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Investigation of precipitation behavior and related hardening in AA 7055 aluminum alloy
Materials Science and Engineering A 500 (2009) 34–42
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Investigation of precipitation behavior and related hardening in AA 7055 aluminum alloy Junzhou Chen a,b , Liang Zhen a,∗ , Shoujie Yang b , Wenzhu Shao a , Shenglong Dai b a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Beijing Institute of Aeronautical Materials, Beijing 100095, PR China
a r t i c l e
i n f o
Article history: Received 8 May 2008 Received in revised form 21 August 2008 Accepted 24 September 2008 Keywords: Precipitation Aluminum alloy Transmission electron microscopy Ageing
a b s t r a c t The precipitation behavior and related hardening in AA 7055 aluminum alloy aged at 120 and 160 ◦ C was investigated in detail. GPI zones were the dominant phase in the alloy upon ageing at 120 ◦ C for 60 min. The metastable phase begins to precipitate in the alloy after being aged at 120 ◦ C for 60 min, and turns to be the main phase after ageing for 300 min. When the alloy was aged at 160 ◦ C, the precipitation was significantly promoted. The results also revealed that the transformation of small GPI zone to phase is the dominant mechanism for formation. Formation and growth of GPI zones and phases led to the increase of the yield strength, while formation and coarsening of resulted in the decrease of the strength. is responsible for the peak hardening of this alloy. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Precipitations from the supersaturated solid solutions have been extensively studied in aluminum alloys [1–10], since they play a very important role in determining the mechanical properties of these alloys. The general precipitation sequence of Al–Zn–Mg (–Cu) alloys can be summarized [6,11] as: Solid solution → GP zones → Metastable → Stable . GP zones are generally accepted as precipitates formed during room temperature ageing and the early stages of artificial ageing. They may lead to the increase of the strength of these alloys. There are two types of GP zones, i.e., GPI and GPII, with different structures [2]. GPI zones are fully coherent with the Al matrix, with internal ordering of Zn and Al or Mg on the {0 0 1}Al planes, based on an AuCu(I)-type sub-unit, and periodic anti-phase boundaries. They could be formed over a wide ageing temperature range, from room temperature to 140–150 ◦ C, independently of quenching temperature. GPII are zn-rich layers on {1 1 1}Al planes, one-to two atoms thick and 3–5 nm wide, with internal order in the form of elongated 1 1 0 domains. They are formed after quenching from temperatures above 450 ◦ C and ageing at temperatures above 70 ◦ C [2,6]. Generally either GPI or GPII zones can form as precursors to the metastable phase. Metastable phase with platelet morphology was thought to be the main hardening phase and responsible for peak hardening in Al–Zn–Mg(–Cu) alloys [6]. It has been concluded that phase has a hexagonal
∗ Corresponding author. Tel.: +86 451 8641 2133; fax: +86 451 8641 3922. E-mail address: [email protected] (L. Zhen). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.065
lattice with a = 0.496 and c = 1.402 nm [5,12], and is semi-coherent with the aluminum matrix [2]. It is well accepted that metastable phase may serve as nuclei for the stable during the artificial ageing. The phase also has a hexagonal structure with lattice dimensions a = 0.521 and c = 0.860 nm, but is incoherent with the aluminum matrix [13,14]. The formation of phase often leads to the decrease of the strength of aluminum alloy. Therefore, in order to optimize and control the mechanic properties, it is necessary to understand which phases are present at different steps of heat treatments. Although there have been lots of work on the evolutions of the precipitates, the mechanism for the formation of metastable phase is still controversial in Al–Zn–Mg(–Cu) alloys. Berg et al. [2] and Li et al. [5] have suggested that GPII zones serve as the precursor to metastable phase and gave the precipitation sequence as: Solid solution → VRC (vacancy-rich clusters) → GPII zones → metastable → stable . Other researchers, however, claim that metastable phase is formed from GPI zones [6,15]. Based on the high-resolution transmission electron microscopy (HRTEM) observations, Mukhopadhyay has suggested that large GPI zones are efficient for heterogeneous nucleation [15]. According to Sha and Cerezo’s work [6], only small GPI zones are important for formation and the transformation of these small GPI zones to phases is the dominant mechanism for formation. Therefore, the precise precipitation sequence remains to be revealed in Al–Zn–Mg(–Cu) alloys. As one kind of newly advanced Al–Zn–Mg–Cu alloys, 7055 aluminum alloy is extensively focused nowadays due to its attractive combined properties, such as high strength, high fracture tough-
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Table 1 Chemical composition (wt%) of AA 7055 aluminum alloy. Type Zn Mg Cu Zr Fe Si Cr Mn Al
AA 7055 7.76 1.94 2.35 0.12 0.061 0.055 0.0051 0.0076 Balance
ness, good resistance to the growth of fatigue cracks and good stress corrosion resistance, and so on [14,16–20]. It is generally believed that the optimum ageing precipitates are responsible for these good properties. However, the detailed information of how to achieve this microstructure is still not very clear in this alloy. The object of the present work is to characterize the precipitation in an AA 7055 aluminum alloy using transmission electron microscope (TEM), high-resolution transmission electron microscope and selected area diffraction pattern (SADP) technique, with the aim of understanding the evolution of precipitates and the related hardening in this alloy. 2. Experimental The chemical compositions of the 7055 aluminum alloy are shown in Table 1. Tensile samples of 15 mm gauge length, 6 mm width and 2 mm thickness were cut out of the plate in the rolling direction. The samples were solid solution treated at 477 ◦ C for 1 h and water quenched to room temperature. Immediately following quenching, the samples were aged at 120 and 160 ◦ C for times up to 3600 and 4320 min, respectively. After ageing treatment, tensile tests were carried out at room temperature using an Instron testing machine. Three samples were prepared for each material condition and a mean yield strength was given. The thin foils for TEM observation were prepared from 3-mm diameter discs punched out from 0.5 mm thick slices, cut from the grip section, mechanically ground and thinned by twin-jet electroplishing in a 30% Nital solution at −20 ◦ C and a potential difference of 15 V. TEM and HRTEM examinations were performed by using a JEM 2010 microscope operating at 200 kV. 3. Results 3.1. Evolution of the properties Fig. 1 shows the strength and elongation evolution of 7055 Al alloy during ageing at 120 and 160 ◦ C for different times. When the alloy was aged at 120 ◦ C, the strength increases significantly with increasing ageing time at the early stage, and then keeps stable at the later stage of the ageing, as shown in Fig. 1(a). When the alloy was aged at 160 ◦ C, the strength also increases significantly with increasing ageing time at the early stage. However, it decreases immediately after a peak value is achieved, as shown in Fig. 1(b). The elongation of the alloy aged at 120 and 160 ◦ C shows a similar trendy with ageing time. 3.2. Precipitation behavior 3.2.1. Precipitation behavior at 120 ◦ C Fig. 2 shows bright-field TEM images near a 0 1 1 zone axis of Al matrix of the samples aged at 120 ◦ C with different ageing times, and HRTEM images of precipitates in some conditions are
Fig. 1. Evolutions of strength and elongation of 7055 aluminum alloy during ageing at (a) 120 ◦ C and (b) 160 ◦ C.
also shown in the insets. The corresponding SADPs in Al 0 0 1 and Al 0 1 1 projections are shown in Figs. 3 and 4, respectively. After ageing for 5 min, darkly imaging precipitates in Al matrix could be observed, as shown in Fig. 2(a). It can be seen that most of precipitates show a round contrast, but some ones are slightly elongated along certain direction. It may be the results from structure factor and strain field. Both of them can contribute to the contrast of these small particles [6]. On the other hand, the uniform distribution of these precipitates in the matrix indicates that they are predominantly formed by homogeneous nucleation in the highly supersaturated solid solution. The HRTEM image in the inset shows that these precipitates are fully coherent with Al matrix lattice, and their average size is about 1–2 nm. From the corresponding SADPs shown in Fig. 3(a and b), weak diffraction spots at {1, (2n + 1)/4, 0} positions in Al 0 0 1 projection, which is corresponding to the diffraction features of GPI zones [6], could be clearly observed. It indicates that GPI zones have been formed at this stage of ageing. Therefore, according to the analysis of SADP these small precipitates observed in Fig. 2(a) can be identified as GPI zones. Although the diffraction spots from Al3 Zr phase could be also observed, we do not concern it hereafter, because Al3 Zr is not a precipitate during artificial ageing. On the other hand, Stiller et al. [9] have been reported that if GPII zones precipitate, weak spots from the GPII zones can be seen near 2/3 {2 2 0}Al in Al 0 0 1 projection. However, in the present work no clear spots near 2/3 {2 2 0} are observed in the diffraction pattern in Al 0 0 1 projection as shown in Fig. 3(a). Additionally, no streaks along {1 1 1} direction are observed in a SADP in Al 0 1 1 projection as shown in Fig. 3(b), which agree with the results of Sha and Cerezo’s work [6]. These all indicates
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Fig. 2. Bright-field TEM images near 0 1 1 zone axis of Al matrix of 7055 aluminum alloy samples aged at 120 ◦ C with different ageing times: (a) 5 min; (b) 60 min; (c) 300 min; (d) 1440 min and (e) 2880 min. HRTEM images of precipitates were shown in the insets in (a, c and e), respectively. The HRTEM images were observed from [0 1 1] direction of Al matrix.
that GPII zones are not the major precipitates at this stage of ageing. After ageing for 60 min, a lot of rod-like precipitates as well as near round ones could be observed in bright-field TEM image, as shown in Fig. 2(b). The size of rod-like precipitates is about 1 nm in width and 3 nm in length. They also uniformly distribute in the matrix. From the corresponding SADPs in Fig. 3(c and d), stronger GPI diffraction spots are observed in Al 0 0 1 SADP compared with that after 5 min ageing, indicating the volume fraction
of GPI zones increased during the 60 min ageing. On the other hand, a diffraction feature corresponding to can be clearly identified, such as weak diffraction spots at 1/3 and 2/3 of {2 2 0} in the Al 0 0 1 SADP, as shown in Fig. 3(c), and diffuse streaks are also observed along {1 1 1} directions in Al 0 1 1 SADP, indicating that metastable precipitates with platelet morphology has formed in the alloy. Therefore, it can be inferred that the rod-like precipitates are metastable precipitates. Again, there is no indication of GPII zone formation in SADPs after 60 min ageing, and since GPII zones
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Fig. 3. Selected area diffraction patterns of 7055 Al alloy samples aged at 120 ◦ C for 5 and 60 min: (a) 0 0 1, (b) 0 1 1 SADPs of a sample aged 5 min, respectively, (c) 0 0 1 and (d) 0 1 1 SADP of a sample aged 60 min, respectively.
forms at early stages, it therefore is unlikely to be present at any later stage of ageing [6]. In this stage, the alloy contains both GPI and metastable , and the former are dominant. After ageing for 300 min, 1440 and 2880 min, rod-like and near round precipitates are still observed in 0 1 1 bright field image, as shown in Figs. 2(c–e), respectively. The small, near round GPI zones decrease with increasing ageing time, while large rod-like precipitates and large near round ones increase with increasing ageing time. According to the previous analysis, the large rod-like precipitates are probably the platelet phases. However, the large ellipsoidal precipitates still remain controversial, whether they are GPI zones [6] or variants of [21]. The HRTEM observation shows that the phases are fully coherent with Al matrix lattice after ageing for 300 min, as shown in the insets of Fig. 2(c), while they are semi-coherent with Al matrix lattice after ageing for 2880 min, as shown in the inset of Fig. 2(e). From the corresponding SADPs in Fig. 4, it can be seen that after 300 min ageing at 120 ◦ C, GPI diffraction spots are still observed but become slightly weaker in Al 0 0 1 SADP, indicating that the GPI density has been decreased at this stage, as shown in Fig. 4(a). However, diffraction spots of phases in Al 0 0 1 SADP as well as streaks along {1 1 1} directions in Al 0 1 1 SADP become slightly stronger, as shown in Fig. 4(a and b). This suggests that the volume fraction of phase increased. After ageing for 1440 min, much stronger diffraction spots at 1/3 and 2/3 of {2 2 0}Al are observed in both Al 0 0 1 SADP and Al 0 1 1 SADP, as shown in Fig. 4(c and d), indicating that volume fraction
increased further in the alloy at this stage. However, a pair of spots very close to 2/3 {2 2 0}Al position in Al 0 0 1 SADP indicate that precipitates have already formed. It agrees with the observation of Sha and Cerezo [6]. Furthermore, the GPI diffraction spots are also observed in Al 0 0 1 SADP. After ageing for 2880 min, as shown in Fig. 4(e and f), GPI diffraction spots become weaker further, and it is still hard to separate the spots from and . 3.2.2. Precipitation behavior at 160 ◦ C Fig. 5 shows bright-field TEM images near 0 1 1 zone axis of Al matrix of samples aged at 160 ◦ C with different ageing times, and HRTEM images of precipitates in some conditions are also shown in the insets. After ageing for 5 min, small particles distributed uniformly in the matrix are observed, as shown in Fig. 5(a). Their size is smaller and the number of particles is larger as compared with that observed in the sample aged at 120 ◦ C for 5 min. These particles should be GPI zones according to the analysis previously. They are also fully coherent with Al matrix lattice according to the HRTEM observation, as shown in the inset of Fig. 5(a). The precipitates grew quickly with increasing ageing time, as shown in Figs. 5(b–f). The rod-like precipitates are dominant after 30 and 300 min ageing, indicating that phases were the main precipitates at these stage. According to the HRTEM observations, phase however becomes semi-coherent with Al matrix lattice after ageing for 300 min, while it is fully coherent with Al matrix lattice after ageing for 30 min, as shown in the insets of Fig. 5(b and d). Further-
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Fig. 4. Selected area diffraction patterns of 7055 Al alloy samples aged at 120 ◦ C for 300, 1440 and 2880 min: (a) 0 0 1, (b) 0 1 1 SAD patterns of a sample aged 300 min, respectively; (c) 0 0 1, (d) 0 1 1 SAD patterns of a sample aged 1440 min, respectively; (e) 0 0 1, (f) 0 1 1 SAD patterns of a sample aged 2880 min, respectively.
more, some extraordinary large precipitates are also observed after ageing for 300 min, as shown in Fig. 5(d). It may indicate that a new phase has been already formed. After ageing for 720 and 2880 min, more extraordinary large precipitates with plate-like morphology are observed and have become the dominant precipitates, although some near round ones are also present. The rod-like phase seems to disappear. After ageing for 2880 min, the precipitates coarsen further. HRTEM observation shows that the coarsen precipitate is incoherent with Al matrix lattice, as shown in the inset of Fig. 5(f).
Fig. 6 shows SADPs in Al 0 0 1, 0 1 1 projections from samples aged at 160 ◦ C for different ageing times after solution treatment. After 30 min ageing at 160 ◦ C, weak spots of GPI and , which have been described above, are observed in Al 0 0 1 SADP, as shown in Fig. 6(a and b). It indicates that the phase precipitates earlier at 160 ◦ C than at 120 ◦ C. After ageing for 300 min, the GPI spots become weaker significantly, whereas that of became slightly stronger, as shown in Fig. 6(c). It suggests that phase become the main phase in this stage, which agree well with the observa-
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Fig. 5. Bright-field TEM images near 0 1 1 zone axis of Al matrix of 7055 aluminum alloy samples aged at 160 ◦ C for different ageing times: (a) 5 min; (b) 30 min; (c) 60 min; (d) 300 min; (e) 720 min and (f) 2880 min. HRTEM images of precipitates were shown in the insets in (a, b, d and e), respectively. The HRTEM images were observed from [0 1 1] direction of Al matrix.
tion in Fig. 5. Simultaneously, two separated diffraction spots near 2/3 {2 2 0}Al positions in Al 011 SADP, as shown in Fig. 6(d), indicates that precipitation of the phase has started after ageing at 160 ◦ C for 300 min. Therefore, the extraordinary large precipitates observed in Fig. 5(d) can be identified as the phase. After ageing for 720 min, although the spots of GPI and still present in Al 011 SADP, the spots of become stronger, as shown in Fig. 6(e). After ageing for 2880 min, the spots of precipitates become stronger
and sharper further, indicating that the precipitates are in dominant at this stage. The streaks along {1 1 1} direction as shown in Fig. 6 h, become diffuse again, suggesting that the volume fraction of precipitates with a platelet morphology have decreased significantly. Additionally, the weak spots of GPI zones are also found as shown in Fig. 6(g). Like ageing at 120 ◦ C, there is no indication of features corresponding to GPII zones formation at any stage of ageing at 160 ◦ C, either.
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Fig. 6. Selected area diffraction patterns of 7055 Al alloy samples aged at 160 ◦ C for 30, 300, 720 and 2880 min: (a) 0 0 1, (b) 0 1 1 SAD patterns of a sample aged 30 min, respectively; (c) 0 0 1, (d) 0 1 1 SAD patterns of a sample aged 300 min, respectively; (e) 0 0 1, (f) 0 1 1 SAD patterns of a sample aged 720 min, respectively; (g) 0 0 1 and (h) 0 1 1 SAD patterns of a sample aged 2880 min, respectively.
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4. Discussion 4.1. The precipitation sequence It is known that two types of GP zones, GPI and GPII, co-exist at the early stage of artificial ageing in Al–Zn–Mg (–Cu) alloys [1–2,6,8]. Their diffraction patterns in Al 0 0 1 projection are different. The diffraction spots from the GPI zones occupy the positions of {1, (2n + 1)/4, 0}, while the spots corresponding to GPII zones are near 2/3 {2 2 0} in Al 0 0 1 projection [9]. In the present work, the spots from GPII zones were not observed at the early stage of ageing. Since GPII zones were not the major precipitates in this study, we will focus on GPI zones only hereafter. Three dimension atom probe (3DAP) investigation [6] has revealed that small clusters will form rapidly immediately after quenching. These clusters grow into large block GP zones while other ones are nucleating and growing at the early stage of ageing, resulting in a spread of size of GPI zones. The large GPI zones are relatively stable and continue to grow with increasing ageing time. It may be the reason that the weak GPI diffraction spots are still observed at the later stage of ageing, such as ageing at 160 ◦ C for 2880 min, although the small precipitates, which are generally considered as GPI zones, are no longer appear (Fig. 5(f)). In fact, it is the large GPI zones with round morphology that provide weak GPI diffraction features, and the smaller GPI zones lead to the formation of platelet in the present work because their surface energy is high and they are therefore much less stable than large GPI zones. In the present work, GPI diffraction spots became weaker with increasing ageing time from 60 to 1440 min at 120 ◦ C, while the diffraction spots of phases became stronger in this period of time. At the same time, it can be also found that the density of uniformly distributed small GPI zones decreased, while the uniformly distributed platelet increased, as shown in Figs. 2(b–d). These all provide direct evidence of the transformation of small GPI zones to platelet . It agrees well with the study of Sha and Cerezo [6]. The mechanism of the transformation of small GPI zones into platelet may proceed as follows. Zn atoms from the matrix diffuse to small GPI zones to feed their growth. In order to reduce the strain produced during the process, the GPI zones grow preferentially in a certain direction, i.e. 1 1 0. Later, another direction of growth can be activated, probably the 1 0 1 or 0 1 1. It therefore results in the small GPI zones developing into a two-dimensional structure on (1 1 1)Al and then transform into the platelet on the (1 1 1)Al . Although it has been reported that the precipitation of via GPII zones appears to be the most efficient mechanism for Al–Zn–Mg alloy, in the present study it is not the major process [22,23]. It is also known that the precipitation driving force increases with increasing the ageing temperature [24]. Therefore, the precipitation is promoted during ageing at 160 ◦ C and the precipitates form rapidly. In the present work, more and smaller GPI zones were observed in the sample aged at 160 ◦ C for 5 min as compared with those observed in the sample aged at 120 ◦ C for 5 min, as shown in Fig. 2(a) and Fig. 5(a). It probably contributes to the higher rate of particle formation than that of particle growth at such a high ageing temperature. After ageing 30–300 min at 160 ◦ C, the strength of the GPI spots decrease, while the strength of spots increase, as shown in Fig. 6. It suggests that some of the small GPI zones dissolute and other transform to metastable rapidly at such high temperature [23]. The metastable was found to form earlier at 160 ◦ C than at 120 ◦ C. Compared with ageing at 120 ◦ C, the significance feature was the clear observation of stable precipitates with a large size at the later stage ageing at 160 ◦ C. It indicates that can be easy to transform to at such higher temperature of 160 ◦ C. Although there are some differences between ageing at 120 and 160 ◦ C, they
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have the same precipitation sequence, which can be summarized as: Solid solution → Small GPI zones → Metastable → Stable , according to the results of the present work. 4.2. The precipitation hardening The yield strength evolution in the present work (Fig. 1) can be reasonably explained on the basis of precipitation shearing and bypassing [25–27]. When the particles are small, coherent and closely spaced, they are cut by moving dislocations. Their contribution to the yield strength, A , of a volume fraction f of shearable particles of radius r follows the equation: A = c3 fm rn , where c3 , m and n are constants [26]. It indicates that the yield strength increases with increasing the volume fraction and radius of the precipitates. As coarsening proceeds the particle strength increases and become incoherent with the matrix. The dislocations bulge between particles and escape by by-passing them. The contribution to the yield strength, B , can be given as B = c4 f1/2 r−1 , where c4 is a constant [26]. In this case, the strength decreases with increasing the radius of precipitates, while the volume fraction is generally constant during the coarsening. In the present work, therefore, the increasing of yield strength of the alloy aged at 120 and 160 ◦ C is attributed to the GPI zones and metastable . The volume fraction and radius of increase with increasing ageing time at the early stage of the ageing, resulting the persistent increasing of the yield strength according to the precipitation shearing mechanism. At the later stage of aging at 120 ◦ C, the main precipitates phases exhibit little changes in volume fraction and radius after ageing for 300 min and later, as shown in Fig. 2, and are still fully coherent with Al matrix. The precipitates can still be cut by the moving dislocations. Therefore, the yield strength keeps stable at a high level without any decreasing, as shown in Fig. 1. However, when the alloy is aged at 160 ◦ C, precipitates, which are incoherent with Al matrix, start to form after ageing for 300 min, and later they grow faster and coarsen further, as shown in Fig. 5. Therefore, the interaction mechanism between dislocation and precipitate transits from shearing to by-passing after the yield strength reaches the peak value, and the strength falls. 5. Conclusions From the results of the present work, the conclusions are drawn as following: 1. GPI zones and metastable platelets were observed in an Al–Zn–Mg–Cu alloy (AA 7055) solution treated for 1 h at 477 ◦ C, quenched and aged at 120 ◦ C for time up to 2880 min. GPI zones were found to be dominant in the alloy upon ageing at 120 ◦ C for 60 min. platelets are the main phase after ageing for 300 min and later. 2. Ageing at 160 ◦ C, the precipitation process is significantly promoted as compared with that at 120 ◦ C. GPI zones are found to decrease rapidly after ageing for 30 min and later. Their size is also smaller than that precipitated at 120 ◦ C during the early stage of ageing. precipitates start to be significant at 30 min, and become dominant when the aging time was between 30 and 300 min. After that, stable phase with large size forms and are the main phase at the later stage of ageing. 3. The precipitation sequence at 120 and 160 ◦ C ageing can be summarized as: Solid solution → small GPI zones → metastable → stable . No GPII zones were clearly identified. Additionally, weak diffraction spots of GPI zones remained at samples after long time ageing, indicating that some stable GPI zones can exist through the ageing process.
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4. Formation and growth of GPI zones and lead to the increase of the yield strength, while formation and coarsening of result in the decease of the strength. is responsible for the peak hardening of this alloy. References
[12] [13] [14] [15] [16] [17] [18]
[1] H. Löffler, I. Kovács, J. Lendvai, J Mater. Sci. 18 (1983) 2215–2240. [2] L.K. Berg, J. Gjønnes, V. Hansen, X.Z. Li, M. Knutson-Wedel, G. Waterloo, D. Schryvers, L.R. Wallenberg, Acta Mater. 49 (2001) 3443–3451. [3] V. Hansen, O.B. Karlsen, Y. Langsrud, J. Gjønnes, Mater. Sci. Technol. 20 (2004) 185–193. [4] D. Godard, P. Archambault, E. Aeby-Gautier, G. Lapasset, Acta Mater. 50 (2002) 2319–2329. [5] X.Z. Li, V. Hansen, J. Gjønnes, L.R. Wallenberg, Acta Mater. 47 (1999) 2651–2659. [6] G. Sha, A. Cerezo, Acta Mater. 52 (2004) 4503–4516. [7] G. Sha, Al Cerezo, Surf. Interf. Anal. 36 (2004) 564–568. [8] T. Engdahl, V. Hansen, P.J. Warren, K. Stiller, Mater. Sci. Eng. A 327 (2002) 59–64. [9] K. Stiller, P.J. Warren, V. Hansen, J. Angenete, J. Gjønnes, Mater. Sci. Eng. A 270 (1999) 55–63. [10] J.C. Werenskiold, A. Deschamps, Y. Bréchet, Mater. Sci. Eng. A 293 (2000) 267–2744. [11] A. Deschamps, F. Livet, Y. Brechet, Acta Mater. 47 (1999) 281–292.
[19] [20] [21] [22] [23] [24] [25] [26] [27]
J.H. Auld, S. McK. Cousland, J. Aust. Inst. Met. 19 (1974) 194–199. P.A. Thackery, J. Inst. Met. 96 (1968) 228–235. T.S. Srivatsan, S. Sriram, J. Mater. Sci. 32 (1997) 2883–3289. A.K. Mukhopadhyay, Q.B. Yang, S.R. Singh, Acta Metall. Mater. 42 (1994) 3083–3091. T.S. Srivatsan, V.K. Vasudevan, JOM 1 (1999) 42–45. T.S. Srivatsan, S. Anand, S. Sriram, V.K. Vasudevan, Mater. Sci. Eng. A 281 (2000) 292–304. J. Champlin, J. Zakrajsek, T.S. Srivatsan, P.C. Lam, M. Manoharan, Mater. Des. 20 (1999) 331–341. K.H. Chen, H.W. Liu, Z. Zhang, S. Li, Richard I. Todd, J. Mater. Proc. Technol. 142 (2003) 190–196. J.Z. Chen, L. Zhen, B.Y. Zhang, Y.X. Cui, S.L. Dai, Mater. Sci. Forum 546–549 (2007) 957–960. M. Dumont, W. Lefebvre, B. Doisneau-Cottignies, A. Deschamps, Acta Mater. 53 (2005) 2881–2892. V. Hansen, K. Stiller, G. Waterloo, J. Gjønnes, X.Z. Li, Mater. Sci. Forum 396–402 (2002) 815–820. X.G. Fan, D.M. Jiang, Q.C. Meng, Z.H. Lai, X.M. Zhang, Mater. Sci. Eng. A 427 (2006) 130–135. A. Deschamps, Y. Brechet, Acta Mater. 47 (1999) 293–305. P. Guyot, L. Cottignies, Acta Mater. 44 (1996) 4161–4167. H.R. Shercliff, M.F. Ashby, Acta Metall. Mater. 38 (1990) 1789–1802. M. Song, Mater. Sci. Eng. A 443 (2007) 172–177.
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Investigation of precipitation behavior and related hardening in AA 7055 aluminum alloy Junzhou Chen a,b , Liang Zhen a,∗ , Shoujie Yang b , Wenzhu Shao a , Shenglong Dai b a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Beijing Institute of Aeronautical Materials, Beijing 100095, PR China
a r t i c l e
i n f o
Article history: Received 8 May 2008 Received in revised form 21 August 2008 Accepted 24 September 2008 Keywords: Precipitation Aluminum alloy Transmission electron microscopy Ageing
a b s t r a c t The precipitation behavior and related hardening in AA 7055 aluminum alloy aged at 120 and 160 ◦ C was investigated in detail. GPI zones were the dominant phase in the alloy upon ageing at 120 ◦ C for 60 min. The metastable phase begins to precipitate in the alloy after being aged at 120 ◦ C for 60 min, and turns to be the main phase after ageing for 300 min. When the alloy was aged at 160 ◦ C, the precipitation was significantly promoted. The results also revealed that the transformation of small GPI zone to phase is the dominant mechanism for formation. Formation and growth of GPI zones and phases led to the increase of the yield strength, while formation and coarsening of resulted in the decrease of the strength. is responsible for the peak hardening of this alloy. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Precipitations from the supersaturated solid solutions have been extensively studied in aluminum alloys [1–10], since they play a very important role in determining the mechanical properties of these alloys. The general precipitation sequence of Al–Zn–Mg (–Cu) alloys can be summarized [6,11] as: Solid solution → GP zones → Metastable → Stable . GP zones are generally accepted as precipitates formed during room temperature ageing and the early stages of artificial ageing. They may lead to the increase of the strength of these alloys. There are two types of GP zones, i.e., GPI and GPII, with different structures [2]. GPI zones are fully coherent with the Al matrix, with internal ordering of Zn and Al or Mg on the {0 0 1}Al planes, based on an AuCu(I)-type sub-unit, and periodic anti-phase boundaries. They could be formed over a wide ageing temperature range, from room temperature to 140–150 ◦ C, independently of quenching temperature. GPII are zn-rich layers on {1 1 1}Al planes, one-to two atoms thick and 3–5 nm wide, with internal order in the form of elongated 1 1 0 domains. They are formed after quenching from temperatures above 450 ◦ C and ageing at temperatures above 70 ◦ C [2,6]. Generally either GPI or GPII zones can form as precursors to the metastable phase. Metastable phase with platelet morphology was thought to be the main hardening phase and responsible for peak hardening in Al–Zn–Mg(–Cu) alloys [6]. It has been concluded that phase has a hexagonal
∗ Corresponding author. Tel.: +86 451 8641 2133; fax: +86 451 8641 3922. E-mail address: [email protected] (L. Zhen). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.065
lattice with a = 0.496 and c = 1.402 nm [5,12], and is semi-coherent with the aluminum matrix [2]. It is well accepted that metastable phase may serve as nuclei for the stable during the artificial ageing. The phase also has a hexagonal structure with lattice dimensions a = 0.521 and c = 0.860 nm, but is incoherent with the aluminum matrix [13,14]. The formation of phase often leads to the decrease of the strength of aluminum alloy. Therefore, in order to optimize and control the mechanic properties, it is necessary to understand which phases are present at different steps of heat treatments. Although there have been lots of work on the evolutions of the precipitates, the mechanism for the formation of metastable phase is still controversial in Al–Zn–Mg(–Cu) alloys. Berg et al. [2] and Li et al. [5] have suggested that GPII zones serve as the precursor to metastable phase and gave the precipitation sequence as: Solid solution → VRC (vacancy-rich clusters) → GPII zones → metastable → stable . Other researchers, however, claim that metastable phase is formed from GPI zones [6,15]. Based on the high-resolution transmission electron microscopy (HRTEM) observations, Mukhopadhyay has suggested that large GPI zones are efficient for heterogeneous nucleation [15]. According to Sha and Cerezo’s work [6], only small GPI zones are important for formation and the transformation of these small GPI zones to phases is the dominant mechanism for formation. Therefore, the precise precipitation sequence remains to be revealed in Al–Zn–Mg(–Cu) alloys. As one kind of newly advanced Al–Zn–Mg–Cu alloys, 7055 aluminum alloy is extensively focused nowadays due to its attractive combined properties, such as high strength, high fracture tough-
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Table 1 Chemical composition (wt%) of AA 7055 aluminum alloy. Type Zn Mg Cu Zr Fe Si Cr Mn Al
AA 7055 7.76 1.94 2.35 0.12 0.061 0.055 0.0051 0.0076 Balance
ness, good resistance to the growth of fatigue cracks and good stress corrosion resistance, and so on [14,16–20]. It is generally believed that the optimum ageing precipitates are responsible for these good properties. However, the detailed information of how to achieve this microstructure is still not very clear in this alloy. The object of the present work is to characterize the precipitation in an AA 7055 aluminum alloy using transmission electron microscope (TEM), high-resolution transmission electron microscope and selected area diffraction pattern (SADP) technique, with the aim of understanding the evolution of precipitates and the related hardening in this alloy. 2. Experimental The chemical compositions of the 7055 aluminum alloy are shown in Table 1. Tensile samples of 15 mm gauge length, 6 mm width and 2 mm thickness were cut out of the plate in the rolling direction. The samples were solid solution treated at 477 ◦ C for 1 h and water quenched to room temperature. Immediately following quenching, the samples were aged at 120 and 160 ◦ C for times up to 3600 and 4320 min, respectively. After ageing treatment, tensile tests were carried out at room temperature using an Instron testing machine. Three samples were prepared for each material condition and a mean yield strength was given. The thin foils for TEM observation were prepared from 3-mm diameter discs punched out from 0.5 mm thick slices, cut from the grip section, mechanically ground and thinned by twin-jet electroplishing in a 30% Nital solution at −20 ◦ C and a potential difference of 15 V. TEM and HRTEM examinations were performed by using a JEM 2010 microscope operating at 200 kV. 3. Results 3.1. Evolution of the properties Fig. 1 shows the strength and elongation evolution of 7055 Al alloy during ageing at 120 and 160 ◦ C for different times. When the alloy was aged at 120 ◦ C, the strength increases significantly with increasing ageing time at the early stage, and then keeps stable at the later stage of the ageing, as shown in Fig. 1(a). When the alloy was aged at 160 ◦ C, the strength also increases significantly with increasing ageing time at the early stage. However, it decreases immediately after a peak value is achieved, as shown in Fig. 1(b). The elongation of the alloy aged at 120 and 160 ◦ C shows a similar trendy with ageing time. 3.2. Precipitation behavior 3.2.1. Precipitation behavior at 120 ◦ C Fig. 2 shows bright-field TEM images near a 0 1 1 zone axis of Al matrix of the samples aged at 120 ◦ C with different ageing times, and HRTEM images of precipitates in some conditions are
Fig. 1. Evolutions of strength and elongation of 7055 aluminum alloy during ageing at (a) 120 ◦ C and (b) 160 ◦ C.
also shown in the insets. The corresponding SADPs in Al 0 0 1 and Al 0 1 1 projections are shown in Figs. 3 and 4, respectively. After ageing for 5 min, darkly imaging precipitates in Al matrix could be observed, as shown in Fig. 2(a). It can be seen that most of precipitates show a round contrast, but some ones are slightly elongated along certain direction. It may be the results from structure factor and strain field. Both of them can contribute to the contrast of these small particles [6]. On the other hand, the uniform distribution of these precipitates in the matrix indicates that they are predominantly formed by homogeneous nucleation in the highly supersaturated solid solution. The HRTEM image in the inset shows that these precipitates are fully coherent with Al matrix lattice, and their average size is about 1–2 nm. From the corresponding SADPs shown in Fig. 3(a and b), weak diffraction spots at {1, (2n + 1)/4, 0} positions in Al 0 0 1 projection, which is corresponding to the diffraction features of GPI zones [6], could be clearly observed. It indicates that GPI zones have been formed at this stage of ageing. Therefore, according to the analysis of SADP these small precipitates observed in Fig. 2(a) can be identified as GPI zones. Although the diffraction spots from Al3 Zr phase could be also observed, we do not concern it hereafter, because Al3 Zr is not a precipitate during artificial ageing. On the other hand, Stiller et al. [9] have been reported that if GPII zones precipitate, weak spots from the GPII zones can be seen near 2/3 {2 2 0}Al in Al 0 0 1 projection. However, in the present work no clear spots near 2/3 {2 2 0} are observed in the diffraction pattern in Al 0 0 1 projection as shown in Fig. 3(a). Additionally, no streaks along {1 1 1} direction are observed in a SADP in Al 0 1 1 projection as shown in Fig. 3(b), which agree with the results of Sha and Cerezo’s work [6]. These all indicates
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Fig. 2. Bright-field TEM images near 0 1 1 zone axis of Al matrix of 7055 aluminum alloy samples aged at 120 ◦ C with different ageing times: (a) 5 min; (b) 60 min; (c) 300 min; (d) 1440 min and (e) 2880 min. HRTEM images of precipitates were shown in the insets in (a, c and e), respectively. The HRTEM images were observed from [0 1 1] direction of Al matrix.
that GPII zones are not the major precipitates at this stage of ageing. After ageing for 60 min, a lot of rod-like precipitates as well as near round ones could be observed in bright-field TEM image, as shown in Fig. 2(b). The size of rod-like precipitates is about 1 nm in width and 3 nm in length. They also uniformly distribute in the matrix. From the corresponding SADPs in Fig. 3(c and d), stronger GPI diffraction spots are observed in Al 0 0 1 SADP compared with that after 5 min ageing, indicating the volume fraction
of GPI zones increased during the 60 min ageing. On the other hand, a diffraction feature corresponding to can be clearly identified, such as weak diffraction spots at 1/3 and 2/3 of {2 2 0} in the Al 0 0 1 SADP, as shown in Fig. 3(c), and diffuse streaks are also observed along {1 1 1} directions in Al 0 1 1 SADP, indicating that metastable precipitates with platelet morphology has formed in the alloy. Therefore, it can be inferred that the rod-like precipitates are metastable precipitates. Again, there is no indication of GPII zone formation in SADPs after 60 min ageing, and since GPII zones
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Fig. 3. Selected area diffraction patterns of 7055 Al alloy samples aged at 120 ◦ C for 5 and 60 min: (a) 0 0 1, (b) 0 1 1 SADPs of a sample aged 5 min, respectively, (c) 0 0 1 and (d) 0 1 1 SADP of a sample aged 60 min, respectively.
forms at early stages, it therefore is unlikely to be present at any later stage of ageing [6]. In this stage, the alloy contains both GPI and metastable , and the former are dominant. After ageing for 300 min, 1440 and 2880 min, rod-like and near round precipitates are still observed in 0 1 1 bright field image, as shown in Figs. 2(c–e), respectively. The small, near round GPI zones decrease with increasing ageing time, while large rod-like precipitates and large near round ones increase with increasing ageing time. According to the previous analysis, the large rod-like precipitates are probably the platelet phases. However, the large ellipsoidal precipitates still remain controversial, whether they are GPI zones [6] or variants of [21]. The HRTEM observation shows that the phases are fully coherent with Al matrix lattice after ageing for 300 min, as shown in the insets of Fig. 2(c), while they are semi-coherent with Al matrix lattice after ageing for 2880 min, as shown in the inset of Fig. 2(e). From the corresponding SADPs in Fig. 4, it can be seen that after 300 min ageing at 120 ◦ C, GPI diffraction spots are still observed but become slightly weaker in Al 0 0 1 SADP, indicating that the GPI density has been decreased at this stage, as shown in Fig. 4(a). However, diffraction spots of phases in Al 0 0 1 SADP as well as streaks along {1 1 1} directions in Al 0 1 1 SADP become slightly stronger, as shown in Fig. 4(a and b). This suggests that the volume fraction of phase increased. After ageing for 1440 min, much stronger diffraction spots at 1/3 and 2/3 of {2 2 0}Al are observed in both Al 0 0 1 SADP and Al 0 1 1 SADP, as shown in Fig. 4(c and d), indicating that volume fraction
increased further in the alloy at this stage. However, a pair of spots very close to 2/3 {2 2 0}Al position in Al 0 0 1 SADP indicate that precipitates have already formed. It agrees with the observation of Sha and Cerezo [6]. Furthermore, the GPI diffraction spots are also observed in Al 0 0 1 SADP. After ageing for 2880 min, as shown in Fig. 4(e and f), GPI diffraction spots become weaker further, and it is still hard to separate the spots from and . 3.2.2. Precipitation behavior at 160 ◦ C Fig. 5 shows bright-field TEM images near 0 1 1 zone axis of Al matrix of samples aged at 160 ◦ C with different ageing times, and HRTEM images of precipitates in some conditions are also shown in the insets. After ageing for 5 min, small particles distributed uniformly in the matrix are observed, as shown in Fig. 5(a). Their size is smaller and the number of particles is larger as compared with that observed in the sample aged at 120 ◦ C for 5 min. These particles should be GPI zones according to the analysis previously. They are also fully coherent with Al matrix lattice according to the HRTEM observation, as shown in the inset of Fig. 5(a). The precipitates grew quickly with increasing ageing time, as shown in Figs. 5(b–f). The rod-like precipitates are dominant after 30 and 300 min ageing, indicating that phases were the main precipitates at these stage. According to the HRTEM observations, phase however becomes semi-coherent with Al matrix lattice after ageing for 300 min, while it is fully coherent with Al matrix lattice after ageing for 30 min, as shown in the insets of Fig. 5(b and d). Further-
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Fig. 4. Selected area diffraction patterns of 7055 Al alloy samples aged at 120 ◦ C for 300, 1440 and 2880 min: (a) 0 0 1, (b) 0 1 1 SAD patterns of a sample aged 300 min, respectively; (c) 0 0 1, (d) 0 1 1 SAD patterns of a sample aged 1440 min, respectively; (e) 0 0 1, (f) 0 1 1 SAD patterns of a sample aged 2880 min, respectively.
more, some extraordinary large precipitates are also observed after ageing for 300 min, as shown in Fig. 5(d). It may indicate that a new phase has been already formed. After ageing for 720 and 2880 min, more extraordinary large precipitates with plate-like morphology are observed and have become the dominant precipitates, although some near round ones are also present. The rod-like phase seems to disappear. After ageing for 2880 min, the precipitates coarsen further. HRTEM observation shows that the coarsen precipitate is incoherent with Al matrix lattice, as shown in the inset of Fig. 5(f).
Fig. 6 shows SADPs in Al 0 0 1, 0 1 1 projections from samples aged at 160 ◦ C for different ageing times after solution treatment. After 30 min ageing at 160 ◦ C, weak spots of GPI and , which have been described above, are observed in Al 0 0 1 SADP, as shown in Fig. 6(a and b). It indicates that the phase precipitates earlier at 160 ◦ C than at 120 ◦ C. After ageing for 300 min, the GPI spots become weaker significantly, whereas that of became slightly stronger, as shown in Fig. 6(c). It suggests that phase become the main phase in this stage, which agree well with the observa-
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Fig. 5. Bright-field TEM images near 0 1 1 zone axis of Al matrix of 7055 aluminum alloy samples aged at 160 ◦ C for different ageing times: (a) 5 min; (b) 30 min; (c) 60 min; (d) 300 min; (e) 720 min and (f) 2880 min. HRTEM images of precipitates were shown in the insets in (a, b, d and e), respectively. The HRTEM images were observed from [0 1 1] direction of Al matrix.
tion in Fig. 5. Simultaneously, two separated diffraction spots near 2/3 {2 2 0}Al positions in Al 011 SADP, as shown in Fig. 6(d), indicates that precipitation of the phase has started after ageing at 160 ◦ C for 300 min. Therefore, the extraordinary large precipitates observed in Fig. 5(d) can be identified as the phase. After ageing for 720 min, although the spots of GPI and still present in Al 011 SADP, the spots of become stronger, as shown in Fig. 6(e). After ageing for 2880 min, the spots of precipitates become stronger
and sharper further, indicating that the precipitates are in dominant at this stage. The streaks along {1 1 1} direction as shown in Fig. 6 h, become diffuse again, suggesting that the volume fraction of precipitates with a platelet morphology have decreased significantly. Additionally, the weak spots of GPI zones are also found as shown in Fig. 6(g). Like ageing at 120 ◦ C, there is no indication of features corresponding to GPII zones formation at any stage of ageing at 160 ◦ C, either.
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Fig. 6. Selected area diffraction patterns of 7055 Al alloy samples aged at 160 ◦ C for 30, 300, 720 and 2880 min: (a) 0 0 1, (b) 0 1 1 SAD patterns of a sample aged 30 min, respectively; (c) 0 0 1, (d) 0 1 1 SAD patterns of a sample aged 300 min, respectively; (e) 0 0 1, (f) 0 1 1 SAD patterns of a sample aged 720 min, respectively; (g) 0 0 1 and (h) 0 1 1 SAD patterns of a sample aged 2880 min, respectively.
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4. Discussion 4.1. The precipitation sequence It is known that two types of GP zones, GPI and GPII, co-exist at the early stage of artificial ageing in Al–Zn–Mg (–Cu) alloys [1–2,6,8]. Their diffraction patterns in Al 0 0 1 projection are different. The diffraction spots from the GPI zones occupy the positions of {1, (2n + 1)/4, 0}, while the spots corresponding to GPII zones are near 2/3 {2 2 0} in Al 0 0 1 projection [9]. In the present work, the spots from GPII zones were not observed at the early stage of ageing. Since GPII zones were not the major precipitates in this study, we will focus on GPI zones only hereafter. Three dimension atom probe (3DAP) investigation [6] has revealed that small clusters will form rapidly immediately after quenching. These clusters grow into large block GP zones while other ones are nucleating and growing at the early stage of ageing, resulting in a spread of size of GPI zones. The large GPI zones are relatively stable and continue to grow with increasing ageing time. It may be the reason that the weak GPI diffraction spots are still observed at the later stage of ageing, such as ageing at 160 ◦ C for 2880 min, although the small precipitates, which are generally considered as GPI zones, are no longer appear (Fig. 5(f)). In fact, it is the large GPI zones with round morphology that provide weak GPI diffraction features, and the smaller GPI zones lead to the formation of platelet in the present work because their surface energy is high and they are therefore much less stable than large GPI zones. In the present work, GPI diffraction spots became weaker with increasing ageing time from 60 to 1440 min at 120 ◦ C, while the diffraction spots of phases became stronger in this period of time. At the same time, it can be also found that the density of uniformly distributed small GPI zones decreased, while the uniformly distributed platelet increased, as shown in Figs. 2(b–d). These all provide direct evidence of the transformation of small GPI zones to platelet . It agrees well with the study of Sha and Cerezo [6]. The mechanism of the transformation of small GPI zones into platelet may proceed as follows. Zn atoms from the matrix diffuse to small GPI zones to feed their growth. In order to reduce the strain produced during the process, the GPI zones grow preferentially in a certain direction, i.e. 1 1 0. Later, another direction of growth can be activated, probably the 1 0 1 or 0 1 1. It therefore results in the small GPI zones developing into a two-dimensional structure on (1 1 1)Al and then transform into the platelet on the (1 1 1)Al . Although it has been reported that the precipitation of via GPII zones appears to be the most efficient mechanism for Al–Zn–Mg alloy, in the present study it is not the major process [22,23]. It is also known that the precipitation driving force increases with increasing the ageing temperature [24]. Therefore, the precipitation is promoted during ageing at 160 ◦ C and the precipitates form rapidly. In the present work, more and smaller GPI zones were observed in the sample aged at 160 ◦ C for 5 min as compared with those observed in the sample aged at 120 ◦ C for 5 min, as shown in Fig. 2(a) and Fig. 5(a). It probably contributes to the higher rate of particle formation than that of particle growth at such a high ageing temperature. After ageing 30–300 min at 160 ◦ C, the strength of the GPI spots decrease, while the strength of spots increase, as shown in Fig. 6. It suggests that some of the small GPI zones dissolute and other transform to metastable rapidly at such high temperature [23]. The metastable was found to form earlier at 160 ◦ C than at 120 ◦ C. Compared with ageing at 120 ◦ C, the significance feature was the clear observation of stable precipitates with a large size at the later stage ageing at 160 ◦ C. It indicates that can be easy to transform to at such higher temperature of 160 ◦ C. Although there are some differences between ageing at 120 and 160 ◦ C, they
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have the same precipitation sequence, which can be summarized as: Solid solution → Small GPI zones → Metastable → Stable , according to the results of the present work. 4.2. The precipitation hardening The yield strength evolution in the present work (Fig. 1) can be reasonably explained on the basis of precipitation shearing and bypassing [25–27]. When the particles are small, coherent and closely spaced, they are cut by moving dislocations. Their contribution to the yield strength, A , of a volume fraction f of shearable particles of radius r follows the equation: A = c3 fm rn , where c3 , m and n are constants [26]. It indicates that the yield strength increases with increasing the volume fraction and radius of the precipitates. As coarsening proceeds the particle strength increases and become incoherent with the matrix. The dislocations bulge between particles and escape by by-passing them. The contribution to the yield strength, B , can be given as B = c4 f1/2 r−1 , where c4 is a constant [26]. In this case, the strength decreases with increasing the radius of precipitates, while the volume fraction is generally constant during the coarsening. In the present work, therefore, the increasing of yield strength of the alloy aged at 120 and 160 ◦ C is attributed to the GPI zones and metastable . The volume fraction and radius of increase with increasing ageing time at the early stage of the ageing, resulting the persistent increasing of the yield strength according to the precipitation shearing mechanism. At the later stage of aging at 120 ◦ C, the main precipitates phases exhibit little changes in volume fraction and radius after ageing for 300 min and later, as shown in Fig. 2, and are still fully coherent with Al matrix. The precipitates can still be cut by the moving dislocations. Therefore, the yield strength keeps stable at a high level without any decreasing, as shown in Fig. 1. However, when the alloy is aged at 160 ◦ C, precipitates, which are incoherent with Al matrix, start to form after ageing for 300 min, and later they grow faster and coarsen further, as shown in Fig. 5. Therefore, the interaction mechanism between dislocation and precipitate transits from shearing to by-passing after the yield strength reaches the peak value, and the strength falls. 5. Conclusions From the results of the present work, the conclusions are drawn as following: 1. GPI zones and metastable platelets were observed in an Al–Zn–Mg–Cu alloy (AA 7055) solution treated for 1 h at 477 ◦ C, quenched and aged at 120 ◦ C for time up to 2880 min. GPI zones were found to be dominant in the alloy upon ageing at 120 ◦ C for 60 min. platelets are the main phase after ageing for 300 min and later. 2. Ageing at 160 ◦ C, the precipitation process is significantly promoted as compared with that at 120 ◦ C. GPI zones are found to decrease rapidly after ageing for 30 min and later. Their size is also smaller than that precipitated at 120 ◦ C during the early stage of ageing. precipitates start to be significant at 30 min, and become dominant when the aging time was between 30 and 300 min. After that, stable phase with large size forms and are the main phase at the later stage of ageing. 3. The precipitation sequence at 120 and 160 ◦ C ageing can be summarized as: Solid solution → small GPI zones → metastable → stable . No GPII zones were clearly identified. Additionally, weak diffraction spots of GPI zones remained at samples after long time ageing, indicating that some stable GPI zones can exist through the ageing process.
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4. Formation and growth of GPI zones and lead to the increase of the yield strength, while formation and coarsening of result in the decease of the strength. is responsible for the peak hardening of this alloy. References
[12] [13] [14] [15] [16] [17] [18]
[1] H. Löffler, I. Kovács, J. Lendvai, J Mater. Sci. 18 (1983) 2215–2240. [2] L.K. Berg, J. Gjønnes, V. Hansen, X.Z. Li, M. Knutson-Wedel, G. Waterloo, D. Schryvers, L.R. Wallenberg, Acta Mater. 49 (2001) 3443–3451. [3] V. Hansen, O.B. Karlsen, Y. Langsrud, J. Gjønnes, Mater. Sci. Technol. 20 (2004) 185–193. [4] D. Godard, P. Archambault, E. Aeby-Gautier, G. Lapasset, Acta Mater. 50 (2002) 2319–2329. [5] X.Z. Li, V. Hansen, J. Gjønnes, L.R. Wallenberg, Acta Mater. 47 (1999) 2651–2659. [6] G. Sha, A. Cerezo, Acta Mater. 52 (2004) 4503–4516. [7] G. Sha, Al Cerezo, Surf. Interf. Anal. 36 (2004) 564–568. [8] T. Engdahl, V. Hansen, P.J. Warren, K. Stiller, Mater. Sci. Eng. A 327 (2002) 59–64. [9] K. Stiller, P.J. Warren, V. Hansen, J. Angenete, J. Gjønnes, Mater. Sci. Eng. A 270 (1999) 55–63. [10] J.C. Werenskiold, A. Deschamps, Y. Bréchet, Mater. Sci. Eng. A 293 (2000) 267–2744. [11] A. Deschamps, F. Livet, Y. Brechet, Acta Mater. 47 (1999) 281–292.
[19] [20] [21] [22] [23] [24] [25] [26] [27]
J.H. Auld, S. McK. Cousland, J. Aust. Inst. Met. 19 (1974) 194–199. P.A. Thackery, J. Inst. Met. 96 (1968) 228–235. T.S. Srivatsan, S. Sriram, J. Mater. Sci. 32 (1997) 2883–3289. A.K. Mukhopadhyay, Q.B. Yang, S.R. Singh, Acta Metall. Mater. 42 (1994) 3083–3091. T.S. Srivatsan, V.K. Vasudevan, JOM 1 (1999) 42–45. T.S. Srivatsan, S. Anand, S. Sriram, V.K. Vasudevan, Mater. Sci. Eng. A 281 (2000) 292–304. J. Champlin, J. Zakrajsek, T.S. Srivatsan, P.C. Lam, M. Manoharan, Mater. Des. 20 (1999) 331–341. K.H. Chen, H.W. Liu, Z. Zhang, S. Li, Richard I. Todd, J. Mater. Proc. Technol. 142 (2003) 190–196. J.Z. Chen, L. Zhen, B.Y. Zhang, Y.X. Cui, S.L. Dai, Mater. Sci. Forum 546–549 (2007) 957–960. M. Dumont, W. Lefebvre, B. Doisneau-Cottignies, A. Deschamps, Acta Mater. 53 (2005) 2881–2892. V. Hansen, K. Stiller, G. Waterloo, J. Gjønnes, X.Z. Li, Mater. Sci. Forum 396–402 (2002) 815–820. X.G. Fan, D.M. Jiang, Q.C. Meng, Z.H. Lai, X.M. Zhang, Mater. Sci. Eng. A 427 (2006) 130–135. A. Deschamps, Y. Brechet, Acta Mater. 47 (1999) 293–305. P. Guyot, L. Cottignies, Acta Mater. 44 (1996) 4161–4167. H.R. Shercliff, M.F. Ashby, Acta Metall. Mater. 38 (1990) 1789–1802. M. Song, Mater. Sci. Eng. A 443 (2007) 172–177.