab casting. Subsequent cold deformation and annealing at a very low heating rate (e.g., 1 °C/min) resulted in precipitation of a Mn bearing phase (i.e., Al6(Mn,Fe)) occurring concomitantly with recrystallization. Therefore, the concurrent precipitation effect took place upon slow heating in the cold-rolled samples without preannealing, which led to the formation of a strong P orientation in these samples. However, this study demonstrated that precipitates of Al6(Mn,Fe) produced in the CC AA3004 Al alloy prior to cold rolling and recrystallization could also give rise to the formation of a strong P orientation, though no concurrent precipitation took place during recrystallization in the alloy. The temperature, 420 °C, used in prior annealing was such that the minimum resistivity was achieved in the alloy after annealing; i.e., the maximum precipitates were produced in the alloy. This also minimized the occurrence of concurrent precipitation during final recrystallization after cold rolling in the alloy. The TEM observations indicated that the Al6(Mn,Fe) precipitates formed before or during recrystallization all strongly interacted with dislocations, thereby retarding the recrystallization process. The results from this study verified that P orientation was indeed developed due to precipitates pinning dislocations and grain boundaries during recrystallization in Al alloys supersaturated with Mn.

The mechanism for the formation of P orientation has not been fully understood so far, though concurrent precipitation is known to be responsible for the P texture formation.[5,6] It has been recognized that recrystallization textures are formed mainly due to rearrangement of dislocations to form low-angle grain boundaries at an elevated temperature. Qualitatively, the driving force for dislocation movement increases with increased temperature, while the resistance to dislocation movement is reduced with the increase in temperature. The onset temperature for recrystallization should be around the temperature at which the driving force equals the resistance. When there are fine precipitates that could pin dislocations, the resistance is enhanced, thereby increasing the onset temperature in the alloy. Recently, it has been found that P orientation is formed at a temperature of 490 °C, 150 °C higher than that for cube orientation in the as-received hot band of CC AA3004 Al alloy.[30] This supports the suggestion that P orientation is formed at a temperature higher than the cube texture, which is formed normally in the alloy without interference of precipitates in Al alloys. Because of the increase in onset temperature for recrystallization, the release of pinned dislocations in a large quantity at relatively higher temperature might not give rise to the formation of the polygonized subgrains, unlike in the case without the precipitates, but it could possibly cause local lattice rotation, which forms the nuclei for P orientation in the alloy studied in this work. For example, one such possible rotation is from a brass orientation to P orientation by a rotation of 19.47 deg around their common zone axis \( \left\langle {011} \right\rangle \).

Figure 10 presents the maximum intensities of the ODF f(g) for a particular angle φ1 along the central line of \( \left\langle {110} \right\rangle \)//ND fiber in the cold-rolled samples without or with prior heat treatment after annealing. The strength of P orientation (which is located at φ 1 = 60 deg) was intensified with increased rolling reduction in the two different samples. It can be seen in Figure 10(a) that the intensity of P orientation was markedly increased when rolling reduction exceeded 50 pct in the as-received samples. As a comparison, the significant increase in the intensity of P orientation took place in the preannealed sample exceeding 75 pct rolling reduction (Figure 10(b)). It is also noted that, as the rolling reduction was greater than 75 pct, the intensity of P orientation in the specimen preannealed at 420 °C was much higher than that in the specimen without prior heat treatment.

Fig. 10
figure 10

Orientation intensity function f(g) along the \( \left\langle {110} \right\rangle \)//ND fiber in the recrystallized alloys after cold rolling at 0, 25, 50, 75, 80, 85, and 90 pct: (a) without the prior heat treatment and (b) with the prior heat treatment

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The observed difference in intensity of P orientation between the two types of samples can be accounted for by the difference in effectiveness of Al6(Mn,Fe) precipitates in pinning dislocations between the two samples. The precipitates produced by preannealing and concurrent precipitation were different in size and distribution in their respective alloys. The quantity of the precipitates produced by preannealing was much more than that by concurrent precipitation, since the resistivity of the preannealed sample was the minimum at 420 °C in the alloy. Both types of precipitates had different distributions, since rolling deformation took place before precipitation in the sample without preannealing and after precipitation in the sample with the preannealing. Rolling deformation changed the distribution if the existing precipitates, e.g., the spacing between precipitates was smaller in the preannealed sample than in the sample without the preannealing in the direction normal to the RD, in addition to the difference in quantity and size of the two types of precipitates between the alloys. It is, therefore, understandable that the effectiveness of precipitates in pinning dislocations was different, leading to the difference in intensity of P orientation in the two alloys. This was consistent with the recent measurement that the critical strain needed for the formation of P orientation was different, –0.8 (65 pct) in the alloy without the preannealing and –1.1 (77 pct) in the preannealed alloy.[30] As mentioned earlier in this article, the grain aspect ratio was increased with increased rolling reduction in the preannealed sample, since the precipitate spacing decreased upon increasing the rolling reduction in the alloy, while in the sample with the preannealed sample, the precipitate spacing was not affected by rolling reduction as significantly as in the preannealed sample. Therefore, the aspect ratio of grains was not affected profoundly in the sample without preannealing, as shown in Figure 7.

As shown in Figures 4(b) and 5, a large number of precipitates (20 to 100 nm in size) and particles (around 0.2 μm in size) existed in the sample preannealed at 420 °C. Strain concentration could take place in the vicinity of the coarse particles in the sample during cold rolling; i.e., a higher energy was stored around these particles. Higher stored energy promoted grain nucleation at these strain concentrated sites, as the driving pressure was enhanced for nucleation during annealing at these sites.[31–36] In the meantime, the fine precipitates had a sound pinning effect on dislocations and grain boundaries (Figures 5), thereby hindering the rearrangement of dislocations and grain boundaries in the alloy at the elevated temperature. Retention of a high dislocation density by precipitate pinning dislocations at an elevated temperature would lead to formation of P-oriented grains at the strain concentrated sites in these alloys. Recent EBSD work by the authors of this paper revealed that most the newly recrystallized grains in the early stage of the recrystallization process were associated with coarse particles and possessed an orientation close to P orientation.[37] It indicated that P orientation was formed at the strain concentration sites, such as particles, in these alloys.

With increasing rolling reduction, more strain concentration sites would be generated at the large particles,[38,39] which gave rise to more nucleation sites for P orientated grains. This was very likely to be responsible for the higher intensity of P orientation in severely deformed samples. Since P-oriented grains grew faster than grains with other orientations (as discussed previously in Reference 6), when the rolling reduction exceeded 70 pct, more nucleation sites led to more P-oriented grains. As a result, the intensity of P orientation was markedly enhanced with a rolling reduction over 70 pct. With further increase in cold rolling reduction over 70 pct, the number of the strain concentration sites for nuclei was likely to be saturated, resulting in only a small increase in the intensity of P orientation in the alloy at the rolling reduction of 85 and 90 pct (Figure 10(b)), since the amount of precipitates produced by concurrent precipitation was limited in the alloy. However, the amount of precipitates was maximized in the preannealed sample; therefore, the intensity of P orientation in the preannealed sample exceeded that in the sample without preannealing, when rolling reduction was over 75 pct. However, it may still be desirable to further study the driving force for and resistance to dislocation movement in the two types of CC 3004 Al alloys, in order to understand the mechanism for the formation of P orientation.

5 Conclusions

  1. 1.

    Annealing at 420 °C led to generation of a large amount of second-phase precipitates in a size ranging from 20 nm to 0.2 μm in the as-received hot band of a CC AA 3004 Al alloy, which had a significant effect on the development of recrystallization microstructures and textures in the cold-rolled alloy.

  2. 2.

    The recrystallization microstructure exhibited a severely elongated grain structure in the alloy with or without prior heat treatment at 420 °C, and the recrystallized grains were refined with an increase in rolling reduction. However, the aspect ratio of grains in the preannealed sample increased with rolling reduction, while that in the sample without the prior heat treatment was not markedly changed.

  3. 3.

    The P orientation was the dominant recrystallization texture in the alloy with or without prior heat treatment, and its volume fraction increased with the increase in rolling reduction in the alloys after recrystallization. When the rolling reduction exceeded 75 pct, the volume fraction of P orientation in the preannealed sample was higher than that in the sample without prior heat treatment.

  4. 4.

    It is found that the concurrent precipitation effect was not the sole reason responsible for the formation of P orientation. The fine second-phase particles generated in preannealing before cold rolling could also give rise to the formation of a strong P orientation. This verified that the formation of P orientation was caused by precipitates pinning dislocations and grain boundaries during recrystallization in CC AA3004 Al alloys.