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Influence of deformation microstructure on the precipitation behaviors of an Al-4Mg-0.3Cu alloy
Accepted Manuscript Influence of deformation microstructure on the precipitation behaviors of an Al– 4Mg-0.3Cu alloy Xiaohui Yang, Kai Li, Xianghai An, Song Ni, Weifeng Wei, Yong Du, Min Song PII:
S0925-8388(16)33542-3
DOI:
10.1016/j.jallcom.2016.11.073
Reference:
JALCOM 39568
To appear in:
Journal of Alloys and Compounds
Received Date: 6 August 2016 Revised Date:
3 November 2016
Accepted Date: 5 November 2016
Please cite this article as: X. Yang, K. Li, X. An, S. Ni, W. Wei, Y. Du, M. Song, Influence of deformation microstructure on the precipitation behaviors of an Al–4Mg-0.3Cu alloy, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.073. 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.
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Influence of deformation microstructure on the precipitation
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behaviors of an Al–4Mg-0.3Cu alloy
Xiaohui Yanga, Kai Lia, Xianghai Anb, Song Nia,*, Weifeng Weia, Yong Dua, Min Songa,c,** a
School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia
Shenzhen Research Institute, Central South University, Shenzhen 518057, China
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c
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b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
Abstract
To provide insight into the effect of deformation microstructure on the precipitation behavior, an
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Al-4Mg-0.3Cu alloy was processed by high-pressure torsion followed by subsequent in-situ heating in a transmission electron microscope. Structural characterization revealed significant differences in the size, composition and spatial distribution of the precipitates in the alloy subjected to different strains.
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Different grain sizes led to various dislocation substructures and grain boundary volume fractions, which in turn governed precipitation behaviors. In addition, the composition of the precipitates along
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the grain boundaries varied with the grain size due to the change of the grain boundary volume fraction and diffusion behavior.
Keywords: Al-Mg; Cu addition; grain boundary; precipitation; in-situ TEM
*Corresponding author: [email protected] (S. Ni) ** Corresponding author: [email protected] (M. Song)
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1. Introduction In the last decade severe plastic deformation (SPD) techniques have been carried out on various
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metallic materials to obtain excellent mechanical properties due to the extensive grain refinement [1-4]. Equal channel angular pressing (ECAP) [2] and high-pressure torsion (HPT) [4] are among the most widely used SPD techniques. In addition to the grain refinement enhanced mechanical properties,
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SPD-induced segregation and/or precipitation of solute elements along the grain boundaries can also
lead to increased strength [5]. For instance, Al-Mg alloy produced by HPT was found to exhibit a very
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high strength, considerably exceeding a prediction based on the Hall–Petch effect for ultrafine grains (UFGs) since the Mg element was segregated along the grain boundaries [6]. It is believed that solute element segregation and/or precipitation during SPD processes are closely related to the formation of a large number of the grain boundaries [7]. As the fast diffusion channels for solute elements, grain
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boundaries play a crucial role in promoting heterogeneous precipitation [8]. Moreover, formation of high-density dislocations during SPD has been verified to enhance the heterogeneous nucleation of precipitates [9]. Thus, understanding the precipitation behaviors of variously deformed microstructures
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after SPD is helpful for further improving the mechanical properties of SPD-processed materials. Recently, abundant studies have been focused on the precipitation behavior of age-hardenable Al
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alloys after SPD processing, including Al-Mg-Si [10], Al-Cu [11], and Al-Zn-Mg-Cu alloys [12]. SPD of non-age-hardenable Al-Mg alloys leads to the segregation of the alloying elements along the grain boundaries that also contributes to the alloy’s hardening [6]. It should be noted that addition of small
amount of Cu to Al-Mg alloys improves the mechanical properties and inhibits the formation of Lüders bands [13]. Therefore, an investigation of the segregation and/or precipitation behaviors of the supersaturated Al-Mg alloys with minor Cu is a key issue for practical applications. Up to now, the
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segregation and precipitation behaviors of Al-Mg alloy with minor Cu after SPD received little attention in comparison to the most studied Al-Mg-Cu alloy [14, 15]. Therefore, in this work an
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Al-4Mg-0.3Cu alloy processed by HPT was used to investigate the influence of SPD on the segregation and precipitation behaviors during heat treatment.
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2. Experimental
The Al-4Mg-0.3Cu (wt.%) alloy used in this study was fabricated by ingot metallurgy, in which
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commercially pure Al, pure Mg (both are 99.9% in purity) and an Al-Cu master alloy were used as the raw materials. The homogenization treatment of the alloy was at 733 K for 10 hours. Specimens for HPT processing were sectioned into disks from the alloy with a diameter slightly less than 20 mm and a thickness of ~2 mm. These disks were polished on both sides using sand papers (400-2000 grades) until
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their thickness reached ~1.5 mm. HPT processing was conducted using a quasi-constrained HPT facility [3]. Disks were subjected to HPT for 5 turns at room temperature under an applied pressure of 4 GPa and a rotation rate of 1 turn per minute. Samples for transmission electron microscopy (TEM)
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investigation were prepared by a twin-jet polishing technique in a mixture of 30% nitric acid and 70% carbinol at -30 oC. Heat treatment was implemented by an in situ heating technique using a model 901
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hot stage controller in a Tecnai G2 20 S-Twin TEM with a temperature ramping rate of 60 oC per min.
After heat treatment, scanning TEM (STEM) observation of the microstructures was carried out in a Titan G2 60-300 microscope operated at 300 kV that provides Z-contrast to exhibit local composition variation. Elemental mapping of the precipitates and/or segregation was performed using energy dispersive X-ray spectroscopy (EDS) with a Super-X detector. To facilitate discussion, the specimen locations for TEM observation are labeled as regions I, II and III, as shown in Figure 1. The effective
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strain calculated from the formula
ε = 2πNr / 3h (where N is the number of HPT turns, h is the
thickness of the disk and r is the distance from the center of the disk) [16] was about 0, 90 and 180 at
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the positions I, II and III, respectively. The average grain sizes are derived by statistical measurements of 50 grains in region I, 100 grains in both regions II and III using TEM. Although the strain of the disk center is theoretically zero, the microstructure can be refined with continuous HPT deformation based
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on strain gradient effects [17, 18].
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3. Results
Optical microscopy observation shows that the average grain size of the Al-4Mg-0.3Cu alloy is about 400 µm. Figure 2 presents typical bright-field TEM images of (a) the alloy and the regions (b) I, (c) II and (d) III of the sample after HPT processing. Few dislocations can be observed in the TEM
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bright field image from the alloy (Fig. 2a). In contrast, many dislocation cells with an average size of ~ 0.6 µm formed in the region I of the deformed sample (Fig. 2b). With increasing the plastic strain, the average grain sizes are further reduced to ~ 200 nm and ~ 130 nm in the regions II and III, respectively,
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as shown in Figs. 2c and 2d. Extensive TEM observations indicated that, with increasing the strain, the average grain size decreased monotonically. Besides, the alloy was initially deformed by dislocations
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multiplication and slip, so the dislocation density increases. As the deformation proceeded, dislocations
slipped towards and annihilated at cell boundaries, making the dislocation cells evolve to subgrains with low angle grain boundaries or even high angle grain boundaries, as shown in Fig.2c and Fig.2d. In this deformation stage, dislocation density decreases. In fact, with the grain being refined to sub-micrometer size or even to nanometer size, the dislocation density in the grain interior decreased
gradually due to two factors. On one hand, the active stress of Frank-Read dislocation source is beyond
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the flow stress. On the other hand, dislocations were effectively dissolved by the grain boundaries. These dislocation density evolution behaviors have also been observed and verified in ECAPed Al and
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Cu [19]. Figure 3 shows typical STEM micrographs of (a) the alloy heated at 200 oC for 30 min and regions
(b) I, (c) II and (d) III of the HPT treated sample after heating at 200 oC for 5 min. The rod-like
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precipitates with various lengths from tens to hundreds of nanometers were observed in the grain interior of the sample (Fig. 3a). However, Fig.3b shows that rod-like precipitates with smaller size
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(arrowed) formed in the grain interior and irregular shaped precipitates were located along the grain boundaries in the region I after in-situ heating. Dislocations are favorable nucleation sites for the precipitates and potentially provide fast diffusion paths for solute atoms [20, 21]. So, precipitates in the grain interior are associated with dislocations formed in the region I. In contrast, only irregular
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precipitates are seen in the regions II and III. In the region II (Fig. 3c), most precipitates were located along the grain boundaries, while a few precipitates presented in the grain interior. In the region III (Fig. 3d), almost all precipitates are located along the grain boundaries. Obviously, the morphology and
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distribution of the precipitates are closely related to the grain size and dislocation density, which is significantly affected by the imposed plastic strain during HPT. It can be seen that precipitates along
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the grain boundaries were heterogeneously distributed. The precipitate size was calculated by averaging the values from measuring the precipitate along two mutually perpendicular directions of the precipitate. By statistical measurement about 150 precipitates along the grain boundaries in region I, II and III using TEM, the size varied in the range from 20 nm to 60 nm. In addition, the average grain sizes in regions I, II and III after in-situ heat treatment are 0.6 µm, 340 nm and 180 nm by statistical measurement of about 50 grains, respectively. Therefore, the fine and uniformly distributed precipitates
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along the grain boundaries and at the triple junctions can slow down the growth of the UFGs. Apart from the shape, size and distribution of the precipitates being affected by different deformed
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microstructures, their chemical composition and corresponding structures are also significantly affected. Figure 4a shows a HRTEM image of the precipitates in the grain interior in the sample after in-situ heat treatment at 200 oC for 30 min. The fast Fourier transform (FFT) pattern from the region marked by a
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white box in Fig.4a demonstrates that the precipitate Al2CuMg (S phase, with the space group of
Cmcm) follows crystallographic orientation relationship [100]S // [100]Al with the Al matrix [22, 23].
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Figure 5 shows a STEM image and the corresponding EDS mapping (b, c and d) of a precipitate along the grain boundary from the region II. The EDS mapping shows the enrichment of Mg and Cu elements in the precipitate. Figs.6a is the HRTEM images of precipitates along the grain boundary from the region II, and Fig.6b is the corresponding FFT pattern from the region marked by a white box in
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Figs.6a, indicating that the precipitates are the Al2CuMg phase. Analysis of ~ 20 precipitates along the grain boundaries from the region II found only the Al2CuMg phase. Figure 7 shows a STEM image and the corresponding EDS mapping (b, c and d) of the
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precipitates at a grain boundary from the region III. EDS mapping of precipitates 1 and 2 showed the lean and enrichment of Mg element, respectively. Figs.8a and 8d show the precipitates along the grain
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boundaries in the region III. The HRTEM images and FFT pattern shows that the precipitate in Fig.8a is the Al2MgCu phase, while the precipitate in Fig.8d is Al2Cu phase (with the space group of I4/mcm) [24]. Extensive analysis of the precipitates along the grain boundaries in the region III indicated two phases: Al2CuMg and Al2Cu. In order to understand the precipitation evolution along the grain
boundaries in detail, the STEM images of the region III of the sample before and after in-situ heating at 100 oC for 5 min are shown in Fig.9a and Fig.9c. To confirm that these bright segregation along the
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grain boundaries are Cu-rich, line scanning were performed along the arrows plotted in Fig.9a and Fig.9c. The EDS line-scanning in Fig.9b and Fig.9d clearly demonstrates that compared to those in the
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grain interior, the Cu content is higher, the Al content is lower and the Mg content remains almost constant at the grain boundary. The above phenomena about precipitates along the grain boundaries can thus be concluded as follows: (i) Cu diffuses to grain boundaries of the UFGs and the average Cu
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concentration along the grain boundaries is higher than that in the grain interiors; (ii) the grain
boundary segregation of Cu induced by HPT processing or in situ heat treatment in UFGs is more rapid
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than that of Mg.
4. Discussion
The results presented in this study indicate that significant difference exists in the size, composition
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and spatial distribution of the precipitates in Al-4Mg-0.3Cu alloys with various microstructures induced by HPT. Difference in grain size, stemmed from different strain, was noted to affect the dislocation
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density and the grain boundary volume, which in turn governed the evolution of the precipitates.
4.1. Effect of grain refinement on the precipitation behavior
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The present TEM results show that dislocation density in the grain interior changes significantly
during the grain refinement process. From Fig.2, the dislocation density in the grain interior increases firstly, and then decreases with decreasing the grain size. A large number of the dislocation lines were
generated in the grain interior in the region I, with an average grain size of about 0.6 µm (see Fig. 2b). Fig.2c shows that only a small number of the dislocations can be observed in the region II. Fig.2d reveals that the almost no dislocation exists in the UFGs with a size of ~130 nm from the region III.
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This observation is in consistent with the results from numerous previous studies [25-27], which showed that conventional dislocation mechanisms, e.g. Frank-Read source and dislocation
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accumulation, are inhibited when the grain size is close to 100 nm [28]. Hayes et al. also reported that there is very limited dislocation activity in the grain interior of cryomilled UFG pure Al and Al 5083 alloy [29]. Due to the relatively small interaction volume in the UFG interiors, dislocations are difficult
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to accumulate within the grains in the order of 100-200 nm. Moreover, in nanocrystalline grains or
UFGs, grain boundaries act as sources and sinks for dislocations, and units or partial dislocations are
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emitted from one grain boundary, traverse the grain and finally are absorbed into the opposite grain boundary without multiplying and interacting with other dislocations. Another feature of Al-4Mg-0.3Cu alloy processed by HPT is that the grain boundary volume increases with the grain refinement. Previous study has demonstrated that short-circuit self-diffusion is greatly enhanced in
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ultrafine-grained Al alloy prepared by ECAP in comparison to that in the well-annealed coarse-grained material [30]. The diffusion enhancement was explained in terms of a specific deformation-induced state of general high-angle grain boundaries. The volume diffusion coefficient of solute in matrix can
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be written as:
D = Do exp(−Q / RT)
D o = 1.49 × 10−5 m 2 ⋅ S−1 and
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where
Q = 120.5kJ ⋅ mol −1 for
(1) Mg
in
Al
[31],
and
D o = 6.54 × 10−5 m 2 ⋅ S−1 and Q = 136 KJ ⋅ mol −1 for Cu in Al [32]. The grain boundary diffusion
coefficients of Mg [33] and Cu [32] in Al can be two orders of magnitude larger than the volume diffusion coefficients of Mg and Cu in Al. In addition, HPT process is known to increase the concentrations of vacancies and grain boundary volume and this consequently increases the diffusion rate of solute atoms within the matrix [34]. Both the contribution of the dislocations in the grain
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interiors and deformation-induced grain boundaries are considered to account for the size and spatial distribution of the precipitates. Firstly, dislocations with a high density provide numerous
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heterogeneous nucleation sites for the precipitates in the region I, which eventually facilitates the formation of a large number of the precipitates near the dislocations. The dislocations serve as fast
diffusion paths for solute atoms in the matrix, thereby permitting rapid diffusion to the precipitates to
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assist their growth. Secondly, the grain boundary diffusion is prominent due to a large number of the grain boundaries volume characterized by UFGs. The fast diffusion of Mg and Cu elements along the
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boundaries in ultrafine-grained Al-4Mg-0.3Cu alloy implies that less solute will be available in the matrix for precipitation. Thus, precipitates inside grains decrease with refining the grain size. In addition, the size of the precipitates along the grain boundaries becomes larger with decreasing the grain size after the same heat treatment due to fast grain boundary diffusion with increasing the grain
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boundary volume.
4.2 The precipitate phase evolution along the grain boundaries in UFGs
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The precipitation may occur preferentially along the grain boundaries rather than in the matrix. Firstly, the free energy barrier for nucleation along the grain boundaries is lower than that in the matrix
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because that the precipitate/matrix interface can be created by a relatively small change in the grain
boundary structure. Secondly, the precipitation nucleation rate along the grain boundaries is higher than that in the matrix because the enhanced solute diffusivity along the grain boundaries increase the critical nucleus rate [35]. Therefore, the number of the precipitates at grain boundaries increases rapidly with decreasing the grain size. Additional factor that could accelerate nucleation includes a higher Cu content at the grain boundaries, since forced migration of boundaries during deformation can result in
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stronger Cu segregation, which would increase the driving force and local availability of solute [36]. Compared to that in regions I and II, the mean spacing of the precipitates along the grain boundaries
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decreases in region III. The “collector plate mechanism” can well account for the precipitate growing process along the grain boundaries [37]. The “collector plate mechanism” is easily understood, based on the fact that the Mg and Cu solute atoms migrate to the grain boundaries through volume diffusion,
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and then swiftly diffuses along the boundaries to the edges of the precipitate [37]. This is due to the fact that the grain boundary diffusivity is in orders of magnitude higher than bulk diffusivity. In this case,
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the mass flow of Mg and Cu to each grain boundary, “collector plate”, can be rapidly involved into the growth of the precipitates. Aaron and Aaronson demonstrated that the grain boundary precipitate lengthening and thickening vary as a function of (time)1/4 and (time)1/2 [37], respectively. However, the precipitates at the grain boundaries in region III show spherical shape and the growth of the precipitates
the precipitates.
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does not follow the rule of Aaron and Aaronson [37], which might be caused by to the coarsening of
Both experiments and atomistic modeling showed that segregation can influence grain boundary
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precipitation by altering the driving force for nucleation and the solute flux arriving at the grain boundaries [36, 38]. In this work, the Cu element shows strongly segregation along the grain
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boundaries than Mg element subjected to HPT processing. The segregation thickness at the grain boundaries of region III after in-situ heat treatment increases obviously, as shown in Fig.9. Similarly, ECAPed Al-Zn-Mg-Cu alloy at 200 oC also demonstrated that Mg and Cu elements segregated strongly
to the grain boundaries in UFGs [39]. The significant segregation of Cu element along the boundaries in UFGs after heat treatment implies that more Cu will be available along the grain boundaries for precipitation. In addition, precipitation in Al–Cu–Mg alloys can be changed by adjusting the Cu/Mg
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atomic ratio [40]. With grain refinement, Cu element diffusing more rapidly to the grain boundaries is attributed to the shorter diffusion paths. So, Al2Cu phase can be formed at the local regions of the grain
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boundaries from region III during heat treatment process due to the high Cu/Mg atomic ratio. In all, the precipitates along the grain boundaries contain Al, Mg and Cu elements in the initial stage. When grain
size is refined to nanoscale, the precipitates along the grain boundaries are composed of Al and Cu
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elements or Al, Mg and Cu elements due to the rapidly diffusion.
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5. Conclusions
Summarizing the results of the precipitation behaviors in supersaturated Al-4Mg alloy with minor Cu element subjected to HPT and in-situ heat treatment in TEM, it is indicated that the dislocations in the grain interior and grain boundaries jointly contribute to the precipitation mechanisms. At low strain,
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precipitates mostly locate in the grain interior due to the generation of high-density dislocations in the grain interior. When grain size is refined to nanoscale at high strain, the precipitates mostly locate along the grain boundaries due to the increase in the grain boundary volume. The size of the precipitates
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along the grain boundaries increases with refining the grain size due to the diffusion enhancement. In addition, the composition of precipitate phases along the grain boundaries changes with grain
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refinement.
Acknowledgements
The financial supports from National Natural Science Foundation of China (51531009 and 51501230), Grants from Shenzhen Science and Technology Project (JCYJ20140509142357196) and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and
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Manufacturing Collaborative Innovation Center are appreciated.
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Figure captions: Figure 1: The regions cut from the disks for TEM observations, i.e. the center region at r=0 marked as I,
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and the regions at r=4 mm and r=8 mm marked as II and III, respectively. Figure 2: TEM images of (a) the alloy, and the alloy subjected to HPT for 5 turns at the (b) region I, (c) region II and (d) region III, respectively.
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Figure 3: HAADF STEM images of (a) the alloy after in situ heat treatment at 200 oC for 30 min, and the (b) region I, (c) region II and (d) region III of the alloy subjected to HPT for 5 turns and
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then in situ heat treated at 200 oC for 5 min. Figure 4: (a) A HRTEM image of the precipitates in the
alloy after in situ heat treatment and (b) the
corresponding FFT pattern from the region marked by a white box in Fig.4a, showing that the precipitate is the Al2CuMg phase.
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Figure 5: (a) A STEM image of a precipitate along the grain boundary from the region II after in situ heat treatment at 200 oC for 5 min and (b-d) EDS mapping of Al, Cu and Mg elements of the precipitate.
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Figure 6: (a) HRTEM images of precipitates along the grain boundary and (b) the corresponding FFT pattern from the region marked by a white box in Fig.6a, showing that the precipitate is the
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Al2CuMg phase.
Figure 7: (a) A STEM image of the precipitates along the grain boundary from the region III and (b-d) EDS mapping of Al, Cu and Mg elements obtained from the white square in Fig.7a, showing that the precipitate 1 has no Mg concentration while the precipitate 2 contains a high concentration of Mg. Figure 8: (a,d) bright field TEM images of grain boundaries from the region III after in situ heat
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treatment at 200 oC for 5 min, (b,e) HRTEM images of the precipitates indicated by black arrow in (a) and (d), respectively, and (c,f) the corresponding FFT patterns from the region
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marked by white boxes in (b) and (e), respectively, showing that the two precipitates are the AlMgCu phase and Al2Cu phase, respectively.
Figure 9: The STEM images of the region III of the sample (a) before and (c) after in-situ heating at
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100 oC for 5 min. (b) and (d) the corresponding EDS line-scanning profiles of Al, Cu and Mg
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elements along the yellow arrow in (a) and (c), respectively.
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Research highlights
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1. Precipitation of an HPT deformed Al-4Mg-0.3Cu alloy was investigated. 2. Precipitation behavior is closely related to the dislocations and grain boundaries.
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3. Precipitates along the grain boundaries change with decreasing the grain size.
S0925-8388(16)33542-3
DOI:
10.1016/j.jallcom.2016.11.073
Reference:
JALCOM 39568
To appear in:
Journal of Alloys and Compounds
Received Date: 6 August 2016 Revised Date:
3 November 2016
Accepted Date: 5 November 2016
Please cite this article as: X. Yang, K. Li, X. An, S. Ni, W. Wei, Y. Du, M. Song, Influence of deformation microstructure on the precipitation behaviors of an Al–4Mg-0.3Cu alloy, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.073. 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.
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Influence of deformation microstructure on the precipitation
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behaviors of an Al–4Mg-0.3Cu alloy
Xiaohui Yanga, Kai Lia, Xianghai Anb, Song Nia,*, Weifeng Weia, Yong Dua, Min Songa,c,** a
School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia
Shenzhen Research Institute, Central South University, Shenzhen 518057, China
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c
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b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
Abstract
To provide insight into the effect of deformation microstructure on the precipitation behavior, an
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Al-4Mg-0.3Cu alloy was processed by high-pressure torsion followed by subsequent in-situ heating in a transmission electron microscope. Structural characterization revealed significant differences in the size, composition and spatial distribution of the precipitates in the alloy subjected to different strains.
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Different grain sizes led to various dislocation substructures and grain boundary volume fractions, which in turn governed precipitation behaviors. In addition, the composition of the precipitates along
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the grain boundaries varied with the grain size due to the change of the grain boundary volume fraction and diffusion behavior.
Keywords: Al-Mg; Cu addition; grain boundary; precipitation; in-situ TEM
*Corresponding author: [email protected] (S. Ni) ** Corresponding author: [email protected] (M. Song)
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1. Introduction In the last decade severe plastic deformation (SPD) techniques have been carried out on various
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metallic materials to obtain excellent mechanical properties due to the extensive grain refinement [1-4]. Equal channel angular pressing (ECAP) [2] and high-pressure torsion (HPT) [4] are among the most widely used SPD techniques. In addition to the grain refinement enhanced mechanical properties,
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SPD-induced segregation and/or precipitation of solute elements along the grain boundaries can also
lead to increased strength [5]. For instance, Al-Mg alloy produced by HPT was found to exhibit a very
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high strength, considerably exceeding a prediction based on the Hall–Petch effect for ultrafine grains (UFGs) since the Mg element was segregated along the grain boundaries [6]. It is believed that solute element segregation and/or precipitation during SPD processes are closely related to the formation of a large number of the grain boundaries [7]. As the fast diffusion channels for solute elements, grain
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boundaries play a crucial role in promoting heterogeneous precipitation [8]. Moreover, formation of high-density dislocations during SPD has been verified to enhance the heterogeneous nucleation of precipitates [9]. Thus, understanding the precipitation behaviors of variously deformed microstructures
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after SPD is helpful for further improving the mechanical properties of SPD-processed materials. Recently, abundant studies have been focused on the precipitation behavior of age-hardenable Al
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alloys after SPD processing, including Al-Mg-Si [10], Al-Cu [11], and Al-Zn-Mg-Cu alloys [12]. SPD of non-age-hardenable Al-Mg alloys leads to the segregation of the alloying elements along the grain boundaries that also contributes to the alloy’s hardening [6]. It should be noted that addition of small
amount of Cu to Al-Mg alloys improves the mechanical properties and inhibits the formation of Lüders bands [13]. Therefore, an investigation of the segregation and/or precipitation behaviors of the supersaturated Al-Mg alloys with minor Cu is a key issue for practical applications. Up to now, the
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segregation and precipitation behaviors of Al-Mg alloy with minor Cu after SPD received little attention in comparison to the most studied Al-Mg-Cu alloy [14, 15]. Therefore, in this work an
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Al-4Mg-0.3Cu alloy processed by HPT was used to investigate the influence of SPD on the segregation and precipitation behaviors during heat treatment.
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2. Experimental
The Al-4Mg-0.3Cu (wt.%) alloy used in this study was fabricated by ingot metallurgy, in which
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commercially pure Al, pure Mg (both are 99.9% in purity) and an Al-Cu master alloy were used as the raw materials. The homogenization treatment of the alloy was at 733 K for 10 hours. Specimens for HPT processing were sectioned into disks from the alloy with a diameter slightly less than 20 mm and a thickness of ~2 mm. These disks were polished on both sides using sand papers (400-2000 grades) until
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their thickness reached ~1.5 mm. HPT processing was conducted using a quasi-constrained HPT facility [3]. Disks were subjected to HPT for 5 turns at room temperature under an applied pressure of 4 GPa and a rotation rate of 1 turn per minute. Samples for transmission electron microscopy (TEM)
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investigation were prepared by a twin-jet polishing technique in a mixture of 30% nitric acid and 70% carbinol at -30 oC. Heat treatment was implemented by an in situ heating technique using a model 901
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hot stage controller in a Tecnai G2 20 S-Twin TEM with a temperature ramping rate of 60 oC per min.
After heat treatment, scanning TEM (STEM) observation of the microstructures was carried out in a Titan G2 60-300 microscope operated at 300 kV that provides Z-contrast to exhibit local composition variation. Elemental mapping of the precipitates and/or segregation was performed using energy dispersive X-ray spectroscopy (EDS) with a Super-X detector. To facilitate discussion, the specimen locations for TEM observation are labeled as regions I, II and III, as shown in Figure 1. The effective
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strain calculated from the formula
ε = 2πNr / 3h (where N is the number of HPT turns, h is the
thickness of the disk and r is the distance from the center of the disk) [16] was about 0, 90 and 180 at
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the positions I, II and III, respectively. The average grain sizes are derived by statistical measurements of 50 grains in region I, 100 grains in both regions II and III using TEM. Although the strain of the disk center is theoretically zero, the microstructure can be refined with continuous HPT deformation based
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on strain gradient effects [17, 18].
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3. Results
Optical microscopy observation shows that the average grain size of the Al-4Mg-0.3Cu alloy is about 400 µm. Figure 2 presents typical bright-field TEM images of (a) the alloy and the regions (b) I, (c) II and (d) III of the sample after HPT processing. Few dislocations can be observed in the TEM
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bright field image from the alloy (Fig. 2a). In contrast, many dislocation cells with an average size of ~ 0.6 µm formed in the region I of the deformed sample (Fig. 2b). With increasing the plastic strain, the average grain sizes are further reduced to ~ 200 nm and ~ 130 nm in the regions II and III, respectively,
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as shown in Figs. 2c and 2d. Extensive TEM observations indicated that, with increasing the strain, the average grain size decreased monotonically. Besides, the alloy was initially deformed by dislocations
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multiplication and slip, so the dislocation density increases. As the deformation proceeded, dislocations
slipped towards and annihilated at cell boundaries, making the dislocation cells evolve to subgrains with low angle grain boundaries or even high angle grain boundaries, as shown in Fig.2c and Fig.2d. In this deformation stage, dislocation density decreases. In fact, with the grain being refined to sub-micrometer size or even to nanometer size, the dislocation density in the grain interior decreased
gradually due to two factors. On one hand, the active stress of Frank-Read dislocation source is beyond
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the flow stress. On the other hand, dislocations were effectively dissolved by the grain boundaries. These dislocation density evolution behaviors have also been observed and verified in ECAPed Al and
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Cu [19]. Figure 3 shows typical STEM micrographs of (a) the alloy heated at 200 oC for 30 min and regions
(b) I, (c) II and (d) III of the HPT treated sample after heating at 200 oC for 5 min. The rod-like
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precipitates with various lengths from tens to hundreds of nanometers were observed in the grain interior of the sample (Fig. 3a). However, Fig.3b shows that rod-like precipitates with smaller size
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(arrowed) formed in the grain interior and irregular shaped precipitates were located along the grain boundaries in the region I after in-situ heating. Dislocations are favorable nucleation sites for the precipitates and potentially provide fast diffusion paths for solute atoms [20, 21]. So, precipitates in the grain interior are associated with dislocations formed in the region I. In contrast, only irregular
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precipitates are seen in the regions II and III. In the region II (Fig. 3c), most precipitates were located along the grain boundaries, while a few precipitates presented in the grain interior. In the region III (Fig. 3d), almost all precipitates are located along the grain boundaries. Obviously, the morphology and
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distribution of the precipitates are closely related to the grain size and dislocation density, which is significantly affected by the imposed plastic strain during HPT. It can be seen that precipitates along
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the grain boundaries were heterogeneously distributed. The precipitate size was calculated by averaging the values from measuring the precipitate along two mutually perpendicular directions of the precipitate. By statistical measurement about 150 precipitates along the grain boundaries in region I, II and III using TEM, the size varied in the range from 20 nm to 60 nm. In addition, the average grain sizes in regions I, II and III after in-situ heat treatment are 0.6 µm, 340 nm and 180 nm by statistical measurement of about 50 grains, respectively. Therefore, the fine and uniformly distributed precipitates
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along the grain boundaries and at the triple junctions can slow down the growth of the UFGs. Apart from the shape, size and distribution of the precipitates being affected by different deformed
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microstructures, their chemical composition and corresponding structures are also significantly affected. Figure 4a shows a HRTEM image of the precipitates in the grain interior in the sample after in-situ heat treatment at 200 oC for 30 min. The fast Fourier transform (FFT) pattern from the region marked by a
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white box in Fig.4a demonstrates that the precipitate Al2CuMg (S phase, with the space group of
Cmcm) follows crystallographic orientation relationship [100]S // [100]Al with the Al matrix [22, 23].
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Figure 5 shows a STEM image and the corresponding EDS mapping (b, c and d) of a precipitate along the grain boundary from the region II. The EDS mapping shows the enrichment of Mg and Cu elements in the precipitate. Figs.6a is the HRTEM images of precipitates along the grain boundary from the region II, and Fig.6b is the corresponding FFT pattern from the region marked by a white box in
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Figs.6a, indicating that the precipitates are the Al2CuMg phase. Analysis of ~ 20 precipitates along the grain boundaries from the region II found only the Al2CuMg phase. Figure 7 shows a STEM image and the corresponding EDS mapping (b, c and d) of the
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precipitates at a grain boundary from the region III. EDS mapping of precipitates 1 and 2 showed the lean and enrichment of Mg element, respectively. Figs.8a and 8d show the precipitates along the grain
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boundaries in the region III. The HRTEM images and FFT pattern shows that the precipitate in Fig.8a is the Al2MgCu phase, while the precipitate in Fig.8d is Al2Cu phase (with the space group of I4/mcm) [24]. Extensive analysis of the precipitates along the grain boundaries in the region III indicated two phases: Al2CuMg and Al2Cu. In order to understand the precipitation evolution along the grain
boundaries in detail, the STEM images of the region III of the sample before and after in-situ heating at 100 oC for 5 min are shown in Fig.9a and Fig.9c. To confirm that these bright segregation along the
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grain boundaries are Cu-rich, line scanning were performed along the arrows plotted in Fig.9a and Fig.9c. The EDS line-scanning in Fig.9b and Fig.9d clearly demonstrates that compared to those in the
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grain interior, the Cu content is higher, the Al content is lower and the Mg content remains almost constant at the grain boundary. The above phenomena about precipitates along the grain boundaries can thus be concluded as follows: (i) Cu diffuses to grain boundaries of the UFGs and the average Cu
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concentration along the grain boundaries is higher than that in the grain interiors; (ii) the grain
boundary segregation of Cu induced by HPT processing or in situ heat treatment in UFGs is more rapid
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than that of Mg.
4. Discussion
The results presented in this study indicate that significant difference exists in the size, composition
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and spatial distribution of the precipitates in Al-4Mg-0.3Cu alloys with various microstructures induced by HPT. Difference in grain size, stemmed from different strain, was noted to affect the dislocation
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density and the grain boundary volume, which in turn governed the evolution of the precipitates.
4.1. Effect of grain refinement on the precipitation behavior
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The present TEM results show that dislocation density in the grain interior changes significantly
during the grain refinement process. From Fig.2, the dislocation density in the grain interior increases firstly, and then decreases with decreasing the grain size. A large number of the dislocation lines were
generated in the grain interior in the region I, with an average grain size of about 0.6 µm (see Fig. 2b). Fig.2c shows that only a small number of the dislocations can be observed in the region II. Fig.2d reveals that the almost no dislocation exists in the UFGs with a size of ~130 nm from the region III.
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This observation is in consistent with the results from numerous previous studies [25-27], which showed that conventional dislocation mechanisms, e.g. Frank-Read source and dislocation
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accumulation, are inhibited when the grain size is close to 100 nm [28]. Hayes et al. also reported that there is very limited dislocation activity in the grain interior of cryomilled UFG pure Al and Al 5083 alloy [29]. Due to the relatively small interaction volume in the UFG interiors, dislocations are difficult
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to accumulate within the grains in the order of 100-200 nm. Moreover, in nanocrystalline grains or
UFGs, grain boundaries act as sources and sinks for dislocations, and units or partial dislocations are
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emitted from one grain boundary, traverse the grain and finally are absorbed into the opposite grain boundary without multiplying and interacting with other dislocations. Another feature of Al-4Mg-0.3Cu alloy processed by HPT is that the grain boundary volume increases with the grain refinement. Previous study has demonstrated that short-circuit self-diffusion is greatly enhanced in
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ultrafine-grained Al alloy prepared by ECAP in comparison to that in the well-annealed coarse-grained material [30]. The diffusion enhancement was explained in terms of a specific deformation-induced state of general high-angle grain boundaries. The volume diffusion coefficient of solute in matrix can
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be written as:
D = Do exp(−Q / RT)
D o = 1.49 × 10−5 m 2 ⋅ S−1 and
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where
Q = 120.5kJ ⋅ mol −1 for
(1) Mg
in
Al
[31],
and
D o = 6.54 × 10−5 m 2 ⋅ S−1 and Q = 136 KJ ⋅ mol −1 for Cu in Al [32]. The grain boundary diffusion
coefficients of Mg [33] and Cu [32] in Al can be two orders of magnitude larger than the volume diffusion coefficients of Mg and Cu in Al. In addition, HPT process is known to increase the concentrations of vacancies and grain boundary volume and this consequently increases the diffusion rate of solute atoms within the matrix [34]. Both the contribution of the dislocations in the grain
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interiors and deformation-induced grain boundaries are considered to account for the size and spatial distribution of the precipitates. Firstly, dislocations with a high density provide numerous
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heterogeneous nucleation sites for the precipitates in the region I, which eventually facilitates the formation of a large number of the precipitates near the dislocations. The dislocations serve as fast
diffusion paths for solute atoms in the matrix, thereby permitting rapid diffusion to the precipitates to
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assist their growth. Secondly, the grain boundary diffusion is prominent due to a large number of the grain boundaries volume characterized by UFGs. The fast diffusion of Mg and Cu elements along the
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boundaries in ultrafine-grained Al-4Mg-0.3Cu alloy implies that less solute will be available in the matrix for precipitation. Thus, precipitates inside grains decrease with refining the grain size. In addition, the size of the precipitates along the grain boundaries becomes larger with decreasing the grain size after the same heat treatment due to fast grain boundary diffusion with increasing the grain
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boundary volume.
4.2 The precipitate phase evolution along the grain boundaries in UFGs
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The precipitation may occur preferentially along the grain boundaries rather than in the matrix. Firstly, the free energy barrier for nucleation along the grain boundaries is lower than that in the matrix
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because that the precipitate/matrix interface can be created by a relatively small change in the grain
boundary structure. Secondly, the precipitation nucleation rate along the grain boundaries is higher than that in the matrix because the enhanced solute diffusivity along the grain boundaries increase the critical nucleus rate [35]. Therefore, the number of the precipitates at grain boundaries increases rapidly with decreasing the grain size. Additional factor that could accelerate nucleation includes a higher Cu content at the grain boundaries, since forced migration of boundaries during deformation can result in
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stronger Cu segregation, which would increase the driving force and local availability of solute [36]. Compared to that in regions I and II, the mean spacing of the precipitates along the grain boundaries
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decreases in region III. The “collector plate mechanism” can well account for the precipitate growing process along the grain boundaries [37]. The “collector plate mechanism” is easily understood, based on the fact that the Mg and Cu solute atoms migrate to the grain boundaries through volume diffusion,
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and then swiftly diffuses along the boundaries to the edges of the precipitate [37]. This is due to the fact that the grain boundary diffusivity is in orders of magnitude higher than bulk diffusivity. In this case,
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the mass flow of Mg and Cu to each grain boundary, “collector plate”, can be rapidly involved into the growth of the precipitates. Aaron and Aaronson demonstrated that the grain boundary precipitate lengthening and thickening vary as a function of (time)1/4 and (time)1/2 [37], respectively. However, the precipitates at the grain boundaries in region III show spherical shape and the growth of the precipitates
the precipitates.
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does not follow the rule of Aaron and Aaronson [37], which might be caused by to the coarsening of
Both experiments and atomistic modeling showed that segregation can influence grain boundary
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precipitation by altering the driving force for nucleation and the solute flux arriving at the grain boundaries [36, 38]. In this work, the Cu element shows strongly segregation along the grain
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boundaries than Mg element subjected to HPT processing. The segregation thickness at the grain boundaries of region III after in-situ heat treatment increases obviously, as shown in Fig.9. Similarly, ECAPed Al-Zn-Mg-Cu alloy at 200 oC also demonstrated that Mg and Cu elements segregated strongly
to the grain boundaries in UFGs [39]. The significant segregation of Cu element along the boundaries in UFGs after heat treatment implies that more Cu will be available along the grain boundaries for precipitation. In addition, precipitation in Al–Cu–Mg alloys can be changed by adjusting the Cu/Mg
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atomic ratio [40]. With grain refinement, Cu element diffusing more rapidly to the grain boundaries is attributed to the shorter diffusion paths. So, Al2Cu phase can be formed at the local regions of the grain
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boundaries from region III during heat treatment process due to the high Cu/Mg atomic ratio. In all, the precipitates along the grain boundaries contain Al, Mg and Cu elements in the initial stage. When grain
size is refined to nanoscale, the precipitates along the grain boundaries are composed of Al and Cu
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elements or Al, Mg and Cu elements due to the rapidly diffusion.
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5. Conclusions
Summarizing the results of the precipitation behaviors in supersaturated Al-4Mg alloy with minor Cu element subjected to HPT and in-situ heat treatment in TEM, it is indicated that the dislocations in the grain interior and grain boundaries jointly contribute to the precipitation mechanisms. At low strain,
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precipitates mostly locate in the grain interior due to the generation of high-density dislocations in the grain interior. When grain size is refined to nanoscale at high strain, the precipitates mostly locate along the grain boundaries due to the increase in the grain boundary volume. The size of the precipitates
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along the grain boundaries increases with refining the grain size due to the diffusion enhancement. In addition, the composition of precipitate phases along the grain boundaries changes with grain
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refinement.
Acknowledgements
The financial supports from National Natural Science Foundation of China (51531009 and 51501230), Grants from Shenzhen Science and Technology Project (JCYJ20140509142357196) and the outstanding graduate project of Advanced Non-ferrous Metal Structural Materials and
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Manufacturing Collaborative Innovation Center are appreciated.
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Figure captions: Figure 1: The regions cut from the disks for TEM observations, i.e. the center region at r=0 marked as I,
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and the regions at r=4 mm and r=8 mm marked as II and III, respectively. Figure 2: TEM images of (a) the alloy, and the alloy subjected to HPT for 5 turns at the (b) region I, (c) region II and (d) region III, respectively.
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Figure 3: HAADF STEM images of (a) the alloy after in situ heat treatment at 200 oC for 30 min, and the (b) region I, (c) region II and (d) region III of the alloy subjected to HPT for 5 turns and
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then in situ heat treated at 200 oC for 5 min. Figure 4: (a) A HRTEM image of the precipitates in the
alloy after in situ heat treatment and (b) the
corresponding FFT pattern from the region marked by a white box in Fig.4a, showing that the precipitate is the Al2CuMg phase.
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Figure 5: (a) A STEM image of a precipitate along the grain boundary from the region II after in situ heat treatment at 200 oC for 5 min and (b-d) EDS mapping of Al, Cu and Mg elements of the precipitate.
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Figure 6: (a) HRTEM images of precipitates along the grain boundary and (b) the corresponding FFT pattern from the region marked by a white box in Fig.6a, showing that the precipitate is the
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Al2CuMg phase.
Figure 7: (a) A STEM image of the precipitates along the grain boundary from the region III and (b-d) EDS mapping of Al, Cu and Mg elements obtained from the white square in Fig.7a, showing that the precipitate 1 has no Mg concentration while the precipitate 2 contains a high concentration of Mg. Figure 8: (a,d) bright field TEM images of grain boundaries from the region III after in situ heat
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treatment at 200 oC for 5 min, (b,e) HRTEM images of the precipitates indicated by black arrow in (a) and (d), respectively, and (c,f) the corresponding FFT patterns from the region
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marked by white boxes in (b) and (e), respectively, showing that the two precipitates are the AlMgCu phase and Al2Cu phase, respectively.
Figure 9: The STEM images of the region III of the sample (a) before and (c) after in-situ heating at
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100 oC for 5 min. (b) and (d) the corresponding EDS line-scanning profiles of Al, Cu and Mg
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elements along the yellow arrow in (a) and (c), respectively.
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Research highlights
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1. Precipitation of an HPT deformed Al-4Mg-0.3Cu alloy was investigated. 2. Precipitation behavior is closely related to the dislocations and grain boundaries.
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3. Precipitates along the grain boundaries change with decreasing the grain size.