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Softening behavior of 8011 alloy produced by accumulative roll bonding process
Scripta Materialia 45 (2001) 597±604
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Softening behavior of 8011 alloy produced by accumulative roll bonding process Z.P. Xingab, S.B. Kangb*, and H.W. Kimb a
Beijing Institute of Aeronautical Materials, Beijing 100095, China Korea Institute of Machinery and Materials, 66 Sangnam-Dong, Changwon, Kyungnam 641-010, South Korea
b
Received 29 January 2001; accepted 9 May 2001
Abstract Ultra-®ne grained Al±Fe±Si 8011 alloy was successfully produced by the accumulative roll bonding (ARB) process. However, softening behavior also happened after 2±3 cycles of ARB. The reason of softening behavior was analyzed preliminarily in this paper. Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Aluminum; Bonding; Ultra-®ne grain; Mechanical properties; Microstructure
Introduction Extensive investigations have demonstrated that a very high plastic strain can produce an ultra-®ne grain size in metallic materials. Several techniques are now available for producing the high strains, including cyclic extrusion compression (CEC) [1], torsion straining under high pressure (TS) [2], equal channel angular pressing (ECAP) [3,4], but one disadvantage of these processes is that they are not applicable to large bulk materials. One of novel intense straining processes for bulk materials using rolling deformation, named accumulative roll bonding (ARB), was developed recently [5±9]. In this process, the achieved strain is theoretically unlimited. The ARB process has been successfully applied to aluminum (1100) [5,6], Al±Mg alloy (5083) [8] and Ti-added interstitial free steel [8,9]. All these several-cycle ARB processed materials have structures with submicron grains and show very high strength at ambient temperature [5±9]. Commercial 8011 alloy is mainly an Al±Fe±Si alloy, which has a wide variety of end applications owing to the fact that it is possible to control the microstructural evolution
*
Corresponding author. Tel.: +82-55-280-3301; fax: +82-55-280-3399. E-mail address: [email protected] (S.B. Kang).
1359-6462/01/$ - see front matter Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 0 6 9 - 7
598
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
of the alloy by means of speci®c thermal and mechanical treatments [10]. However, in this study we ®nd that the strength at ambient temperature of the ultra-®ne grains formed in 8011 alloy by the ARB process did not increase with increasing number of cycles and was lower than that of 1100 alloy in the same processing condition while the solute and particle content in the latter are lower than those in the former [5,6]. For analyzing the particular softening behavior of 8011 alloy, the changes in mechanical properties and microstructural evaluation during ARB process are presented in this paper. Some complementary experiments on 1100 alloy processed with the same condition as 8011 alloy are also performed for purpose of comparison. Experimental The material used in this study was 8011 alloy whose chemical composition is given in Table 1. For comparison, the chemical composition of 1100 alloy is also given. The initial dimension of the materials is 4.0 mm in thickness, and the mean grain size of the 8011 alloy is about 26 lm. For ARB process, the initial materials are cold rolled to the sheets with 1 mm thick (0 cycle, equivalent strain of 1.6). The repeated cycle of ARB process can be simpli®ed as follows: Cutting ! Surface treatment ! Stacking ! Heating ! Roll bonding: Firstly, the cold-rolled materials are reduced to sheets with dimensions of 1.0 mm in thickness, 30 mm in width and 300 mm in length. Then the interface between two sheets is degreased by acetone and wire-brushed. After that the two sheets are layered in order to put brushed surfaces in contact and ®xed to each other closely by wires. At last, the stacked sheets are kept at 200°C for 5 min and then rolled. ARB process implies a reduction in thickness per cycle of 50% (equivalent strain, e 0:8/cycle). The rolling diameter is 175 mm and the roll peripheral speed is about 1 m/min. The above procedures are repeated up to 9 cycles (equivalent total strain of 8.8). The values reported for Hv represent the average of seven separate measurements taken at randomly selected points using a load of 200 g for 10 s. The tensile specimens were machined from the rolled sheets according to the ASTM E8M standard, oriented along the rolling direction. The gauge length is 25 mm. Tensile tests at room temperature are conducted on a standard universal testing machine at a strain rate of 8:3 10 4 s 1 . The optical examination of the samples is conducted under condition of polarized light. Etching is carried out electrolytically using Baker's solution of 5 ml HBF4 and Table 1 Chemical composition of the 8011 and 1100 alloy (mass%) Materials
Si
Fe
Cu
Mn
Mg
Ti
Cr
Zn
Al
A8011 A1100
0.625 0.17
0.725 0.57
0.007 0.11
0.003 0.004
0.002 0.001
0.013 0.017
0.001 0.002
0.002 0.002
Balance Balance
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
599
200 ml distilled water and a stainless steel cathode. To observe the particle size and distribution, some samples were electro-polished in a solution of methanol and perchloric acid with a ratio of 9:1. These samples are examined by a JSM-5800 scanning electron microscope (SEM) with an X-ray energy dispersive spectrometer (EDS). All optical and SEM microstructure were observed along the transverse direction of the rolled samples. Specimens are also examined using a JEM-2000 FX II transmission electron microscope (TEM) operating at 200 kV. Thin foils parallel to the rolling plane are prepared by a twin-jet electro-polisher using a solution of 80 ml HClO4 and 320 ml CH3 OH. Selected area electron diraction (SAD) patterns are taken from regions having diameters of 3 lm. Measurements of the sub-grain and grain sizes are made quantitatively from the TEM photomicrographs using an image analysis program. Results Mechanical properties Fig. 1(a) presents the Vickers microhardness of specimens with the plastic strain. It shows that the Vickers microhardness of 8011 alloy produced by ARB increased signi®cantly with strain up to e 1:6 (0 cycle). Then it kept nearly no change until a strain of 4.0 (3 cycles). After that, it decreased gradually with strain. However, in the same processing condition, the hardness of 1100 alloy increased with strain until strain e 4:0, and then leveled out. The ambient tensile properties of the ARB processed 8011 alloy are presented in Fig. 1(b). It can be seen that yield strength (YS) increased with the strain only up to e 2:4 (1 cycle), and tensile strength (UTS) increased only to the strain of 3.2 (2 cycles). After that, the strength decreased with the strain. In contrast, the elongation decreased with strain to e 2:4, and then increased. This result is very dierent from those of 1100 alloy (Fig. 1b) produced with the same route. In 1100 alloy, UTS and YS increased with
Fig. 1. (a) Vickers microhardness and (b) ambient tensile properties of the 8011 and 1100 alloy produced by ARB at 200°C.
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Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
the strain up to the strain of 4.8, and then kept a constant value. Its elongation decreased with the strain. The results of 1100 alloy in our experiments were similar to those in other works [5,6]. For a given strain, especially in the higher strain range, the hardness and strength of 8011 alloy were obviously lower than those of 1100 alloy. However, the elongation of the 8011 alloy was 2±3 times higher than that of 1100 alloy in the high strain range (Fig. 1b). Microstructure Optical micrographs of as-ARBed samples are presented in Fig. 2. The interface introduced at each cycle was observed occasionally. The bonding of the interfaces was quite good except after one cycle. Fig. 3 (a)±(c) shows SEM micrographs of as-ARBed 8011 samples. In 8011 alloy, there were two kinds of particles. During ARB process, the gray particles and big white particles did not change greatly. However, the number of the small white particles increased with increasing number of cycles. It is believed that these white particles result from precipitation at the interfaces. Indeed, it is clearly seen from Fig. 3(b) that the small white particles are formed in the interfaces. The composition of the two particles was analyzed by EDS using SEM. The white particles contained only Si and Al, and the weight ratio of Si and Al elements in some bigger particles (1.5±2.0 lm in diameter) was
Fig. 2. Optical microstructure of 8011 alloy after ARB at 200°C: (a) 0 cycle (strain of 1.6), (b) 1 cycle (strain of 2.4), (c) 3 cycles (strain of 4.0), (d) 6 cycles (strain of 6.4).
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
601
Fig. 3. SEM micrographs of 8011 and 1100 alloys after ARB at 200°C: (a) 0 cycle in 8011, (b) 3 cycles in 8011, (c) 9 cycles in 8011, (d) 3 cycles in 1100.
found in the range of (25±50):(75±50). The content of Al decreased with increasing particle size. The Al content measured in the particles is probably related to the base composition, as the beam size of the probe is about 2±3 lm, which is bigger than most of the white particles. So the white particles were supposed to be Si phases. The gray particles contained Al, Fe and Si, and the weight content of some bigger particles (P3 lm) was determined to be about 7±10% Si and 28±32% Fe. Because the size of the gray phases is quite big, the measurement of their composition is more accurate. The composition of the gray particles is in the range of the a-AlFeSi phases [11], so they were most probably a-AlFeSi phases. In addition, the composition in the interface was nearly the same as that in the base material. The small white particles in the interface were probably the Si phases that precipitated after annealing process in every cycle or during ARB process. However, they were too small to be veri®ed using the SEM±EDS technique. For comparison, SEM micrograph of 1100 alloy at the same strain condition as Fig. 3(b) is presented in Fig. 3(d). There was only one kind of particles in the 1100 alloy and the particles changed slightly with the processing cycles. Their shape and distribution was similar to the gray particles observed in 8011 alloy, but their number was less important than in the 8011 alloy. The particles in the 1100 alloy contained only Fe and
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Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
Fig. 4. TEM micrographs and corresponding SAD patterns of the 8011 alloy after ARB at 200°C: (a) 0 cycle, (b) 1 cycle, (c) 3 cycles, (d) 9 cycles.
Al, and the weight content of Fe in some bigger particles (2±3 lm) was about 15±22%. They were probably FeAl3 or FeAl6 particles, and need further study. It can be seen from Fig. 3 that the particle type and number observed in the two alloys are very dierent. In particular there was no Si particles in 1100 alloy while the Si particles in 8011 alloy increased with increasing number of cycles. Fig. 4 shows TEM microstructures and their corresponding SAD patterns of the 8011 alloy after various cycles of ARB. At 0 cycle (strain of 1.6), the ®ne grain (or subgrain) structures whose average diameters were less than 1 lm can be seen in many regions (Fig. 4(a)). However, SAD patterns show that large misorientation grains existed only in some regions. Most regions were ®lled with the small misorientation subgrains. The fraction of the regions with large misorientation grains increased with the increasing cycles in some extent. Furthermore, image analysis shows that the average grain (or subgrain) sizes of this material were about 800±900 nm, and they did not change dramatically during the ARB cycles. This means that ultra-®ne grains have been formed in most regions of this alloy after several cycles, but the grain (or subgrain) size was bigger than those in other alloys made by the similar process [1±5].
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
603
Discussion Strengthening in 8011 alloy may be from solid±solution hardening, second-phase precipitation hardening, grain re®nement hardening, and/or strain hardening. Fig. 5 presents the eect of the annealing at 200°C/5 min on the hardness of ARB processed 8011 alloy. It can be seen that the annealing resulted in the decrease of hardness at each cycle in some extent. The ARB process increased the hardness in the ®rst 3 cycles (strain of 4), then had less and less eect on the hardness of the material. This means strain hardening has some eect only before 3 cycles for this material, then has nearly no eect to the hardness of the material. It was reported that dynamic recovery and static recovery takes place easily in the Al±Fe alloy [12]. From our result, the grain (or subgrain) sizes in 8011 alloy were about 800±900 nm, which are quite bigger than those in other alloys such as 1100 and 5083 alloys [5±9]. This means the contribution of the grain re®nement to the hardness of the material is not so large as for other alloys. So the grain re®nement contributes to the increase of the strength of 8011 alloy only during the ®rst cycle. The Si precipitation was con®rmed to start above 30°C and continue up to 300°C [10]. Hot rolling at 200°C in this experiment certainly promotes the precipitation of Si particles. This results in the increasing number of Si particles with increasing number of cycles. The increasing precipitation of Si particles may enhance the strength and hardness for some extent, but meanwhile it results in the decrease of concentration of Si in solid solution. The rates of strength increase by second-phase precipitation may be lower than those by solid solution hardening of Si [13,14]. This may be another reason for the decrease of hardness and strength of 8011 alloy with increasing the cycles. Furthermore, the main composition dierence between 8011 and 1100 alloys was that there are 0.625% Si and more higher Fe contents in 8011 alloy. This results in the presence of two kind of particles (Si and a-AlFeSi) in 8011 alloy and only one kind of
Fig. 5. Eect of the annealing at 200°C/5 min on the hardness of ARB processed 8011 alloy.
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Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
particles (Alx Fe) in the 1100 alloy. The particles in the 8011 alloy cause dynamic and static recovery more easily than other alloys [12]. Therefore, the existence of Si particles in 8011 alloy is probably the most important contribution to the softening of the 8011 alloy. Not only the decrease of solid solution hardening is caused by the precipitation of Si particles, but also the dynamic and static recovery is probably caused by it. In one word, dynamic recovery and static recovery caused by precipitation during ARB process and/or after annealing result in the decrease of grain re®nement strengthening and/or strain hardening. The precipitation of Si particles with the cycles also results in the decrease of solid solution hardening. These were probably the main reason for the softening behavior in 8011 alloy after 2±3 cycles.
Conclusion Ultra-®ne grained 8011 alloy whose mean grain (or subgrain) size was 800±900 nm was successfully produced by the ARB process. However, softening behavior also happened after 2±3 cycles of ARB. Dynamic recovery and static recovery caused by particles and the decrease of solid solution hardening during ARB process may play the main role on the softening behavior of the 8011 alloy.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
Richert, J., & Richert, M. (1986). Aluminium 62, 604. Valiev, R. Z., Mulyukov, R. R., Ovchinnikov, V. V., & Shabashov, V. A. (1991). Scripta Mater 25, 2717. Valiev, R. Z., Krasilnikov, N. A., & Tsenev, N. K. (1991). Mater Sci Eng A 137, 35. Iwahashi, Y., Horita, Z., Nemoto, M., & Langdon, T. G. (1997). Acta Mater 45, 4733. Saito, Y., Tsuji, N., Utsunomiya, H., Sakai, T., & Hong, R. G. (1998). Scripta Mater 39, 1221. Saito, Y., Utsunomiya, H., Tsuji, N., Sakai, T., & Hong, R. G. (1999). J Jpn Inst Metals 63, 790. Tsuji, N., Murakami, T., & Saito, Y. (1999). J Jpn Inst Metals 63, 243. Saito, Y., Utsunomiya, H., Tsuji, N., & Sakai, T. (1999). Acta Mater 47, 579. Tsuji, N., Saito, Y., Utsunomiya, H., & Tanigawa, S. (1999). Scripta Mater 40, 795. Luiggi, A. N. (1998). J Metal Mater Trans A 29, 2669. Mondolf, L. F. (1976). Aluminum Alloys: Structure and Properties. London: Butterworths, 534. Tsuchida, S., & Baba, Y. In J. G. Morris & H. D. Merchant, et al. (Eds.). Al Alloys for Packaging. Chicago: The Minerals, Metals and Materials Society. [13] Alicante, R. M. (1992). Aluminium 68, 1081. [14] Davis J. R. (1993). ASM Specialty Handbook: Aluminium and Aluminium Alloys. Materials Park, OH: ASM International, 33.
www.elsevier.com/locate/scriptamat
Softening behavior of 8011 alloy produced by accumulative roll bonding process Z.P. Xingab, S.B. Kangb*, and H.W. Kimb a
Beijing Institute of Aeronautical Materials, Beijing 100095, China Korea Institute of Machinery and Materials, 66 Sangnam-Dong, Changwon, Kyungnam 641-010, South Korea
b
Received 29 January 2001; accepted 9 May 2001
Abstract Ultra-®ne grained Al±Fe±Si 8011 alloy was successfully produced by the accumulative roll bonding (ARB) process. However, softening behavior also happened after 2±3 cycles of ARB. The reason of softening behavior was analyzed preliminarily in this paper. Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Aluminum; Bonding; Ultra-®ne grain; Mechanical properties; Microstructure
Introduction Extensive investigations have demonstrated that a very high plastic strain can produce an ultra-®ne grain size in metallic materials. Several techniques are now available for producing the high strains, including cyclic extrusion compression (CEC) [1], torsion straining under high pressure (TS) [2], equal channel angular pressing (ECAP) [3,4], but one disadvantage of these processes is that they are not applicable to large bulk materials. One of novel intense straining processes for bulk materials using rolling deformation, named accumulative roll bonding (ARB), was developed recently [5±9]. In this process, the achieved strain is theoretically unlimited. The ARB process has been successfully applied to aluminum (1100) [5,6], Al±Mg alloy (5083) [8] and Ti-added interstitial free steel [8,9]. All these several-cycle ARB processed materials have structures with submicron grains and show very high strength at ambient temperature [5±9]. Commercial 8011 alloy is mainly an Al±Fe±Si alloy, which has a wide variety of end applications owing to the fact that it is possible to control the microstructural evolution
*
Corresponding author. Tel.: +82-55-280-3301; fax: +82-55-280-3399. E-mail address: [email protected] (S.B. Kang).
1359-6462/01/$ - see front matter Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 0 6 9 - 7
598
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
of the alloy by means of speci®c thermal and mechanical treatments [10]. However, in this study we ®nd that the strength at ambient temperature of the ultra-®ne grains formed in 8011 alloy by the ARB process did not increase with increasing number of cycles and was lower than that of 1100 alloy in the same processing condition while the solute and particle content in the latter are lower than those in the former [5,6]. For analyzing the particular softening behavior of 8011 alloy, the changes in mechanical properties and microstructural evaluation during ARB process are presented in this paper. Some complementary experiments on 1100 alloy processed with the same condition as 8011 alloy are also performed for purpose of comparison. Experimental The material used in this study was 8011 alloy whose chemical composition is given in Table 1. For comparison, the chemical composition of 1100 alloy is also given. The initial dimension of the materials is 4.0 mm in thickness, and the mean grain size of the 8011 alloy is about 26 lm. For ARB process, the initial materials are cold rolled to the sheets with 1 mm thick (0 cycle, equivalent strain of 1.6). The repeated cycle of ARB process can be simpli®ed as follows: Cutting ! Surface treatment ! Stacking ! Heating ! Roll bonding: Firstly, the cold-rolled materials are reduced to sheets with dimensions of 1.0 mm in thickness, 30 mm in width and 300 mm in length. Then the interface between two sheets is degreased by acetone and wire-brushed. After that the two sheets are layered in order to put brushed surfaces in contact and ®xed to each other closely by wires. At last, the stacked sheets are kept at 200°C for 5 min and then rolled. ARB process implies a reduction in thickness per cycle of 50% (equivalent strain, e 0:8/cycle). The rolling diameter is 175 mm and the roll peripheral speed is about 1 m/min. The above procedures are repeated up to 9 cycles (equivalent total strain of 8.8). The values reported for Hv represent the average of seven separate measurements taken at randomly selected points using a load of 200 g for 10 s. The tensile specimens were machined from the rolled sheets according to the ASTM E8M standard, oriented along the rolling direction. The gauge length is 25 mm. Tensile tests at room temperature are conducted on a standard universal testing machine at a strain rate of 8:3 10 4 s 1 . The optical examination of the samples is conducted under condition of polarized light. Etching is carried out electrolytically using Baker's solution of 5 ml HBF4 and Table 1 Chemical composition of the 8011 and 1100 alloy (mass%) Materials
Si
Fe
Cu
Mn
Mg
Ti
Cr
Zn
Al
A8011 A1100
0.625 0.17
0.725 0.57
0.007 0.11
0.003 0.004
0.002 0.001
0.013 0.017
0.001 0.002
0.002 0.002
Balance Balance
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
599
200 ml distilled water and a stainless steel cathode. To observe the particle size and distribution, some samples were electro-polished in a solution of methanol and perchloric acid with a ratio of 9:1. These samples are examined by a JSM-5800 scanning electron microscope (SEM) with an X-ray energy dispersive spectrometer (EDS). All optical and SEM microstructure were observed along the transverse direction of the rolled samples. Specimens are also examined using a JEM-2000 FX II transmission electron microscope (TEM) operating at 200 kV. Thin foils parallel to the rolling plane are prepared by a twin-jet electro-polisher using a solution of 80 ml HClO4 and 320 ml CH3 OH. Selected area electron diraction (SAD) patterns are taken from regions having diameters of 3 lm. Measurements of the sub-grain and grain sizes are made quantitatively from the TEM photomicrographs using an image analysis program. Results Mechanical properties Fig. 1(a) presents the Vickers microhardness of specimens with the plastic strain. It shows that the Vickers microhardness of 8011 alloy produced by ARB increased signi®cantly with strain up to e 1:6 (0 cycle). Then it kept nearly no change until a strain of 4.0 (3 cycles). After that, it decreased gradually with strain. However, in the same processing condition, the hardness of 1100 alloy increased with strain until strain e 4:0, and then leveled out. The ambient tensile properties of the ARB processed 8011 alloy are presented in Fig. 1(b). It can be seen that yield strength (YS) increased with the strain only up to e 2:4 (1 cycle), and tensile strength (UTS) increased only to the strain of 3.2 (2 cycles). After that, the strength decreased with the strain. In contrast, the elongation decreased with strain to e 2:4, and then increased. This result is very dierent from those of 1100 alloy (Fig. 1b) produced with the same route. In 1100 alloy, UTS and YS increased with
Fig. 1. (a) Vickers microhardness and (b) ambient tensile properties of the 8011 and 1100 alloy produced by ARB at 200°C.
600
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
the strain up to the strain of 4.8, and then kept a constant value. Its elongation decreased with the strain. The results of 1100 alloy in our experiments were similar to those in other works [5,6]. For a given strain, especially in the higher strain range, the hardness and strength of 8011 alloy were obviously lower than those of 1100 alloy. However, the elongation of the 8011 alloy was 2±3 times higher than that of 1100 alloy in the high strain range (Fig. 1b). Microstructure Optical micrographs of as-ARBed samples are presented in Fig. 2. The interface introduced at each cycle was observed occasionally. The bonding of the interfaces was quite good except after one cycle. Fig. 3 (a)±(c) shows SEM micrographs of as-ARBed 8011 samples. In 8011 alloy, there were two kinds of particles. During ARB process, the gray particles and big white particles did not change greatly. However, the number of the small white particles increased with increasing number of cycles. It is believed that these white particles result from precipitation at the interfaces. Indeed, it is clearly seen from Fig. 3(b) that the small white particles are formed in the interfaces. The composition of the two particles was analyzed by EDS using SEM. The white particles contained only Si and Al, and the weight ratio of Si and Al elements in some bigger particles (1.5±2.0 lm in diameter) was
Fig. 2. Optical microstructure of 8011 alloy after ARB at 200°C: (a) 0 cycle (strain of 1.6), (b) 1 cycle (strain of 2.4), (c) 3 cycles (strain of 4.0), (d) 6 cycles (strain of 6.4).
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
601
Fig. 3. SEM micrographs of 8011 and 1100 alloys after ARB at 200°C: (a) 0 cycle in 8011, (b) 3 cycles in 8011, (c) 9 cycles in 8011, (d) 3 cycles in 1100.
found in the range of (25±50):(75±50). The content of Al decreased with increasing particle size. The Al content measured in the particles is probably related to the base composition, as the beam size of the probe is about 2±3 lm, which is bigger than most of the white particles. So the white particles were supposed to be Si phases. The gray particles contained Al, Fe and Si, and the weight content of some bigger particles (P3 lm) was determined to be about 7±10% Si and 28±32% Fe. Because the size of the gray phases is quite big, the measurement of their composition is more accurate. The composition of the gray particles is in the range of the a-AlFeSi phases [11], so they were most probably a-AlFeSi phases. In addition, the composition in the interface was nearly the same as that in the base material. The small white particles in the interface were probably the Si phases that precipitated after annealing process in every cycle or during ARB process. However, they were too small to be veri®ed using the SEM±EDS technique. For comparison, SEM micrograph of 1100 alloy at the same strain condition as Fig. 3(b) is presented in Fig. 3(d). There was only one kind of particles in the 1100 alloy and the particles changed slightly with the processing cycles. Their shape and distribution was similar to the gray particles observed in 8011 alloy, but their number was less important than in the 8011 alloy. The particles in the 1100 alloy contained only Fe and
602
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
Fig. 4. TEM micrographs and corresponding SAD patterns of the 8011 alloy after ARB at 200°C: (a) 0 cycle, (b) 1 cycle, (c) 3 cycles, (d) 9 cycles.
Al, and the weight content of Fe in some bigger particles (2±3 lm) was about 15±22%. They were probably FeAl3 or FeAl6 particles, and need further study. It can be seen from Fig. 3 that the particle type and number observed in the two alloys are very dierent. In particular there was no Si particles in 1100 alloy while the Si particles in 8011 alloy increased with increasing number of cycles. Fig. 4 shows TEM microstructures and their corresponding SAD patterns of the 8011 alloy after various cycles of ARB. At 0 cycle (strain of 1.6), the ®ne grain (or subgrain) structures whose average diameters were less than 1 lm can be seen in many regions (Fig. 4(a)). However, SAD patterns show that large misorientation grains existed only in some regions. Most regions were ®lled with the small misorientation subgrains. The fraction of the regions with large misorientation grains increased with the increasing cycles in some extent. Furthermore, image analysis shows that the average grain (or subgrain) sizes of this material were about 800±900 nm, and they did not change dramatically during the ARB cycles. This means that ultra-®ne grains have been formed in most regions of this alloy after several cycles, but the grain (or subgrain) size was bigger than those in other alloys made by the similar process [1±5].
Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
603
Discussion Strengthening in 8011 alloy may be from solid±solution hardening, second-phase precipitation hardening, grain re®nement hardening, and/or strain hardening. Fig. 5 presents the eect of the annealing at 200°C/5 min on the hardness of ARB processed 8011 alloy. It can be seen that the annealing resulted in the decrease of hardness at each cycle in some extent. The ARB process increased the hardness in the ®rst 3 cycles (strain of 4), then had less and less eect on the hardness of the material. This means strain hardening has some eect only before 3 cycles for this material, then has nearly no eect to the hardness of the material. It was reported that dynamic recovery and static recovery takes place easily in the Al±Fe alloy [12]. From our result, the grain (or subgrain) sizes in 8011 alloy were about 800±900 nm, which are quite bigger than those in other alloys such as 1100 and 5083 alloys [5±9]. This means the contribution of the grain re®nement to the hardness of the material is not so large as for other alloys. So the grain re®nement contributes to the increase of the strength of 8011 alloy only during the ®rst cycle. The Si precipitation was con®rmed to start above 30°C and continue up to 300°C [10]. Hot rolling at 200°C in this experiment certainly promotes the precipitation of Si particles. This results in the increasing number of Si particles with increasing number of cycles. The increasing precipitation of Si particles may enhance the strength and hardness for some extent, but meanwhile it results in the decrease of concentration of Si in solid solution. The rates of strength increase by second-phase precipitation may be lower than those by solid solution hardening of Si [13,14]. This may be another reason for the decrease of hardness and strength of 8011 alloy with increasing the cycles. Furthermore, the main composition dierence between 8011 and 1100 alloys was that there are 0.625% Si and more higher Fe contents in 8011 alloy. This results in the presence of two kind of particles (Si and a-AlFeSi) in 8011 alloy and only one kind of
Fig. 5. Eect of the annealing at 200°C/5 min on the hardness of ARB processed 8011 alloy.
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Z.P. Xing et al./Scripta Materialia 45 (2001) 597±604
particles (Alx Fe) in the 1100 alloy. The particles in the 8011 alloy cause dynamic and static recovery more easily than other alloys [12]. Therefore, the existence of Si particles in 8011 alloy is probably the most important contribution to the softening of the 8011 alloy. Not only the decrease of solid solution hardening is caused by the precipitation of Si particles, but also the dynamic and static recovery is probably caused by it. In one word, dynamic recovery and static recovery caused by precipitation during ARB process and/or after annealing result in the decrease of grain re®nement strengthening and/or strain hardening. The precipitation of Si particles with the cycles also results in the decrease of solid solution hardening. These were probably the main reason for the softening behavior in 8011 alloy after 2±3 cycles.
Conclusion Ultra-®ne grained 8011 alloy whose mean grain (or subgrain) size was 800±900 nm was successfully produced by the ARB process. However, softening behavior also happened after 2±3 cycles of ARB. Dynamic recovery and static recovery caused by particles and the decrease of solid solution hardening during ARB process may play the main role on the softening behavior of the 8011 alloy.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
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