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Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy
Accepted Manuscript Title: Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy Author: Hongying Li, Desheng Huang, Wei Kang, Jiaojiao Liu, Yangxun Ou, Dewang Li PII: DOI: Reference:
S1005-0302(16)00023-2 http://dx.doi.org/doi: 10.1016/j.jmst.2016.01.018 JMST 649
To appear in:
Journal of Materials Science & Technology
Received date: Revised date: Accepted date:
7-5-2015 23-7-2015 24-8-2015
Please cite this article as: Hongying Li, Desheng Huang, Wei Kang, Jiaojiao Liu, Yangxun Ou, Dewang Li, Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy, Journal of Materials Science & Technology (2016), http://dx.doi.org/doi: 10.1016/j.jmst.2016.01.018. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy Hongying Li 1,2,, Desheng Huang 1,Wei Kang 1,Jiaojiao Liu 1, Yangxun Ou 1, Dewang Li 1 1
School of Materials Science and Engineering, Central South University, Changsha 410083,
China 2
Key Laboratory of Non-ferrous Materials, Ministry of Education, Central South University,
Changsha 410083, China [Received 7 May 2015; Received in revised form 23 July 2015; Accepted 24 August 2015]
Corresponding author. Prof., Ph.D.; Tel.: +86 731 88836328.
E-mail address: [email protected] (Hongying Li). The effects of different aging processes on the microstructure and mechanical properties of a novel Al–Cu–Li alloy have been investigated by X-ray diffraction, scanning electron microscopy and transmission electron microscopy. It is found that tensile properties of 2198 alloy are sensitive to aging processes, which corresponds to different microstructures. σ (Al5Cu6Mg2) and T1 (Al2CuLi) phases are the major precipitates for the alloy in T6 aging condition (165 °C/60 h). After duplex aging condition (150 °C/24 h + 180 °C/12 h), σ, θ´ (Al2Cu) and T1 phases are detected. Only the T1 phases can be found in the T8 state alloy (6% pre-strain+135 °C/60 h). The failure modes of alloy in T6 and duplex aging conditions are dimple-intergranular fracture, while typical quasi-cleavage fracture in T8 condition. Key words: Al–Cu–Li alloy; Aging process; Microstructure; Mechanical properties 1. Introduction Al–Cu–Li alloy has extensive application in the field of aerospace, especially in rocket fuel tank and large passenger aircraft due to its low density, good comprehensive mechanical properties, good corrosion resistance, low fatigue crack propagation rate and other characteristics. Previous investigations using methods, like differential scanning calorimetry (DSC), transmission electron microscopy (TEM), X-ray diffraction (XRD), etc. have thoroughly researched the
1 Page 1 of 12
competition of precipitation kinetics between the precipitates, aged precipitation behavior, ageing strengthening mechanism and precipitated phase structure of Al–Cu–Li alloys (such as 2195, 2197)[1,2]. Recently, Jia et al.[3] and Lin et al.[4] studied the microstructural evolution of the Al–Cu–Li alloys during homogenization and solution treatment. Zhang et al.[5] studied the effect of naturally-aged and artificial aging conditions on the microstructure and property of 2198 alloy. However, there is no literature to make a comparison between different aging processes on the microstructure and properties of the Al–Cu–Li alloy. The strengthening phase includes δ´(Al3Li), T1(Al2CuLi) and θ´(Al2Cu), which have been founded in Al–Cu–Li alloys[6–12]. In addition to these phases mentioned above, a cubic σ (Al5Cu6Mg2) phase has been observed in Al–Cu–Mg system where minimum concentrations of Si atoms act as cores for nucleation[13,14]. Alloy performance as a result of the combined action of precipitated phase characteristics, and features of the precipitated phase are closely related to the aging process. It is necessary to understand the effect of precipitated phases on the mechanical properties. The objective of the present study is to find the relationship between the kind of precipitated phases and mechanical properties under the peak aging of various heat treatment conditions. 2. Experimental High purity ingot of Al, pure Li and master alloys of Al–Cu, Al–Mg, Al–Zn, Al–Mn and Al–Zr were melted and cast in a crucible furnace. Chemical composition of the ingot was measured to be Al–3.7wt%Cu–1.5wt%Li–0.50wt%Zn–0.37wt%Mg–0.30wt%Mn–0.14wt%Zr. The ingot was homogenized at 520 °C for 24 h before rolling, then hot rolling and cold rolling process were used to produce a 2 mm-thick sheet. The
at 520 °C for 2 h
2 Page 2 of 12
and then quenched into water prior to artificial aging treatment. The different aging treatment procedures of the studied alloy are shown in Table 1. The tensile samples were cut along through-thickness direction of the cold rolled sheet. The size of tensile samples is shown in Fig. 1. Tensile test was performed at a MTS 810 machine and the strain was measured using an extensometer. Tensile test was carried out at room temperature. A strain rate of 1.0 × 10−3 s−1 was applied. The tensile value was derived from the average of five tensile testing samples. The fracture morphology was observed by the Sirion200 type scanning electron microscopy (SEM). The XRD measurements were carried out in Rigaku D/Max2500 X-ray diffractometer using CuKα radiation. The TEM samples were cut from the testing samples along the rolling sheet, then mechanically thinned to 0.1 mm, and punched into disks of 3 mm in diameter. The disks were finally thinned by jet electro polishing in a 75% methanol and 25% nitric acid solution cooled down to approximately 243 K by liquid nitrogen. TEM observation was carried out using a TecnaiG220 instrument operating at 160 kV. 3. Results Fig. 2 illustrates the mechanical properties of a novel Al–Cu–Li alloy samples with various heat treatments. As can be seen, the yield strength (YS), ultimate tensile strength (UTS) and elongation (El) of the studied alloy are 453 MPa, 493 MPa and 12.5% in the T6 condition, respectively. Compared with the data in T6 state, the yield and ultimate tensile strength in duplex aging get slightly decreased, but the elongation rate gets appreciably increased. The yield strength and ultimate tensile strength of T8 samples increase significantly, indicating a strong aging response to pre-deformation of this alloy. The yield strength and tensile strength is up to 539 MPa and 567 MPa, respectively. However, the elongation of T8 aged samples decreases to 10.8%. 3 Page 3 of 12
Fig. 3 presents the typical tensile fracture surface of specimens in the T6, T8 and duplex aging conditions. A comparison between the morphology of these three individual states reveals distinct differences in the underlying mechanisms in each state. There are many small dimples covering the fracture surface, and the dimples are shallow, so the fracture pattern of T6-aged studied alloy is the intergranular fracture(Fig. 3(a) and (b)). Considering the T8 state (Fig. 3(c) and (d)), the dimples are much less observed. In addition, the fracture surface has an obvious delamination features. Smooth surfaces and tear ridge emerge on the fracture surface. This is the typical quasi-cleavage fracture. After the duplex aging treatment of the alloy (Fig. 3(e) and (f)), it is also a dimple-fracture morphology, but the dimples are much deep, which coresponds to coarser precipitates at higher aging temperature. This feature shows the typical ductile fracture resulting from void growth and coalescence. Fig. 4 shows the TEM images and corresponding selected area electron diffraction (SAED) patterns of the studied alloy in solution state and different aging conditions, viewed along <110> and <001> directions. It can be observed that the different heat treatments will precipitate different phases. From the SAED pattern, the diffraction spots 1/2 {220}Al and the streaks through the {000}Al spot indicate the δ´ phases along <200> and <
> direction. Meanwhile, it can be seen
that the spots and streaks are also very weak. From the bright field (BF) image, it can be observed that a few δ´ phases were found in the alloy matrix after solution treatment. The main phases in T6-aged studied alloy are a large number of dispersed cubic σ phases and a small quantity of plate-shaped T1 phases (Fig. 4(b)). Two types of intersections are generally observed for T1 plates: one plate completely penetrates through the other plate (intersection points A and B) or one plate impinges upon the other plate (intersection points C and D). Fig. 4(c) shows that the T1 phase uniformly disperses in the Al matrix, and the size of T1 phase is small and thin. As can be seen, the σ phase, θ´ phase and T1 phase exist in the Al matrix (Fig. 4(d)). The θ´ phase, σ phase and T1 phase hold the orientations of
,
, and
,
respectively[15–17]. So the orientation of θ´ phase is parallel to σ phase, but the orientation of T1 4 Page 4 of 12
phase and σ phase has a certain angle, which can be verified in the Fig. 4(c). To make better identification of the phase, the XRD patterns of the sample are shown in Fig. 5. It shows that the main kinds of the intermetallic phases were σ, T1 and θ´ phases, which was in accordance with designations reported in literature[5]. The σ and T1 phases are the main phases in the specimens under T6 state. In the specimens under T8 condition, only the T1 phase was found. There exists the combinative precipitation of σ, T1 and θ´ phases in the duplex aging sample. The existence of θ´ phase does not obviously improve the strength of the alloy, which suggests that T1 and σ are the main hardening phases. 4. Discussion It can be seen from the Fig. 4 that the emergence of the σ phase often coexists with that of T1 and θ´ phases. However, the presence of θ´ and T1 phases consumes the amount of Cu atoms, so the number of σ phases reduces or even disappears, accompanied by a fine dispersion of T1 or θ´ phases precipitate, indicating a competitive precipitating relationship between σ and T1 phase, or σ and θ´ phase. Schueller et al.[18] reported that the σ phase, once nucleated, is a relatively stable phase which resists dissolution to high temperature, so the σ phase will not change to the θ´ and T1 phases during the aging process. In contrast to the duplex aging alloy, in the T6 alloy σ phase in the process of heat treatment is much more, but there is no significant increase in the strength, indicating that the σ phase is not the main strengthing phase. The misfit of lattice parameter between the σ phase and the aluminum matrix was calculated as 2.8% (assuming aAl=0.404 nm). Therefore, according to Van der Merwe’s investigation, the σ phase should be semicoherent[19]. Fig. 6 is the morphology of σ phase. It can be seen that the σ phase exists in the Moire fringe. The major precipitated phases in the peak aging state of three type aging processes are the σ,
5 Page 5 of 12
T1 and θ´ phases. From the Fig. 4(a), a few δ´ phases exist in the alloy matrix after solution treatment, but it is hard to find the δ´ phase under the different peak aging conditions, it is possible that the growing up of T1 phases consume many Li atoms during the aging process, so that the δ´ phase decreased. There are two ways to dissolve the δ´ phase in the process of growing the T1 phase[20]: the T1 phase is closely tangent to the δ´ phase, and it makes the Li atoms of the δ´ phase diffuse to the boundary layer of T1 phase. The boundary layer of T1 phase grows along the coherent surface of T1/α and crosses the δ´ phase, so that the δ´ phase is absorbed into the T1 phase. The pre-deformation before aging, as is well known, has an evident effect on the precipitation process, which can be applied to improve the age hardening[21]. The presence of dislocations often accelerate precipitation and coarse the precipitates[22,23]. Cassada et al.[24] found that matrix dislocations play a significant role in the nucleation and growth of T1 plate. A detailed study[25] believed that the degree of pre-deformation has an effect on the length and thickness of T1 plate. The yield strength will increase 100 MPa if the number density increased two orders of magnitude. At the same time, the time to reach the peak strength is 20% shorter than that the aging process without plastic deformation. The number of strengthening phase (like T1 phase) is only a few under the T6 and duplex aging conditions. Numerous tiny dimples distribute in the fracture surface, and the dimples seem unusually shallow. There also been found a few scattered secondary cracks in the fracture surface. The mechanism of ductile fracture in metals is always explained by the classical void growth model[26], which divides the fracture process into three stages: void initiation, growth and coalescence. As is shown in Fig. 3, coalescence voids can be clearly seen at the precipitates, which 6 Page 6 of 12
mean that the precipitates separate from matrix in the process of deformation and the micro-voids generate there. So the fracture mechanisms of T6 and duplex aging alloy are dimple-intergranular fracture. Many researchers concluded that the grain-boundary precipitates and the precipitate free zone (PFZ) adjacent to grain boundaries can attribute to intergranular fracture[27–30]. Strain localization in soft PFZs results in nucleation and growth of voids around grain-boundary precipitates so that intergranular fracture surfaces are produced. The size and depth of dimples decrease with decreasing PFZ width, as would be expected. In the present study, the grain-boundary precipitates and PFZ adjacent to the grain boundaries have been found in the studied alloy under the T6 and duplex aging conditions, as can be seen in Fig. 7(a) and (b). Thus, the dimples would be expected on intergranular fracture surfaces, and are indeed observed on same facets. Large amount of dislocation would be produced during predeformation before aging, which can promote the decomposition of supersaturated solid solution, and avoid the occurrence of PFZ, as can be seen in Fig. 7(c). The T8 state alloy has the obvious layering feature. The layering feature has the crack divider effect, hindering the dislocation glide, which has restricted to the plastic zone at the crack tip. The stratification characteristics decrease the three dimensional tensile stress at the crack tip, and the plane strain state transforms into the plane stress state. These cause the improvement of macro-fracture toughness and therefore the elongation of the T8-aged alloy is slightly lower than that of the T6 and duplex aging alloy. 5. Conclusions （1） The novel Al–Cu–Li alloy exhibits a strong aging response in different aging processes. The ultimate tensile strength of studied alloy can reach 567 MPa under the T8 condition, which is much higher than that in T6 and duplex aging processes. 7 Page 7 of 12
（2） The failure modes of T6 and duplex aging condition are dimple-intergranular fracture, while typical quasi-cleavage fracture in T8 condition. （3） Different heat treatments will cause different precipitated phases. T1, θ´ and σ phases are the three major strengthening phases in the novel Al–Cu–Li alloy. In the T6 state alloy, σ and T1 phases are the main precipitating phases. The alloy precipitates many T1 phases under T8 condition. By duplex aging process, the alloy precipitates σ, T1 and θ´ phases. Acknowledgement Authors thank the National High Technology Research and Development Program of China (Grant No. 2013AA032401) for the financial support. References [1] T. Honma, S. Yanagita, K. Hono,
52 (2004) 1997–2003.
[2] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 299–306.
[3] M. Jia, Z.Q. Zheng, Z. Gong, J. Alloy. Compd. 614 (2014) 131–139. [4] Y. Lin, Z.Q. Zheng, S.C. Li, Arch. Civ. Mech. Eng. 14 (2014) 61–71. [5] S.F. Zhang, W.D. Zeng, W.H. Yang, C.L. Shi, H.J. Wang, Mater. Des. 63 (2014) 368–374. [6] B.P. Huang, Z.Q. Zheng,
46 (1998) 4381–4393.
[7] S.P. Ringer, B.C. Muddle, I.J. Polmear,
42 (1994) 1715–1725.
[8] B.M. Gable, A.W. Zhu, A.A. Csontos, E.A. Starke, J. Light Metals 1 (2001) 1–14. [9] J.I. Pérez-Landazábal, M.L. Nó, G. Madariaga, V. Recarte, J.S. Juan,
48 (2000)
1283–1296. [10] C.S. Tsao, C.Y. Chen, T.Y. Kuo, T.L. Lin, M.S. Yu, [11] J.C. Huang, A.J. Ardell, [12] F.W. Gayle, B. Vandersande,
363 (2003) 228–233.
36 (1988) 2995–3066. 37 (1989) 1033–1046. 8 Page 8 of 12
[13] K.S. Kumar, S.A. Brown, J.R. Pickens,
44 (1996) 1899–1951.
[14] A.K. Mukhopadhyay, A. Kumar, S. Raveendra, I. Samajdar,
64 (2011)
386–389. [15] B.C. Muddle, I.J. Polmear, Acta Mater. 37 (1989) 777–789. [16] J.C. Huang, A. Ardell, J.
3 (1987) 176–188.
[17] R.D. Schueller, A.K. Sachdev, F.E. Wawner,
27 (1992) 617–622.
[18] R.D. Schueller, F.E. Wawner,A.K. Sachdev, J. Mater. Sci. 29 (1994) 239–249. [19] F.C. Frank, J.H. Van der Merwe, Proc. Roy. Soc. London A 198 (1949) 205. [20] J.C. Huang, A.J. Ardell,
36 (1988) 2995–3006.
[21] S.P. Ringer, B.C. Muddle, I.J. Polmear, [22] A. Deschamps, F. Livet, Y. Brechet, [23] Y. Birol,
. 26 (1995) 1659–1671. 47 (1998) 281–292.
52 ( 2005) 169–173.
[24] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 287–297.
[25] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 299–306.
[26] F. Thomason, Ductile Fracture of Metals, Ergamon Press, Oford, 1990. 2 (1986) 1205–1209.
[27] A.K. Vasudevan, E.A. Ludwiczak, S.F. Baumann, [28] S. Suresh, A.K. Vasudevan, M. Tosten, [29] T.H. Sanders Jr., E.A. Ludwiczak, R.R. Sawtell, [30] S.C. Jha, T.H. Sanders Jr., M.A. Dayananda,
35 (1987) 25–46. 43 (1980) 247–260. 35 (1987) 473–482.
Figure Captions Fig. 1. Dimensions of the tensile sample. Fig. 2. Tensile properties of a novel Al–Cu–Li alloy sample with various heat treatments. 9 Page 9 of 12
Fig. 3. Fractographs of the tensile specimens under conditions of T6 (a, b), T8 (c, d), and duplex aging (e, f). Fig. 4. TEM images of specimens under different aging conditions: (a) BF image of solid solution state taken along [001]Al; (b) BF image of T6 taken along [110]Al; (c) BF image of T8 taken along [110]Al; (d) BF image of duplex aging taken along [110]Al. Fig. 5. XRD patterns of a novel Al–Cu–Li alloy at various aged processes. Fig. 6. Morphology of σ phase: (a) Moire fringe; (b) HRTEM. Fig. 7. Grain-boundary precipitates and PFZ observed in studied alloy: (a) T6 aging; (b) duplex aging; (c) T8 aging. Table Captions Table 1 Aging treatment procedures used for a novel Al–Li alloy in this study Number
Condition
Heat treatment procedure
1 2 3
T6 T8 Duplex aging
165 °C/60 h 6% pre-strain + 135 °C/60 h 150 °C /24 h + 180 °C /12 h
Figures
Fig. 1.
Fig. 2.
10 Page 10 of 12
Fig. 3.
Fig. 4. 11 Page 11 of 12
Fig. 5.
Fig. 6.
Fig. 7.
12 Page 12 of 12
S1005-0302(16)00023-2 http://dx.doi.org/doi: 10.1016/j.jmst.2016.01.018 JMST 649
To appear in:
Journal of Materials Science & Technology
Received date: Revised date: Accepted date:
7-5-2015 23-7-2015 24-8-2015
Please cite this article as: Hongying Li, Desheng Huang, Wei Kang, Jiaojiao Liu, Yangxun Ou, Dewang Li, Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy, Journal of Materials Science & Technology (2016), http://dx.doi.org/doi: 10.1016/j.jmst.2016.01.018. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effect of Different Aging Processes on the Microstructure and Mechanical Properties of a Novel Al–Cu–Li Alloy Hongying Li 1,2,, Desheng Huang 1,Wei Kang 1,Jiaojiao Liu 1, Yangxun Ou 1, Dewang Li 1 1
School of Materials Science and Engineering, Central South University, Changsha 410083,
China 2
Key Laboratory of Non-ferrous Materials, Ministry of Education, Central South University,
Changsha 410083, China [Received 7 May 2015; Received in revised form 23 July 2015; Accepted 24 August 2015]
Corresponding author. Prof., Ph.D.; Tel.: +86 731 88836328.
E-mail address: [email protected] (Hongying Li). The effects of different aging processes on the microstructure and mechanical properties of a novel Al–Cu–Li alloy have been investigated by X-ray diffraction, scanning electron microscopy and transmission electron microscopy. It is found that tensile properties of 2198 alloy are sensitive to aging processes, which corresponds to different microstructures. σ (Al5Cu6Mg2) and T1 (Al2CuLi) phases are the major precipitates for the alloy in T6 aging condition (165 °C/60 h). After duplex aging condition (150 °C/24 h + 180 °C/12 h), σ, θ´ (Al2Cu) and T1 phases are detected. Only the T1 phases can be found in the T8 state alloy (6% pre-strain+135 °C/60 h). The failure modes of alloy in T6 and duplex aging conditions are dimple-intergranular fracture, while typical quasi-cleavage fracture in T8 condition. Key words: Al–Cu–Li alloy; Aging process; Microstructure; Mechanical properties 1. Introduction Al–Cu–Li alloy has extensive application in the field of aerospace, especially in rocket fuel tank and large passenger aircraft due to its low density, good comprehensive mechanical properties, good corrosion resistance, low fatigue crack propagation rate and other characteristics. Previous investigations using methods, like differential scanning calorimetry (DSC), transmission electron microscopy (TEM), X-ray diffraction (XRD), etc. have thoroughly researched the
1 Page 1 of 12
competition of precipitation kinetics between the precipitates, aged precipitation behavior, ageing strengthening mechanism and precipitated phase structure of Al–Cu–Li alloys (such as 2195, 2197)[1,2]. Recently, Jia et al.[3] and Lin et al.[4] studied the microstructural evolution of the Al–Cu–Li alloys during homogenization and solution treatment. Zhang et al.[5] studied the effect of naturally-aged and artificial aging conditions on the microstructure and property of 2198 alloy. However, there is no literature to make a comparison between different aging processes on the microstructure and properties of the Al–Cu–Li alloy. The strengthening phase includes δ´(Al3Li), T1(Al2CuLi) and θ´(Al2Cu), which have been founded in Al–Cu–Li alloys[6–12]. In addition to these phases mentioned above, a cubic σ (Al5Cu6Mg2) phase has been observed in Al–Cu–Mg system where minimum concentrations of Si atoms act as cores for nucleation[13,14]. Alloy performance as a result of the combined action of precipitated phase characteristics, and features of the precipitated phase are closely related to the aging process. It is necessary to understand the effect of precipitated phases on the mechanical properties. The objective of the present study is to find the relationship between the kind of precipitated phases and mechanical properties under the peak aging of various heat treatment conditions. 2. Experimental High purity ingot of Al, pure Li and master alloys of Al–Cu, Al–Mg, Al–Zn, Al–Mn and Al–Zr were melted and cast in a crucible furnace. Chemical composition of the ingot was measured to be Al–3.7wt%Cu–1.5wt%Li–0.50wt%Zn–0.37wt%Mg–0.30wt%Mn–0.14wt%Zr. The ingot was homogenized at 520 °C for 24 h before rolling, then hot rolling and cold rolling process were used to produce a 2 mm-thick sheet. The
at 520 °C for 2 h
2 Page 2 of 12
and then quenched into water prior to artificial aging treatment. The different aging treatment procedures of the studied alloy are shown in Table 1. The tensile samples were cut along through-thickness direction of the cold rolled sheet. The size of tensile samples is shown in Fig. 1. Tensile test was performed at a MTS 810 machine and the strain was measured using an extensometer. Tensile test was carried out at room temperature. A strain rate of 1.0 × 10−3 s−1 was applied. The tensile value was derived from the average of five tensile testing samples. The fracture morphology was observed by the Sirion200 type scanning electron microscopy (SEM). The XRD measurements were carried out in Rigaku D/Max2500 X-ray diffractometer using CuKα radiation. The TEM samples were cut from the testing samples along the rolling sheet, then mechanically thinned to 0.1 mm, and punched into disks of 3 mm in diameter. The disks were finally thinned by jet electro polishing in a 75% methanol and 25% nitric acid solution cooled down to approximately 243 K by liquid nitrogen. TEM observation was carried out using a TecnaiG220 instrument operating at 160 kV. 3. Results Fig. 2 illustrates the mechanical properties of a novel Al–Cu–Li alloy samples with various heat treatments. As can be seen, the yield strength (YS), ultimate tensile strength (UTS) and elongation (El) of the studied alloy are 453 MPa, 493 MPa and 12.5% in the T6 condition, respectively. Compared with the data in T6 state, the yield and ultimate tensile strength in duplex aging get slightly decreased, but the elongation rate gets appreciably increased. The yield strength and ultimate tensile strength of T8 samples increase significantly, indicating a strong aging response to pre-deformation of this alloy. The yield strength and tensile strength is up to 539 MPa and 567 MPa, respectively. However, the elongation of T8 aged samples decreases to 10.8%. 3 Page 3 of 12
Fig. 3 presents the typical tensile fracture surface of specimens in the T6, T8 and duplex aging conditions. A comparison between the morphology of these three individual states reveals distinct differences in the underlying mechanisms in each state. There are many small dimples covering the fracture surface, and the dimples are shallow, so the fracture pattern of T6-aged studied alloy is the intergranular fracture(Fig. 3(a) and (b)). Considering the T8 state (Fig. 3(c) and (d)), the dimples are much less observed. In addition, the fracture surface has an obvious delamination features. Smooth surfaces and tear ridge emerge on the fracture surface. This is the typical quasi-cleavage fracture. After the duplex aging treatment of the alloy (Fig. 3(e) and (f)), it is also a dimple-fracture morphology, but the dimples are much deep, which coresponds to coarser precipitates at higher aging temperature. This feature shows the typical ductile fracture resulting from void growth and coalescence. Fig. 4 shows the TEM images and corresponding selected area electron diffraction (SAED) patterns of the studied alloy in solution state and different aging conditions, viewed along <110> and <001> directions. It can be observed that the different heat treatments will precipitate different phases. From the SAED pattern, the diffraction spots 1/2 {220}Al and the streaks through the {000}Al spot indicate the δ´ phases along <200> and <
> direction. Meanwhile, it can be seen
that the spots and streaks are also very weak. From the bright field (BF) image, it can be observed that a few δ´ phases were found in the alloy matrix after solution treatment. The main phases in T6-aged studied alloy are a large number of dispersed cubic σ phases and a small quantity of plate-shaped T1 phases (Fig. 4(b)). Two types of intersections are generally observed for T1 plates: one plate completely penetrates through the other plate (intersection points A and B) or one plate impinges upon the other plate (intersection points C and D). Fig. 4(c) shows that the T1 phase uniformly disperses in the Al matrix, and the size of T1 phase is small and thin. As can be seen, the σ phase, θ´ phase and T1 phase exist in the Al matrix (Fig. 4(d)). The θ´ phase, σ phase and T1 phase hold the orientations of
,
, and
,
respectively[15–17]. So the orientation of θ´ phase is parallel to σ phase, but the orientation of T1 4 Page 4 of 12
phase and σ phase has a certain angle, which can be verified in the Fig. 4(c). To make better identification of the phase, the XRD patterns of the sample are shown in Fig. 5. It shows that the main kinds of the intermetallic phases were σ, T1 and θ´ phases, which was in accordance with designations reported in literature[5]. The σ and T1 phases are the main phases in the specimens under T6 state. In the specimens under T8 condition, only the T1 phase was found. There exists the combinative precipitation of σ, T1 and θ´ phases in the duplex aging sample. The existence of θ´ phase does not obviously improve the strength of the alloy, which suggests that T1 and σ are the main hardening phases. 4. Discussion It can be seen from the Fig. 4 that the emergence of the σ phase often coexists with that of T1 and θ´ phases. However, the presence of θ´ and T1 phases consumes the amount of Cu atoms, so the number of σ phases reduces or even disappears, accompanied by a fine dispersion of T1 or θ´ phases precipitate, indicating a competitive precipitating relationship between σ and T1 phase, or σ and θ´ phase. Schueller et al.[18] reported that the σ phase, once nucleated, is a relatively stable phase which resists dissolution to high temperature, so the σ phase will not change to the θ´ and T1 phases during the aging process. In contrast to the duplex aging alloy, in the T6 alloy σ phase in the process of heat treatment is much more, but there is no significant increase in the strength, indicating that the σ phase is not the main strengthing phase. The misfit of lattice parameter between the σ phase and the aluminum matrix was calculated as 2.8% (assuming aAl=0.404 nm). Therefore, according to Van der Merwe’s investigation, the σ phase should be semicoherent[19]. Fig. 6 is the morphology of σ phase. It can be seen that the σ phase exists in the Moire fringe. The major precipitated phases in the peak aging state of three type aging processes are the σ,
5 Page 5 of 12
T1 and θ´ phases. From the Fig. 4(a), a few δ´ phases exist in the alloy matrix after solution treatment, but it is hard to find the δ´ phase under the different peak aging conditions, it is possible that the growing up of T1 phases consume many Li atoms during the aging process, so that the δ´ phase decreased. There are two ways to dissolve the δ´ phase in the process of growing the T1 phase[20]: the T1 phase is closely tangent to the δ´ phase, and it makes the Li atoms of the δ´ phase diffuse to the boundary layer of T1 phase. The boundary layer of T1 phase grows along the coherent surface of T1/α and crosses the δ´ phase, so that the δ´ phase is absorbed into the T1 phase. The pre-deformation before aging, as is well known, has an evident effect on the precipitation process, which can be applied to improve the age hardening[21]. The presence of dislocations often accelerate precipitation and coarse the precipitates[22,23]. Cassada et al.[24] found that matrix dislocations play a significant role in the nucleation and growth of T1 plate. A detailed study[25] believed that the degree of pre-deformation has an effect on the length and thickness of T1 plate. The yield strength will increase 100 MPa if the number density increased two orders of magnitude. At the same time, the time to reach the peak strength is 20% shorter than that the aging process without plastic deformation. The number of strengthening phase (like T1 phase) is only a few under the T6 and duplex aging conditions. Numerous tiny dimples distribute in the fracture surface, and the dimples seem unusually shallow. There also been found a few scattered secondary cracks in the fracture surface. The mechanism of ductile fracture in metals is always explained by the classical void growth model[26], which divides the fracture process into three stages: void initiation, growth and coalescence. As is shown in Fig. 3, coalescence voids can be clearly seen at the precipitates, which 6 Page 6 of 12
mean that the precipitates separate from matrix in the process of deformation and the micro-voids generate there. So the fracture mechanisms of T6 and duplex aging alloy are dimple-intergranular fracture. Many researchers concluded that the grain-boundary precipitates and the precipitate free zone (PFZ) adjacent to grain boundaries can attribute to intergranular fracture[27–30]. Strain localization in soft PFZs results in nucleation and growth of voids around grain-boundary precipitates so that intergranular fracture surfaces are produced. The size and depth of dimples decrease with decreasing PFZ width, as would be expected. In the present study, the grain-boundary precipitates and PFZ adjacent to the grain boundaries have been found in the studied alloy under the T6 and duplex aging conditions, as can be seen in Fig. 7(a) and (b). Thus, the dimples would be expected on intergranular fracture surfaces, and are indeed observed on same facets. Large amount of dislocation would be produced during predeformation before aging, which can promote the decomposition of supersaturated solid solution, and avoid the occurrence of PFZ, as can be seen in Fig. 7(c). The T8 state alloy has the obvious layering feature. The layering feature has the crack divider effect, hindering the dislocation glide, which has restricted to the plastic zone at the crack tip. The stratification characteristics decrease the three dimensional tensile stress at the crack tip, and the plane strain state transforms into the plane stress state. These cause the improvement of macro-fracture toughness and therefore the elongation of the T8-aged alloy is slightly lower than that of the T6 and duplex aging alloy. 5. Conclusions （1） The novel Al–Cu–Li alloy exhibits a strong aging response in different aging processes. The ultimate tensile strength of studied alloy can reach 567 MPa under the T8 condition, which is much higher than that in T6 and duplex aging processes. 7 Page 7 of 12
（2） The failure modes of T6 and duplex aging condition are dimple-intergranular fracture, while typical quasi-cleavage fracture in T8 condition. （3） Different heat treatments will cause different precipitated phases. T1, θ´ and σ phases are the three major strengthening phases in the novel Al–Cu–Li alloy. In the T6 state alloy, σ and T1 phases are the main precipitating phases. The alloy precipitates many T1 phases under T8 condition. By duplex aging process, the alloy precipitates σ, T1 and θ´ phases. Acknowledgement Authors thank the National High Technology Research and Development Program of China (Grant No. 2013AA032401) for the financial support. References [1] T. Honma, S. Yanagita, K. Hono,
52 (2004) 1997–2003.
[2] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 299–306.
[3] M. Jia, Z.Q. Zheng, Z. Gong, J. Alloy. Compd. 614 (2014) 131–139. [4] Y. Lin, Z.Q. Zheng, S.C. Li, Arch. Civ. Mech. Eng. 14 (2014) 61–71. [5] S.F. Zhang, W.D. Zeng, W.H. Yang, C.L. Shi, H.J. Wang, Mater. Des. 63 (2014) 368–374. [6] B.P. Huang, Z.Q. Zheng,
46 (1998) 4381–4393.
[7] S.P. Ringer, B.C. Muddle, I.J. Polmear,
42 (1994) 1715–1725.
[8] B.M. Gable, A.W. Zhu, A.A. Csontos, E.A. Starke, J. Light Metals 1 (2001) 1–14. [9] J.I. Pérez-Landazábal, M.L. Nó, G. Madariaga, V. Recarte, J.S. Juan,
48 (2000)
1283–1296. [10] C.S. Tsao, C.Y. Chen, T.Y. Kuo, T.L. Lin, M.S. Yu, [11] J.C. Huang, A.J. Ardell, [12] F.W. Gayle, B. Vandersande,
363 (2003) 228–233.
36 (1988) 2995–3066. 37 (1989) 1033–1046. 8 Page 8 of 12
[13] K.S. Kumar, S.A. Brown, J.R. Pickens,
44 (1996) 1899–1951.
[14] A.K. Mukhopadhyay, A. Kumar, S. Raveendra, I. Samajdar,
64 (2011)
386–389. [15] B.C. Muddle, I.J. Polmear, Acta Mater. 37 (1989) 777–789. [16] J.C. Huang, A. Ardell, J.
3 (1987) 176–188.
[17] R.D. Schueller, A.K. Sachdev, F.E. Wawner,
27 (1992) 617–622.
[18] R.D. Schueller, F.E. Wawner,A.K. Sachdev, J. Mater. Sci. 29 (1994) 239–249. [19] F.C. Frank, J.H. Van der Merwe, Proc. Roy. Soc. London A 198 (1949) 205. [20] J.C. Huang, A.J. Ardell,
36 (1988) 2995–3006.
[21] S.P. Ringer, B.C. Muddle, I.J. Polmear, [22] A. Deschamps, F. Livet, Y. Brechet, [23] Y. Birol,
. 26 (1995) 1659–1671. 47 (1998) 281–292.
52 ( 2005) 169–173.
[24] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 287–297.
[25] W.A. Cassada, G.J. Shiflet, E.A. Starke,
22 (1991) 299–306.
[26] F. Thomason, Ductile Fracture of Metals, Ergamon Press, Oford, 1990. 2 (1986) 1205–1209.
[27] A.K. Vasudevan, E.A. Ludwiczak, S.F. Baumann, [28] S. Suresh, A.K. Vasudevan, M. Tosten, [29] T.H. Sanders Jr., E.A. Ludwiczak, R.R. Sawtell, [30] S.C. Jha, T.H. Sanders Jr., M.A. Dayananda,
35 (1987) 25–46. 43 (1980) 247–260. 35 (1987) 473–482.
Figure Captions Fig. 1. Dimensions of the tensile sample. Fig. 2. Tensile properties of a novel Al–Cu–Li alloy sample with various heat treatments. 9 Page 9 of 12
Fig. 3. Fractographs of the tensile specimens under conditions of T6 (a, b), T8 (c, d), and duplex aging (e, f). Fig. 4. TEM images of specimens under different aging conditions: (a) BF image of solid solution state taken along [001]Al; (b) BF image of T6 taken along [110]Al; (c) BF image of T8 taken along [110]Al; (d) BF image of duplex aging taken along [110]Al. Fig. 5. XRD patterns of a novel Al–Cu–Li alloy at various aged processes. Fig. 6. Morphology of σ phase: (a) Moire fringe; (b) HRTEM. Fig. 7. Grain-boundary precipitates and PFZ observed in studied alloy: (a) T6 aging; (b) duplex aging; (c) T8 aging. Table Captions Table 1 Aging treatment procedures used for a novel Al–Li alloy in this study Number
Condition
Heat treatment procedure
1 2 3
T6 T8 Duplex aging
165 °C/60 h 6% pre-strain + 135 °C/60 h 150 °C /24 h + 180 °C /12 h
Figures
Fig. 1.
Fig. 2.
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Fig. 3.
Fig. 4. 11 Page 11 of 12
Fig. 5.
Fig. 6.
Fig. 7.
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