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Microstructural evolution, mechanical property and thermal stability of Al–Li 2198-T8 alloy processed by high pressure torsion
Author’s Accepted Manuscript Microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy processed by high pressure torsion Jian Han, Zhixiong Zhu, Huijun Li, Chong Gao www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(15)30575-X http://dx.doi.org/10.1016/j.msea.2015.10.112 MSA32963
To appear in: Materials Science & Engineering A Received date: 19 September 2015 Revised date: 13 October 2015 Accepted date: 29 October 2015 Cite this article as: Jian Han, Zhixiong Zhu, Huijun Li and Chong Gao, Microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy processed by high pressure torsion, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.112 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 galley proof before it is published in its final citable 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.
Microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy processed by high pressure torsion Jian Hana, Zhixiong Zhua*, Huijun Lia, Chong Gaob*
a
School of Mechanical, Materials and Mechatronic Engineering, University of Wollongong,
Wollongong NSW 2522, Australia b
School of Materials Science and Engineering, Beihang University, Beijing 100191, P.R.
China
Keywords: Al-Li alloy, 2198-T8, high-pressure torsion, microstructural evolution, mechanical property, thermal stability.
Corresponding author: Zhixiong Zhu at School of Mechanical, Materials and Mechatronic Engineering,
University
of
Wollongong,
Wollongong
NSW
2522,
Australia,
[email protected], phone number: +61 24223592, fax number: +61 242213238. Chong Gao at School of Materials Science and Engineering, Beihang University, Beijing 100191, China, [email protected], phone number: +86 18618153725.
ABSTRACT
In this study, the Al-Li 2198-T8 (T8: under artificially peak-aged condition) alloy was fabricated by high pressure torsion (HPT) and then subjected to heat treatment at 175 °C for different durations. The microstructural evolution, mechanical property and thermal stability of the HPT-processed alloy were studied. The results demonstrate that the precipitates appearing in the as-received alloy were dissolved into the matrix after HPT processing; and the grains were significantly refined, accompanying by the modification of dislocations. The
dislocation density gradually increased along the radial direction from the centre to the edge of the obtained disk. The improvement of tensile strength of HPT-processed 2198-T8 alloy was dominantly associated with grain refinement and dislocation strengthening. The subsequent heat treatment of the HPT-processed alloy led to slight grain growth and reduction of dislocation density. Also, no strengthening phases precipitated out during the post heat treatment. These factors accounted for the decrease of tensile property for HPT 2198-T8 alloy aged at 175 °C for up to 12 h, suggesting a low thermal stability.
1. Introduction
The history of age-hardening aluminum-lithium (Al-Li) alloys can trace back to the 1920s. As one type of promising materials, the development of Al-Li alloys has experienced three generations [1]. The previous two generations of Al-Li alloys, e.g. 2020 alloy (Al-1.21Li-4.45Cu-0.51Mn-0.20Cd,
in
wt.%)
and
8090
alloy
(Al-2.5Li-1.11Cu-1.16Mn-0.16Zr, in wt.%), had ductility problems or exhibited significant anisotropy in mechanical properties [2, 3]. With characteristics of weight-saving and damage-tolerance, the third generation of Al-Li alloys has been considered as an ideal candidate of the structural materials applied in the modern aerospace industries due to their low density and excellent mechanical properties. The desirable combinations of the mechanical properties are mainly related to precipitation strengthening, which is related to the secondary phases, such as T1 (Al2CuLi), θ' (Al2Cu), S' (Al2CuMg) and δ' (Al3Li) [4, 5]. However, with booming and applying of innovative materials such as carbon-fiber composites, it is imperative to explore the potential of Al-Li alloys with more attractive properties [6-8].
Severe plastic deformation (SPD) has been identified as an effective processing technique to improve mechanical properties of Al-Li alloys. Among of all SPD techniques, equal-channel angular pressing (ECAP) and high pressure torsion (HPT) are the most commonly used for producing bulk nano- / ultrafine-grained materials [9]. In terms of HPT, an extremely high
pressure is applied on the specimen and thus exceptional grain refinement can be delivered. For the HPT-processed materials, the strength can be improved via different mechanisms: grain refinement based on the Hall-Petch theory [10, 11], dislocation accumulation based on Bailey-Hirsch theory, solid solution hardening based on Fleischer and Labusch theory, and precipitation hardening based on Orowan theory [12]. In the last decade, some work has been carried out on the HPT technique to study the microstructural evolution, mechanical properties and strengthening mechanisms. Ghosh et al. [13] concluded that HPT processing refined the grain size up to 120 nm for 7150 alloy (Al-Zn-Mg-Cu) and the grains near the edge of the disk were finer than that near the centre. The hardness of the HPT-processed disk was influenced by not only grain size but also remaining strain near the edge of the disk. Das et al. [9] found that the microstructure of the disk edge for 6063 (Al-Si-Mg) alloy after HPT was severely deformed and consisted of array ultra fine sub-grains, whereas the grain shape was irregular. The hardness gradually increased with distance from the centre to the edge. Lee et al. [14] reported an ultrafine-grained structure with an average grain size of ~140 nm was achieved in 2091 alloy (Al-Li-Cu-Mg) through HPT processing. The microhardness increased with strain, and saturated to a constant level. A further increase in hardness was achieved by ageing the HPT-processed alloy at specified temperatures.
In the present study, the microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy subjected to HPT processing and / or post-HPT ageing were studied. The microstructures and mechanical properties were characterised using optical microscope (OM), transmission electron microscope (TEM), X-ray diffractometer (XRD) and relative mechanical testers. The main objective of this paper is to study the grain refinement and strengthening mechanisms of the obtained materials.
2. Experimental procedures
The studied material is a 1.5 mm cold-rolled sheet of Al-Li 2198 alloy under T8 condition (i.e. solution treatment, quenching, pre-deformation under tension state, and aging at 175 °C). The chemical composition is listed in Table 1. The as-received material was cut into the disks with a diameter of 10 mm, and then the disks were ground to 1.2 mm. Afterwards, the prepared disks were processed between two anvils and strained by concurrent rotation of the two anvils with respect to each other for 2 revolutions at room temperature using a facility operating under quasi-constrained conditions with an applied pressure of 3.0 GPa and a rotation speed of 0.5 rpm. During the process, some limited material is outflow around the periphery of the disks between both anvils [9], and the thickness of the specimen was reduced to 1.1 mm (Fig. 1a). After analysing the corresponding microstructures and properties, the HPT-processed disks (designated as ‘HPT 2198-T8’) were subjected to 175 °C for 6 h and 12 h respectively to study the influence of artificial ageing on the thermal stability.
Table 1 Chemical composition of the studied Al-Li 2198-T8 alloy (in wt.%). Li
Cu
Mg
Mn
Zn
0.98
3.29
0.36
0.05
0.35
Zr
Ag
Si
Fe
Al
0.16 0.30 0.12 0.15 Bal.
The specimens after HPT processing were sectioned and prepared according to standard metallographic procedures. Then the surfaces of the specimens were etched in Keller’s solution (2 ml HF + 3 ml HCl + 5 ml HNO3 + 190 ml H2O) for 15 s. The microstructures of the specimens were studied using the OM. To evaluate the grain size and dislocation density of the HPT 2198-T8, three regions, i.e. area-A (-1.5 ~ 1.5 mm), area-B (0.5 ~ 3.5 mm) and area-C (2.0 ~ 5.0 mm) in Fig. 1a, were separately cut from the specimens and analysed. The
thin foils for TEM study were prepared using a twin-jet electrolytic polishing technique with the electrolyte of 25 vol.% nitric acid and 75 vol.% methanol at the temperature of -35 °C, voltage of 20 V and current range of 60 ~ 90 mA. The precipitates were observed on JEM 2100F TEM. Quantitative XRD measurements were carried out on a D/max 2200PC X-ray diffractometer equipped with a Cu target at the accelerating voltage of 40 kV and a current of 40 mA, together with a scanning step of 0.02° and scanning speed of 6 °/min.
The Vickers hardness testing was performed on an HXZ-1000 Hardness tester using a load of 0.2 kg, dwelling time of 10 s and indentation spacing of 0.5 mm. The tensile testing was performed on an Instron-8801 at room temperature with a strain rate of 6.7 × 10-3 s-1. The dimension of non-standard tensile specimen used in this study was shown in Fig. 1b. Since the dimension of the tensile specimen has to meet the requirements of tensile instrument, area-A was selected for cutting tensile specimen due to a comparatively small size of the HPT-processed disk.
Fig. 1. (a) Disk of HPT 2198-T8 indicating the location of TEM and XRD specimens; (b) schematic representation of tensile specimen.
3. Results and discussion
3.1. Microstructure of 2198-T8 before and after HPT processing
Figs. 2a and b show the optical micrograph of the as-received cold rolled 2198-T8 sheet and the statistical distribution of grain size respectively. It is evident that the typical grains were dominantly polygon-shaped with average grain size of 159.0 μm. The deformed grains were elongated along the rotation direction during HPT processing, and the grain boundaries turned indistinct on the surface of torsion plane (Fig 3). Furthermore, grain structures were invisible from the magnified optical micrographs of area-A, area-B and area-C, due to the grain
refinement after HPT processing.
Fig. 2. (a) Optical micrograph of 2198-T8 alloy; (b) statistical distribution of grain size.
Fig. 3. Optical micrographs of HPT 2198-T8 alloy.
Fig. 4 displays the TEM images of as-received 2198-T8 alloy viewed along [ 12]Al axis. It can be identified from the selected area electron diffraction (SAED) patterns in Fig. 4a that there were three sets of diffraction spots: the relatively bright diffraction spots from Al matrix; the 1/3 [02 ] and 1/3 [31 ] diffraction spots indexed as T1 (Al2CuLi) phases (indicated by yellow circles) precipitated out on the {111}Al plane of Al matrix; and the 1/2 [02 ] and 1/2 [ 00] diffraction spots indexed as super-lattice of L12 structure δ' (Al3Li) precipitates (indicated by blue circles) [1]. The corresponding bright field (BF) and dark field (DF) images (Figs. 4b and c) show that large amounts of needle-shaped T1 (Al2CuLi) phase and some spherical δ' (Al3Li) phase formed within the Al matrix together with dislocations. For the major strengthening T1 phase, the aspect ratio was estimated to be ~10:1 (Fig. 4c) and the length ranged from 27.5 nm to 60.9 nm (Fig. 4d).
Fig. 4. TEM images of Al-Li 2198-T8 alloy viewed along [ 12]Al axis: (a) SAED patterns; (b) BF image; (c) DF image; (d) dimension distribution of T1 phase.
Fig. 5 exhibits the TEM images of the HPT 2198-T8 at different distances away from the centre of the disk. It can be observed from the SAED patterns that the concentric rings formed at the three regions (area-A, area-B and area-C) along the radial direction. Among the three areas, the most uniform ring-shape diffraction pattern appeared in area-B, indicating that the region including finest grains formed in this area. From the BF and DF images of the HPT
2198-T8 alloy, it is apparent that the grains were significantly refined into nano-structures and plenty of dislocations tangled in the grains. The statistical values of the grain size D for area-A, area-B and area-C are 190.1 ± 28.5, 130.2 ± 36.3, and 128.2 ± 27.8 nm, respectively. It is well acceptable that the HPT processing is capable of producing fine grains, and the grains were not homogeneously distributed along the radial direction. Additionally, based on the microstructure (Fig. 5c), some heating is supposed to occur during HPT in the edge area. The principle of HPT is that the disk deformed by shear strain to ensure deformation processes under a quasi-hydrostatic pressure. Assuming that the thickness of the disk is independent of the rotation angle, the shear strain of the disk subjected to HPT processing can be predicted by Equation (1). And the equivalent von Mises strain (ε) can be calculated using Equation (2) [15, 16]. However, for large shear strain (> 0.8), the equivalent strain should be calculated from Equation (3). Based on the distances to the centre (0 mm, 1.5 mm, 3.5 mm and 5 mm), the corresponding shear strains obtained from Equation (1) are 0, 17.1, 40.0 and 57.1, respectively. The equivalent strain can be overlooked if shear strain is small (£ 0.8).
(1) (2) (3)
where γ and ε are the shear strain and equivalent strain respectively, N is the number of rotations, r is the disk radius and h is the thickness of the disk. The equivalent strain ε determined by Equations (1) ~ (3) were 0 ~ 9.8% for the area-A, 6.1 ~ 12.8% for the area-B and 10.8 ~ 14.0% for the area-C, respectively.
It has been reported that the grain size decreased with increasing strain and then reached a steady-state [17, 18]. However, in our study the grain size in HPT 2198-T8 was firstly
decreased, then increased and leveled off along the radial direction. The finest grain size, i.e. 93.9 nm, was achieved within area-B. Xie et al. [19] concluded that the dislocation cell structures formed in the original grains at low strain levels, with increasing strain, the dislocation cells were continuously transformed into the equiaxed sub-grains and finally into fine equiaxed grains.
However, no precipitates were observed in the HPT-processed alloy. One explanation is that the precipitates in the as-received alloy are firstly broken up into small particles due to large shear strain during the HPT process. Murayama et al. [20] reported that large surface free energy was released from those broken particles, which induced small particles re-dissolved into the matrix. On the other hand, Xu et al. [21] believed that the main reason was related to the strain energy from the lattice distortion, the increase of the dislocation density and other crystal defects.
Fig. 5. TEM images of HPT 2198-T8: (a) area-A; (b) area-B; (c) area-C.
Apart from the grain size and precipitation, dislocations can greatly influence the mechanical properties, especially strength. As far as dislocation density is concerned, the equation of Williamson-Hall is available for the estimation of dislocation density by means of XRD peak broadening [22, 23]. Fig. 6 shows the XRD results of the HPT 2198-T8 in different regions. It is evident that the full width at half maximum (FWHM) of the diffraction peaks decreased gradually from the centre to the edge. The dislocation density ρ can be assessed and calculated by the values of FWHM. The peak broadening usually results from grain refinement and / or formation of dislocations. The relationship between dislocation density, grain size and peak broadening can be expressed by the following equations:
(4) (5)
where β is the peak width at half maximum in rad, λ is the wave length of X-ray beam (0.1542 nm for Cu Kα), D is the crystallite size, ε is the lattice strain, θ is the Bragg angle, ρ is the dislocation density and b is the Burgers vector. It is reported that the effect of grain size on the peak broadening is significant when the grain size is finer than 100 nm [24]. In this study, the grains are much coarser than 100 nm, thus the influence of grain size on peak broadening can be neglected. Hence only the influence of dislocation needs to be taken into account. The dislocation density ρ was evaluated by Equations (4) and (5) to be 20.6, 27.8 and 40.8 × 1014 m−2 for area-A, area-B and area-C, respectively. Compared to the original 2198-T8 alloy (14.8 × 1014 m−2) [25], the dislocation density gradually increased with the distance from the centre to the edge.
Fig. 6. XRD results of different areas of HPT 2198-T8 alloy (area-A, area-B and area-C).
In addition, the evolution of the dislocation density in fine-grained alloys is different from that in coarse-grained materials [26]. Wang et al. [27] revealed that the deformation-induced dislocation density in nano-crystalline grains initially decreased and then increased with strain. For area-A, area-B and area-C of HPT 2198-T8, the dislocation density is proportional to the equivalent strain that is related to the distortion of the diffraction peaks. The correlation between the equivalent strain ε and microstructural evolution of HPT 2198-T8 was summarised in Fig. 7. It is clear that the grain size has a linear relationship with the equivalent strain and / or dislocation density.
Fig. 7. Correlation between the equivalent strain ε and microstructural evolution of area-A,
area-B and area-C.
3.2. Thermal stability of HPT 2198-T8
Post-HPT ageing was carried out for the evaluation of the Thermal stability of HPT 2198-T8. The TEM images of HPT 2198-T8 aged at 175 °C for 6 h and 12 h are shown in Fig. 8. Based on the polycrystalline rings of the SAED patterns, finer grains were obtained in the sample aged for 6 h than that for 12 h. As can be seen from Fig. 8, grain growth and formation of sub-grains were occurred. After ageing for 6 h and 12 h, the grains grew up to 305.1 ± 95.4 nm and 350.5 ± 36.8 nm respectively. Furthermore, some sub-grains were observed within grains (marked by yellow arrows).
Fig. 8. TEM images in area-A of HPT 2198-T8 aged at 175 °C for: (a) 6 h and (b) 12 h.
Fig. 9 displays the XRD results of area-A of the HPT 2198-T8 aged at 175 °C for 6 h and 12 h. It is evident that the FWHM decreased gradually with the extension of holding time. The dislocation density ρ was calculated by Equations (4) and (5) to be approximately 6.3 and 4.6 × 1014 m−2 for 6 h and 12 h, respectively. Compared to the HPT 2198-T8 without any post-HPT ageing, the values of dislocation density were gradually reduced with the ageing time when the duration is less than 12 h.
Fig. 9. XRD patterns of area-A of HPT 2198-T8 aged at 175 °C up to 12 h.
Fig. 10 displays the hardness distribution of the HPT 2198-T8 before and after post-HPT ageing. For the HPT specimen without ageing, the hardness gradually increased from 200 HV0.2 at the centre (area-A) to peak value (approximately 255 HV0.2 in area-B), which was related to the equivalent strain ε and also the variations of the grain size and dislocation density induced from the equivalent strain ε. When approaching the periphery of the disk, the
hardness slightly went down. Compared to the hardness of as-received alloy under cold-rolled condition (180 HV0.2), the hardness of HPT 2198-T8 significantly increased (by 11.1% to 41.7%). The post-HPT ageing led to considerable reduction of hardness. When the ageing time was 6 h, the hardness gradually increased from 165 HV0.2 in the centre (area-A) to 197 HV0.2 on the edge (area-C). When the ageing time was prolonged to 12 h, the hardness further decreased and leveled off at ~90 HV0.2 along the radial direction.
Fig. 10. Hardness distribution of HPT 2198-T8 aged at 175 °C up to 12 h compared to HPT 2198-T8 without ageing.
Fig. 11 shows the tensile properties of the HPT 2198-T8 before and after post heat treatment compared to the as-received alloy. For the as-received alloy, the ultimate tensile strength (UTS) and yield strength (YS) are 518.7 MPa and 470.4 MPa, and the elongation (EL) is 15.5%; whereas the UTS and YS of the HPT 2198-T8 were increased to 710.9 MPa and 700.1 MPa, and the EL was only less than 5%. After ageing at 175 °C for 6 h and 12 h, the UTS and YS were decreased to 482.7 / 473.2 MPa and 339.9 / 322.4 MPa respectively, and the EL was lower than 5%. The variations of tensile strength were consistent with the hardness. Furthermore, since the grains significantly grew up to more than 300 nm, it was understandable that the EL did not greatly change with the variations of tensile strength.
As mentioned above, several strengthening mechanisms determine the final properties of ultrafine-grained materials. The contribution of grain refinement is supposed to be the dominant strengthening mechanism since the grains after HPT turn from micro to nano scale. The importance of dislocation accumulation should be considered, due to its large contribution to the strength [1]. The precipitation hardening can be ignored since no precipitation behaviour is observed after HPT and post-HPT. Additionally, the contribution of solid-solution hardening to the total strength is small [12]. It is reported that the microstructures of some Al alloys processed by HPT were metastable due to high-angle grain
boundaries and dislocations accumulated in grains [19, 27]. This is because in these Al alloys the main factors determining the strength are grain size and dislocation density. In this study, the dominant strengthening mechanisms of the HPT 2198-T8 are grain refinement and dislocation strengthening [23, 28]. The average grain size in area-A was 190 nm with the dislocation density of 20.6 × 1014 m−2 in the grains after HPT processing (Fig. 5a), which led to a significant increase in the mechanical properties. After ageing at 175 °C for 6 h, the grains grew up to 305 nm and the dislocation density decreased to 6.3 × 1014 m−2, which gave rise to a lower strength. When the duration further increased to 12 h, the grains continuously grew to 351 nm and the dislocation density further reduced to 4.6 × 1014 m−2, which undoubtedly further lowered the strength. It can be concluded that the post-HPT ageing has a negative influence on the strength of the HPT 2198-T8. In other words, the HPT 2198-T8 has a low thermal stability.
Fig. 11. Tensile properties of HPT 2198-T8 alloy before and after post-HPT ageing compared to the as-received alloy.
Conclusions
The microstructures of the Al-Li 2198-T8 alloy processed by HPT were non-uniform. With the increase of the strain (0 ~ 14.0%) in the HPT-processed specimen, no precipitates were observed and the average grain sizes varied from 190 nm to 128 nm. Grain refinement is the most important characteristic during HPT processing. Also, the dislocation density gradually increased along the radial direction from the centre to the edge. The mechanical properties (hardness and strength) of the HPT 2198-T8 significantly increased due to grain refinement and dislocation strengthening.
The thermal stability of HPT 2198-T8 was low, which means that the post-HPT ageing led to a large variation of the mechanical properties. After ageing at 175 °C for 6 h and 12h, the
grains grew up, and the dislocation density decreased. Meanwhile, no strengthening phases precipitated out. These factors, including grain size, dislocation density and precipitation, led to a lower strength of HPT 2198-T8 after post-HPT ageing.
Acknowledgements The authors are grateful to the cooperative partners X.Z. Liao and X.H. An (University of Sydney) for their help in discussion and HPT experiments.
References
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PII: DOI: Reference:
S0921-5093(15)30575-X http://dx.doi.org/10.1016/j.msea.2015.10.112 MSA32963
To appear in: Materials Science & Engineering A Received date: 19 September 2015 Revised date: 13 October 2015 Accepted date: 29 October 2015 Cite this article as: Jian Han, Zhixiong Zhu, Huijun Li and Chong Gao, Microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy processed by high pressure torsion, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.112 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 galley proof before it is published in its final citable 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.
Microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy processed by high pressure torsion Jian Hana, Zhixiong Zhua*, Huijun Lia, Chong Gaob*
a
School of Mechanical, Materials and Mechatronic Engineering, University of Wollongong,
Wollongong NSW 2522, Australia b
School of Materials Science and Engineering, Beihang University, Beijing 100191, P.R.
China
Keywords: Al-Li alloy, 2198-T8, high-pressure torsion, microstructural evolution, mechanical property, thermal stability.
Corresponding author: Zhixiong Zhu at School of Mechanical, Materials and Mechatronic Engineering,
University
of
Wollongong,
Wollongong
NSW
2522,
Australia,
[email protected], phone number: +61 24223592, fax number: +61 242213238. Chong Gao at School of Materials Science and Engineering, Beihang University, Beijing 100191, China, [email protected], phone number: +86 18618153725.
ABSTRACT
In this study, the Al-Li 2198-T8 (T8: under artificially peak-aged condition) alloy was fabricated by high pressure torsion (HPT) and then subjected to heat treatment at 175 °C for different durations. The microstructural evolution, mechanical property and thermal stability of the HPT-processed alloy were studied. The results demonstrate that the precipitates appearing in the as-received alloy were dissolved into the matrix after HPT processing; and the grains were significantly refined, accompanying by the modification of dislocations. The
dislocation density gradually increased along the radial direction from the centre to the edge of the obtained disk. The improvement of tensile strength of HPT-processed 2198-T8 alloy was dominantly associated with grain refinement and dislocation strengthening. The subsequent heat treatment of the HPT-processed alloy led to slight grain growth and reduction of dislocation density. Also, no strengthening phases precipitated out during the post heat treatment. These factors accounted for the decrease of tensile property for HPT 2198-T8 alloy aged at 175 °C for up to 12 h, suggesting a low thermal stability.
1. Introduction
The history of age-hardening aluminum-lithium (Al-Li) alloys can trace back to the 1920s. As one type of promising materials, the development of Al-Li alloys has experienced three generations [1]. The previous two generations of Al-Li alloys, e.g. 2020 alloy (Al-1.21Li-4.45Cu-0.51Mn-0.20Cd,
in
wt.%)
and
8090
alloy
(Al-2.5Li-1.11Cu-1.16Mn-0.16Zr, in wt.%), had ductility problems or exhibited significant anisotropy in mechanical properties [2, 3]. With characteristics of weight-saving and damage-tolerance, the third generation of Al-Li alloys has been considered as an ideal candidate of the structural materials applied in the modern aerospace industries due to their low density and excellent mechanical properties. The desirable combinations of the mechanical properties are mainly related to precipitation strengthening, which is related to the secondary phases, such as T1 (Al2CuLi), θ' (Al2Cu), S' (Al2CuMg) and δ' (Al3Li) [4, 5]. However, with booming and applying of innovative materials such as carbon-fiber composites, it is imperative to explore the potential of Al-Li alloys with more attractive properties [6-8].
Severe plastic deformation (SPD) has been identified as an effective processing technique to improve mechanical properties of Al-Li alloys. Among of all SPD techniques, equal-channel angular pressing (ECAP) and high pressure torsion (HPT) are the most commonly used for producing bulk nano- / ultrafine-grained materials [9]. In terms of HPT, an extremely high
pressure is applied on the specimen and thus exceptional grain refinement can be delivered. For the HPT-processed materials, the strength can be improved via different mechanisms: grain refinement based on the Hall-Petch theory [10, 11], dislocation accumulation based on Bailey-Hirsch theory, solid solution hardening based on Fleischer and Labusch theory, and precipitation hardening based on Orowan theory [12]. In the last decade, some work has been carried out on the HPT technique to study the microstructural evolution, mechanical properties and strengthening mechanisms. Ghosh et al. [13] concluded that HPT processing refined the grain size up to 120 nm for 7150 alloy (Al-Zn-Mg-Cu) and the grains near the edge of the disk were finer than that near the centre. The hardness of the HPT-processed disk was influenced by not only grain size but also remaining strain near the edge of the disk. Das et al. [9] found that the microstructure of the disk edge for 6063 (Al-Si-Mg) alloy after HPT was severely deformed and consisted of array ultra fine sub-grains, whereas the grain shape was irregular. The hardness gradually increased with distance from the centre to the edge. Lee et al. [14] reported an ultrafine-grained structure with an average grain size of ~140 nm was achieved in 2091 alloy (Al-Li-Cu-Mg) through HPT processing. The microhardness increased with strain, and saturated to a constant level. A further increase in hardness was achieved by ageing the HPT-processed alloy at specified temperatures.
In the present study, the microstructural evolution, mechanical property and thermal stability of Al-Li 2198-T8 alloy subjected to HPT processing and / or post-HPT ageing were studied. The microstructures and mechanical properties were characterised using optical microscope (OM), transmission electron microscope (TEM), X-ray diffractometer (XRD) and relative mechanical testers. The main objective of this paper is to study the grain refinement and strengthening mechanisms of the obtained materials.
2. Experimental procedures
The studied material is a 1.5 mm cold-rolled sheet of Al-Li 2198 alloy under T8 condition (i.e. solution treatment, quenching, pre-deformation under tension state, and aging at 175 °C). The chemical composition is listed in Table 1. The as-received material was cut into the disks with a diameter of 10 mm, and then the disks were ground to 1.2 mm. Afterwards, the prepared disks were processed between two anvils and strained by concurrent rotation of the two anvils with respect to each other for 2 revolutions at room temperature using a facility operating under quasi-constrained conditions with an applied pressure of 3.0 GPa and a rotation speed of 0.5 rpm. During the process, some limited material is outflow around the periphery of the disks between both anvils [9], and the thickness of the specimen was reduced to 1.1 mm (Fig. 1a). After analysing the corresponding microstructures and properties, the HPT-processed disks (designated as ‘HPT 2198-T8’) were subjected to 175 °C for 6 h and 12 h respectively to study the influence of artificial ageing on the thermal stability.
Table 1 Chemical composition of the studied Al-Li 2198-T8 alloy (in wt.%). Li
Cu
Mg
Mn
Zn
0.98
3.29
0.36
0.05
0.35
Zr
Ag
Si
Fe
Al
0.16 0.30 0.12 0.15 Bal.
The specimens after HPT processing were sectioned and prepared according to standard metallographic procedures. Then the surfaces of the specimens were etched in Keller’s solution (2 ml HF + 3 ml HCl + 5 ml HNO3 + 190 ml H2O) for 15 s. The microstructures of the specimens were studied using the OM. To evaluate the grain size and dislocation density of the HPT 2198-T8, three regions, i.e. area-A (-1.5 ~ 1.5 mm), area-B (0.5 ~ 3.5 mm) and area-C (2.0 ~ 5.0 mm) in Fig. 1a, were separately cut from the specimens and analysed. The
thin foils for TEM study were prepared using a twin-jet electrolytic polishing technique with the electrolyte of 25 vol.% nitric acid and 75 vol.% methanol at the temperature of -35 °C, voltage of 20 V and current range of 60 ~ 90 mA. The precipitates were observed on JEM 2100F TEM. Quantitative XRD measurements were carried out on a D/max 2200PC X-ray diffractometer equipped with a Cu target at the accelerating voltage of 40 kV and a current of 40 mA, together with a scanning step of 0.02° and scanning speed of 6 °/min.
The Vickers hardness testing was performed on an HXZ-1000 Hardness tester using a load of 0.2 kg, dwelling time of 10 s and indentation spacing of 0.5 mm. The tensile testing was performed on an Instron-8801 at room temperature with a strain rate of 6.7 × 10-3 s-1. The dimension of non-standard tensile specimen used in this study was shown in Fig. 1b. Since the dimension of the tensile specimen has to meet the requirements of tensile instrument, area-A was selected for cutting tensile specimen due to a comparatively small size of the HPT-processed disk.
Fig. 1. (a) Disk of HPT 2198-T8 indicating the location of TEM and XRD specimens; (b) schematic representation of tensile specimen.
3. Results and discussion
3.1. Microstructure of 2198-T8 before and after HPT processing
Figs. 2a and b show the optical micrograph of the as-received cold rolled 2198-T8 sheet and the statistical distribution of grain size respectively. It is evident that the typical grains were dominantly polygon-shaped with average grain size of 159.0 μm. The deformed grains were elongated along the rotation direction during HPT processing, and the grain boundaries turned indistinct on the surface of torsion plane (Fig 3). Furthermore, grain structures were invisible from the magnified optical micrographs of area-A, area-B and area-C, due to the grain
refinement after HPT processing.
Fig. 2. (a) Optical micrograph of 2198-T8 alloy; (b) statistical distribution of grain size.
Fig. 3. Optical micrographs of HPT 2198-T8 alloy.
Fig. 4 displays the TEM images of as-received 2198-T8 alloy viewed along [ 12]Al axis. It can be identified from the selected area electron diffraction (SAED) patterns in Fig. 4a that there were three sets of diffraction spots: the relatively bright diffraction spots from Al matrix; the 1/3 [02 ] and 1/3 [31 ] diffraction spots indexed as T1 (Al2CuLi) phases (indicated by yellow circles) precipitated out on the {111}Al plane of Al matrix; and the 1/2 [02 ] and 1/2 [ 00] diffraction spots indexed as super-lattice of L12 structure δ' (Al3Li) precipitates (indicated by blue circles) [1]. The corresponding bright field (BF) and dark field (DF) images (Figs. 4b and c) show that large amounts of needle-shaped T1 (Al2CuLi) phase and some spherical δ' (Al3Li) phase formed within the Al matrix together with dislocations. For the major strengthening T1 phase, the aspect ratio was estimated to be ~10:1 (Fig. 4c) and the length ranged from 27.5 nm to 60.9 nm (Fig. 4d).
Fig. 4. TEM images of Al-Li 2198-T8 alloy viewed along [ 12]Al axis: (a) SAED patterns; (b) BF image; (c) DF image; (d) dimension distribution of T1 phase.
Fig. 5 exhibits the TEM images of the HPT 2198-T8 at different distances away from the centre of the disk. It can be observed from the SAED patterns that the concentric rings formed at the three regions (area-A, area-B and area-C) along the radial direction. Among the three areas, the most uniform ring-shape diffraction pattern appeared in area-B, indicating that the region including finest grains formed in this area. From the BF and DF images of the HPT
2198-T8 alloy, it is apparent that the grains were significantly refined into nano-structures and plenty of dislocations tangled in the grains. The statistical values of the grain size D for area-A, area-B and area-C are 190.1 ± 28.5, 130.2 ± 36.3, and 128.2 ± 27.8 nm, respectively. It is well acceptable that the HPT processing is capable of producing fine grains, and the grains were not homogeneously distributed along the radial direction. Additionally, based on the microstructure (Fig. 5c), some heating is supposed to occur during HPT in the edge area. The principle of HPT is that the disk deformed by shear strain to ensure deformation processes under a quasi-hydrostatic pressure. Assuming that the thickness of the disk is independent of the rotation angle, the shear strain of the disk subjected to HPT processing can be predicted by Equation (1). And the equivalent von Mises strain (ε) can be calculated using Equation (2) [15, 16]. However, for large shear strain (> 0.8), the equivalent strain should be calculated from Equation (3). Based on the distances to the centre (0 mm, 1.5 mm, 3.5 mm and 5 mm), the corresponding shear strains obtained from Equation (1) are 0, 17.1, 40.0 and 57.1, respectively. The equivalent strain can be overlooked if shear strain is small (£ 0.8).
(1) (2) (3)
where γ and ε are the shear strain and equivalent strain respectively, N is the number of rotations, r is the disk radius and h is the thickness of the disk. The equivalent strain ε determined by Equations (1) ~ (3) were 0 ~ 9.8% for the area-A, 6.1 ~ 12.8% for the area-B and 10.8 ~ 14.0% for the area-C, respectively.
It has been reported that the grain size decreased with increasing strain and then reached a steady-state [17, 18]. However, in our study the grain size in HPT 2198-T8 was firstly
decreased, then increased and leveled off along the radial direction. The finest grain size, i.e. 93.9 nm, was achieved within area-B. Xie et al. [19] concluded that the dislocation cell structures formed in the original grains at low strain levels, with increasing strain, the dislocation cells were continuously transformed into the equiaxed sub-grains and finally into fine equiaxed grains.
However, no precipitates were observed in the HPT-processed alloy. One explanation is that the precipitates in the as-received alloy are firstly broken up into small particles due to large shear strain during the HPT process. Murayama et al. [20] reported that large surface free energy was released from those broken particles, which induced small particles re-dissolved into the matrix. On the other hand, Xu et al. [21] believed that the main reason was related to the strain energy from the lattice distortion, the increase of the dislocation density and other crystal defects.
Fig. 5. TEM images of HPT 2198-T8: (a) area-A; (b) area-B; (c) area-C.
Apart from the grain size and precipitation, dislocations can greatly influence the mechanical properties, especially strength. As far as dislocation density is concerned, the equation of Williamson-Hall is available for the estimation of dislocation density by means of XRD peak broadening [22, 23]. Fig. 6 shows the XRD results of the HPT 2198-T8 in different regions. It is evident that the full width at half maximum (FWHM) of the diffraction peaks decreased gradually from the centre to the edge. The dislocation density ρ can be assessed and calculated by the values of FWHM. The peak broadening usually results from grain refinement and / or formation of dislocations. The relationship between dislocation density, grain size and peak broadening can be expressed by the following equations:
(4) (5)
where β is the peak width at half maximum in rad, λ is the wave length of X-ray beam (0.1542 nm for Cu Kα), D is the crystallite size, ε is the lattice strain, θ is the Bragg angle, ρ is the dislocation density and b is the Burgers vector. It is reported that the effect of grain size on the peak broadening is significant when the grain size is finer than 100 nm [24]. In this study, the grains are much coarser than 100 nm, thus the influence of grain size on peak broadening can be neglected. Hence only the influence of dislocation needs to be taken into account. The dislocation density ρ was evaluated by Equations (4) and (5) to be 20.6, 27.8 and 40.8 × 1014 m−2 for area-A, area-B and area-C, respectively. Compared to the original 2198-T8 alloy (14.8 × 1014 m−2) [25], the dislocation density gradually increased with the distance from the centre to the edge.
Fig. 6. XRD results of different areas of HPT 2198-T8 alloy (area-A, area-B and area-C).
In addition, the evolution of the dislocation density in fine-grained alloys is different from that in coarse-grained materials [26]. Wang et al. [27] revealed that the deformation-induced dislocation density in nano-crystalline grains initially decreased and then increased with strain. For area-A, area-B and area-C of HPT 2198-T8, the dislocation density is proportional to the equivalent strain that is related to the distortion of the diffraction peaks. The correlation between the equivalent strain ε and microstructural evolution of HPT 2198-T8 was summarised in Fig. 7. It is clear that the grain size has a linear relationship with the equivalent strain and / or dislocation density.
Fig. 7. Correlation between the equivalent strain ε and microstructural evolution of area-A,
area-B and area-C.
3.2. Thermal stability of HPT 2198-T8
Post-HPT ageing was carried out for the evaluation of the Thermal stability of HPT 2198-T8. The TEM images of HPT 2198-T8 aged at 175 °C for 6 h and 12 h are shown in Fig. 8. Based on the polycrystalline rings of the SAED patterns, finer grains were obtained in the sample aged for 6 h than that for 12 h. As can be seen from Fig. 8, grain growth and formation of sub-grains were occurred. After ageing for 6 h and 12 h, the grains grew up to 305.1 ± 95.4 nm and 350.5 ± 36.8 nm respectively. Furthermore, some sub-grains were observed within grains (marked by yellow arrows).
Fig. 8. TEM images in area-A of HPT 2198-T8 aged at 175 °C for: (a) 6 h and (b) 12 h.
Fig. 9 displays the XRD results of area-A of the HPT 2198-T8 aged at 175 °C for 6 h and 12 h. It is evident that the FWHM decreased gradually with the extension of holding time. The dislocation density ρ was calculated by Equations (4) and (5) to be approximately 6.3 and 4.6 × 1014 m−2 for 6 h and 12 h, respectively. Compared to the HPT 2198-T8 without any post-HPT ageing, the values of dislocation density were gradually reduced with the ageing time when the duration is less than 12 h.
Fig. 9. XRD patterns of area-A of HPT 2198-T8 aged at 175 °C up to 12 h.
Fig. 10 displays the hardness distribution of the HPT 2198-T8 before and after post-HPT ageing. For the HPT specimen without ageing, the hardness gradually increased from 200 HV0.2 at the centre (area-A) to peak value (approximately 255 HV0.2 in area-B), which was related to the equivalent strain ε and also the variations of the grain size and dislocation density induced from the equivalent strain ε. When approaching the periphery of the disk, the
hardness slightly went down. Compared to the hardness of as-received alloy under cold-rolled condition (180 HV0.2), the hardness of HPT 2198-T8 significantly increased (by 11.1% to 41.7%). The post-HPT ageing led to considerable reduction of hardness. When the ageing time was 6 h, the hardness gradually increased from 165 HV0.2 in the centre (area-A) to 197 HV0.2 on the edge (area-C). When the ageing time was prolonged to 12 h, the hardness further decreased and leveled off at ~90 HV0.2 along the radial direction.
Fig. 10. Hardness distribution of HPT 2198-T8 aged at 175 °C up to 12 h compared to HPT 2198-T8 without ageing.
Fig. 11 shows the tensile properties of the HPT 2198-T8 before and after post heat treatment compared to the as-received alloy. For the as-received alloy, the ultimate tensile strength (UTS) and yield strength (YS) are 518.7 MPa and 470.4 MPa, and the elongation (EL) is 15.5%; whereas the UTS and YS of the HPT 2198-T8 were increased to 710.9 MPa and 700.1 MPa, and the EL was only less than 5%. After ageing at 175 °C for 6 h and 12 h, the UTS and YS were decreased to 482.7 / 473.2 MPa and 339.9 / 322.4 MPa respectively, and the EL was lower than 5%. The variations of tensile strength were consistent with the hardness. Furthermore, since the grains significantly grew up to more than 300 nm, it was understandable that the EL did not greatly change with the variations of tensile strength.
As mentioned above, several strengthening mechanisms determine the final properties of ultrafine-grained materials. The contribution of grain refinement is supposed to be the dominant strengthening mechanism since the grains after HPT turn from micro to nano scale. The importance of dislocation accumulation should be considered, due to its large contribution to the strength [1]. The precipitation hardening can be ignored since no precipitation behaviour is observed after HPT and post-HPT. Additionally, the contribution of solid-solution hardening to the total strength is small [12]. It is reported that the microstructures of some Al alloys processed by HPT were metastable due to high-angle grain
boundaries and dislocations accumulated in grains [19, 27]. This is because in these Al alloys the main factors determining the strength are grain size and dislocation density. In this study, the dominant strengthening mechanisms of the HPT 2198-T8 are grain refinement and dislocation strengthening [23, 28]. The average grain size in area-A was 190 nm with the dislocation density of 20.6 × 1014 m−2 in the grains after HPT processing (Fig. 5a), which led to a significant increase in the mechanical properties. After ageing at 175 °C for 6 h, the grains grew up to 305 nm and the dislocation density decreased to 6.3 × 1014 m−2, which gave rise to a lower strength. When the duration further increased to 12 h, the grains continuously grew to 351 nm and the dislocation density further reduced to 4.6 × 1014 m−2, which undoubtedly further lowered the strength. It can be concluded that the post-HPT ageing has a negative influence on the strength of the HPT 2198-T8. In other words, the HPT 2198-T8 has a low thermal stability.
Fig. 11. Tensile properties of HPT 2198-T8 alloy before and after post-HPT ageing compared to the as-received alloy.
Conclusions
The microstructures of the Al-Li 2198-T8 alloy processed by HPT were non-uniform. With the increase of the strain (0 ~ 14.0%) in the HPT-processed specimen, no precipitates were observed and the average grain sizes varied from 190 nm to 128 nm. Grain refinement is the most important characteristic during HPT processing. Also, the dislocation density gradually increased along the radial direction from the centre to the edge. The mechanical properties (hardness and strength) of the HPT 2198-T8 significantly increased due to grain refinement and dislocation strengthening.
The thermal stability of HPT 2198-T8 was low, which means that the post-HPT ageing led to a large variation of the mechanical properties. After ageing at 175 °C for 6 h and 12h, the
grains grew up, and the dislocation density decreased. Meanwhile, no strengthening phases precipitated out. These factors, including grain size, dislocation density and precipitation, led to a lower strength of HPT 2198-T8 after post-HPT ageing.
Acknowledgements The authors are grateful to the cooperative partners X.Z. Liao and X.H. An (University of Sydney) for their help in discussion and HPT experiments.
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