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Influence of aging treatment on mechanical properties of 6061 aluminum alloy
Materials and Design 31 (2010) 972–975
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Short Communication
Influence of aging treatment on mechanical properties of 6061 aluminum alloy F. Ozturk a,*, A. Sisman b, S. Toros a, S. Kilic a, R.C. Picu c a
Department of Mechanical Engineering, Nigde University, Nigde, Turkey Department of Mechanical Engineering, Kahramanmaras Sutcu Imam University, Kahramanmaras, Turkey c Department of Mechanical, Aerospace and Nuclear Engineering, Rensselaer Polytechnic Institute, Troy, NY, USA b
a r t i c l e
i n f o
Article history: Received 19 June 2009 Accepted 11 August 2009 Available online 14 August 2009
a b s t r a c t Aluminum–magnesium–silicon (Al–Mg–Si) alloys show medium strength, excellent formability, good corrosion resistance and are widely used in extruded products and automotive body panels. The major advantage of these alloys is their age hardening response during the paint baking process as well as the fact that they exhibit no yield point phenomenon and Lüdering. In this study, the mechanical properties of a commercially available AA6061 alloy aged to various levels were studied. Peak-aged conditions were reached in this particular alloy after a 2 h heat treatment at 200 °C. The variation of the yield stress, ultimate tensile strength, ductility and strain hardening rate with aging time is measured and discussed in relation to the microstructural changes induced by the heat treatment. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Aluminum and its alloys are desirable in many industries such as aircraft, automotive, appliances and food packaging due to their high strength to weight ratio and corrosion resistance. Aluminum– magnesium–silicon (Al–Mg–Si) denoted as 6XXX series alloys are medium strength heat treatable alloys and have excellent formability and good corrosion resistance characteristics [1]. Mg and Si are the major solutes; they increase the strength of the alloy by precipitation hardening. There has been a considerable industrial interest in these alloys because two-thirds of all extruded products are made of aluminum and 90% of those are made from 6XXX series alloys [2]. In this series, AA6061 is one of the most widely used alloys [3]. These materials can be heat treated to produce precipitation to various degrees. The T6 treatment involving solution heat treatment and subsequent artificial aging and quenching is a common method to increase the strength of the alloy [4]. The solution heat treatment is first performed at 500 °C to obtain the supersaturated a solid solution. Artificial aging is obtained by heating to about 200 °C for various amounts of time and leads to precipitation of various phases (leading to the stable b phase). The hardness and strength are determined by the precipitate type, density and size [5]. The precipitation sequence was studied by several groups [6–9]. The formation of the stable b phase (Mg2Si) is preceded by a chain of transformations involving various coherent and semi-coherent metastable phases. Specifically, the sequence is: formation of independent clusters of Mg and Si atoms, formation of co-clusters that contain both Si and Mg, formation and growth of b needle-shaped * Corresponding author. Tel.: +90 388 225 2254; fax: +90 388 225 0112. E-mail address: [email protected] (F. Ozturk). 0261-3069/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2009.08.017
precipitates, the transformation of the b precipitates in B lathshaped precipitates and b rod-shaped precipitates, followed at long annealing times by the formation of the incoherent, stable b. Two other post-b phases, denoted as U1 and U2, have been reported. The pre-b phases are coherent, while b is semi-coherent (coherent with the Al matrix along the axis of the needle). Al–Mg–Si series alloys generally have lower formability than Al–Mg alloys (5XXX series), but provide higher strength after paint baking [10]. Age hardening during paint baking is the major advantage of these alloys over the 5XXX series. Demir and Gunduz [11] investigated the effect of artificial aging on the machinability of AA6061 in as-received, solution heat treated and aged conditions. Their results revealed that aging at 180 °C for various times significantly affects the surface roughness of the workpiece. Kim et al. [12] investigated the mechanical and tribological properties of rheo-formed AA6061 wrought alloy. Peak hardness and surface roughness were determined after a 530 °C solution heat treatment for 10 h. Surface roughness increases with the aging time. Preferential precipitation at grain boundaries, the chemistry and crystallography of these precipitates were discussed by de Haas et al. [13]. The presence of grain boundary precipitates, in particular in Si-rich alloys, promotes intergranular fracture and reduced overall ductility. Doward and Bouvier [14] compared several alloys of 6061-type investigating, in particular, the effect of the departure of Si content from the nominal one (excess Si/excess Mg) on the mechanical properties and the nature of fracture. They concluded that optimal combinations of strength and toughness are obtained from Si-lean and balanced chemistries. In the present investigation, the mechanical properties of AA6061 subjected to solution heat treatment followed by aging to various degrees are investigated and discussed in relation to the respective microstructural changes during heat treatment. This
973
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
adds to existing literature data that refer mostly to hardness and the evolution of the microstructure during aging.
Aging time (min) 0.0
15
30
90 120
1.5
2.0
1080
2880
3.0
3.5
90
2. Materials and experimental procedure Commercial AA6061 alloy in the nominal T6 condition in the form of rods of 25 mm diameter were used in this study. The chemical composition was determined by spectral analysis and is given in Table 1. Tensile test specimens were prepared following the ASTM E8 standard. The aging treatment procedure is shown in Fig. 1. Samples in the as-received condition which were solutionized without aging were also tested for reference (denoted as NoHT). The hardness was measured after aging and before mechanical testing using a Mitutoya HV-112 Vickers Hardness Testing Machines with a 10 kg load. Tensile tests were performed by Shimadzu Autograph 100 kN testing machine and deformation was measured with a video-extensometer.
Hardness (HV)
80
70
60
50
0.0
0.5
1.0
2.5
Log aging time (min) Fig. 2. Hardness measurements for various aging conditions.
3. Results and discussion
Table 1 Chemical composition of 6061 alloy (wt.%). Mg
Si
Cu
Mn
Fe
Cr
Zn
0.91
0.80
0.321
0.212
0.456
0.021
0.178
350
300
250
Stress (MPa)
The hardness obtained after aging to various degrees is shown in Fig. 2. In average, 20 hardness measurements were performed for each aging condition. Maximum hardness was reached after approximately 200 min of aging. It should be noted that increasing Mg content slows down aging [9], therefore identifying the proper composition (Table 1) is important. An incubation period of about 10–15 min is observed, in which the hardness does not change significantly. Further aging leads to the increase of hardness, with a maximum obtained after about 200 min. This variation is associated with microstructural evolution [7]. The formation of Mg and Si co-clusters (as well as the formation of nm scale precipitates whose nature was not yet clarified [7]) takes place initially. These contribute marginally to increasing hardness and yield stress. Peak-aging is associated with a dense population of b needleshaped precipitates aligned in the h1 0 0iAl crystal directions, which appear to be optimal barriers for dislocations [7,9]. Part of these precipitates remain in the microstructure during subsequent aging (over-aging) as the other metastable phases, such as B0 and b0 , form. The relationship established in [7] between hardness and the microstructure is used here to identify the state of precipitation in our samples and relate to the other mechanical properties, such as strength and ductility. Specifically, we select the
200
NoHT 15min 30min 90min 120min 1080min 2880min
150
100
50
0 0.00
0.05
0.10
0.15
0.20
0.25
Strain (mm/ mm) Fig. 3. True stress vs. true strain curves.
aging conditions of the samples to be tested such to cover the entire range of microstructures described above, and identify the respective microstructure based on the measured hardness. Stress–strain curves for all aging conditions are presented in Fig. 3. The formation of solute clusters and the subsequent precipitation leads to higher yield and flow stress. In over-aged conditions, the flow stress decreases, as expected. It is interesting to
Fig. 1. Aging process diagram.
974
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
observe that the strain hardening behavior of samples aged up to 30 min is identical to that of the reference NoHT sample. This indicates that formation of solute clusters has some effect on the yield stress, but does not modify significantly the evolution of the dislocation structure (hardening). However, the behavior changes dramatically at peak age, when the microstructure is dominated by small (nanoscale) b” precipitates [9]. The yield stress and hardness increase rapidly with aging time and the overall shape of the stress–strain curve changes. Samples aged longer, including those over-aged, exhibit the same reduced strain hardening rate observed for the b” dominated microstructure. This is associated with a shift from precipitate sharing to Orowan looping, as the particles loose coherency upon continued annealing. Fig. 4 shows the yield stress vs. hardness. The data points organize in two groups, one at low hardness, corresponding to underaged microstructures, and the other at large hardness, corresponding to peak-aged and over-aged samples. The data points from the second group can be described by the equation HV ¼ 13 ry , while those from the first group are well below this line. This correlates with the observed differences between the mechanical behavior of samples with and without b00 precipitates. The variation of the yield stress and ultimate tensile strength (UTS) with aging time is shown in Fig. 5. The yield stress follows largely the evolution of the hardness (Fig. 4), however the UTS ex-
hibit a weaker variation with the aging time. The difference between the yield point and ultimate tensile strength decreases with increasing aging time. This is due both to reduced strain hardening rate and to the reduction of the strain to failure. In peak-aged conditions, the ultimate tensile strength is 321 MPa, which is 73% higher than the corresponding value for the NoHT material. The ductility follows an inverse relation with the strength, as expected. The formability of the material for each aging condition was evaluated by measuring the uniform and total elongation, Fig. 6. These parameters are almost constant for over-aged materials. The uniform elongation is initially (NoHT) 16.4%, 6.6% at peakaged, and 6.6% at over-aged conditions. The corresponding total tensile elongation, which is an important parameter in terms of formability, is 24.6% for NoHT material, 13.2% for peak-aged, and 14.2% for the over-aged materials. These results reveal that two different points of view to the manufacturers are considered. The first one is to form the material first and then aged. The second point is to age the material first and then formed. This decision is generally made based on the application. For instance, if the application requires high strengthening on some region which means the material needs to be aged first in order to get high resistance to thinning and then formed. Otherwise, the desired shape cannot be obtained. 0.26
300
Total and uniform elongation (mm/mm)
y=3x
250
Yield strength (MPa)
Total elongation Uniform elongation
0.24
200
150
100
0.22 0.20 0.18 0.16 0.14 0.12 0.10 0.08 0.06 0.0
50 50
60
70
80
90
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Log aging time (min)
100
Hardness HV Fig. 6. Elongations for various aging conditions. Fig. 4. Yield strengths vs. hardness for various aging conditions.
4000
No HT 15min 30min 90min 120min 1080min 2880min
320
3500
σy
3000
240
dσ/dε
and UTS (MPa)
280
200
2500 2000 1500
160
1000 120
UTS Yield strength
80
500 0
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Log aging time (min) Fig. 5. Yield and ultimate tensile strengths for various aging conditions.
100
150
200
250
300
σ (MPa) Fig. 7. Strain hardening rate vs. stress for various aging conditions.
350
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
Fig. 7 shows strain hardening ddr vs. the flow stress and indicates, as observed above, that peak and over-aged samples have a relatively reduced capability to harden during deformation. This has implications for formability since strain hardening is an important factor promoting stable plastic flow. 4. Conclusions In this study, as-received 6061 alloy was aged and the effect of aging time on mechanical properties was determined. The following conclusions are obtained: (1) Peak-aging conditions are reached after 200 min of aging at 200 °C. Slightly different Mg/Si atomic ratio would lead to a different time to peak age. (2) The presence of b” precipitates leads to significant changes in the mechanical behavior of the material: the yield stress increases significantly, the hardness varies linearly with the yield stress and the strain hardening capability is reduced. Comparatively, all other microstructural changes occurring during aging have a much smaller effect on the mechanical response of the material.
References [1] Murtha SJ. New 6XXX aluminum alloy for automotive body sheet applications. SAE Int J Mater Manuf 1995;104:657–66.
975
[2] Parson NC, Yiu HL. In: Campbell PG, editor. Light metals. Warrendale (PA, USA): TMS; 1989. [3] Buha J, Lumley RN, Crosky AG. Microstructural development and mechanical properties of interrupted aged Al–Mg–Si–Cu alloy. Metall Mater Trans A 2006;37A:3119–30. [4] Abis S, Boeuf A, Caciuffo R, Fiorini P, Magnani M, Melone S, et al. Investigation of Mg2Si precipitation in an Al–Mg–Si alloy by small angle neutron scattering. J Nucl Mater 1985;135:181–9. [5] Buha J, Lumley RN, Crosky AG, Hono K. Secondary precipitation in an Al–Mg– Si–Cu alloy. Acta Mater 2007;55:3015–24. [6] Marioara CD, Nordmark H, Andersen SJ, Holmestad R. Post-b phases and their influence on microstructure ad hardness in 6xxx Al–Mg–Si alloys. J Mater Sci 2006;41:471–8. [7] Edwards GA, Stiller K, Dunlop GL, Couper MJ. The precipitation sequence in Al– Mg–Si alloys. Acta Mater 1998;46:3893–904. [8] Marioara CD, Andersen SJ, Jansen J, Zandbergen HW. The influence of temperature and storage time at RT on nucleation of the beta phase in a 6082 Al–Mg–Si alloy. Acta Mater 2003;51:789–96. [9] Yassar SR, Field PD. Transmission electron microscopy and differential scanning calorimetry studies on the precipitation sequence in an Al–Mg–Si alloy: AA6022. J Mater Res 2005;20:2705–11. [10] Burger GB, Gupta AK, Jeffrey PW, Llyod DJ. Microstructural control of aluminum sheet used in automotive applications. Mater Character 1995;35:23–39. [11] Demir H, Gunduz S. The effects of aging on machinability of 6061 aluminum alloy. Mater Des 2009;30:1480–3. [12] Kim HH, Cho SH, Kang CG. Evaluation of microstructure and mechanical properties by using nano/micro-indentation and nanoscratch during aging treatment of rheo-forged Al 6061 alloy. Mater Sci Eng A 2008;485: 272–81. [13] de Haas M, van Scherpenzeel SM, de Hosson JTM. Grain boundary segregation and precipitation in Al alloy AA6061. Mat Sci Forum 2006;519–521:467–72. [14] Doward RC, Bouvier C. A rationalization of factors affecting strength, ductility and toughness of AA6061-type Al–Mg–Si–(Cu) alloys. Mater Sci Eng A 1998;254:33–44.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Short Communication
Influence of aging treatment on mechanical properties of 6061 aluminum alloy F. Ozturk a,*, A. Sisman b, S. Toros a, S. Kilic a, R.C. Picu c a
Department of Mechanical Engineering, Nigde University, Nigde, Turkey Department of Mechanical Engineering, Kahramanmaras Sutcu Imam University, Kahramanmaras, Turkey c Department of Mechanical, Aerospace and Nuclear Engineering, Rensselaer Polytechnic Institute, Troy, NY, USA b
a r t i c l e
i n f o
Article history: Received 19 June 2009 Accepted 11 August 2009 Available online 14 August 2009
a b s t r a c t Aluminum–magnesium–silicon (Al–Mg–Si) alloys show medium strength, excellent formability, good corrosion resistance and are widely used in extruded products and automotive body panels. The major advantage of these alloys is their age hardening response during the paint baking process as well as the fact that they exhibit no yield point phenomenon and Lüdering. In this study, the mechanical properties of a commercially available AA6061 alloy aged to various levels were studied. Peak-aged conditions were reached in this particular alloy after a 2 h heat treatment at 200 °C. The variation of the yield stress, ultimate tensile strength, ductility and strain hardening rate with aging time is measured and discussed in relation to the microstructural changes induced by the heat treatment. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Aluminum and its alloys are desirable in many industries such as aircraft, automotive, appliances and food packaging due to their high strength to weight ratio and corrosion resistance. Aluminum– magnesium–silicon (Al–Mg–Si) denoted as 6XXX series alloys are medium strength heat treatable alloys and have excellent formability and good corrosion resistance characteristics [1]. Mg and Si are the major solutes; they increase the strength of the alloy by precipitation hardening. There has been a considerable industrial interest in these alloys because two-thirds of all extruded products are made of aluminum and 90% of those are made from 6XXX series alloys [2]. In this series, AA6061 is one of the most widely used alloys [3]. These materials can be heat treated to produce precipitation to various degrees. The T6 treatment involving solution heat treatment and subsequent artificial aging and quenching is a common method to increase the strength of the alloy [4]. The solution heat treatment is first performed at 500 °C to obtain the supersaturated a solid solution. Artificial aging is obtained by heating to about 200 °C for various amounts of time and leads to precipitation of various phases (leading to the stable b phase). The hardness and strength are determined by the precipitate type, density and size [5]. The precipitation sequence was studied by several groups [6–9]. The formation of the stable b phase (Mg2Si) is preceded by a chain of transformations involving various coherent and semi-coherent metastable phases. Specifically, the sequence is: formation of independent clusters of Mg and Si atoms, formation of co-clusters that contain both Si and Mg, formation and growth of b needle-shaped * Corresponding author. Tel.: +90 388 225 2254; fax: +90 388 225 0112. E-mail address: [email protected] (F. Ozturk). 0261-3069/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2009.08.017
precipitates, the transformation of the b precipitates in B lathshaped precipitates and b rod-shaped precipitates, followed at long annealing times by the formation of the incoherent, stable b. Two other post-b phases, denoted as U1 and U2, have been reported. The pre-b phases are coherent, while b is semi-coherent (coherent with the Al matrix along the axis of the needle). Al–Mg–Si series alloys generally have lower formability than Al–Mg alloys (5XXX series), but provide higher strength after paint baking [10]. Age hardening during paint baking is the major advantage of these alloys over the 5XXX series. Demir and Gunduz [11] investigated the effect of artificial aging on the machinability of AA6061 in as-received, solution heat treated and aged conditions. Their results revealed that aging at 180 °C for various times significantly affects the surface roughness of the workpiece. Kim et al. [12] investigated the mechanical and tribological properties of rheo-formed AA6061 wrought alloy. Peak hardness and surface roughness were determined after a 530 °C solution heat treatment for 10 h. Surface roughness increases with the aging time. Preferential precipitation at grain boundaries, the chemistry and crystallography of these precipitates were discussed by de Haas et al. [13]. The presence of grain boundary precipitates, in particular in Si-rich alloys, promotes intergranular fracture and reduced overall ductility. Doward and Bouvier [14] compared several alloys of 6061-type investigating, in particular, the effect of the departure of Si content from the nominal one (excess Si/excess Mg) on the mechanical properties and the nature of fracture. They concluded that optimal combinations of strength and toughness are obtained from Si-lean and balanced chemistries. In the present investigation, the mechanical properties of AA6061 subjected to solution heat treatment followed by aging to various degrees are investigated and discussed in relation to the respective microstructural changes during heat treatment. This
973
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
adds to existing literature data that refer mostly to hardness and the evolution of the microstructure during aging.
Aging time (min) 0.0
15
30
90 120
1.5
2.0
1080
2880
3.0
3.5
90
2. Materials and experimental procedure Commercial AA6061 alloy in the nominal T6 condition in the form of rods of 25 mm diameter were used in this study. The chemical composition was determined by spectral analysis and is given in Table 1. Tensile test specimens were prepared following the ASTM E8 standard. The aging treatment procedure is shown in Fig. 1. Samples in the as-received condition which were solutionized without aging were also tested for reference (denoted as NoHT). The hardness was measured after aging and before mechanical testing using a Mitutoya HV-112 Vickers Hardness Testing Machines with a 10 kg load. Tensile tests were performed by Shimadzu Autograph 100 kN testing machine and deformation was measured with a video-extensometer.
Hardness (HV)
80
70
60
50
0.0
0.5
1.0
2.5
Log aging time (min) Fig. 2. Hardness measurements for various aging conditions.
3. Results and discussion
Table 1 Chemical composition of 6061 alloy (wt.%). Mg
Si
Cu
Mn
Fe
Cr
Zn
0.91
0.80
0.321
0.212
0.456
0.021
0.178
350
300
250
Stress (MPa)
The hardness obtained after aging to various degrees is shown in Fig. 2. In average, 20 hardness measurements were performed for each aging condition. Maximum hardness was reached after approximately 200 min of aging. It should be noted that increasing Mg content slows down aging [9], therefore identifying the proper composition (Table 1) is important. An incubation period of about 10–15 min is observed, in which the hardness does not change significantly. Further aging leads to the increase of hardness, with a maximum obtained after about 200 min. This variation is associated with microstructural evolution [7]. The formation of Mg and Si co-clusters (as well as the formation of nm scale precipitates whose nature was not yet clarified [7]) takes place initially. These contribute marginally to increasing hardness and yield stress. Peak-aging is associated with a dense population of b needleshaped precipitates aligned in the h1 0 0iAl crystal directions, which appear to be optimal barriers for dislocations [7,9]. Part of these precipitates remain in the microstructure during subsequent aging (over-aging) as the other metastable phases, such as B0 and b0 , form. The relationship established in [7] between hardness and the microstructure is used here to identify the state of precipitation in our samples and relate to the other mechanical properties, such as strength and ductility. Specifically, we select the
200
NoHT 15min 30min 90min 120min 1080min 2880min
150
100
50
0 0.00
0.05
0.10
0.15
0.20
0.25
Strain (mm/ mm) Fig. 3. True stress vs. true strain curves.
aging conditions of the samples to be tested such to cover the entire range of microstructures described above, and identify the respective microstructure based on the measured hardness. Stress–strain curves for all aging conditions are presented in Fig. 3. The formation of solute clusters and the subsequent precipitation leads to higher yield and flow stress. In over-aged conditions, the flow stress decreases, as expected. It is interesting to
Fig. 1. Aging process diagram.
974
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
observe that the strain hardening behavior of samples aged up to 30 min is identical to that of the reference NoHT sample. This indicates that formation of solute clusters has some effect on the yield stress, but does not modify significantly the evolution of the dislocation structure (hardening). However, the behavior changes dramatically at peak age, when the microstructure is dominated by small (nanoscale) b” precipitates [9]. The yield stress and hardness increase rapidly with aging time and the overall shape of the stress–strain curve changes. Samples aged longer, including those over-aged, exhibit the same reduced strain hardening rate observed for the b” dominated microstructure. This is associated with a shift from precipitate sharing to Orowan looping, as the particles loose coherency upon continued annealing. Fig. 4 shows the yield stress vs. hardness. The data points organize in two groups, one at low hardness, corresponding to underaged microstructures, and the other at large hardness, corresponding to peak-aged and over-aged samples. The data points from the second group can be described by the equation HV ¼ 13 ry , while those from the first group are well below this line. This correlates with the observed differences between the mechanical behavior of samples with and without b00 precipitates. The variation of the yield stress and ultimate tensile strength (UTS) with aging time is shown in Fig. 5. The yield stress follows largely the evolution of the hardness (Fig. 4), however the UTS ex-
hibit a weaker variation with the aging time. The difference between the yield point and ultimate tensile strength decreases with increasing aging time. This is due both to reduced strain hardening rate and to the reduction of the strain to failure. In peak-aged conditions, the ultimate tensile strength is 321 MPa, which is 73% higher than the corresponding value for the NoHT material. The ductility follows an inverse relation with the strength, as expected. The formability of the material for each aging condition was evaluated by measuring the uniform and total elongation, Fig. 6. These parameters are almost constant for over-aged materials. The uniform elongation is initially (NoHT) 16.4%, 6.6% at peakaged, and 6.6% at over-aged conditions. The corresponding total tensile elongation, which is an important parameter in terms of formability, is 24.6% for NoHT material, 13.2% for peak-aged, and 14.2% for the over-aged materials. These results reveal that two different points of view to the manufacturers are considered. The first one is to form the material first and then aged. The second point is to age the material first and then formed. This decision is generally made based on the application. For instance, if the application requires high strengthening on some region which means the material needs to be aged first in order to get high resistance to thinning and then formed. Otherwise, the desired shape cannot be obtained. 0.26
300
Total and uniform elongation (mm/mm)
y=3x
250
Yield strength (MPa)
Total elongation Uniform elongation
0.24
200
150
100
0.22 0.20 0.18 0.16 0.14 0.12 0.10 0.08 0.06 0.0
50 50
60
70
80
90
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Log aging time (min)
100
Hardness HV Fig. 6. Elongations for various aging conditions. Fig. 4. Yield strengths vs. hardness for various aging conditions.
4000
No HT 15min 30min 90min 120min 1080min 2880min
320
3500
σy
3000
240
dσ/dε
and UTS (MPa)
280
200
2500 2000 1500
160
1000 120
UTS Yield strength
80
500 0
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
Log aging time (min) Fig. 5. Yield and ultimate tensile strengths for various aging conditions.
100
150
200
250
300
σ (MPa) Fig. 7. Strain hardening rate vs. stress for various aging conditions.
350
F. Ozturk et al. / Materials and Design 31 (2010) 972–975
Fig. 7 shows strain hardening ddr vs. the flow stress and indicates, as observed above, that peak and over-aged samples have a relatively reduced capability to harden during deformation. This has implications for formability since strain hardening is an important factor promoting stable plastic flow. 4. Conclusions In this study, as-received 6061 alloy was aged and the effect of aging time on mechanical properties was determined. The following conclusions are obtained: (1) Peak-aging conditions are reached after 200 min of aging at 200 °C. Slightly different Mg/Si atomic ratio would lead to a different time to peak age. (2) The presence of b” precipitates leads to significant changes in the mechanical behavior of the material: the yield stress increases significantly, the hardness varies linearly with the yield stress and the strain hardening capability is reduced. Comparatively, all other microstructural changes occurring during aging have a much smaller effect on the mechanical response of the material.
References [1] Murtha SJ. New 6XXX aluminum alloy for automotive body sheet applications. SAE Int J Mater Manuf 1995;104:657–66.
975
[2] Parson NC, Yiu HL. In: Campbell PG, editor. Light metals. Warrendale (PA, USA): TMS; 1989. [3] Buha J, Lumley RN, Crosky AG. Microstructural development and mechanical properties of interrupted aged Al–Mg–Si–Cu alloy. Metall Mater Trans A 2006;37A:3119–30. [4] Abis S, Boeuf A, Caciuffo R, Fiorini P, Magnani M, Melone S, et al. Investigation of Mg2Si precipitation in an Al–Mg–Si alloy by small angle neutron scattering. J Nucl Mater 1985;135:181–9. [5] Buha J, Lumley RN, Crosky AG, Hono K. Secondary precipitation in an Al–Mg– Si–Cu alloy. Acta Mater 2007;55:3015–24. [6] Marioara CD, Nordmark H, Andersen SJ, Holmestad R. Post-b phases and their influence on microstructure ad hardness in 6xxx Al–Mg–Si alloys. J Mater Sci 2006;41:471–8. [7] Edwards GA, Stiller K, Dunlop GL, Couper MJ. The precipitation sequence in Al– Mg–Si alloys. Acta Mater 1998;46:3893–904. [8] Marioara CD, Andersen SJ, Jansen J, Zandbergen HW. The influence of temperature and storage time at RT on nucleation of the beta phase in a 6082 Al–Mg–Si alloy. Acta Mater 2003;51:789–96. [9] Yassar SR, Field PD. Transmission electron microscopy and differential scanning calorimetry studies on the precipitation sequence in an Al–Mg–Si alloy: AA6022. J Mater Res 2005;20:2705–11. [10] Burger GB, Gupta AK, Jeffrey PW, Llyod DJ. Microstructural control of aluminum sheet used in automotive applications. Mater Character 1995;35:23–39. [11] Demir H, Gunduz S. The effects of aging on machinability of 6061 aluminum alloy. Mater Des 2009;30:1480–3. [12] Kim HH, Cho SH, Kang CG. Evaluation of microstructure and mechanical properties by using nano/micro-indentation and nanoscratch during aging treatment of rheo-forged Al 6061 alloy. Mater Sci Eng A 2008;485: 272–81. [13] de Haas M, van Scherpenzeel SM, de Hosson JTM. Grain boundary segregation and precipitation in Al alloy AA6061. Mat Sci Forum 2006;519–521:467–72. [14] Doward RC, Bouvier C. A rationalization of factors affecting strength, ductility and toughness of AA6061-type Al–Mg–Si–(Cu) alloys. Mater Sci Eng A 1998;254:33–44.