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Precipitation in Al–Cu–Mg alloy during creep exposure
Materials Science & Engineering A 556 (2012) 796–800
Contents lists available at SciVerse ScienceDirect
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Precipitation in Al–Cu–Mg alloy during creep exposure Y.C. Lin a,b,c,n, Yu-Chi Xia a,b, Yu-Qiang Jiang a,b, Lei-Ting Li a,b a b c
School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China State Key Laboratory of High Performance Complex Manufacturing, Changsha 410083, China State Key Laboratory of Material Processing and Die & Mould Technology, Wuhan, 430074, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 May 2012 Received in revised form 14 July 2012 Accepted 16 July 2012 Available online 24 July 2012
The uniaxial tensile creep tests of Al–Cu–Mg alloy were carried out over a wide range of temperature and applied stress. The effects of applied stress and creep aging temperature on the precipitation in Al–Cu–Mg alloy were discussed. It is found that the precipitation process is very sensitive to the applied stress and creep aging temperature, and S phase (CuMgAl2) is the main precipitate under the tested creep conditions. With the increase of applied stress and creep aging temperature, S phase easily grows up and becomes sparse. Meanwhile, the creep aging leads to the discontinuous distribution of precipitation phase in grain boundary, which can improve the corrosion-resistance of Al–Cu–Mg alloy. Because the lattice misfit of acicular exudation phase is far less than that of disk-shaped phase, the applied stress/strain fields will alert the competitive precipitation equilibrium between different strengthening phases. So, the large applied stress and high aging temperature easily make the preferential precipitation process as SSS-GPB-S00 -S0 -S. & 2012 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloy Aging Precipitation
1. Introduction Due to the excellent properties of aluminum, the aluminum and its alloys are widely used for aerospace and aviation [1–6]. Al–Cu–Mg alloy, a kind of deformation aluminum alloy, is developed primarily for elevated temperature applications requiring improved short transverse properties (e.g., fracture toughness and ductility), relatively high strength, good corrosion resistance, and creep resistance at elevated temperatures. For example, the primary use of 2124-T851 aluminum alloy thick plate is for aerospace vehicles structure, such as machined fuselage bulkheads and wing skins in high-performance military aircraft. The excellent strength properties of the aluminum alloys depend on the precipitation of metastable and stable phases during the age hardening treatment. Despite some efforts invested into the hot deformation behaviors of Al–Cu–Mg alloy [7–12], the precipitation process during creep aging needs to be further investigated. Some researchers have investigated the effects of the aging processing parameters on the precipitation process of strengthening phases for different metals or alloys [13–19]. Ringer et al. [13] studied the clustering and precipitation reactions during the early stages of aging at 180 1C in an Al–1.7Cu–0.3Mg (at%) alloy and equivalent alloys containing 0.1 and 0.2 at% Ag. Raviprasad
n Corresponding author at: Central South University, School of Mechanical and Electrical Engineering, Changsha 410083, China. Tel.: þ86 15200817337. E-mail addresses: [email protected], [email protected] (Y.C. Lin).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.07.069
et al. [14] investigated the influence of combined microalloying additions of Ag and Si on the evolution of microstructure, and the age hardening behavior of an Al–2.5Cu–1.5Mg (at%) base alloy produce a fine scale microstructure, which involves precipitation processes across the Al–Cu–Mg phase diagram. Wang and Starink [15] gave a critical review on the formation of precipitates and intermetallic phases in dilute precipitation hardening Al–Cu–Mg alloys and without Li additions. Zhu and Starke [16] investigated the aging process of Al–Cu alloy and found that the stressorienting effect is relative with the applied stress, temperature, alloying components and aging time. Some researchers investigated the relationship between the evolution of the hardening precipitation during creep and the creep behavior of Al–Cu–Mg alloy. Przydatek [17] confirmed the coarsening of the hardening precipitation during creep at 175 1C for 2650 Al alloy. Majimel et al. [18] showed that S00 tends to disappear during thermal exposure presumably because S00 is consumed by S0 precipitates. The disappearance of S00 is responsible for the softening of 2650 Al alloy. Feng et al. [19] investigated the heterogeneous nucleation and growth of precipitates at dislocations in Al–Cu–Mg alloy by examining samples aged at 195 1C for various times from 10 min to 9 h, and found that the precipitation sequence of S (Al2CuMg) phase along dislocations is SSS-GPB zones-S (Type I) -S (Type I) þS (Type II). The objective of this study is to investigate the precipitation processes in Al–Cu–Mg alloy during creep aging. The uniaxial tensile creep tests of Al–Cu–Mg alloy were carried out over a wide range of temperature and applied stress. Also, the effects of the
Y.C. Lin et al. / Materials Science & Engineering A 556 (2012) 796–800
applied stress and creep temperature on the precipitation processes of Al–Cu–Mg alloy were discussed.
2. Material and experiments 2124-T851 aluminum alloy, which is one typical commercial Al–Cu–Mg alloy, was used in this investigation. This material was made by hot rolling, pre-stretching (stretch leveling ratio 1.8%) and heat treatment [20]. The compositions (wt%) of this material is 4.67Cu–1.46Mg–0.63Mn–0.18Fe–0.04Zn–0.01Ti–0.12Si–(bal.) Al. The specimens have been machined out from the rolled Al–Cu–Mg alloy plate with gauge length of 40 mm and diameter of 10 mm. The uniaxial tensile creep experiments were carried out in MTS-GWT2105 creep test machine. The controlling accuracies of the applied stress and temperature are 70.1 MPa and 72 K, respectively. The grating ruler is used to measure the creep displacement and it can be recorded by computer automatically, and the accuracy of grating ruler is 0.1 mm. Three different temperatures (473 K, 503 K and 533 K) and three different applied stresses (120 MPa, 140 MPa and 160 MPa) were used in the creep aging experiments. Firstly, the specimens were heated to the test temperature at a heating rate of 10 K/min, and hold for 0.5 h to eliminate thermal gradients before loading. After test, the fractured samples were naturally cooled to the room temperature in the furnace. In order to investigate the precipitation processes of Al–Cu–Mg alloy during creep aging, transmission electron microscopy (TEM) was carried out. The thin foil samples for TEM experiments were prepared in the following steps. Firstly, the block specimens were cut into 0.5–1.0 mm thick foil. Then, several disks with 3 mm in diameter and 0.25 mm in thickness were punched out from these slices, and subsequently electropolished using a solution of HNO3 and methanol (1:3 in volume). TEM foils were examined using a TecnaiG2-20 microscope operating at 200 KV.
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Fig. 1 shows the isothermal section of ternary Al–Cu–Mg phase diagram at 473 K. Some investigations [27–29] indicate that (1) when the ratio of Cu/Mg is between 8 and 4, the main precipitate are y and S phases at the same time. (2) When the ratio of Cu/Mg ranges from 4 to 1.5, the main precipitate is S phase. 2124-T851 aluminum alloy used in this study is one typical commercial Al–Cu–Mg alloy. For 2124-T851 aluminum alloy, the contents of Cu and Mg are 4.67% and 1.46% (wt%), respectively. So, the ratio of Cu/Mg is about 3.198, and then it can be concluded that the main precipitate is S phase during creep aging. Fig. 2 shows the TEM micrographs of as-received and creep aged samples. Here, the creep aging conditions are the temperature of 533 K and the applied stress of 140 MPa. From Fig. 2, it can be found that the acicular exudation phase is the main precipitate within the grains for both samples. For the samples without creep aging, as shown in Fig. 2(a), the acicular exudation phases homogeneously distribute within the grains, and the corresponding direction parallel to [1 0 0]Al or[0 1 1]Al. For the samples subjected to the creep aging, as shown in Fig. 2(b), the acicular exudation phase grows up and the quantity of acicular exudation phase decreases. Still, the distribution of precipitate phase is uniform within the grain. Fig. 3 shows the TEM micrographs and corresponding diffraction patterns for the sample without creep aging. The [1 1 2]Al selected area diffraction (SAD) pattern (shown in Fig. 3b) reveals light spots, and the light spots existing at 1/3o2004 and 2/3o2004 directions corresponding to S phase. So, it can be found that the main precipitate is S phase in 2124-T851 aluminum alloy. This result is consistent with others’ report [30]. Fig. 4 shows the TEM micrographs and corresponding diffraction patterns for the sample subjected to the creep aging under the temperature of 473 K and applied stress of 160 MPa for 100 h. Obviously, the [110]Al selected area diffraction (SAD) pattern (shown in Fig. 4b) reveals cross reflections [31], which match all spots of S phase.
3. Results and discussion 3.1. Precipitation hardening processes in Al–Cu–Mg alloys Precipitation strengthening is the main strengthening mechanism for ultra-high strength alloys, and the types of precipitation phase depend on the contents of alloy elements. For Al–Cu–Mg alloy, the types of precipitates are relative with the ratio of Cu/Mg. When the content of Mg element is low, the main precipitate is y phase (CuAl2), and the quantities of y phase decrease with the increase of the content of Mg element. When the Mg element is increased to some extent, the main precipitate will be S phase (CuMgAl2), which has the excellent high temperature strength. However, if the content of Mg element is continuously increased, T phase (A16CuMg4) and b phase (Al3Mg2) appear, which will deteriorate the mechanical properties of Al–Cu–Mg alloy. So, the content of Mg element is too high to strengthen the alloys [21]. Some investigators [22–25] proposed the following precipitation sequence of y and S. 00
Fig.1. Isothermal section of ternary Al–Cu–Mg phase diagram at 473 K (a ¼Al, y ¼ CuAl2, S ¼Al2CuMg, T¼ Al6CuMg4 and b ¼ Al12CuMg17).
0
SSS-GP-y -y -y SSS-GPB-S00 -S0 -S where SSS represents the supersaturated solid solution. The term GPB (Guinier–Preston–Bagaryatsky) zones was firstly reported by Silcock who suggested GPB zones (Cu and Mg ‘‘atomic clusters’’) are different from GP (Guinier–Preston) zone in Al–Cu alloys discovered earlier [26]. This process is ‘‘competitive balance between precipitations’’.
Fig. 2. TEM micrographs of (a) as-received condition; and (b) 533 K 140 MPa.
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d2 ¼0.0188, d3 ¼0.0149 [33]. Obviously, the misfit of acicular
Fig. 3. TEM micrographs and corresponding diffraction patterns for the as-received sample.
exudation phase is far less than the misfit of disk-shaped phase. According to Eshelby’s theory [34–35], firstly, the larger the misfit of precipitation phase is, the more the nucleation-growth stage of precipitation phase is sensitive to the applied stress. Secondly, the shape of precipitation phase depends on the elastic strain energy and interface energy, and the elastic strain energy is the main resistance for the growth of precipitation phase. Generally, the elastic strain energy of acicular and sphere phases is lower than that of disk-shaped phase. So, the elastic strain energy of diskshaped phase will strongly influence the behavior of precipitation phase. i.e., comparing to S0 phase, y0 phase is more sensitive to the applied stress field. So, the applied stress will break the precipitation equilibrium and alert the precipitation processes. For the studied Al–Cu–Mg alloy, the large applied stress and high aging temperature easily make the preferential precipitation process as 00 0 SSS-GPB-S -S -S, in order to decrease the active energy. This precipitation sequence will take in more solute atoms in the metallic matrix. Meanwhile, this process will decrease the content of y0 phase. In other words, the applied stress field will hinder the precipitation of y phase and promote the growth of S phase for the studied Al–Cu–Mg alloy.
3.2. Effect of creep test parameters on precipitation hardening
Fig. 4. TEM micrographs and corresponding diffraction patterns for the sample aged at 473 K 160 MPa.
From the above experimental results and analysis, it can be found that the preferential growth orientation of precipitation phase is sensitive to the applied stress for the studied 2124-T851 aluminum alloy. The acicular exudation precipitation phases uniformly distribute in the metallic matrix and the precipitation phases grow up in a preferential orientation. This is because that the stress/strain field induced by the creep aging conditions changes the system’s elastic energy, which makes the precipitation phases grow up in a preferential orientation, namely, stressorientation effect. The structures of GP or GPB zone is extremely asymmetric in the space, which results in the different misfits between the strengthening phase and metallic matrix in different orientations. Therefore, there is great difference of strain energy in different directions, and then the stress oriented effect is obvious at the nucleation stage of precipitation phase [32]. When alloys are affected by the applied stress, precipitation phases will grow up in a preferential orientation to balance the difference of strain energy and decrease the system’s energy. Additionally, the effects of stress on the penetration rates in various orientations are strongly linked to the anisotropic microstructure. Generally, the status (coherent or incoherent) of the strengthening phases is affected by the changes of stress and strain distortion fields, which are induced by the misfits between aluminum matrix and precipitation particles during the precipitation processes. Different precipitation phases are of different shapes, structures and physical properties in Al–Cu–Mg alloy. On the one hand, GP zone, y00 phase and y0 phase is formed by diskshaped Cu atoms segregated at {001} surface of aluminum matrix, and these precipitation phases is of tetragonal misfit strain field and low face misfit. These precipitation phases’ marginal misfit is d3,[GP] ¼ 0.10, d3y00 ¼ 0.052, d3y0 ¼ 0.045. On the other hand, the S0 phase is acicular, and GPB zone is solute cluster which grow along with the direction of o001 4 in aluminum matrix. So, S0 phase is of lower lattice misfit, and the misfit is d1 ¼ 0.0123,
3.2.1. Effect of applied stress Fig. 5 shows the TEM micrographs for the sample at the creep aging temperature of 473 K and applied stress of 120, 140, and 160 MPa. It can be found that, for the sample without creep aging (Fig. 2(a)), the length of acicular exudation phase is about 140 nm. With the increase of applied stress, the acicular exudation phase grows up and the quantity of acicular exudation phase decreases. For example, the length of S phase is near 200 nm after creep aging with the applied stress of 140 MPa. This is because the applied stress changes the competition of precipitation processes
Fig. 5. TEM micrographs for the sample aged at the creep aging temperature of 473 K and applied stress of (a) 120 MPa; (b) 140 MPa; and (c) 160 MPa.
Y.C. Lin et al. / Materials Science & Engineering A 556 (2012) 796–800
of the studied Al–Cu–Mg alloy, and the precipitation phases S is easy to grow up and coarsen under high stress. 3.2.2. Effect of temperature In this section, the effects of creep aging temperature on the precipitation process of Al–Cu–Mg alloys will be discussed. Fig. 2(b) and Fig. 5(b) show the TEM micrographs under the applied stress of 140 MPa and creep aging temperature of 533 K and 473 K, respectively. It can be found that with the increase of creep aging temperature, the scale of S phase quickly increased. Because the growth of S phase absorbs the solute atoms around it, the density of S phase decreases. The lengths of acicular exudation phase for the samples at the creep aging temperature of 533 K and 473 K are near and over 200 nm, respectively. However, for the samples without creep aging, the length of S phase is only about 140 nm. 3.2.3. Effect of creep exposure on precipitate free zones Grain boundaries act as sinks for solute and vacancies. This has the effect of increasing the precipitation on the grain boundary relative to the surrounding region, often leaving a precipitate free zone (PFZ) close to the boundary. This phenomenon is detrimental to the material since there is a reduction in precipitation hardening in the PFZ. Fig. 6 shows the precipitation phases on the grain boundaries of as-received and creep aged samples. Here, the creep aging conditions is the temperature of 473 K and applied stress of 120, 140, and 160 MPa. From Fig. 6, it can be found that the precipitates almost continuously distribute on the grain boundaries for the as-received samples, and the mechanical properties will be greatly deteriorated. However, for those creep aged ones, the grain boundary precipitates merge and concentrate, which result in a discontinuous distribution of precipitates on the grain boundaries. This is because the growth of precipitate needs to take in a lot of supersaturated solute atoms around the grain boundary. In other
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words, it suppresses the growth of precipitates around the grain boundaries. Therefore, the precipitates around grain boundaries are smaller than those in metallic matrix. Moreover, the discontinuous distribution of grain boundary precipitates can improve the material’s corrosion-resistance.
4. Conclusions The effects of the applied stress and creep aging temperature on the precipitation processes of 2124-T851 aluminum alloy were investigated by uniaxial tensile creep experiments. It is found that the precipitation process is very sensitive to the applied stress and creep aging temperature, and the dominant precipita00 0 tion process is SSS-GPB-S -S -S. So, the main precipitate is S phases under the tested creep aging conditions. With the increase of the applied stress and creep aging temperature, the S phase easily grows up and coarsens, and then the density of it decreases. Meanwhile, the creep aging will result in the discontinuous distribution of precipitation phase in the grain boundaries, which is useful to improve the corrosion-resistance of Al–Cu–Mg alloy.
Acknowledgments This work was supported by the National Key Basic Research Program (Grant no. 2010CB731702), the Program for New Century Excellent Talents in University (No. NCET-10-0838), the Sheng-hua Yu-ying Program and the Graduate Degree Thesis Innovation Foundation of Central South University (no. 2011ssxt094), State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology (no. 2012P04), China. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]
Fig. 6. Precipitation phases on grain boundary for the sample aged at different creep aging conditions. (a) as-received condition; (b) 473 K–120 MPa; (c) 473 K– 140 MPa; and (d) 473 K 160 MPa.
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[25] S.P. Ringer, K. Hono, I.J. Polmear, T. Sakurai, Appl. Surf. Sci. 94–95 (1996) 253–260. [26] J.M. Silcock, J. Inst. Met. 89 (1960) 203–210. [27] L.F. Mondolfo, Aluminum Alloys: Structure and Properties, Butterworths, London, 1976. [28] C.G. Cordovilla, E. Louis, J. Mater. Sci. 19 (1984) 279–290. [29] S.P. Ringer, B.T. Sofyan, K.S. Prasad, G.C. Quan, Acta Mater. 56 (2008) 2147–2160.
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Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Precipitation in Al–Cu–Mg alloy during creep exposure Y.C. Lin a,b,c,n, Yu-Chi Xia a,b, Yu-Qiang Jiang a,b, Lei-Ting Li a,b a b c
School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China State Key Laboratory of High Performance Complex Manufacturing, Changsha 410083, China State Key Laboratory of Material Processing and Die & Mould Technology, Wuhan, 430074, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 May 2012 Received in revised form 14 July 2012 Accepted 16 July 2012 Available online 24 July 2012
The uniaxial tensile creep tests of Al–Cu–Mg alloy were carried out over a wide range of temperature and applied stress. The effects of applied stress and creep aging temperature on the precipitation in Al–Cu–Mg alloy were discussed. It is found that the precipitation process is very sensitive to the applied stress and creep aging temperature, and S phase (CuMgAl2) is the main precipitate under the tested creep conditions. With the increase of applied stress and creep aging temperature, S phase easily grows up and becomes sparse. Meanwhile, the creep aging leads to the discontinuous distribution of precipitation phase in grain boundary, which can improve the corrosion-resistance of Al–Cu–Mg alloy. Because the lattice misfit of acicular exudation phase is far less than that of disk-shaped phase, the applied stress/strain fields will alert the competitive precipitation equilibrium between different strengthening phases. So, the large applied stress and high aging temperature easily make the preferential precipitation process as SSS-GPB-S00 -S0 -S. & 2012 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloy Aging Precipitation
1. Introduction Due to the excellent properties of aluminum, the aluminum and its alloys are widely used for aerospace and aviation [1–6]. Al–Cu–Mg alloy, a kind of deformation aluminum alloy, is developed primarily for elevated temperature applications requiring improved short transverse properties (e.g., fracture toughness and ductility), relatively high strength, good corrosion resistance, and creep resistance at elevated temperatures. For example, the primary use of 2124-T851 aluminum alloy thick plate is for aerospace vehicles structure, such as machined fuselage bulkheads and wing skins in high-performance military aircraft. The excellent strength properties of the aluminum alloys depend on the precipitation of metastable and stable phases during the age hardening treatment. Despite some efforts invested into the hot deformation behaviors of Al–Cu–Mg alloy [7–12], the precipitation process during creep aging needs to be further investigated. Some researchers have investigated the effects of the aging processing parameters on the precipitation process of strengthening phases for different metals or alloys [13–19]. Ringer et al. [13] studied the clustering and precipitation reactions during the early stages of aging at 180 1C in an Al–1.7Cu–0.3Mg (at%) alloy and equivalent alloys containing 0.1 and 0.2 at% Ag. Raviprasad
n Corresponding author at: Central South University, School of Mechanical and Electrical Engineering, Changsha 410083, China. Tel.: þ86 15200817337. E-mail addresses: [email protected], [email protected] (Y.C. Lin).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.07.069
et al. [14] investigated the influence of combined microalloying additions of Ag and Si on the evolution of microstructure, and the age hardening behavior of an Al–2.5Cu–1.5Mg (at%) base alloy produce a fine scale microstructure, which involves precipitation processes across the Al–Cu–Mg phase diagram. Wang and Starink [15] gave a critical review on the formation of precipitates and intermetallic phases in dilute precipitation hardening Al–Cu–Mg alloys and without Li additions. Zhu and Starke [16] investigated the aging process of Al–Cu alloy and found that the stressorienting effect is relative with the applied stress, temperature, alloying components and aging time. Some researchers investigated the relationship between the evolution of the hardening precipitation during creep and the creep behavior of Al–Cu–Mg alloy. Przydatek [17] confirmed the coarsening of the hardening precipitation during creep at 175 1C for 2650 Al alloy. Majimel et al. [18] showed that S00 tends to disappear during thermal exposure presumably because S00 is consumed by S0 precipitates. The disappearance of S00 is responsible for the softening of 2650 Al alloy. Feng et al. [19] investigated the heterogeneous nucleation and growth of precipitates at dislocations in Al–Cu–Mg alloy by examining samples aged at 195 1C for various times from 10 min to 9 h, and found that the precipitation sequence of S (Al2CuMg) phase along dislocations is SSS-GPB zones-S (Type I) -S (Type I) þS (Type II). The objective of this study is to investigate the precipitation processes in Al–Cu–Mg alloy during creep aging. The uniaxial tensile creep tests of Al–Cu–Mg alloy were carried out over a wide range of temperature and applied stress. Also, the effects of the
Y.C. Lin et al. / Materials Science & Engineering A 556 (2012) 796–800
applied stress and creep temperature on the precipitation processes of Al–Cu–Mg alloy were discussed.
2. Material and experiments 2124-T851 aluminum alloy, which is one typical commercial Al–Cu–Mg alloy, was used in this investigation. This material was made by hot rolling, pre-stretching (stretch leveling ratio 1.8%) and heat treatment [20]. The compositions (wt%) of this material is 4.67Cu–1.46Mg–0.63Mn–0.18Fe–0.04Zn–0.01Ti–0.12Si–(bal.) Al. The specimens have been machined out from the rolled Al–Cu–Mg alloy plate with gauge length of 40 mm and diameter of 10 mm. The uniaxial tensile creep experiments were carried out in MTS-GWT2105 creep test machine. The controlling accuracies of the applied stress and temperature are 70.1 MPa and 72 K, respectively. The grating ruler is used to measure the creep displacement and it can be recorded by computer automatically, and the accuracy of grating ruler is 0.1 mm. Three different temperatures (473 K, 503 K and 533 K) and three different applied stresses (120 MPa, 140 MPa and 160 MPa) were used in the creep aging experiments. Firstly, the specimens were heated to the test temperature at a heating rate of 10 K/min, and hold for 0.5 h to eliminate thermal gradients before loading. After test, the fractured samples were naturally cooled to the room temperature in the furnace. In order to investigate the precipitation processes of Al–Cu–Mg alloy during creep aging, transmission electron microscopy (TEM) was carried out. The thin foil samples for TEM experiments were prepared in the following steps. Firstly, the block specimens were cut into 0.5–1.0 mm thick foil. Then, several disks with 3 mm in diameter and 0.25 mm in thickness were punched out from these slices, and subsequently electropolished using a solution of HNO3 and methanol (1:3 in volume). TEM foils were examined using a TecnaiG2-20 microscope operating at 200 KV.
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Fig. 1 shows the isothermal section of ternary Al–Cu–Mg phase diagram at 473 K. Some investigations [27–29] indicate that (1) when the ratio of Cu/Mg is between 8 and 4, the main precipitate are y and S phases at the same time. (2) When the ratio of Cu/Mg ranges from 4 to 1.5, the main precipitate is S phase. 2124-T851 aluminum alloy used in this study is one typical commercial Al–Cu–Mg alloy. For 2124-T851 aluminum alloy, the contents of Cu and Mg are 4.67% and 1.46% (wt%), respectively. So, the ratio of Cu/Mg is about 3.198, and then it can be concluded that the main precipitate is S phase during creep aging. Fig. 2 shows the TEM micrographs of as-received and creep aged samples. Here, the creep aging conditions are the temperature of 533 K and the applied stress of 140 MPa. From Fig. 2, it can be found that the acicular exudation phase is the main precipitate within the grains for both samples. For the samples without creep aging, as shown in Fig. 2(a), the acicular exudation phases homogeneously distribute within the grains, and the corresponding direction parallel to [1 0 0]Al or[0 1 1]Al. For the samples subjected to the creep aging, as shown in Fig. 2(b), the acicular exudation phase grows up and the quantity of acicular exudation phase decreases. Still, the distribution of precipitate phase is uniform within the grain. Fig. 3 shows the TEM micrographs and corresponding diffraction patterns for the sample without creep aging. The [1 1 2]Al selected area diffraction (SAD) pattern (shown in Fig. 3b) reveals light spots, and the light spots existing at 1/3o2004 and 2/3o2004 directions corresponding to S phase. So, it can be found that the main precipitate is S phase in 2124-T851 aluminum alloy. This result is consistent with others’ report [30]. Fig. 4 shows the TEM micrographs and corresponding diffraction patterns for the sample subjected to the creep aging under the temperature of 473 K and applied stress of 160 MPa for 100 h. Obviously, the [110]Al selected area diffraction (SAD) pattern (shown in Fig. 4b) reveals cross reflections [31], which match all spots of S phase.
3. Results and discussion 3.1. Precipitation hardening processes in Al–Cu–Mg alloys Precipitation strengthening is the main strengthening mechanism for ultra-high strength alloys, and the types of precipitation phase depend on the contents of alloy elements. For Al–Cu–Mg alloy, the types of precipitates are relative with the ratio of Cu/Mg. When the content of Mg element is low, the main precipitate is y phase (CuAl2), and the quantities of y phase decrease with the increase of the content of Mg element. When the Mg element is increased to some extent, the main precipitate will be S phase (CuMgAl2), which has the excellent high temperature strength. However, if the content of Mg element is continuously increased, T phase (A16CuMg4) and b phase (Al3Mg2) appear, which will deteriorate the mechanical properties of Al–Cu–Mg alloy. So, the content of Mg element is too high to strengthen the alloys [21]. Some investigators [22–25] proposed the following precipitation sequence of y and S. 00
Fig.1. Isothermal section of ternary Al–Cu–Mg phase diagram at 473 K (a ¼Al, y ¼ CuAl2, S ¼Al2CuMg, T¼ Al6CuMg4 and b ¼ Al12CuMg17).
0
SSS-GP-y -y -y SSS-GPB-S00 -S0 -S where SSS represents the supersaturated solid solution. The term GPB (Guinier–Preston–Bagaryatsky) zones was firstly reported by Silcock who suggested GPB zones (Cu and Mg ‘‘atomic clusters’’) are different from GP (Guinier–Preston) zone in Al–Cu alloys discovered earlier [26]. This process is ‘‘competitive balance between precipitations’’.
Fig. 2. TEM micrographs of (a) as-received condition; and (b) 533 K 140 MPa.
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d2 ¼0.0188, d3 ¼0.0149 [33]. Obviously, the misfit of acicular
Fig. 3. TEM micrographs and corresponding diffraction patterns for the as-received sample.
exudation phase is far less than the misfit of disk-shaped phase. According to Eshelby’s theory [34–35], firstly, the larger the misfit of precipitation phase is, the more the nucleation-growth stage of precipitation phase is sensitive to the applied stress. Secondly, the shape of precipitation phase depends on the elastic strain energy and interface energy, and the elastic strain energy is the main resistance for the growth of precipitation phase. Generally, the elastic strain energy of acicular and sphere phases is lower than that of disk-shaped phase. So, the elastic strain energy of diskshaped phase will strongly influence the behavior of precipitation phase. i.e., comparing to S0 phase, y0 phase is more sensitive to the applied stress field. So, the applied stress will break the precipitation equilibrium and alert the precipitation processes. For the studied Al–Cu–Mg alloy, the large applied stress and high aging temperature easily make the preferential precipitation process as 00 0 SSS-GPB-S -S -S, in order to decrease the active energy. This precipitation sequence will take in more solute atoms in the metallic matrix. Meanwhile, this process will decrease the content of y0 phase. In other words, the applied stress field will hinder the precipitation of y phase and promote the growth of S phase for the studied Al–Cu–Mg alloy.
3.2. Effect of creep test parameters on precipitation hardening
Fig. 4. TEM micrographs and corresponding diffraction patterns for the sample aged at 473 K 160 MPa.
From the above experimental results and analysis, it can be found that the preferential growth orientation of precipitation phase is sensitive to the applied stress for the studied 2124-T851 aluminum alloy. The acicular exudation precipitation phases uniformly distribute in the metallic matrix and the precipitation phases grow up in a preferential orientation. This is because that the stress/strain field induced by the creep aging conditions changes the system’s elastic energy, which makes the precipitation phases grow up in a preferential orientation, namely, stressorientation effect. The structures of GP or GPB zone is extremely asymmetric in the space, which results in the different misfits between the strengthening phase and metallic matrix in different orientations. Therefore, there is great difference of strain energy in different directions, and then the stress oriented effect is obvious at the nucleation stage of precipitation phase [32]. When alloys are affected by the applied stress, precipitation phases will grow up in a preferential orientation to balance the difference of strain energy and decrease the system’s energy. Additionally, the effects of stress on the penetration rates in various orientations are strongly linked to the anisotropic microstructure. Generally, the status (coherent or incoherent) of the strengthening phases is affected by the changes of stress and strain distortion fields, which are induced by the misfits between aluminum matrix and precipitation particles during the precipitation processes. Different precipitation phases are of different shapes, structures and physical properties in Al–Cu–Mg alloy. On the one hand, GP zone, y00 phase and y0 phase is formed by diskshaped Cu atoms segregated at {001} surface of aluminum matrix, and these precipitation phases is of tetragonal misfit strain field and low face misfit. These precipitation phases’ marginal misfit is d3,[GP] ¼ 0.10, d3y00 ¼ 0.052, d3y0 ¼ 0.045. On the other hand, the S0 phase is acicular, and GPB zone is solute cluster which grow along with the direction of o001 4 in aluminum matrix. So, S0 phase is of lower lattice misfit, and the misfit is d1 ¼ 0.0123,
3.2.1. Effect of applied stress Fig. 5 shows the TEM micrographs for the sample at the creep aging temperature of 473 K and applied stress of 120, 140, and 160 MPa. It can be found that, for the sample without creep aging (Fig. 2(a)), the length of acicular exudation phase is about 140 nm. With the increase of applied stress, the acicular exudation phase grows up and the quantity of acicular exudation phase decreases. For example, the length of S phase is near 200 nm after creep aging with the applied stress of 140 MPa. This is because the applied stress changes the competition of precipitation processes
Fig. 5. TEM micrographs for the sample aged at the creep aging temperature of 473 K and applied stress of (a) 120 MPa; (b) 140 MPa; and (c) 160 MPa.
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of the studied Al–Cu–Mg alloy, and the precipitation phases S is easy to grow up and coarsen under high stress. 3.2.2. Effect of temperature In this section, the effects of creep aging temperature on the precipitation process of Al–Cu–Mg alloys will be discussed. Fig. 2(b) and Fig. 5(b) show the TEM micrographs under the applied stress of 140 MPa and creep aging temperature of 533 K and 473 K, respectively. It can be found that with the increase of creep aging temperature, the scale of S phase quickly increased. Because the growth of S phase absorbs the solute atoms around it, the density of S phase decreases. The lengths of acicular exudation phase for the samples at the creep aging temperature of 533 K and 473 K are near and over 200 nm, respectively. However, for the samples without creep aging, the length of S phase is only about 140 nm. 3.2.3. Effect of creep exposure on precipitate free zones Grain boundaries act as sinks for solute and vacancies. This has the effect of increasing the precipitation on the grain boundary relative to the surrounding region, often leaving a precipitate free zone (PFZ) close to the boundary. This phenomenon is detrimental to the material since there is a reduction in precipitation hardening in the PFZ. Fig. 6 shows the precipitation phases on the grain boundaries of as-received and creep aged samples. Here, the creep aging conditions is the temperature of 473 K and applied stress of 120, 140, and 160 MPa. From Fig. 6, it can be found that the precipitates almost continuously distribute on the grain boundaries for the as-received samples, and the mechanical properties will be greatly deteriorated. However, for those creep aged ones, the grain boundary precipitates merge and concentrate, which result in a discontinuous distribution of precipitates on the grain boundaries. This is because the growth of precipitate needs to take in a lot of supersaturated solute atoms around the grain boundary. In other
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words, it suppresses the growth of precipitates around the grain boundaries. Therefore, the precipitates around grain boundaries are smaller than those in metallic matrix. Moreover, the discontinuous distribution of grain boundary precipitates can improve the material’s corrosion-resistance.
4. Conclusions The effects of the applied stress and creep aging temperature on the precipitation processes of 2124-T851 aluminum alloy were investigated by uniaxial tensile creep experiments. It is found that the precipitation process is very sensitive to the applied stress and creep aging temperature, and the dominant precipita00 0 tion process is SSS-GPB-S -S -S. So, the main precipitate is S phases under the tested creep aging conditions. With the increase of the applied stress and creep aging temperature, the S phase easily grows up and coarsens, and then the density of it decreases. Meanwhile, the creep aging will result in the discontinuous distribution of precipitation phase in the grain boundaries, which is useful to improve the corrosion-resistance of Al–Cu–Mg alloy.
Acknowledgments This work was supported by the National Key Basic Research Program (Grant no. 2010CB731702), the Program for New Century Excellent Talents in University (No. NCET-10-0838), the Sheng-hua Yu-ying Program and the Graduate Degree Thesis Innovation Foundation of Central South University (no. 2011ssxt094), State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology (no. 2012P04), China. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]
Fig. 6. Precipitation phases on grain boundary for the sample aged at different creep aging conditions. (a) as-received condition; (b) 473 K–120 MPa; (c) 473 K– 140 MPa; and (d) 473 K 160 MPa.
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