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The effect of plastic deformation on structure and properties of chosen 6000 series aluminium alloys
Materials Science and Engineering A324 (2002) 239– 243 www.elsevier.com/locate/msea
The effect of plastic deformation on structure and properties of chosen 6000 series aluminium alloys J. Dutkiewicz *, L. Litynska Institute of Metallurgy and Materials Science, Polish Academy of Sciences, 25 Reymonta Street, 30 -059 Krakow, Poland
Abstract In order to determine the effect of increased copper addition on the strengthening behaviour during deformation and subsequent ageing, two types of Al alloys were continuously cast and extruded: 6013 containing 1.15% Mg, 1.08% Si, 0.7% Mn, 0.3% Fe and 1.1% Cu, and 6XXX alloy containing 1.09% Mg, 0.9% Si, 0.1% Fe and 1.6%Cu (in wt.%). The 6XXX alloy aged at 165°C showed a hardness maximum of 150 HV, while alloy 6013 revealed lower precipitation hardening at higher strain hardening. A similar maximum hardness of about 170 HV was observed for both alloys after ageing in a deformed state. Transmission electron microscopy performed after deformation of 60% and 90% by cold rolling showed the presence of large, elongated subgrains. Narrow deformation bands consisting of fine subgrains of large misorientation were observed in the alloys deformed 60% by rolling, and their density increased in the alloys deformed up to 90%. Fine precipitates were observed after an ageing time corresponding to the maximum of hardness for alloy 6XXX deformed in the as quenched state. They formed predominantly on dislocations as confirmed by dark field imaging. In the alloys deformed by 60% the precipitates were identified as belonging to the metastable B2 phase. In the alloy deformed by 90%, larger needle-like precipitates were observed. They seemed to possess a quaternary Q-precipitate structure. A frequent rotational moire pattern observed in this state allowed a misorientation as 2–3° to be determined. The presence of fine subgrains formed after ageing indicated the activation of recovery process at the ageing temperature. © 2002 Published by Elsevier Science B.V. Keywords: Al– Mg– Si–Cu alloys; Plastic deformation; Age-hardening; Precipitation; Transmission electron microscopy
1. Introduction The 6000 series alloys have recently found increased application in automotive and construction industry. Therefore, several research works have been undertaken to strengthen the alloys either by small copper additions [1–5,13] or by a predeformation treatment [6 –10,14 –16]. The copper addition increases the peak hardness and yield strength during ageing [1 – 5,13], and it increases even more the peak hardness than that of the solid solution of the as quenched alloys. This indicates that copper concentrates in the precipitates and increases the volume of the precipitates formed [2]. It does not change the electron diffraction pattern from
This paper is dedicated to Professor Pavel Luka´cˇ on the occasion of his 65th birthday. * Corresponding author. Tel.: + 48-12-6374200; fax: + 48-126372192. E-mail address: [email protected] (J. Dutkiewicz).
the b¦ phase characteristic for Al –Mg –Si alloys [2,11], but shows very strong streaks due to the needle-like shape. In earlier works it was suggested that the strengthening due to copper addition was caused by the additional S% and u% precipitation [3,13]. Another type of precipitates reported in Al –Mg –Si–Cu alloys is the quaternary Q phase [12] with a lath morphology formed at later ageing stages. Two types of precipitates were observed on dislocations in the deformed Al –Mg –Si alloy at an early stage of ageing [14]. Some consisted of small precipitates that looked like string beads on dislocation lines. Others were elongated and resembled end-on of the needle- or rod-shaped precipitates having elongated cross section. There were also random precipitates and the b’’ phase in the matrix, which formed additionally to those on dislocations. It is concluded that the plastically deformed solution treated alloy has higher wear resistance than the undeformed one [15]. The results of [16] show that the strength of Al –Mg –Si alloys in the underaged
0921-5093/02/$ - see front matter © 2002 Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 3 1 8 - 1
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
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Table 1 The chemical compositions of investigated alloys (in wt.%) Alloy
Al
Mg
Si
Cu
Mn
Cr
Fe
Zn
Ti
6013 6XXX
bal. bal.
1.15 1.09
1.08 0.90
1.1 1.60
0.73 0.001
0.09 0.098
0.26 0.14
0.012 0.005
0.009 0.06
condition is greatly increased by pre-stretching immediately after water quench. The b¦ precipitates are observed to form directly on dislocations with a larger size than those which form in the dislocation free areas. When the prestretching is less than 5%, the increase in pre-deformation results in the rise of the b¦ precipitate density, whereas the pre-stretching by 10% leads to larger sizes on account of density of the b¦ precipitates at the same time. In high-resolution transmission microscopy (HRTEM) studies of an Al– Mg – Si alloy [17] artificially aged after mechanical deformation, the observation of precipitates cross sections revealed that the transverse orientation of the precipitate deviated 10 degrees from the [1 0 0] or [0 1 0] orientation of the Al matrix. With longer ageing time, the specific precipitates disappeared and the b% phase particles became dominant. The above mentioned results indicate that structure of precipitates formed during ageing of the deformed Al – Mg –Si solid solution can be different from that of aged directly after quenching. Some samples to be aged were deformed in tensile tests [7,8,14– 17]. The effect of higher deformation was studied mainly on texture and recrystallization [5], however, then higher annealing temperatures than those for artificial ageing were used; furthermore, no information was found about the effect of deformation on precipitation and mechanical properties of Al– Si –Mg alloys with copper addition. Therefore in the present paper there have been investigated the effect of deformation of the as quenched Al–Mg – Si – Cu alloys with various copper additions on the structure of matrix and precipitates during artificial ageing as well as the effect of increasing deformation by rolling up to 90% on hardness changes during ageing.
The structure of alloys after deformation and ageing was studied using transmission analytical electron microscope Philips CM20. Thin foils were prepared parallel to the rolling direction for 60% deformed material and perpendicular to the rolling plane, while for 90% of deformation they were parallel to the rolling plane. The Fishione double jet electropoliser was used for thin foil preparation in electrolyte, consisting of 1/3 nitric acid and 2/3 methanol at temperature − 20°C and voltage 15 V.
3. Results and discussion The hardness measurements performed during ageing at room temperatures, 165 and 250°C are shown in Fig. 1. The hardness after quenching was about 70 HV for both alloys. The higher hardness was observed for the 6XXX alloy due to its higher copper content [4,11] for all temperatures of ageing. The maximum value of HV = 143 and HV= 152 was attained for the 6013 and 6XXX alloys, respectively after 8 h of ageing at 165°C. The hardness measurements were carried out during ageing at 165°C for alloys deformed by rolling after quenching. The curves obtained for 6013 and 6XXX, alloy after 30%, 60% and 90% of reduction are presented in Fig. 2. The increase of hardness occurred
2. Experimental procedure Two alloys of the composition given in Table 1 were continuously cast at the rate of 30 mm/min from the elements of 99.9% purity and extruded at 400°C. The heat treatment consisted of annealing at 530°C, followed by water quench and rolling directly after quenching. The alloys were aged at 165°C after having been cold rolled by 60% and 90%. The changes of alloy hardness were measured using Vickers method directly after quenching and after rolling and ageing at 165°C.
Fig. 1. Hardness versus ageing time at room temperatures, 165 and 250°C of quenched 6013 and 6XXX alloys.
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
Fig. 2. Hardness versus ageing time at 165°C of 6013 and 6XXX alloys deformed after quenching 30%, 60% and 90% by rolling.
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respectively) and up to HV= 145 (for both alloys) after 90% degree of deformation. The hardness maximum HV = 170 was reached after 2 h of ageing for both alloys deformed by 90%. The structure studies of quenched and deformed 6013 and 6XXX alloys were carried out using transmission electron microscopy. The presence of large elongated subgrains following the traces {1 1 1} were observed for both alloys deformed by 60%. Fig. 3 shows the tilt experiment made for the 60% deformed alloy in order to measure the subgrain misorientation. The B subgrain visible in Fig. 3(b) after tilting is in the same diffraction condition as subgrain A in Fig. 3(a) (the corresponding diffraction pattern is inserted). The determined subgrain misorientation was estimated to be equal to 2° from the double tilt goniometer readout. The deformation bands consisting of fine crystallites of large misorientation were observed in both alloys. Their density was much higher in alloys deformed by 90%. In Fig. 4 the microstructure of deformation band and the corresponding diffraction patterns obtained from the band area and neighbouring grains are shown for the 6XXX alloy deformed 90%. Both grains are only slightly tilted, but grains within a deformation band show high misorientation. This fact suggests the presence of a regular deformation band or shear band [18]. Changes in the precipitation structure in the 6XXX alloy deformed 60% and 90% and aged for 2 h at 165°C (corresponding to maximum of hardness) are visible in the series of bright and dark field images formed for the
Fig. 3. Transmission electron micrographs of 6XXX alloy quenched and deformed 60% by rolling illustrating the tilt experiment to measure the subgrains misorientation equal to 2°.
during ageing for both alloys in spite of the increase of the initial hardness after 30% of deformation up to HV =105 and HV=115 (for the 6013 and 6XXX,
Fig. 4. Transmission electron micrograph of 6XXX alloy deformed 90% after quenching showing deformation band and corresponding diffraction patterns from the left grain, the deformation band and the right grain.
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Fig. 5. The bright and dark field TEM images and corresponding diffraction patterns from 6XXX alloy quenched, deformed 60% (a,c,e) and 90% (b,d,f) and aged for 2 h at 165°C.
6013 alloy shows lower density of precipitates as compared to 6XXX alloy. It can be concluded that increase of hardness maximum with increasing Cu content in the investigated alloys is connected with the increase of precipitates density. The previous results on the effect of predeformation on ageing characteristic of ternary Al –Mg –Si alloys [7,8] indicate fewer precipitates on dislocations when compared to our results. The reason may be the lower degree of deformation, but it could be also the effect of copper, causing higher strains around precipitates and changing the precipitation behaviour. The structure of precipitates seems to be similar to that observed in the case of the quenched and aged alloys, as far as it can be judged from the diffraction pattern analysis; however, one cannot exclude the formation of new phases like those found after deformation and ageing in ternary Al–Mg –Si using HREM method [17]. Rotational moire patterns frequently observed in deformed and aged alloys are shown in Fig. 6 (for the 6XXX alloy). The measured distance between the fringes allowed to calculate the misorientation of 2°–3° degress for fine subgrains formed after ageing indicating the presence of a recovery process occurring already at 165°C. The frequently observed presence of moire pattern suggests that at least two subgrains of very small thickness of several nanometers are present within the foil thickness.
4. Conclusions
Fig. 6. Transmission electron micrograph of 6XXX alloy quenched, deformed 90% and aged for 2 h with rotational moire fringes formed by overlapping subgrains (calculated misorienation is about 3°); corresponding diffraction pattern.
same diffraction conditions with [0 0 1] zone axis (Fig. 5). The dislocation structure and precipitates are well visible in dark field images, which were obtained using precipitates reflections in the positions corresponding to 0 0 1 matrix reflection. The precipitates were formed mainly on dislocations and their density increased with increasing deformation. In the 60% deformed alloy the structure of the precipitates can be identified as a metastable B2 phase. A similar observation made for the deformed and aged
(1) The hardness measurements of the deformed and aged alloys 6013 and 6XXX show that the increase of the degree of deformation above 60% of reduction by rolling increases the maximum hardness up to 30 HV above the maximum attained during artificial ageing at 165°C. At 30% of reduction, the maximum hardness was not changed. The hardness maximum is reached after 2-h ageing instead of 24 h without deformation. (2) The precipitates in the deformed and aged alloys form predominantly on dislocations after the ageing time corresponding to the hardness maximum. Their density increases with the degree of deformation. The diffraction patterns at this stage of ageing are similar for both alloys and indicate formation of a transition cubic phase (at lower deformations) and the Q phase at higher deformations and longer ageing time. (3) The deformation by rolling causes formation of elongated subgrain structure with misorientation increasing with the degree of rolling. At 60% of deformation, the deformation bands begin to form. Their density and thickness increase with the degree of reduction. They consist of very small grains of high misorientation. (4) The recovery process occurs after ageing at 165°C. It is manifested by formation of moire fringes
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
indicating tilting type of misorientation (typically 2–3°) through the thickness of the foil. No recrystallization effects can be seen at this ageing temperature. Acknowledgements The financial support of the Research Project (Zamawiany 15–15) of the State Committee for Scientific Research is gratefully acknowledged. References [1] T. Moons, P. Ratchev, P. De Smeet, B. Verlinden, P. Van Houtte, Scripta Mater. 35 (1996) 939. [2] M. Tamizifar, G.W. Lorimer, in: L. Arnberg, O. Lohne, E. Nes, N. Ryum (Eds.), Proceedings of the Third International Conference on Aluminium Alloys, NTH, Trondheim, 1992, p. 220. [3] D.G. Eskin, Z. Metallkd. 83 (1992) 10. [4] H. Uchida, H. Yoshida, H. Hira, T. Amano, Mater. Sci. Forum, vols. 217 – 222, p. 1753. [5] S.C. Bergsma, M.E. Kassner, Mater. Sci. Forum, vols. 217 – 222, p. 1801. [6] G. Burger, A.K. Gupta, L. Sutak, D.J. Lloyd, Mater. Sci Forum, vols. 217 –222, p. 471.
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[7] L. Zhen, S.B. Kang, Mater. Sci. Technol. 14 (1998) 317. [8] Z.H. Ismail, Scripta Metal. Mater. 32 (1995) 457. [9] B. Davis, F.J. Humphreys, K.M. Gatenby, C.P. Johnson, in: T. Ginza, C. Ku (Eds.), Sixth International Conference on Aluminium Alloys, Japan Institute of Light Metals, Tokyo, Japan, 1998. [10] S. Ikeno, K. Matsuda, K Nakajima, S. Rengakuji, Y. Uetani, J. Jpn. Ins. Light Metals 48 (211) (1998) 207. [11] I. Okumura, K. Matsuda, Y. Uetani, F. Shinagawa, S. Ikeno, in: M. Koiwa, K. Otsuka, T. Miyazaki (Eds.), Proceedings of the International Conference on Solid-Solid Phase Transformations, The Japan Institute of Metals, 1999, pp. 205. [12] L. Sagalowicz, G. Hug, D. Bechet, P. Sainfort, G. Lapasset, in: G. Sanders, E.A. Starke (Eds.), Fourth International Conference on Aluminium Alloys, Georgia Institute of Technology, Atlanta, 1994, pp. 636. [13] T. Sakurai, T. Eto, Proceedings of the Third International Conference on Aluminium Alloys, NTH, Trondheim, 1992, pp. 208. [14] K. Matsuda, S. Shimizu, H. Gamada, Y. Uetani, F. Shinagawa, S. Ikeno, J. Soc. Mater. Sci., Japan 48 (1999) 10. [15] H. Cimenoglu, Y. Sun, M. Baydogan, Mater. Lett. 38 (1999) 221. [16] L. Zhen, S.B. Kang, Mater. Sci. Technol. 14 (1998) 317. [17] K. Matsuda, H. Gamada, Y. Uetani, S. Rengakuji, F. Shinagawa, S. Ikeno, J. Jpn. Inst. Light Metals 48 (1998) 471. [18] D. Kuhlman-Wilsdorf, Acta Mater. 47 (1999) 1697.
The effect of plastic deformation on structure and properties of chosen 6000 series aluminium alloys J. Dutkiewicz *, L. Litynska Institute of Metallurgy and Materials Science, Polish Academy of Sciences, 25 Reymonta Street, 30 -059 Krakow, Poland
Abstract In order to determine the effect of increased copper addition on the strengthening behaviour during deformation and subsequent ageing, two types of Al alloys were continuously cast and extruded: 6013 containing 1.15% Mg, 1.08% Si, 0.7% Mn, 0.3% Fe and 1.1% Cu, and 6XXX alloy containing 1.09% Mg, 0.9% Si, 0.1% Fe and 1.6%Cu (in wt.%). The 6XXX alloy aged at 165°C showed a hardness maximum of 150 HV, while alloy 6013 revealed lower precipitation hardening at higher strain hardening. A similar maximum hardness of about 170 HV was observed for both alloys after ageing in a deformed state. Transmission electron microscopy performed after deformation of 60% and 90% by cold rolling showed the presence of large, elongated subgrains. Narrow deformation bands consisting of fine subgrains of large misorientation were observed in the alloys deformed 60% by rolling, and their density increased in the alloys deformed up to 90%. Fine precipitates were observed after an ageing time corresponding to the maximum of hardness for alloy 6XXX deformed in the as quenched state. They formed predominantly on dislocations as confirmed by dark field imaging. In the alloys deformed by 60% the precipitates were identified as belonging to the metastable B2 phase. In the alloy deformed by 90%, larger needle-like precipitates were observed. They seemed to possess a quaternary Q-precipitate structure. A frequent rotational moire pattern observed in this state allowed a misorientation as 2–3° to be determined. The presence of fine subgrains formed after ageing indicated the activation of recovery process at the ageing temperature. © 2002 Published by Elsevier Science B.V. Keywords: Al– Mg– Si–Cu alloys; Plastic deformation; Age-hardening; Precipitation; Transmission electron microscopy
1. Introduction The 6000 series alloys have recently found increased application in automotive and construction industry. Therefore, several research works have been undertaken to strengthen the alloys either by small copper additions [1–5,13] or by a predeformation treatment [6 –10,14 –16]. The copper addition increases the peak hardness and yield strength during ageing [1 – 5,13], and it increases even more the peak hardness than that of the solid solution of the as quenched alloys. This indicates that copper concentrates in the precipitates and increases the volume of the precipitates formed [2]. It does not change the electron diffraction pattern from
This paper is dedicated to Professor Pavel Luka´cˇ on the occasion of his 65th birthday. * Corresponding author. Tel.: + 48-12-6374200; fax: + 48-126372192. E-mail address: [email protected] (J. Dutkiewicz).
the b¦ phase characteristic for Al –Mg –Si alloys [2,11], but shows very strong streaks due to the needle-like shape. In earlier works it was suggested that the strengthening due to copper addition was caused by the additional S% and u% precipitation [3,13]. Another type of precipitates reported in Al –Mg –Si–Cu alloys is the quaternary Q phase [12] with a lath morphology formed at later ageing stages. Two types of precipitates were observed on dislocations in the deformed Al –Mg –Si alloy at an early stage of ageing [14]. Some consisted of small precipitates that looked like string beads on dislocation lines. Others were elongated and resembled end-on of the needle- or rod-shaped precipitates having elongated cross section. There were also random precipitates and the b’’ phase in the matrix, which formed additionally to those on dislocations. It is concluded that the plastically deformed solution treated alloy has higher wear resistance than the undeformed one [15]. The results of [16] show that the strength of Al –Mg –Si alloys in the underaged
0921-5093/02/$ - see front matter © 2002 Published by Elsevier Science B.V. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 3 1 8 - 1
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
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Table 1 The chemical compositions of investigated alloys (in wt.%) Alloy
Al
Mg
Si
Cu
Mn
Cr
Fe
Zn
Ti
6013 6XXX
bal. bal.
1.15 1.09
1.08 0.90
1.1 1.60
0.73 0.001
0.09 0.098
0.26 0.14
0.012 0.005
0.009 0.06
condition is greatly increased by pre-stretching immediately after water quench. The b¦ precipitates are observed to form directly on dislocations with a larger size than those which form in the dislocation free areas. When the prestretching is less than 5%, the increase in pre-deformation results in the rise of the b¦ precipitate density, whereas the pre-stretching by 10% leads to larger sizes on account of density of the b¦ precipitates at the same time. In high-resolution transmission microscopy (HRTEM) studies of an Al– Mg – Si alloy [17] artificially aged after mechanical deformation, the observation of precipitates cross sections revealed that the transverse orientation of the precipitate deviated 10 degrees from the [1 0 0] or [0 1 0] orientation of the Al matrix. With longer ageing time, the specific precipitates disappeared and the b% phase particles became dominant. The above mentioned results indicate that structure of precipitates formed during ageing of the deformed Al – Mg –Si solid solution can be different from that of aged directly after quenching. Some samples to be aged were deformed in tensile tests [7,8,14– 17]. The effect of higher deformation was studied mainly on texture and recrystallization [5], however, then higher annealing temperatures than those for artificial ageing were used; furthermore, no information was found about the effect of deformation on precipitation and mechanical properties of Al– Si –Mg alloys with copper addition. Therefore in the present paper there have been investigated the effect of deformation of the as quenched Al–Mg – Si – Cu alloys with various copper additions on the structure of matrix and precipitates during artificial ageing as well as the effect of increasing deformation by rolling up to 90% on hardness changes during ageing.
The structure of alloys after deformation and ageing was studied using transmission analytical electron microscope Philips CM20. Thin foils were prepared parallel to the rolling direction for 60% deformed material and perpendicular to the rolling plane, while for 90% of deformation they were parallel to the rolling plane. The Fishione double jet electropoliser was used for thin foil preparation in electrolyte, consisting of 1/3 nitric acid and 2/3 methanol at temperature − 20°C and voltage 15 V.
3. Results and discussion The hardness measurements performed during ageing at room temperatures, 165 and 250°C are shown in Fig. 1. The hardness after quenching was about 70 HV for both alloys. The higher hardness was observed for the 6XXX alloy due to its higher copper content [4,11] for all temperatures of ageing. The maximum value of HV = 143 and HV= 152 was attained for the 6013 and 6XXX alloys, respectively after 8 h of ageing at 165°C. The hardness measurements were carried out during ageing at 165°C for alloys deformed by rolling after quenching. The curves obtained for 6013 and 6XXX, alloy after 30%, 60% and 90% of reduction are presented in Fig. 2. The increase of hardness occurred
2. Experimental procedure Two alloys of the composition given in Table 1 were continuously cast at the rate of 30 mm/min from the elements of 99.9% purity and extruded at 400°C. The heat treatment consisted of annealing at 530°C, followed by water quench and rolling directly after quenching. The alloys were aged at 165°C after having been cold rolled by 60% and 90%. The changes of alloy hardness were measured using Vickers method directly after quenching and after rolling and ageing at 165°C.
Fig. 1. Hardness versus ageing time at room temperatures, 165 and 250°C of quenched 6013 and 6XXX alloys.
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
Fig. 2. Hardness versus ageing time at 165°C of 6013 and 6XXX alloys deformed after quenching 30%, 60% and 90% by rolling.
241
respectively) and up to HV= 145 (for both alloys) after 90% degree of deformation. The hardness maximum HV = 170 was reached after 2 h of ageing for both alloys deformed by 90%. The structure studies of quenched and deformed 6013 and 6XXX alloys were carried out using transmission electron microscopy. The presence of large elongated subgrains following the traces {1 1 1} were observed for both alloys deformed by 60%. Fig. 3 shows the tilt experiment made for the 60% deformed alloy in order to measure the subgrain misorientation. The B subgrain visible in Fig. 3(b) after tilting is in the same diffraction condition as subgrain A in Fig. 3(a) (the corresponding diffraction pattern is inserted). The determined subgrain misorientation was estimated to be equal to 2° from the double tilt goniometer readout. The deformation bands consisting of fine crystallites of large misorientation were observed in both alloys. Their density was much higher in alloys deformed by 90%. In Fig. 4 the microstructure of deformation band and the corresponding diffraction patterns obtained from the band area and neighbouring grains are shown for the 6XXX alloy deformed 90%. Both grains are only slightly tilted, but grains within a deformation band show high misorientation. This fact suggests the presence of a regular deformation band or shear band [18]. Changes in the precipitation structure in the 6XXX alloy deformed 60% and 90% and aged for 2 h at 165°C (corresponding to maximum of hardness) are visible in the series of bright and dark field images formed for the
Fig. 3. Transmission electron micrographs of 6XXX alloy quenched and deformed 60% by rolling illustrating the tilt experiment to measure the subgrains misorientation equal to 2°.
during ageing for both alloys in spite of the increase of the initial hardness after 30% of deformation up to HV =105 and HV=115 (for the 6013 and 6XXX,
Fig. 4. Transmission electron micrograph of 6XXX alloy deformed 90% after quenching showing deformation band and corresponding diffraction patterns from the left grain, the deformation band and the right grain.
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Fig. 5. The bright and dark field TEM images and corresponding diffraction patterns from 6XXX alloy quenched, deformed 60% (a,c,e) and 90% (b,d,f) and aged for 2 h at 165°C.
6013 alloy shows lower density of precipitates as compared to 6XXX alloy. It can be concluded that increase of hardness maximum with increasing Cu content in the investigated alloys is connected with the increase of precipitates density. The previous results on the effect of predeformation on ageing characteristic of ternary Al –Mg –Si alloys [7,8] indicate fewer precipitates on dislocations when compared to our results. The reason may be the lower degree of deformation, but it could be also the effect of copper, causing higher strains around precipitates and changing the precipitation behaviour. The structure of precipitates seems to be similar to that observed in the case of the quenched and aged alloys, as far as it can be judged from the diffraction pattern analysis; however, one cannot exclude the formation of new phases like those found after deformation and ageing in ternary Al–Mg –Si using HREM method [17]. Rotational moire patterns frequently observed in deformed and aged alloys are shown in Fig. 6 (for the 6XXX alloy). The measured distance between the fringes allowed to calculate the misorientation of 2°–3° degress for fine subgrains formed after ageing indicating the presence of a recovery process occurring already at 165°C. The frequently observed presence of moire pattern suggests that at least two subgrains of very small thickness of several nanometers are present within the foil thickness.
4. Conclusions
Fig. 6. Transmission electron micrograph of 6XXX alloy quenched, deformed 90% and aged for 2 h with rotational moire fringes formed by overlapping subgrains (calculated misorienation is about 3°); corresponding diffraction pattern.
same diffraction conditions with [0 0 1] zone axis (Fig. 5). The dislocation structure and precipitates are well visible in dark field images, which were obtained using precipitates reflections in the positions corresponding to 0 0 1 matrix reflection. The precipitates were formed mainly on dislocations and their density increased with increasing deformation. In the 60% deformed alloy the structure of the precipitates can be identified as a metastable B2 phase. A similar observation made for the deformed and aged
(1) The hardness measurements of the deformed and aged alloys 6013 and 6XXX show that the increase of the degree of deformation above 60% of reduction by rolling increases the maximum hardness up to 30 HV above the maximum attained during artificial ageing at 165°C. At 30% of reduction, the maximum hardness was not changed. The hardness maximum is reached after 2-h ageing instead of 24 h without deformation. (2) The precipitates in the deformed and aged alloys form predominantly on dislocations after the ageing time corresponding to the hardness maximum. Their density increases with the degree of deformation. The diffraction patterns at this stage of ageing are similar for both alloys and indicate formation of a transition cubic phase (at lower deformations) and the Q phase at higher deformations and longer ageing time. (3) The deformation by rolling causes formation of elongated subgrain structure with misorientation increasing with the degree of rolling. At 60% of deformation, the deformation bands begin to form. Their density and thickness increase with the degree of reduction. They consist of very small grains of high misorientation. (4) The recovery process occurs after ageing at 165°C. It is manifested by formation of moire fringes
J. Dutkiewicz, L. Litynska / Materials Science and Engineering A324 (2002) 239–243
indicating tilting type of misorientation (typically 2–3°) through the thickness of the foil. No recrystallization effects can be seen at this ageing temperature. Acknowledgements The financial support of the Research Project (Zamawiany 15–15) of the State Committee for Scientific Research is gratefully acknowledged. References [1] T. Moons, P. Ratchev, P. De Smeet, B. Verlinden, P. Van Houtte, Scripta Mater. 35 (1996) 939. [2] M. Tamizifar, G.W. Lorimer, in: L. Arnberg, O. Lohne, E. Nes, N. Ryum (Eds.), Proceedings of the Third International Conference on Aluminium Alloys, NTH, Trondheim, 1992, p. 220. [3] D.G. Eskin, Z. Metallkd. 83 (1992) 10. [4] H. Uchida, H. Yoshida, H. Hira, T. Amano, Mater. Sci. Forum, vols. 217 – 222, p. 1753. [5] S.C. Bergsma, M.E. Kassner, Mater. Sci. Forum, vols. 217 – 222, p. 1801. [6] G. Burger, A.K. Gupta, L. Sutak, D.J. Lloyd, Mater. Sci Forum, vols. 217 –222, p. 471.
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[7] L. Zhen, S.B. Kang, Mater. Sci. Technol. 14 (1998) 317. [8] Z.H. Ismail, Scripta Metal. Mater. 32 (1995) 457. [9] B. Davis, F.J. Humphreys, K.M. Gatenby, C.P. Johnson, in: T. Ginza, C. Ku (Eds.), Sixth International Conference on Aluminium Alloys, Japan Institute of Light Metals, Tokyo, Japan, 1998. [10] S. Ikeno, K. Matsuda, K Nakajima, S. Rengakuji, Y. Uetani, J. Jpn. Ins. Light Metals 48 (211) (1998) 207. [11] I. Okumura, K. Matsuda, Y. Uetani, F. Shinagawa, S. Ikeno, in: M. Koiwa, K. Otsuka, T. Miyazaki (Eds.), Proceedings of the International Conference on Solid-Solid Phase Transformations, The Japan Institute of Metals, 1999, pp. 205. [12] L. Sagalowicz, G. Hug, D. Bechet, P. Sainfort, G. Lapasset, in: G. Sanders, E.A. Starke (Eds.), Fourth International Conference on Aluminium Alloys, Georgia Institute of Technology, Atlanta, 1994, pp. 636. [13] T. Sakurai, T. Eto, Proceedings of the Third International Conference on Aluminium Alloys, NTH, Trondheim, 1992, pp. 208. [14] K. Matsuda, S. Shimizu, H. Gamada, Y. Uetani, F. Shinagawa, S. Ikeno, J. Soc. Mater. Sci., Japan 48 (1999) 10. [15] H. Cimenoglu, Y. Sun, M. Baydogan, Mater. Lett. 38 (1999) 221. [16] L. Zhen, S.B. Kang, Mater. Sci. Technol. 14 (1998) 317. [17] K. Matsuda, H. Gamada, Y. Uetani, S. Rengakuji, F. Shinagawa, S. Ikeno, J. Jpn. Inst. Light Metals 48 (1998) 471. [18] D. Kuhlman-Wilsdorf, Acta Mater. 47 (1999) 1697.