Fine precipitation scenarios of AlZnMg(Cu) alloys revealed by advanced atomic-resolution electron microscopy study Part I: Structure determination of the precipitates in AlZnMg(Cu) alloys

Materials Characterization 99 (2015) 277–286

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Materials Characterization journal homepage: www.elsevier.com/locate/matchar

Fine precipitation scenarios of AlZnMg(Cu) alloys revealed by advanced atomic-resolution electron microscopy study Part I: Structure determination of the precipitates in AlZnMg(Cu) alloys☆ J.Z. Liu a,b,c, J.H. Chen a,b,⁎, D.W. Yuan a, C.L. Wu a, J. Zhu c, Z.Y. Cheng c a b c

Center for High-Resolution Electron Microscopy, College of Materials Science & Engineering, Hunan University, Changsha, Hunan 410082, China Hunan Province Key Laboratory for Spray Deposition Technology and Application, Hunan University, Changsha, Hunan 410082, China Beijing National Center for Electron Microscopy, Department of Materials Science & Engineering, Tsinghua University, Beijing 100084, China

a r t i c l e

i n f o

Article history: Received 4 October 2014 Accepted 21 November 2014 Available online 14 January 2015 Keywords: Electron microscopy Aluminum alloys Al–Zn–Mg–Cu alloys Precipitation Aging Phase transformation

a b s t r a c t Although they are among the most important precipitation-hardened materials for industry applications, the high-strength AlZnMg(Cu) alloys have thus far not yet been understood adequately about their underlying precipitation scenarios in relation with the properties. This is partly due to the fact that the structures of a number of different precipitates involved in electron microscopy in association with quantitative image simulations have to be employed; a systematic study of these hardening precipitates in different alloys is also necessary. In Part I of the present study, it is shown that there are five types of structurally different precipitates including the equilibrium η-phase precipitate. Using two state-of-the-art atomic-resolution imaging techniques in electron microscopy in association with quantitative image simulations, we have determined and clarified all the unknown precipitate structures. It is demonstrated that atomic-resolution imaging can directly suggest approximate structure models, whereas quantitative image analysis can refine the structure details that are much smaller than the resolution of the microscope. This combination is crucially important for solving the difficult structure problems of the strengthening precipitates in AlZnMg(Cu) alloys. © 2015 Elsevier Inc. All rights reserved.

1. Introduction Since the 1930s, many structural components of aircrafts and spacevehicles have been made of aluminum alloys that can provide the strength of steel at less than half the weight (i.e., a high strength-toweight ratio) [1,2]. With proper thermal processes the AlZnMg(Cu) alloys with high Zn(Cu)-contents (up to 8%) can achieve an ultrahigh strength (650 MPa to 800 MPa) and other combined properties required for aerospace applications [3], largely owing to the formation of a large number of very fine disc-shaped aggregates or precipitates of the alloying elements in the Al-grains [4–8]. These precipitates concordantly resist the grain deformation upon loaded stresses, therefore strengthening the material. This phenomenon is known as precipitation hardening [9]. It has been demonstrated for an AlZnMg(Cu) alloy that precipitate hardening contributes 66% to 74% of the total strength in comparison with the sum of solid solution strengthening and grain-size strengthening [10]. Hence understanding as precise as possible of the natures of these

☆ The Part 2 of this article entitled "Fine precipitation scenarios of AlZnMg(Cu) alloys revealed by advanced atomic-resolution electron microscopy study Part II: Fine precipitation scenarios in AlZnMg(Cu) alloys" can be referred in the same issue on pages 142 to 149. ⁎ Corresponding author at: Center for High-Resolution Electron Microscopy, College of Materials Science & Engineering, Hunan University, Changsha, Hunan 410082, China. E-mail address: [email protected] (J.H. Chen).

http://dx.doi.org/10.1016/j.matchar.2014.11.028 1044-5803/© 2015 Elsevier Inc. All rights reserved.

precipitates has long been a serious study in physical metallurgy, in order to improve the properties of AlZnMg(Cu) alloys through proper manufacturing processes. However, due to their small size in all three dimensions and generally the limitation of suitable characterization tools [11,12], in association with their complexities in formation and in structure evolution, the hardening precipitates in AlZnMg(Cu) alloys have long been very difficult to assess adequately in composition and structure, leading to significant controversies and misleading conclusions about the precipitate types and the precipitation behaviors in this alloy system [7,13–20]. There are fundamental reasons why the current understandings about these hardening precipitates are ambiguous and inconsistent. The key reason is that lacking of powerful characterization tools and quantitative structure analysis methods has long been a technique hamper to understand precisely the natures of the early-stage precipitates in aluminum alloys [11,12]. Even now when state-of-the-art aberrationcorrected high-resolution transmission electron microscopy (HRTEM) and aberration-corrected scanning transmission electron microscopy (STEM) become available, to precisely determine the structures of such small precipitates is still a difficult task, since this requires quantitative image analysis and refinement. Characterized by X-ray diffraction (XRD), conventional transmission electron microscopy (TEM) associated with selected-area electron diffraction (SAED), and atom-probe tomography (APT), the

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precipitates in AlZnMg(Cu) alloys have long been believed to form in the following sequence upon thermal aging: super-saturated solid solution (SSSS) → GP → η′ → (ηp→) η [4,6,7,13–16,18–21], where the GP stands for the Guinier–Preston zones and the ηp-precipitates denote a recently identified hardening phase [7]. However, this precipitation sequence is still in debates [13,15,19,20]. Actually, examining the precipitates without using quantitative image analysis could be insufficient to determine their structures and to distinguish the difference between these precipitates, as shown in detail in the present study. The current technical situation is that a quantitative image simulation refinement procedure for determining an unknown structure is much less time-consuming for HRTEM than for STEM. Hence making use of the advantages of both atomic-resolution imaging techniques in HRTEM and in STEM and performing quantitative image simulation refinements can be more effective to solve the difficult and complex problems that we are facing about the atomic-scale metallurgy of AlZnMg(Cu) alloys. In the present study, in order to have a complete view about the hardening precipitates in the high-strength AlZnMg(Cu) alloys with high Zn(Cu)-contents, a series of alloys with varying Zn(Cu)-to-Mg atomic ratios have been systematically investigated. Two state-of-theart atomic-resolution imaging techniques in advanced HR(S)TEM in association with quantitative image simulations and first principles energy calculations were employed to classify the precipitates and to determine all their structures. It is shown that a consistent understanding of the complex precipitation behaviors in AlZnMg(Cu) alloys with high Zn(Cu)-to-Mg atomic ratios can be achieved. To accomplish such a goal, the present work consists of two parts (or two steps), i.e., Part I: Structure determination of the precipitates in AlZnMg(Cu) alloys, and Part II: Fine precipitation scenarios in AlZnMg(Cu) alloys. In Part I, we focus on structure determination for all the precipitates observed in these alloys using advanced HR(S)TEM with emphasizing the methods employed. This part of the study demonstrates the following major results. (1) The quantitatively determined ηp-structure indicates that the η p -precipitates are structurally distinguishable from the equilibrium η-phase and from the η′-precipitates. (2) By determined their structures, the GP zones in the alloys are classified accurately. (3) Atomic-resolution imaging in STEM, on the one hand, can provide straightforward chemistry and structure information about the hardening-precipitates in the first step to well distinguish their subtle differences and to propose their reasonable structure models at the atomicscale. Atomic-resolution imaging in HRTEM with rapid quantitative image simulation analysis, on the other hand, can provide the refined structures of the precipitates with more precise structure information about the precipitates in the second step to quantitatively explain the contrast variations, which may not be intuitive but can reflect fine defective and compositional changes in the atomic columns and to determine structure models with high precision beyond the resolution limitation of the microscopes.

2. Materials and methods 2.1. Alloys and thermal processes A typical commercial 7150 alloy with the composition of Al–6.48Zn– 1.57Mg–2.17Cu (wt.%) has been prepared for this part of the study. The alloy plates of about 20 mm in thickness were homogenized at 475 °C before cutting into suitable samples for age-hardening heattreatments and for mechanical property measurements. Prior to the hardening age, the alloys were solution heat-treated at 470 °C for 1 h and then water quenched to room temperature. Subsequent thermal aging was carried out in an oil bath. Series of the alloy samples being aged for periods of 0–120 h at 120 °C or 150 °C were prepared for investigations. Specimens for STEM and HRTEM observations were prepared first by mechanical polishing and then by electro-polishing until perforation.

2.2. Atomic-resolution imaging and quantitative image analysis in HRTEM 2.2.1. Electron microscopy instruments Three modern TEM instruments were employed in the present study in order to comprehensively investigate the precipitation behaviors of these widely used important alloys. Firstly, a FEI Tecnai F20 HRTEM instrument at Hunan University was employed for the general routine observations of various precipitates in the alloys. This microscope, operating at 200 kV and in the high angle annular dark field STEM (HAADFSTEM) imaging mode with a half-angle of 12 mrad for probe convergence and a collection inner semi-angle of 36 mrad, can achieve a point resolution of 0.16 nm at best. In the HRTEM mode the microscope has an information limit (resolution) of 0.14 nm [22,23]. Secondly, in order to resolve the fine structure details smaller than 0.14 nm in the precipitate structures, a FEI Titan 80–300 Cs-corrected HRTEM instrument at Tsinghua University was used to record through-focus (TF) series of HRTEM images of precipitates for exit-wavefunction reconstruction (TFEWR). This state-of-the-art microscope, with an aberration-corrector implemented below its objective lens for imaging, enables an (resolution) information transfer of down to 0.07 nm at 300 kV in the TEM mode. Thirdly, a FEI Titan ChemiSTEM instrument at Zhejiang University was also used to resolve the structure details smaller than 0.14 nm in the precipitates. The microscopy is Cs-corrected for the probe-forming lens and has a resolution of 0.08 nm in the HAADF-STEM imaging mode with a half-angle of 21.4 mrad for probe convergence and a collection inner semi-angle of 53 mrad. 2.2.2. HAAD-STEM imaging In the HAADF-STEM imaging mode, the obtained image shows a Zcontrast (where Z represents the atomic numbers of the atoms in the specimen) proportional to Z1.7 ∼ 2.0 [24–26]. Hence the HAADF image contrast is relatively easy to interpret in terms of the structure, so long as the composition of the observing material is known, though the light atoms are more difficult to be visualized than the heavy atoms when they are largely different in Z and co-exist in the structure. 2.2.3. TF-EWR imaging in HRTEM For atomic-resolution imaging by TF-EWR in the HRTEM imaging mode, three necessary steps are required to obtain the right images. Firstly each focus-variation series of about 20 images of the observing object can be recorded automatically with the FEI Titan 80–300 Cscorrected HRTEM instrument. Secondly the image series is processed by a FEI software called TrueImage [27,28] to recover the complex wavefunction at the image-plane. Thirdly, the aberrations in the obtained wavefunction such as defocus, coma, two-fold astigmatism as well as three-fold astigmatism, should be corrected in order to retrieve the final right wavefunction below the specimen, which normally is a difficult step in the procedure and requires quite a lot of experiences in practice [27–30]. This step will become much easier if an aberration-corrected HRTEM instrument is used for recording the image series [31]. After removing the lens aberrations, the phase-map of the obtained wavefunction shows an atomic-resolution image, whose contrast reflects the projected atomic potential of the observing object, so long as the specimen thickness is sufficiently thin (to roughly meet the phase-object approximation) [12,32]. 2.2.4. Quantitative image simulation analysis Image simulations were performed to verify the atomic-resolution images obtained by TF-EWR, and to refine the obtained structure models, using the MacTempas image simulation package [33]. If the structure model proposed by atomic-resolution imaging was right, the simulated wavefunction would well match the one experimentally retrieved by TF-EWR. In image simulations, super-cells were used in order to include both the Al-matrix and the precipitates. For decades simulations of images, exit-wave functions and electron diffraction patterns have been a standard tool in HRTEM [34–38]. Especially for

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determining the defect structure inside a known crystal (i.e., an internal standard exists), the image simulation analysis can be performed quantitatively and therefore fine structural details much smaller than the point resolution of the microscope can be revealed [12,22,23,31,39].

plane. It has been difficult for conventional HRTEM and other techniques such as XRD, SAED and 3DAP to answer the following questions: How many different GP zones exist and how do they transform from one to another?

3. Results

3.1.2. HAADF-STEM observations Fig. 2 shows the HAADF-STEM images of typical hardening precipitates in the 7150 alloy in comparison with their HRTEM images, viewed along a b112NAl direction. It is seen form Fig. 2e–h that due to the image delocalization effect in HRTEM [27,28], the exact thicknesses (or numbers of the atomic layers) of these disc-like precipitates cannot be measured for sure in non-aberration-corrected HRTEM. In contrast, the differences among these precipitates are easily seen in HAADF-STEM (even without aberration-correction for the scanning electron probe), as shown in Fig. 2a–d. It can be further characterized by HAADF-STEM observations that in the 7150 alloy there are characteristically 4 types of strengthening precipitates upon thermal aging, apart from the large equilibrium η-phase particles (Fig. 1 g). The first type is the ηp-precipitates with three features easily seen in HAADF-STEM: (i) with a thickness of at least 11-atomic-layers in the {111}Al planes (Fig. 2d), (ii) with its lattice parameter c = 4dAl b 111N approximately, and (iii) always with an interface dislocation loop, i.e., 11 of precipitate atomic-layers are matched by 10 of {111}Al atomic-layers, as shown in Fig. 3a (and more clearly shown in Fig. 4). The second type is the long-time well-known η′-precipitates distinguishable easily by their lattice parameter c = 6dAl b 111N (Fig. 2c and g). The third type of these disc-like precipitates is the first type of GP zones distinguished by the following features (Figs. 2b and 3c and d): (i) with a thickness of 7-atomic-layers in the {111}Al planes, (ii) with a variable concentration of alloying elements, and (iii) with two atomic “panels” as their stable skeleton for changing continuously in composition and structure. In Part II of the present study, it will further be shown convincingly that these GP zones are the early-stage precipitates or the precursors of the ηp-precipitates and therefore they are referred to the GP-ηp zones hereafter. The fourth type of precipitates (the second type of GP zones) observed is the GP zones or the “cloudy” clusters

3.1. Morphology and crystallography of the precipitates 3.1.1. HRTEM observations Fig. 1a–b shows the overviews of the disc-like hardening precipitates in the 7150 AlZnMg(Cu) alloy, observed in TEM from a b112NAl direction and a b110NAl direction, respectively, which are the two best viewing directions for these precipitates in order to obtain their edgeon images, as schematically illustrated in Fig. 1c. Fig. 1d–g are the HRTEM images of 4 types of different particles that are distinguishable in conventional HRTEM for their characteristic features in crystallography. Fig. 1e shows the well-known η′-precipitate distinguishable from the others by its lattice parameter c = 6dAl b 111N approximately [15], whereas Fig. 1f demonstrates the less known ηp, or η-precursor precipitate distinguishable from the others by its lattice parameter c = 4dAl b 111N approximately [7]. The typical equilibrium η-phase particle of large size is shown in Fig. 1g, which is formed typically in the very later stage of thermal aging and usually do not have a fixed orientation relationship with the matrix when grown into a large 3D nano-crystal. As mentioned above, two of the confusion points in the current understandings about these precipitates are the following. On the one hand, the difference between the ηp-precipitates and the η′-precipitates is not recognized [19], and their relation has then not been discussed [7, 15]. On the other hand, without using quantitative image simulations to precisely determine the ηp-structure, the difference between the ηp-precipitates and the equilibrium η-phase precipitates has not been realized [20]. The reasons for these are objectively complex and shall be elucidated in the present study in steps. Fig. 1d shows the typical early-stage precipitates in the alloys, the GP zones [16], formed prior to the others and featured by lacking of periodicity in the third direction, i.e., the direction normal to their lying {111}Al

Fig. 1. The morphology and crystallography of precipitates in Al–Zn–Mg–(Cu) alloys. a and b, the bright-field images of the precipitates in 7150 alloys viewed along a b112NAl and a b110NAl directions, respectively. c, a schematic three-dimensional drawing illustrating that all the disc-like precipitates shown in a and b lie on the {111}Al planes in the Al-matrix, and some appear edge-on and the others appear about of round-shape in the [110]Al and [112]Al projections. d–g, HRTEM images of the GP-zones (d), η′-precipitate (e), ηp-precipitate (f) and equilibrium η-phase precipitate (g) in the alloy. The inset in g shows the Fourier transformation pattern of the η-particle.

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Fig. 2. The HAADF-STEM images (top) and HRTEM images (bottom) of 4 types of hardening precipitates observed in a 7150 alloy, illustrated in comparison, a vs. e, etc.

with a variable thickness and without clear features in comparison with the GP-ηp zones (Fig. 2a and e). It will be shown that these GP zones are related to the η′-precipitates and therefore they are referred to the GP-η′ zones hereafter. According to the morphological and crystallographic features of the observed precipitates, it is clear that upon thermal aging 5 types of precipitates may form in the 7150 alloy, i.e., the GP-η′ zones, the GP-ηp zones, the η′-precipitates, the ηp-precipitates and the large equilibrium η-phase particles. Nonetheless, to understand the precipitate evolution and transformation occurring among these different particles, their structures have to be known as precise as possible. This is still a difficult

task and requires not only state-of-the-art atomic-resolution imaging techniques but also quantitative image analysis, as well as first principle energy calculations, as shown in the present study by steps. 3.2. Structure determinations of the precipitates by atomic-resolution imaging and quantitative image simulation analysis Since the equilibrium η-phase (MgZn2) has a well-known Laves phase structure [40], we began with structure determination for the ηp-precipitates, using both atomic-resolution imaging in aberrationcorrected HAADF-STEM and atomic-resolution imaging in aberration-

Fig. 3. HAADF-STEM images of precipitates. a and b, a ηP-precipitates viewed along the [112]Al direction and the [110]Al direction, respectively. c and d, the 7-atomic-layers GP-zones in different stages viewed along the [112]Al direction.

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Fig. 4. Atomic-resolution HAADF-STEM images of two 11-atomic-layer ηp-particles, each with an interface dislocation loop, viewed along the [112]Al direction, a, with a perfect stacking order, b, with a disordered stacking order. It is marked in (a) that there are two types of Zn-columns, ZnI and ZnII, distinguishable in brightness in the structures.

corrected HRTEM. The former imaging technique is more sensitive for detecting the Zn(Cu)-atoms in the structure, and the later one is more sensitive for detecting the Mg/Al atoms in the structure and more suitable for structure refinement by repeated image simulations. 3.2.1. Structure of the ηp-precipitates Fig. 4 shows aberration-corrected HAADF-STEM images of two typical ηp-precipitates, both of which have a thickness of 11-atomic-layers and an interface dislocation loop. Analyzing the two atomic-columnresolved images, the following points appear. (i) The ηp-precipitate shown in Fig. 4a is a single crystalline particle with two identical unit cells in the thickness direction, whereas the one shown in Fig. 4b has a disordered stacking sequence in its thickness direction. Nonetheless,

the basic “atomic building blocks” in the two structures are the same. (ii) The positions of Zn(Cu) atoms in the structures are clearly revealed, whereas Al/Mg atoms in the structures are not visualized clearly. (iii) There are two types of Zn(Cu)-containing atom columns in each “building block”, and one (marked as ZnI) appear brighter and broader than another (marked as ZnII). Without identifying the precise positions of Mg/Al atoms and analyzing the image contrast quantitatively, the ηp-structure can easily be misinterpreted as the equilibrium η-structure [20], i.e., the Laves phase MgZn2 structure [40]. To precisely locate the Mg/Al positions and therefore determine the ηp-structure quantitatively by image simulation refinement, aberrationcorrected HRTEM was employed. Fig. 5a–b shows two atomic-resolution images of the ηp-precipitates viewed in two independent directions.

Fig. 5. The atomic-resolution TF-EWR phase-images matched with image simulations to reveal the structure of the ηp-precipitate. a and b, the atomic-resolution images of the ηp-precipitates viewed along the [112]Al direction (a) and the [110]Al direction (b), which both well match with the simulated images (the insets) calculated from the refined ηp-structure model. c and d, the atomic-resolution images of the ηp-precipitates viewed along the [112]Al direction (a) and the [110]Al direction (b), which both cannot match with the simulated images (the insets) calculated from the equilibrium η-phase structure model. The inserted line-scan profile in c indicates that the ZnI-columns in the ηp-structure actually consist of two sub Zn-columns with a small distance of about 0.1 nm or less.

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These images were obtained by performing TF-EWR over the TF-HRTEM image series recorded in an aberration-corrected FEI Titan microscope with a highest resolution of 0.8 nm [12,41]. All the atomic-columns in the structure are clearly resolved as bright dots. From the two atomiccolumn-resolved images a three-step procedure was followed to deduce the structure quantitatively. (1) Based on the weak-phase object approximation [31,42], the images themselves have already suggested a rough starting structure model for image simulation refinement [12,29]. The HAADF-STEM images (Fig. 4) also provide important messages for the initial structure model. (2) The model was repeatedly adjusted such that the images calculated from the model match well with the experimentally reconstructed ones (the insets in Fig. 5a–b). The image calculations were performed using the MacTempas image simulation software [33] over a supercell model to include the matrix (as an internal standard). (3) The model is further analyzed by energy calculations to reach an energy minimum, using density functional theory implemented in the Vienna Ab initio Simulation Package (VASP) (see Part II of the present study for the results and analysis of energy calculations). In addition, by image simulations we have also tested whether or not the η p -precipitates would have the equilibrium η-phase structure. Fig. 5c–d clearly demonstrates that different from the previous speculations [19,20], the ηp-precipitates do not have the η-phase (MgZn 2 ) structure but have their own structure, since the simulated images of the η-structure do not match the experimental images at some atomic positions. Through such quantitative analysis of image contrast, a proper atomic structure model was finally obtained for the ηp-structure, as shown by its projections in the [112]Al direction (Fig. 6a) and in the [110]Al direction (Fig. 6b) respectively, as well as by its 3D structure model (Fig. 6c). Analyzing the obtained atomic structure of the ηp-precipitate, the following points appear: (i) the ηp-precipitate still contains a high percentage of Al atoms in its structure; (ii) in the TF-EWR image, opposite to the observation in HAADF-STEM (Fig. 4), the ZnII-columns are brighter than the ZnI-columns, since the later actually consists of two separated sub Zn-columns with a small distance of about 0.107 nm or less, as marked in Fig. 5c; (iii) in the [110] Al projection of the

ηp -structure, different atomic-columns are densely projected and cannot be resolved directly by both the two imaging techniques employed in the present study. Without quantitative simulation analysis of the image contrast, the closely located atomic-columns in the [110] Al projection could not be identified precisely. Table 1 lists the obtained atomic coordinates of the ηp-structure through matching the calculated phase images with the experimentally reconstructed ones. Furthermore, Table 1 lists the crystallographic structure data of the ideal MgZn4Al ηp-crystal, further refined after minimizing the formation energy of the structure by VASP calculations that allowing a full relaxation of both the cell dimension and atomic coordinates from the experimentally refined structure data (see Part II for details of the VASP calculations). Since quantitative image matching analysis was performed through repeatedly adjusting specimen parameters that may influence the image contrast, including specimen thickness and tilt angle, as well as Debye–Waller (DW) factors, inter-atomic distances and compositions of all atomic-columns in the structure, as shown in Table 1. Another important message about the obtained ηp-structure is that when classified by their DW factors and compositions there are three kinds of Zn-atomic-columns in the structure, as marked in Fig. 6a. The Zn1columns are well aligned (indicated by a small DW factor), whereas the Zn2-columns are not well aligned (indicated by a large DW factor); and furthermore the Zn3-columns not only have a large DW factor but also contain vacancies. Similar information is also obtained about the Mg and Al atomic columns. Quantitative analysis shows that there could be up to 50% of Al atoms in the Mg columns and up to 30% vacancies in some of the Zn columns. The large DW factors and mixed atomic-columns suggested by quantitative structure refinement imply that these atomic-columns are arranged slightly in a (randomly) zigzag manner, indicating that the ηp-precipitates were still in change before the alloy was quenched to room temperature for investigation [12]. Hence, the ηp -structure actually has a dynamic composition of Mg1 − xZn4 − yAl1 + x (0 b x b 0.5, 0 b y b 0.3), evolving towards its metastable composition of MgZn4Al (Fig. 6). This also explains why the same atomic-columns in Fig. 5a may locate at slightly

Fig. 6. The refined structure model of the ηp-precipitate. a, projected along the b112NAl direction, in which the marked ZnI and ZnII columns are classified by their brightness in Fig. 5c, whereas the Zn1, Zn2 and Zn3 columns are classified by their Debye–Waller (DW) factors and occupancy compositions in quantitative image simulation analysis; b, projected along the b110NAl direction; c, the 3D structure model.

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Table 1 The atomic coordinates (left) of the ηp-structure obtained through matching the calculated phase images with the experimentally reconstructed ones, and those (right) fully relaxed away from the constraint of aluminum matrix after energy minimization by VASP calculations. The VASP relaxed structure has a hexagonal lattice with a = 5.00 Å, c = 9.27 Å, belonging to the space group P6 (No. 174). Before relaxation the model has a hexagonal unit cell with a = b = 0.496 nm, c = 0.886 nm and α = β = 90°, γ = 120°. In the [112]Al image (Fig. 5a) the optimum matching was obtained for a specimen thickness of 3.5 nm and for a crystal tilt of 1.14° and in the [110]Al image (Fig. 5b) that was achieved for a specimen thickness of 4.3 nm and a crystal tilt of 0°. #Represents that there are 0.2 vacancies in the atomic columns. Atom

Zn1 Zn2 Zn3 Zn4# Zn5 Zn6 Zn7 Zn8# Mg9(Al) Mg10(Al) Al11 Al12

Experiment

VASP relaxed

x/a

y/b

z/c

DW

Occupancy

Site

x/a

y/b

z/c

0.93401 0.72124 0.34475 0.00000 0.27906 0.06735 0.65359 0.00000 2/3 2/3 1/3 1/3

0.27876 0.65525 0.06599 0.00000 0.93265 0.34641 0.72049 0.00000 1/3 1/3 2/3 2/3

0.99600 0.99600 0.99600 0.21146 0.49400 0.49400 0.49400 0.73278 0.23756 0.73995 0.25105 0.72653

10.00 10.00 0.10 1.50 10.00 10.00 0.10 1.50 0.25 0.25 1.00 1.00

1.0 1.0 1.0 0.8 1.0 1.0 1.0 0.8 2/3(1/3) 2/3(1/3) 1.0 1.0

3j 3j 3j 2g 3k 3k 3k 2g 2i 2i 2h 2h

0.921 0.746 1/3 0 0.254 0.079 2/3 0 2/3 2/3 1/3 1/3

0.254 2/3 0.079 0 0.921 1/3 0.746 0 1/3 1/3 2/3 2/3

0 0 0 0.24918 0.500 0.500 0.500 0.75082 0.29097 0.70903 0.20546 0.79436

different position and do not have the same brightness, which approximately reflects the projected potential of the atomic-columns. It can be seen that the Zn-to-Mg atomic ratio for an ideal ηp-precipitate is 4.0. 3.2.2. Structure of the GP-ηp zones Fig. 7 shows atomic-resolution HAADF-STEM images of several typical GP-ηp zones in different stages of evolution. From these observations, the following are clear: (i) the GP-ηp zones are dynamic precipitates during thermal aging, i.e., every GP-ηp zone is different from another, since upon initiation a GP-ηp zone shall proceed to evolve its dimension and structure/composition with its own evolution speed in a path towards a relatively stable or matured structure; (ii) the common feature of these dynamic precipitates is that they all have a “doubleatomic-panel”, which not only has a stable atomic structure but also defines their characteristic thickness of 7-atomic-layers; (iii) the “doubleatomic-panel” serves as a stable skeleton for a GP-ηp zone to evolve and the Al-atoms between the two atomic panels shall be replaced gradually by Zn and Mg atoms in the evolution. Hence, determination of all possible structures of a dynamic GP-ηp zone is endless. Nonetheless, the atomic structure of their “double-atomic-panel” skeleton is well defined. So is its matured structure well defined (Fig. 7d). These two structures can readily be determined by atomic-resolution imaging and quantitative image analysis. From any image in Fig. 7a–d, a structure model for the two atomicpanels of a GP-ηp zone can be obtained. To precisely determine the double-atomic-panel structure of the GP-ηp zones and its matured

structure (Fig. 7c–d) as well quantitatively by image simulation refinement, aberration-corrected HRTEM was also employed. Fig. 8a shows the atomic-resolution image of a matured GP-ηp zone obtained by performing TF-EWR over the HRTEM image series recorded in an aberration-corrected FEI Titan microscope. An initial structure model “read” from the HAADF-STEM image (Fig. 7d) was used for the image simulation refinement on Fig. 8a until good image matching has been achieved, as shown by the inserted simulated image in Fig. 8a. Fig. 8b demonstrates the determined structure of the double-atomic-panel of the GP-ηp zones without visualizing its 3 internal atomic-layers between the two panels, whereas Fig. 8c shows the relatively stable structure of a matured GP-ηp zone. The obtained results indicate that the Znto-Mg atomic ratio of a GP-ηp zone can be as high as 5.0. Fig. 8b reveals that each panel with a composition of MgZn5 consists of one outmost Zn–Mg-layer and one inner Zn-layer. Due to the intrinsic symmetry of the FCC Al-lattice, the two formed atomic-panels are symmetric to each other through an inversion center. Interesting is that such a structure configuration of one outmost Zn–Mg-layer plus one inner Zn-layer for each panel of the GP-ηp zones is also adopted by the ηpprecipitates with 11-atomic-layers serving as their outmost “atomicskin” structure (Fig. 4). This indicates that such an atomic configuration is an energetically favored interface structure between the Al-lattice and the ηp/GP-ηp precipitates. Fig. 8c shows the final or matured structure of the GP-ηp zones proposed by quantitative image analysis. In comparison with the ηpstructure (Fig. 6), it can be seen that (i) the basic “atomic-building

Fig. 7. The atomic-resolution HAADF-STEM images of the 7-atomic-layer GP-ηp zones in their different stages of development with a double-panel structure, viewed along the b112NAl direction. a and b, early-stage GP-ηp zones containing a large portion of Al-atoms; c and d, late-stage GP-ηp zones with more and more alloying element atoms replacing the prior Al-atoms.

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Fig. 8. a, the atomic-resolution TF-EWR phase-image matched with image simulation (the inset) to reveal the atomic structure of the GP-ηp zones viewed along the b112NAl direction; b, the refined atomic structure of the out “panels” of the GP-ηp zones (left) in comparison of the Al matrix (right); c, the atomic structure of a developed late-stage GP-ηp zone.

blocks” of the ηp-structure can also be found in the matured GP-ηp structure, and (ii) the matured GP-ηp structure still includes an Allattice layer, which consists of mostly Al-atoms at the FCC Al-lattice sites but can be mixed with Zn or Mg atoms at some of the Al sites. However, to accomplish the nucleation of a ηp-precipitate, a matured GP-ηp, has to grow from 7 atomic-layers into 11 atomic-layers in a specific manner, as observed experimentally and discussed in Part II of the present study. 3.2.3. Structures of the η′-precipitates and the GP-η′ zones Fig. 9a shows the atomic-resolution image of a η′-precipitate obtained by performing TF-EWR over the HRTEM image series recorded in an aberration-corrected FEI Titan microscope. Since the structure of the η′precipitates has been studied extensively and different structure models have thus far been suggested by different authors [15,18,21,43–45], we firstly examined these models through image simulations. It was found that none of them can yield a simulated image that exactly matches with the experimentally obtained image of the η′-precipitate shown in Fig. 9a. Our further study suggests that a η′-structure model (Mg2Zn5 − xAl2 + x with x N 0) shown in Fig. 9b can yield a simulated image (the inset in Fig. 9a) that well matches with the experimentally obtained image. Comparing our model with the one previously proposed for the η′-precipitates by Li et al. using HRTEM image simulations [15], it can be seen that the two models are rather similar, except that the Zn-atoms in the pure Zn-layers are arranged slightly different, as shown in Fig. 9b. Compared with the GP-ηp zones and the ηp-precipitates, a η′-precipitate requires a much lower Zn-to-Mg atomic ratio (b2.5) to form.

In comparison with a typical GP-ηp zone, Fig. 10 shows the atomicresolution image of a GP-η′ zone obtained by performing TF-EWR over the HRTEM image series recorded in an aberration-corrected FEI Titan microscope. It is seen that rather different from the structure of the GP-ηp zone that contains 7-atomic-layers without exception, the structure of the GP-η′ zone is more similar to that of the η′-precipitate (Fig. 9a). However, the long range ordering of atoms in the GP-η′ structure is much less well-defined and no proper structure models can be suggested in the present study. 4. Discussion Before drawing conclusions from this part of study, the following point appears worth to be further discussed. We have seen that the two state-of-the-art atomic-resolution imaging techniques have played the key rules in solving the structure problems of various precipitates in these important alloys. Fig. 11 illustrates the key problem that we have faced about the ηp-precipitates. Due to the small atomic distances in both the b112NAl and the b110NAl projections of the ηp-structure, some atomic-columns may not be directly resolved. This can easily lead to a conclusion that ηp-precipitates are just the small η-precipitates [19,20]. Atomic-resolution imaging in HAADF-STEM, on the one hand, can provide straightforward chemistry and structure information about the hardening-precipitates in the first crucial step to well distinguish their subtle differences and to propose their reasonable structure models at the atomic-scale. Atomic-resolution imaging in HRTEM with rapid quantitative image simulation analysis, on the other hand, can provide the refined structures of the precipitates with more precise structure

Fig. 9. a, the atomic-resolution TF-EWR phase-image matched with simulated image (the inset) to reveal the atomic structure of the η′-precipitate viewed along the b112NAl direction; b, the refined structure model suggested by the present work (right) in comparison with that proposed by Li et al. [17].

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the structures is the first crucial step in order to reveal the fine precipitation scenarios in the alloys. The next step has to be such that the relations among these different precipitates are well understood, including their phase or structure transformations, their relative stabilities and their formation sequences with respect to thermal aging and to the variation of alloy composition. These shall be presented systematically in Part II of the present study. 5. Conclusions

Fig. 10. The atomic-resolution TF-EWR phase-image of a GP-η′ zone (left) in comparison with that of a GP-ηp zone (right), viewed along the b112NAl direction.

information about the precipitates for the second step to quantitatively explain the contrast variations, which may not be intuitive but can reflect fine defective and compositional changes in the atomic columns, and to determine structure models with high precision beyond the resolution limitation of the microscopes. The image interpretation in HAADF-STEM is generally more straightforward than that in HRTEM, but is still very time-consuming to be quantitative by image simulations. The combination of the two techniques can be more powerful in solving difficult structure problems in material science, such as the ones that we are facing in AlZn(Cu)Mg alloys. We have seen that a few different types of hardening precipitates may form in the AlZnMg(Cu) alloys upon thermal aging. Determining

By employing two state-of-the-art atomic-resolution imaging techniques in electron microscopy in association with quantitative image simulations, all the hardening precipitates formed in a typical 7150 AlZnMg(Cu) alloys have been investigated systematically for their morphology, crystallography and structure. Based on our experimental observations and quantitative data analysis, the following can be concluded. (1). In the 7150 AlZnMg(Cu) alloys, 5 types of structurally different precipitates can be observed: (i) The large sized equilibrium η-phase (MgZn2) precipitates with a well-known Laves phase structure; (ii) the ηp-precipitates with its lattice parameter c = 4dAl b 111N, many of which have a typical thickness of 11 atomic-layers and an on-side interface dislocation loop; (iii) the early-stage GP-ηp zones featured by their characteristic thickness of 7 atomic-layers, structurally stable doubleatomic-panel and variable content of alloying elements; (iv)

Fig. 11. The schematic side-views and top-view showing the difficulties in solving the structure problem of the ηp-precipitate. a, the b110NAl side-view of the typical quadruple atomic layer in the ηp-structure (right) in comparison within the η-structure (left); b, the b111NAl top-view of the same for a; c, the b112NAl side-view of the same for a and b.

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the η′-precipitates with its lattice parameter c = 6dAl b 111N; (v) the early-stage GP-η′ zones without featured thickness and without obvious periodic spacing in comparison with other precipitates. (2). Due to their small sizes, disc-like shape and structure/composition deviations from the Al-lattice and the η-phase structure, these hardening precipitates have to be studied with state-of-the-art atomic-resolution imaging techniques in association with quantitative image simulation analysis, in order to determine precisely their structures. Atomic-resolution imaging in HAADF-STEM can provide straightforward structure models at the atomic-scale, whereas atomic-resolution imaging in HRTEM with rapid quantitative image simulation analysis can provide the refined structures with high precision beyond the resolution limitation of the microscope. The combination of the two techniques can be more powerful in solving difficult structure problems in material science. (3). Structure determinations through quantitative image analysis have been conducted for four of the five types of observed precipitates. The ηp-structure has a composition of Mg1 − xZn4 − yAl1 + x (0 b x b 0.5, 0 b y b 0.3), evolving towards its metastable composition of MgZn4Al. The GP-ηp zones have a well-defined double-panel structure with a stable composition of MgZn5 for their skeleton that guides their continuous structure evolution towards a matured structure. It is confirmed that the η′-precipitates with a composition of Mg2Zn5 − xAl2 + x with x N 0, have a structure that is very similar to the model previously proposed by Li et al. [15]. For the GP-η′ zones, it is shown that they do not have a well defined structure because of lacking a periodic unit cell in the structure.

Acknowledgments This work is supported by the National Basic Research (973) Program of China (No. 2009CB623704), the National Natural Science Foundation of China (Nos. 51301064, 51371081, 51171063), the Instrumental Innovation Foundation of Hunan Province (No. 2011TT1003), the Aid Program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province, the China Scholarship Council, the Beijing National Center for Electron Microscopy at Tsinghua University and the Center of Electron Microscopy at Zhejiang University. We are grateful to Prof. Z. Zhang, Prof. Y. Wang and Miss Z.J. Shen for their technical supports. References [1] J.C. Williams, E.A. Starke, Progress in structural materials for aerospace systems, Acta Mater. 51 (2003) 5775–5799. [2] E.A. Starke Jr., J.T. Staley, Application of modern aluminum alloys to aircraft, Prog. Aerosp. Sci. 32 (1996) 131–172. [3] K. Chen, H. Liu, Z. Zhang, S. Li, R.I. Todd, The improvement of constituent dissolution and mechanical properties of 7055 aluminum alloy by stepped heat treatments, J. Mater. Process. Technol. 142 (2003) 190–196. [4] G. Sha, A. Cerezo, Early-stage precipitation in Al–Zn–Mg–Cu alloy (7050), Acta Mater. 52 (2004) 4503–4516. [5] R. Ferragut, A. Somoza, A. Tolley, Microstructural evolution of 7012 alloy during the early stages of artificial ageing, Acta Mater. 47 (1999) 4355–4364. [6] M. Liu, B. Klobes, K. Maier, On the age-hardening of an Al–Zn–Mg–Cu alloy: a vacancy perspective, Scr. Mater. 64 (2011) 21–24. [7] J.Z. Liu, J.H. Chen, X.B. Yang, S. Ren, C.L. Wu, H.Y. Xu, J. Zou, Revisiting the precipitation sequence in Al–Zn–Mg-based alloys by high-resolution transimission electron microscopy, Scr. Mater. 63 (2010) 1061–1064. [8] A. Deschamps, G. Fribourg, Y. Bréchet, J.L. Chemin, C.R. Hutchinson, In situ evaluation of dynamic precipitation during plastic straining of an Al–Zn–Mg–Cu alloy, Acta Mater. 60 (2012) 1905–1916. [9] E. Hornbogen, Hundred years of precipitation hardening, J. Light. Met. 1 (2001) 127–132. [10] J.T. Staley, Microstructure and toughness of high-strength aluminium alloys, in: W.R. Warke (Ed.), Properties Related to Fracture Toughness, American Society for Testing and Materials, Baltimore, 1976, pp. 71–96.

[11] H.W. Zandbergen, S.J. Andersen, J. Jansen, Structure determination of Mg5Si6 particles in Al by dynamic electron diffraction studies, Science 277 (1997) 1221–1225. [12] J.H. Chen, E. Costan, M.A. van Huis, Q. Xu, H.W. Zandbergen, Atomic pillar-based nanoprecipitates strengthen AlMgSi alloys, Science 312 (2006) 416–419. [13] H. Löffler, I. Kovacs, J. Lendvai, Decomposition processes in Al–Zn–Mg Alloys, J. Mater. Sci. 18 (1983) 2215–2240. [14] J. Lendvai, Precipitation and strengthening in aluminium alloys, in: J.H. Driver, B. Dubost, F. Durand, R. Fougeres, P. Guyot, P. Sainfort, M. Suery (Eds.),Materials Science Forum, 1996, pp. 43–56. [15] X.Z. Li, V. Hansen, J. Gjonnes, L.R. Wallenberg, HREM study and structure modeling of the eta ' phase, the hardening precipitates in commercial Al–Zn–Mg alloys, Acta Mater. 47 (1999) 2651–2659. [16] L.K. Berg, J. Gjonnes, V. Hansen, X.Z. Li, M. Knutson-Wedel, G. Waterloo, D. Schryvers, L.R. Wallenberg, GP-zones in Al–Zn–Mg alloys and their role in artificial aging, Acta Mater. 49 (2001) 3443–3451. [17] G. Waterloo, V. Hansen, J. Gjonnes, S.R. Skjervold, Effect of predeformation and preaging at room temperature in Al–Zn–Mg-(Cu, Zr) alloys, Mater. Sci. Eng. A 303 (2001) 226–233. [18] A. Kverneland, V. Hansen, R. Vincent, K. Gjonnes, J. Gjonnes, Structure analysis of embedded nano-sized particles by precision electron diffraction eta '-precipitate in an Al–Zn–Mg alloy as example, Ultramicroscopy 106 (2006) 492–502. [19] Y.Y. Li, L. Kovarik, P.J. Phillips, Y.F. Hsu, W.H. Wang, M.J. Mills, High-resolution characterization of the precipitation behavior of an Al–Zn–Mg–Cu alloy, Philos. Mag. Lett. 92 (2012) 166–178. [20] C.D. Marioara, W. Lefebvre, S.J. Andersen, J. Friis, Atomic structure of hardening precipitates in an Al–Mg–Zn–Cu alloy determined by HAADF-STEM and first-principles calculations: relation to eta-MgZn2, J. Mater. Sci. 48 (2013) 3638–3651. [21] C. Wolverton, Crystal structure and stability of complex precipitate phases in Al–Cu– Mg–(Si) and Al–Zn–Mg alloys, Acta Mater. 49 (2001) 3129–3142. [22] Z.R. Liu, J.H. Chen, S.B. Wang, D.W. Yuan, M.J. Yin, C.L. Wu, The structure and the properties of S-phase in AlCuMg alloys, Acta Mater. 59 (2011) 7396–7405. [23] S.B. Wang, J.H. Chen, M.J. Yin, Z.R. Liu, D.W. Yuan, J.Z. Liu, C.H. Liu, C.L. Wu, Doubleatomic-wall-based dynamic precipitates of the early-stage S-phase in AlCuMg alloys, Acta Mater. 60 (2012) 6573–6580. [24] E.J. Kirkland, R.F. Loane, J. Silcox, Simulation of annular dark field stem images using a modified multislice method, Ultramicroscopy 23 (1987) 77–96. [25] S. Hillyard, J. Silcox, Detector geometry, thermal diffuse scattering and strain effects in ADF STEM imaging, Ultramicroscopy 58 (1995) 6–17. [26] P.D. Nellist, S.J. Pennycook, Incoherent imaging using dynamically scattered coherent electrons, Ultramicroscopy 78 (1999) 111–124. [27] W.M.J. Coene, A. Thust, M. Op de Beeck, D. Van Dyck, Maximum-likelihood method for focus-variation image reconstruction in high resolution transmission electron microscopy, Ultramicroscopy 64 (1996) 109–135. [28] A. Thust, W.M.J. Coene, M. Op de Beeck, D. Van Dyck, Focal-series reconstruction in HRTEM: simulation studies on non-periodic objects, Ultramicroscopy 64 (1996) 211–230. [29] W.M.J. Coene, T.J.J. Denteneer, Improved methods for the determination of the spherical aberration coefficient in high-resolution electron microscopy from micrographs of an amorphous object, Ultramicroscopy 38 (1991) 225–233. [30] J.H. Chen, D. Van Dyck, M. Op de Beeck, J. Van Landuyt, Computational comparisons between the conventional multislice method and the third-order multislice method for calculating high-energy electron diffraction and imaging, Ultramicroscopy 69 (1997) 219–240. [31] J.H. Chen, H.W. Zandbergen, D.V. Dyck, Atomic imaging in aberration-corrected high-resolution transmission electron microscopy, Ultramicroscopy 98 (2004) 81–97. [32] J.M. Cowley, Twenty forms of electron holography, Ultramicroscopy 41 (1992) 335–348. [33] M.A. O'Keefe, R. Kilaas, Advances in High-resolution Image Simulation, 1988. [34] M.J. Hÿtch, W.M. Stobbs, Quantitative comparison of high resolution TEM images with image simulations, Ultramicroscopy 53 (1994) 191–203. [35] M.A. O'Keefe, P.R. Buseck, S. Iijima, Computed crystal structure images for high resolution electron microscopy, Nature 274 (1978) 322–324. [36] J.H. Chen, D. Van Dyck, Accurate multislice theory for elastic electron scattering in transmission electron microscopy, Ultramicroscopy 70 (1997) 29–44. [37] J.C.H. Spence, J.M. Zuo, Electron Microdiffraction, Plenum Press, New York, 1992. [38] C. Boothroyd, Why don't high-resolution simulations and images match? J. Microsc. 190 (1998) 99–108. [39] C.L. Jia, K. Urban, Atomic-resolution measurement of oxygen concentration in oxide materials, Science 303 (2004) 2001–2004. [40] Y. Komura, K. Tokunaga, Structural studies of stacking variants in Mg-Base FriaufLaves Phases, Acta Crystallogr. B Struct. Sci. 36 (1980) 1548–1554. [41] M. Haider, S. Uhlemann, E. Schwan, H. Rose, B. Kabius, K. Urban, Electron microscopy image enhanced, Nature 392 (1998) 768–769. [42] J.M. Cowley, Diffraction Physics, North-Holland Publishing, New York, 1981. [43] J.H. Auld, S.M. Cousland, The structure of the metastable η′ phase in aluminium– zinc–magnesium Alloys, J. Aust. Inst. Metal. 19 (1974) 194–199. [44] L.F. Mondolfo, N.A. Gjostein, D.W. Levinson, Structural changes during the aging in an Al–Mg–Zn alloy, Trans. Am. Inst. Min. Metall. Eng. 206 (1956) 1378–1385. [45] J. Gjønnes, C.J. Simensen, An electron microscope investigation of the microstructure in an aluminium–zinc–magnesium alloy, Acta Metall. 18 (1970) 881–890.