The enhanced microhardness in a rapidly solidified Al alloy

Author’s Accepted Manuscript The enhanced microhardness in a rapidly solidified Al alloy Yaojun Lin, Shuaiying Mao, Zhigang Yan, Yaqi Zhang, Limin Wang www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(17)30350-7 http://dx.doi.org/10.1016/j.msea.2017.03.052 MSA34835

To appear in: Materials Science & Engineering A Received date: 4 January 2017 Revised date: 9 March 2017 Accepted date: 13 March 2017 Cite this article as: Yaojun Lin, Shuaiying Mao, Zhigang Yan, Yaqi Zhang and Limin Wang, The enhanced microhardness in a rapidly solidified Al alloy, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.03.052 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.

The enhanced microhardness in a rapidly solidified Al alloy Yaojun Lina,b,*, Shuaiying Maob, Zhigang Yanb, Yaqi Zhangb, and Limin Wangb a

School of Materials Science and Engineering, Wuhan University of Technology, Wuhan, Hubei 430070, China

b

State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, Hebei 066004, China *

Corresponding author: [email protected] (Y.J. Lin)

Abstract The objective of the present study is to explore the feasibility of producing 7075 Al alloy ribbons with the increased microhardness via rapid solidification processing, namely melt spinning. Our results show that the melt spun + aged 7075 Al alloy ribbons exhibit approximately 17% (307 MPa) higher microhardness than that of the coarse-grained (CG) counterpart solid-solution treated + aged under the same conditions. The higher microhardness of the melt spun + aged 7075 Al alloy stems from its microstructural features, including (i) sub-micrometric/micrometric sized grains, (ii) intragranular subgrains, and (iii) nanoscale precipitates with inter-precipitate distance slightly smaller than that in the CG counterpart. In an effort to understand the mcirostructure in melt spun + aged 7075 Al alloy, the microstructural evolution during aging of the as-melt spun 7075 Al alloy was analyzed and discussed. By quantifying the contributions of precipitation, grain boundary, dislocation, solid solution and subgrain boundary strengthening to the yield strength based on the microstructural analysis, the higher microhardness of the melt spun + aged 7075 Al alloy ribbons was rationalized. Keywords: Aluminum alloys; Rapid solidification; Melt spinning; Microstructure; Microhardness.

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1. Introduction Over the past few decades, rapid solidification techniques have attracted considerable interest due to their ability to significantly refine microstructures and even to create amorphous structures, and thus the potential to achieve excellent mechanical properties [1-3]. Among the various rapid solidification techniques developed, including atomization methods (e.g., gas, water and centrifugal atomization), chilling methods (e.g., melt spinning and melt drag), etc., melt spinning has received particular attention as a result of a relatively higher cooling rate (105 to 107 K/s) during melt spinning and its ability of mass production [1-3]. Inspection of the published studies indicates that melt spinning has been widely used to produce ribbons of two classes of Al-based alloys: (i) Al-based amorphous alloys, including Al-EM-LM (EM: early transition metals, including Ti, Zr, Hf, V, Nb, Ta, Cr, Mo or W; LM: late transition metals, including Fe, Co, Ni or Cu) [4,5], Al-R (R: rare earth metals, including Y, La, Ce, Pr, Nd, Sm, Gd, Tb, Dy, Ho, Er or Yb) [6,7], Al-R-M (M: transition metals) [7-10], etc., and (ii) crystalline Al alloys containing transition elements that exhibit extremely low equilibrium solubility in the Al matrix (e.g., Cr, Fe, Ti, Mn, Ni, Zr, etc.), such as Al-Fe-Cr alloys [11-13], Al-Fe-Cr-Ti alloys [14], Al-Mn-Fe alloys [15], Al-Si-Zr alloys [16], Al-20Si-5Fe-2X (X=Ni or Cr) [17], etc. Usually, the melt spun ribbons of the two classes of Al-based alloys were implemented as precursors to prepare bulk Al-based alloys with nanogranular structures for the amorphous Al-based alloy ribbons and bulk Al-based alloys strengthened by quasicrystalline or nanoscale particles for the crystalline Al-based alloy ribbons. In doing so, the melt spun ribbons were first fragmented into powders and then powders were consolidated at elevated temperatures. In addition to the application as precursors to prepare bulk Al-based alloys, the melt spun ribbons of the two classes of Al-based alloys may be implemented as raw materials to fabricate high-strength micro-sheet-metal components with a few tens of micrometers in

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thickness via micro-manufacturing techniques, such as micro-machining, micro-forming, micro-joining, etc. [18]. Despite the high strength, the brittleness of the melt spun ribbons of the two classes of Al-based alloys, however, presents a stringent challenge. The above discussion invites an interesting question: given the good ductility of commercial Al alloys, can their melt spun ribbons be used to make high-strength micro-sheet-metal components? Inspection of the published studies shows that the ribbons of commercial Al alloys, such as 6061 Al [19], 8006 Al [20], 5083 Al [21], A413.1 Al [22], Al-Zn-Mg-Cu alloys [23], 7010 Al [24], etc., have been prepared via melt spinning. To accomplish this goal, one of the critical issues is whether melt spinning can enhance the strength of commercial Al alloys or not. In view of the aforementioned discussion, the objective of the present study is to explore the feasibility of producing commercial Al alloy ribbons with the enhanced strength using melt spinning. Given that 7000 series Al alloys generally exhibit the highest strength of various commercial Al alloys [25-27], the present study concentrated on this class of commercial Al alloys. Herein we selected a widely used 7075 Al alloy as the model material. It has been well documented [25-33] that in the solid-solution treated + aged coarse-grained (CG) 7000 series Al alloys the primary strengthening mechanism is precipitation strengthening imparted by nano-scale second-phase particles, including Guinier-Preston (G-P) zones,  and  phases, which precipitate during aging. Generally, the other strengthening mechanisms, such as grain boundary strengthening, dislocation strengthening and solid-solution strengthening, make insignificant contributions to the overall strength, since high temperatures during solid-solution treatment yield coarse grains (a few ten to hundred micrometers in grain size [25-27,34,35]) and a low density of dislocations (on the order of 1012 m-2) [34,35], and most alloying elements precipitate to form second-phase particles [36,37]. In order to enhance the strength, ultra-fine grained (UFG) 7000 series Al alloys have been developed via various severe plastic deformation approaches, such as equal channel

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angular pressing (ECAP) [34-36,38,39], high-pressure torsion [40-43], multi-directional compression [44], cryogenic rolling [45,46], high-pressure sliding [47], accumulative-rolling bonding [48], friction stir welding [49], ball milling + consolidation [37], etc. In addition to precipitation strengthening of nano-scale second-phase particles, ultra-fine grains and the high density of dislocations provide grain boundary strengthening and dislocation strengthening comparable to precipitation strengthening. Interestingly, in a related study [40], Liddicoat et al. created a hierarchically nanostructured 7075 Al alloy solid solution that features nano-sized grains containing a high density of dislocations, intragranular subnanometric solute clusters and intergranular nanometric solute structures. Accordingly, the major strengthening mechanisms involve precipitation strengthening of the subnanometric solute clusters, grain boundary strengthening and dislocation strengthening. Similar to the CG 7000 series Al alloys, in UFG 7000 series Al alloys solid-solution strengthening also plays a minor role due to the precipitation of most alloying elements [36,37]. As-melt spun 7075 Al alloy has been investigated in our most recently published study [50]. Our results show that the as-melt spun 7075 Al alloy is characterized by a fine-grained (FG, average grain size 1.6 m) nearly single-phase supersaturated solid solution with a medium density of dislocations (5.71013 m-2). The high cooling rate on the order of 105-107 K/s during melt spinning induced a high undercooling degree and thus the formation of a high density of nuclei and a FG structure. Moreover, the high cooling rate resulted in an extremely short timeframe, which is responsible for the formation of nearly single-phase supersaturated solid solution and of a medium density of dislocations: the alloying elements do not have a sufficiently long timeframe for precipitation from the matrix Al; it is difficult for those completely disordered atoms in the molten state to be rearranged into perfect crystal lattice and consequently, a medium density of crystal defects were left predominantly in the form of dislocations. In contrast, the solid-solution treated commercial 7075 Al alloy features

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a CG structure (a few ten to a few hundred micrometers in average grain size [25-27,34,35]) and a low density of dislocations (on the order of 1012 m-2 [34,35]), which are attributed to grain growth and recovery/recrystallization caused by a high temperature during solid-solution treatment. Accordingly, the FG structure and the higher density of dislocations present a potential to enhance the strength of the melt spun 7075 Al alloy relative to the CG counterpart by the additional grain boundary strengthening and dislocation strengthening. However, as an age-hardened alloy, precipitation strengthening, which is attained by aging to form precipitates, makes a significant contribution to the strength of 7075 Al alloy as discussed above. Hence, the answers to the following questions associated with microstructural evolution during aging of the as-melt spun 7075 Al alloy critically influence the

strength

of

melt

spun

+

aged

7075

Al

alloy

ribbons.

(i)

How

do

sub-micrometric/micrometric sizes of grains with a medium density of dislocations (5.71013 m-2) change during aging of the as-melt spun 7075 Al alloy? The answer to this question is actually complicated, depending on the following processes: (1) Whether may grain growth occur with aging at 120 C for a prolonged period 24 h or not? (2) Given such a medium density of dislocations, whether may recrystallization occur with aging at 120 C for 24 h or not? Moreover, given such a medium density of dislocations, even though recrystallization occurs, its effect on the grain sizes is complex. It may further refine or increase the grain sizes in the as-melt spun 7075 Al alloy, depending on the density of recrystallized nuclei [51]. (ii) By what mechanism, recovery or recrystallization, is the dislocation density reduced? On the one hand, if recrystallization occurs, almost all dislocations are annihilated as a result of sink at the migrating GBs [51]. On the other hand, if only recovery occurs in absence of recrystallization, what microstructure are the dislocations evolved into and what is the remaining dislocation density? (iii) More interestingly, with a medium density of dislocations and sub-micrometric/micrometric grain sizes in the as-melt spun 7075 Al alloy, can the 5

nucleation of precipitates be enhanced during aging of the as-melt spun 7075 Al alloy such that the precipitate sizes and inter-precipitate distances are effectively reduced? If so, the microhardness in the melt spun + aged 7075 Al can be significantly improved. Based on the above discussion associated with the microstructural evolution during aging of the as-melt spun 7075 Al alloy, it can be suggested that the studies on the microstructural evolution are both of technological importance and of scientific interest. From a technological standpoint, the studies enable the understanding of the microhardness in the melt spun + aged 7075 Al alloy. From a scientific perspective, the studies reveal the mechanisms underlying the evolution of grains, dislocations and precipitates during aging of a sub-micrometric/micrometric-grained supersaturated solid solution with a medium density of dislocations, which hitherto remain poorly understood. Apparently, the studies on microstructural evolution in the present study are totally different from that reported in Ref. [50], involving the formation of submicrometric/micrometric grain size and of supersaturated solid solution via rapid solidification, and of dislocations as a result of uncompleted rearrangement of completely disordered atoms in molten state into perfect crystal lattice, as described above. Moreover, given the complexity of the microstructural evolution during aging of the as-melt spun 7075 Al alloy as discussed above, it is necessary to conduct in-depth studies of microstructural evolution. In view of the above discussion, the present study investigated the microstructural evolution during aging of the as-melt spun 7075 Al alloy, and the resultant microstructure and microhardness in the melt spun + aged 7075 Al alloy ribbons.

2. Experimental 2.1. Materials preparation The CG 7075 Al alloy plates used in the present study were commercially produced via

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chill casting followed by homogenization at 450 C to 470 C for 20 hours and hot rolling at 380 C to 410 C. The composition of the CG 7075 Al alloy plates is as follows (wt.%): Al-5.60Zn-2.50Mg-1.50Cu-0.40Fe-0.35Si-0.30Mn-0.25Cr-0.15Ti. The 7075 Al alloy ribbons were processed via melt spinning in an apparatus shown in Fig. 1. The melt spinning procedure is described as follows. First, the CG 7075 Al alloy pieces were sectioned from the aforementioned plates, put into a quartz tube and melted via induction heating under argon protection. Then, the molten alloy was ejected onto the surface of a water-cooled copper wheel with 220 mm in diameter using a pressure of 0.05 MPa through a hole with 0.5 mm in diameter at the bottom of the quartz tube, which is positioned 10 mm above the impinging point of the molten alloy onto the wheel surface; the copper wheel was rotated at a speed of 40 m/s on its surface. By doing so, the ribbons with 40 m in thickness and 2-3 mm in width were produced. Finally, the as-melt spun ribbons were aged at 120 C for 24 h (i.e., the aging recipe of T6 temper [25-27]). For comparison, the aforementioned CG 7075 Al alloy plates were solid-solution treated at 480 C for 5 h, followed by quenching and immediate aging at 120 C for 24 h, namely T6 temper [25-27].

2.2. Microhardness measurements and microstructural characterizations Microhardness of the melt spun + aged 7075 Al alloy ribbons and of the solid-solution treated + aged CG 7075 Al alloy plates was measured using a FM-ARS9000 Vickers microhardness tester with a load of 10 g and a loading time of 10 s. In order to prepare specimens for microhardness measurements, the melt spun + aged 7075 Al alloy ribbons were first ground to 20 m in thickness (i.e., nearly a half thickness of the ribbons) from either the air or the wheel sides, then mechanically polished on the ground surface. Microhardness measurements were performed in 20 - 25 randomly selected points on the aforementioned polished surface. For the solid-solution treated + aged CG 7075 Al alloy 7

plates, three planes perpendicular to each other, i.e., the planes determined by rolling direction (RD) and normal direction (ND), RD and transverse direction (TD), and ND and TD, were mechanically ground and polished. Microhardness measurements were carried out in 20 - 25 randomly selected points on each of the three polished surfaces. The microstructures of both materials were studied using optical microscopy (OM, model: Axiovert 200 MAT). For the melt spun + aged 7075 Al alloy ribbons, planar-view OM observations were conducted on the mechanically polished surface where microhardness measurements were performed as described above. The microstructure was revealed using Keller’s reagent (2 ml HF + 3 ml HCl + 5 ml HNO3 + 190 ml H2O). For the solid-solution treated + aged CG 7075 Al alloy plates, OM analysis was carried out under polarized light on the aforementioned three mechanically polished surfaces where microhardness measurements were done. In doing so, the three polished surfaces were electro-chemically etched by 60 ml fluoroboric acid in 400 ml deionized water for anodic oxidation. Based on the OM micrographs, grain sizes were measured using linear intercept method per ASTM E112 [52] to obtain the statistical distribution of grain sizes. The two materials were also characterized using transmission electron microscopy (TEM, model: JEOL JEM 2010) equipped with electron-energy loss spectroscopy (EELS) system, operated at 200 kV. For the melt spun + aged 7075 Al alloy ribbons, planar-view TEM observations were performed. The preparation of the corresponding TEM specimens is described as follows. First, the ribbons were mechanically ground to 20 m in thickness by equally grinding both the wheel and the air sides. Then, the specimens were thinned to electron transparency via twin-jet polishing using a solution of 25 vol.% nitric acid and 75 vol.% methanol at -30 C. For the solid-solution treated + aged CG 7075 Al alloy, TEM observations were carried out along RD, ND and TD. When the TEM specimens were prepared, the materials were also mechanically ground to 20 m in thickness at first, and 8

then the same subsequent operations as those for the melt spun + aged 7075 Al alloy ribbons were implemented for final thinning. The TEM studies involved microstructural observations via bright field (BF) and dark field (DF) imaging, selected area electron diffraction (SAED), and EELS analysis on the view fields observed. Herein EELS analysis was implemented to in situ measure the specimen thickness corresponding to the view fields by using Gatan Digitalmicrograph software [53] that was developed based on the Kramers–Kronig sum-rule and log-ratio methodology [54]. X-ray diffraction (XRD) analysis was conducted on the two materials using a Rigaku D/MAX-2500 diffractometer equipped with a Cu target at the speed of 0.02 and the count time of 3 s per step. For the melt spun + aged 7075 Al alloy ribbons, XRD analysis was conducted on the mechanically polished surface where microhardness measurements were performed. For the solid-solution treated + aged CG 7075 Al alloy, XRD analysis was carried out on the polished plane determined by RD and ND. It is noted that the purpose of the strategies used to mechanically grind the melt spun + aged 7075 Al alloy ribbons for the preparation of microhardness, OM, TEM and XRD specimens as described above is to measure microhardness and to study the microstructure located in the nearly half thickness of the ribbons. It is well documented that the microstructure progressively changes along the thickness of a melt spun ribbon from the wheel to air sides. The microstructure and microhardness in the nearly half thickness may represent the average scenario of microstructure and microhardness of the ribbon.

3. Results 3.1. Microhardness Fig. 2a shows the distribution of the measured Vickers microhardness of the melt spun + aged 7075 Al alloy. Figs. 2b, 2c, and 2d display that of the solid-solution treated + aged CG

9

7075 Al alloy on the planes determined by RD and ND, RD and TD, and ND and TD, respectively. By fitting these distributions using a normal probability function, the Vickers microhardness was evaluated to be 206873 MPa, 1896136 MPa, 164153 MPa, and 174751 MPa (average value  standard deviation) corresponding to Figs. 2a through 2d, respectively. The difference in microhardness on the different planes in the solid-solution treated + aged CG 7075 Al alloy can be attributed to the presence of crystallographic textures in the rolled materials. In the present study, we used the arithmetic average value of microhardness on the three planes perpendicular to each other, 1761 MPa, to represent the microhardness of the solid-solution treated + aged CG 7075 Al alloy. Comparison of the average microhardness of the two materials indicates that Vickers microhardness of the melt spun + aged 7075 Al alloy is 17% (307 MPa) higher than that of the solid-solution treated + aged CG counterpart.

3.2. Microstructure 3.2.1. The melt spun + aged 7075 Al alloy ribbons As shown in Fig. 3a, the planar-view OM microstructure of the melt spun + aged 7075 Al alloy ribbons consists of sub-micrometric and micrometric sized grains. By measuring randomly selected 447 grains from a series of OM micrographs, the statistical distribution of grain sizes was determined and presented in Fig. 3b. It is shown that the grain sizes fall in the range from 0.4 m to 5.8 m and the sub-micrometric sized grains account for 7% in number fraction. By fitting the statistical distribution using a lognormal probability function 0.9 [34,44], the grain sizes were evaluated as D= 1.8 -0.6 m (average grain size D  standard

deviation). Fig. 4 demonstrates the typical planar-view TEM microstructure at grain interiors. The TEM BF and DF images as shown in Figs. 4a and 4b, respectively, reveal the presence of 10

block-like crystallites at grain interiors. The boundaries between these crystallites comprise dense dislocation walls, whereas the dislocation densities at crystallite interiors are much lower than those at crystallite boundaries; these microstructural features can be observed more clearly in the TEM DF image (Fig. 4b). Fig. 4c presents a SAED pattern from a circled region of 0.7 m in diameter (i.e., diffraction aperture size) at the grain interior in Fig. 4a, which includes five crystallites. The SAED pattern shows less than 5 of misorientations between these crystallites, indicating that these crystallites are subgrains. The Study of misorientations between subgrains inside a large number of grains reveal that the misorientations fall in the range of 1 – 5 with a distribution of 2.50.7 (average misorientation  standard deviation). In addition, the measurements of the lengths of a large number of subgrain boundaries show 0.450.06 m-1 of the length of subgrain boundaries per unit area. The higher magnification TEM BF image in Fig. 4d, which is the enlarged micrograph of the region marked by a white rectangle in Fig. 4a, displays nanoscale precipitates at the grain interior. Based on the TEM BF images from a large number of view fields, the average size of precipitates, d , was measured to be 7.5 nm. The thickness of these view fields was in situ determined on the basis of EELS analysis and it varies from 90 nm to 120 nm. By doing so, the volumetric density of precipitates can be estimated as 3.6×1022 m-3, and the average inter-precipitate distance, s , is calculated to be 30 nm.

3.2.2. Solid-solution treated + aged CG 7075 Al Due to the occurrence of recrystallization during hot rolling and sometimes solid-solution treatment, the solid-solution treated + aged CG 7075 Al alloy exhibits the similar equiaxed grains on the three planes determined by RD and ND, RD and TD, and ND and TD, respectively. Therefore, only the typical OM microstructure on the plane determined by RD 11

and ND is shown as a representative in Fig. 5a for purpose of brevity. In a lot of randomly selected OM micrographs, 150 grains were analyzed and the statistical distribution of grain sizes was obtained, as reported in Fig. 5b. By fitting the statistical distribution with a 159.5 lognormal probability function [34,44], the grain sizes were assessed to be D= 215.0 -91.6

m (average grain size D  standard deviation). Again, given the similarity of TEM microstructure on the three planes, only the TEM BF image on the plane determined by RD and ND is shown in Fig. 6a, revealing a relatively low density of dislocations at grain interiors. Fig. 6b is the enlarged image of the region included by a white rectangle in Fig. 6a, which also shows nanoscale precipitates in the CG 7075 Al alloy. By measuring the precipitates from many TEM view fields, their average size, d , is evaluated to be 8.3 nm. Based on thickness of these view fields (90 nm - 110 nm) in situ measured via EELS analysis, the volumetric density of precipitates is estimated to be 2.8×1022 m-3 and thus the average inter-precipitate distance, s , is computed to be 33 nm.

3.3. XRD results Fig. 7a shows the XRD patterns of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG 7075 Al alloy. In the XRD pattern of the solid-solution treated + aged CG 7075 Al alloy, the intensity of 220, 331 and 420 peaks is much lower relative to that of 111, 200 and 311 peaks, suggesting the presence of crystallographic texture, which results in different magnitudes of microhardness along different crystallographic directions as presented in section 3.1. The dislocation densities in the two materials were estimated based on peak broadening after subtraction of instrumental broadening via the following procedure. First, the average dimension of coherent domains ( d CD ) and microstrain (   2 1 / 2 ) are calculated using the following equation [55]:

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 / tan 2   ( / d CD )( / tan sin  )  16   2  2

(1)

where  is the X-ray wavelength,  is the angle corresponding to the peak maximum position (i.e., Bragg angle) and  is the integral breadth corresponding to  after subtraction of instrumental broadening. Then, the dislocation density  can be estimated via the following equation [56,57]:

  2 3   2 1 / 2 / bd CD

(2)

where b = 0.286 nm is the dislocation Burgers vector in an Al alloy. For the peaks acquired in the melt spun + aged 7075 Al alloy,  / tan2  as a function 2

of  / tan sin is plotted in Fig. 7b. We approximated the melt spun + aged 7075 Al alloy as an isotropic materials. By performing a least-square fitting to  / tan2  against 2

 / tan sin , a straight line is obtained. However, when all eight peaks in the diffraction pattern of the melt-spun + aged 7075 Al alloy were used for least-square fitting to obtain the straight line, the linearity of fitting is as low as 0.2; in contrast, when the two peaks, 331 and 420 peaks are removed and the rest six peaks are used for least-square fitting, the linearity of fitting is as high as 0.98. Thus, we argue that the data in the 331 and 420 peaks are abnormal. We thus removed the two peaks during the least-square fitting. Based on the fitted straight line (Fig. 7b), d CD and   2 1 / 2 can be determined. Using Eq. 2, the dislocation density in the melt spun + aged 7075 Al alloy was calculated to be (3.70.4)1013 m-2. Since the solid-solution treated + aged CG 7075 Al alloy underwent rolling as described in Section 2.1, the crystallographic textures should be developed in the material. To that effect, peak broadening in both 111-222 and 200-400 diffraction peak pairs was utilized to evaluate d CD and   2 1 / 2 along the <111> and <100> directions, respectively [58-60]. In doing so,  / tan2  as a function of  / tan sin is plotted for the two diffraction 2

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peak pairs as shown in Fig. 7b. d CD and   2 1 / 2 were accessed based on the two straight lines connecting 111 to 222, and 200 to 400. Then, the dislocation densities corresponding to <111> and <100> can be calculated using Eq. 2. In the present study, the arithmetic average value of the dislocation densities along the two directions is considered as the overall dislocation density in the solid-solution treated + aged CG 7075 Al alloy, which is determined to be (1.30.3)1013 m-2. The XRD patterns are also used to determine the lattice constants of the two materials based on the Nelson-Riley function [61], i.e., (cos 2  / sin   cos 2  /  ) / 2 . For each of the two materials, the apparent lattice constant a (θ) calculated using the Bragg equation against (cos 2  / sin   cos 2  /  ) / 2 is plotted, as shown in Fig. 7c; by fitting these data points using

least-square method, a straight line is obtained and the lattice constant is calculated by extrapolation at (cos 2  / sin   cos 2  /  ) / 2 =0 (i.e., =90). By using the Nelson-Riley function, the lattice constants a are determined to be 4.0580.002 Å and 4.0590.001 Å for the melt-spun + aged 7075 Al alloy and the solid-solution treated + aged CG 7075 Al alloy, respectively.

4. Discussion 4.1. Microstructural evolution during aging of the as-melt spun 7075 Al alloy ribbons The microstructure of the melt spun + aged 7075 Al alloy evolved from that of the as-melt spun 7075 Al alloy. As reported in our most recently published study [50], the microstructure of the as-melt spun 7075 Al alloy comprises sub-micrometric/micrometric sized equiaxed grains (0.2 m - 5.5 m of grain size range and 1.6 m of average grain size) with a medium density of tangled dislocations (5.7×1013 m-2) at grain interiors. After aging, equiaxed grains were retained and no essential changes occurred for grain sizes (0.4 14

m - 5.8 m of grain size range and 1.8 m of average grain size); however, at grain interiors subgrain structure replaced the tangled dislocation structure, as presented in Section 3.2.1, and moreover, the dislocation density was reduced from 5.7×1013 m-2 to 3.7×1013 m-2, as reported in Section 3.3. Based on these microstructural observations, we suggest that recovery rather than recrystallization occurred during aging of the as-melt spun 7075 Al alloy ribbons [51]. Absence of recrystallization may be attributable to the stored energy (primary in dislocations) lower than the critical magnitude to trigger recrystallization under the aging temperature 120 C [51]. Indeed, the stored energy in dislocations, E, can be estimated using the following equation [51]: E  Gb 2 / 2

(3)

where G=26 GPa is the shear modulus in an Al alloy. The stored energy is then assessed as 6  104 J/m3 (herein =5.7×1013 m-2 as calculated in Section 3.3). Inspection of the published scientific literature shows that the stored energy that can activate recrystallization generally ranges 106 – 107 J/m3 [51,62], and the stored energy that enables recovery is generally on the order of 104 – 105 J/m3 [51]. Thus, recrystallization cannot be triggered and recovery occurred during aging of the as-melt spun 7075 Al alloy. In the present study, the mechanisms underlying the decrease in the dislocation density during aging (recovery process) may involve dislocation annihilation via counteracting of opposite-sign dislocations and via sink at grain boundaries, and dislocation rearrangement to form subgrain boundaries [51]. Given the fact that grain sizes essentially remained unchanged during aging, grain growth has been retarded for the following three reasons. (i) GBs exhibit a relatively low mobility due to the aging temperature as low as 120 C. (ii) At the early stage of aging, solute atoms at GBs impeded GB migration [51,63]. In the as-melt spun 7075 Al alloy, GBs contain the solutes with the concentration higher than the overall concentration 11 wt.% due to solute 15

segregation. (iii) At the later stage of aging, precipitates can effectively retard grain growth by pinning GBs. In the melt spun + aged 7075 Al alloy, the volume fraction of precipitates can be calculated to be f = 2.26% using the equation [44]: f  (2 π/ 3)[2s / d  (8 / 3)1 / 2 ]2 . The limit average grain size for grain growth is then estimated as 0.44 m using the equation [51]: D lim  4d / 3 f . Another microstructural evolution during aging is the occurrence of precipitation. The melt spun + aged 7075 Al alloy ribbons present a higher volumetric density of precipitates (i.e., a smaller inter-precipitate distance) and a smaller average precipitate size than those in the solid-solution treated + aged CG counterpart. The formation of this microstructural feature can be explained as follows. In a 7075 Al alloy, precipitation during aging follows the sequence [25-27]: saturated solid solution  Guinier-Preston (G-P) zones   phase   phase. In other words,  and  precipitates evolve from the G-P zones. Hence, the density of nuclei of G-P zones critically affects the volumetric density of precipitates and thus the inter-precipitate distance in the aged alloy. The difference in microstructural characteristics between the as-melt spun 7075 Al alloy and the solid-solution treated CG 7075 Al alloy, including smaller grain sizes and higher dislocation density in the former [50], may influence the densities of nuclei of G-P zones during aging of the two materials. However, based on the published studies [35,37], it appears that precipitates are uniformly distributed at grain boundaries and interiors in the aged both CG and UFG 7075 Al alloys, suggesting a limited influence of grain boundaries (i.e., grain size) on the density of nuclei of G-P zones. Inspection of the published scientific literature [64] indicates that dislocations do not significantly enhance the nucleation of G-P zones. On the basis of the preceding discussion, other factors than grain sizes and dislocations may be responsible for the higher volumetric density of precipitates in the melt spun + aged 7075 Al alloy. 16

In fact, it has been reported by the published literature [35,65] that vacancies may effectively promote the nucleation of G-P zones during aging of 7000 series Al alloys. An early study by Embury and Nicholson [65] shows that in an Al-Zn-Mg alloy the concentration and distribution of vacant lattice sites critically affected the nucleation of G-P zones. In another related work, Zhao et al. [35] observed a higher volumetric density of G-P zones in the UFG 7075 Al alloy produced by solid-solution treatment of CG counterpart + ECAP + natural aging than that in the CG counterpart processed by solid-solution treatment + natural aging. The higher volumetric density of G-P zones in the UFG 7075 Al alloy may be attributable to vacancies created during ECAP. In the present study, as-melt spun 7075 Al alloy and solid-solution treated CG 7075 Al alloy were obtained by quenching from the temperature higher than the liquidus temperature of 7075 Al alloy 635 C [25], and from the solid-solution temperature 480 C, respectively, inducing a higher concentration of vacancies in the former. Despite annihilation of a proportion of excess vacancies at room temperature and during heating to the aging temperature [51], the former may still contain a higher concentration of vacancies, leading to the higher density of nuclei of G-P zones and thus the higher volumetric density of precipitates in the melt spun + aged 7075 Al alloy. Nevertheless, inspection of the published scientific literature reveals that the concentrations of equilibrium vacancies are substantially low [66,67] in materials, e.g., on the order of 10-4 to 10-3 atomic fraction in Al even at a temperature close to the melting point. Consequently, the difference in the concentrations of vacancies between the melt spun + aged 7075 Al alloy and the solid-solution treated + aged CG 7075 Al alloy may be quite small, i.e., a slightly higher volumetric density of precipitates in the melt spun + aged 7075 Al alloy. As presented in section 3.3, the melt spun + aged 7075 Al alloy (a = 4.0580.002 Å) exhibits almost the same lattice constant as that of the solid-solution treated + aged CG counterpart (a = 4.0590.001 Å). Thus, the dissolved and precipitated amounts of alloying elements in the melt spun + 17

aged 7075 Al alloy should be almost the same as those in the solid-solution treated + aged CG 7075 Al alloy. In other words, the melt spun + aged 7075 Al alloy contains almost the same volume fraction of precipitates as that in the solid-solution treated + aged CG counterpart. Given the slightly higher volumetric density of precipitates in the melt spun + aged 7075 Al alloy, the dimensions of precipitates in the melt spun + aged 7075 Al alloy are slightly smaller.

4.2. The increased microhardness of the melt spun + aged 7075 Al alloy ribbons As presented in Section 3.1, the melt spun + aged 7075 Al alloy exhibits 17% (307 MPa) higher microhardness than that of the solid-solution treated + aged CG counterpart. Given the fact that the equations correlating microhardness with strengthening mechanisms are unavailable in the published scientific literature, in order to rationalize the enhanced microhardness, we rely on the calculated contributions of various strengthening mechanisms to the improvement in yield strength and the Tabor equation connecting yield strength to microhradness. In both melt spun + aged 7075 Al alloy and the solid-solution treated + aged CG counterpart, the following strengthening mechanisms may be operative: precipitation strengthening, grain boundary strengthening, dislocation strengthening and solid solution strengthening. Due to the presence of subgrains in the melt spun + aged 7075 Al alloy, subgrain boundaries may provide additional strengthening. The contributions of various strengthening mechanisms to the improvements in yield strengths of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG counterpart are analyzed as follows. The contribution of precipitation strengthening to the improvement in the yield strength can be calculated using the following equation [68]:

18

 p  M

0.4Gb ln( d / b) π(1   )1 / 2 s

(4)

where  p is the contribution of precipitation strengthening to the improvement in the yield strength, M=3.06 is the Taylor factor, d is the average diameter of circular cross-sections in random planes for a spherical precipitate, d  2 / 3 d , with an assumption of spherical precipitates in the present study, and =0.33 is the poisson ratio. Despite the presence of crystallographic texture in the solid solution treated + aged CG 7075 Al alloy, leading to different M values along different crystallographic directions, we still take M=3.06 corresponding to isotropic materials to calculate the contribution of precipitation strengthening in this material. By doing so, we can calculate the average contribution of precipitation strengthening over different crystallographic directions. As presented in Section 3.1, the measured microhardness of the solid-solution treated + aged CG 7075 Al alloy is an average one over different crystallographic directions. Hence, the calculated value can still be used in the comparison with the measured one. The contributions of precipitation strengthening to the improvements in the yield strengths of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG counterpart are evaluated as 361 MPa and 340 MPa, respectively ( d =7.5 nm and s =30 nm for the melt spun + aged 7075 Al alloy; d =8.3 nm and s =33 nm for the solid-solution treated + aged CG 7075 Al alloy, as presented

in Sections 3.2.1 and 3.2.2). The contribution of grain boundary strengthening to the improvement in yield strength can be assessed using the Hall-Petch equation:

 g  k D

1 / 2

(5)

where  g is the contribution of grain boundary strengthening to the improvement in the yield strength, and k is the Hall-Petch slope (k=120 MPa m1/2 for 7075 Al alloy [25,37]). 19

The contributions of grain boundary strengthening to the improvements in yield strengths of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG counterpart are then assessed as 89 MPa and 8 MPa, respectively ( D =1.8 m for the melt spun + aged 7075 Al alloy and D =215.0 m for the solid-solution treated + aged CG 7075 Al alloy, as presented in Sections 3.2.1 and 3.2.2). The contribution of dislocation strengthening to the improvement in the yield strength can be determined by the following equation [69]:  dis  MGb 1 / 2

(6)

where  dis is the contribution of dislocation strengthening to the improvement in the yield strength,  =0.24 is a constant. Again, we still take M=3.06 corresponding to isotropic materials to calculate the contribution of dislocation strengthening in the the solid solution treated + aged CG 7075 Al alloy. The contributions of dislocation strengthening to the improvements in yield strengths of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG counterpart are computed to be 33 MPa and 20 MPa, respectively (=3.71013 m-2 for the melt spun + aged 7075 Al alloy and =1.31013 m-2 for the solid-solution treated + aged CG 7075 Al alloy, as presented in Section 3.3). Given almost the same lattice constants (4.0580.002 Å and 4.0590.001 Å as reported in section 3.3) and thus concentrations of solutes for the two materials, the contributions of solid solution strengthening to the improvements in yield strengths of the melt spun + aged 7075 Al alloy and of the solid-solution treated + aged CG counterpart can be considered to be similar to each other. Given the unavailability of concentrations of solutes, the contributions of solid-solution strengthening cannot be quantitatively evaluated. However, the upper bound of the contributions of solid-solution strengthening in a T6 treated 7075 Al alloy can be estimated as 82 MPa [37], assuming the dissolution of all alloying elements in the Al matrix.

20

Precipitation of all alloying elements corresponds to the lower bound of the contributions of solid-solution strengthening (0 MPa). The contribution of subgrain boundary strengthening to the improvement in the yield strength, sg, can be evaluated using the following equation [70]:  sg  MG(1.5bS v  sg )1 / 2

(7)

where Sv is the subgrain boundary area per unit volume, and sg is the average misorientation between subgrains. Sv can be correlated with the subgrain boundary length per unit area, La, using the equation of stereology [71]: Sv=(4/) La. Based on the measured sg = 2.5 and La = 0.45 m-1 as reported in section 3.2.1, sg can be calculated to be 62 MPa in the melt-spun + aged 7075 Al alloy. Apparently, no subgrain boundary strengthening is provided (sg=0 MPa) in the solid-solution treated + aged CG 7075 Al alloy due to absence of subgrain boundaries. Based on the preceding calculations, the total improvements in yield strengths of the melt spun + aged 7075 Al alloy and of the solid solution treated + aged CG 7075 Al alloy imparted by various strengthening mechanisms are 545 - 627 MPa and 368 - 450 MPa, respectively. Herein the upper and lower bounds in the total improvements in yield strengths correspond to the upper and lower bounds of the contributions of solid-solution strengthening as analyzed above. In order to correlate the improvement in yield strength (y ) with that in microhardness (H), we implement the modified Tabor relation [72]:  y  (H / 3)(0.1) n

(8)

where n is the work hardening index. It can be derived that n equals to the true strain corresponding to the ultimate tensile strength [69]. Inspection of the published studies [37,39,40] shows that the engineering strains corresponding to the ultimate tensile strengths of the T6 treated CG 7075 Al alloy range from 9% to 11%. By taking the arithmetic average 21

value of the engineering strains, 10%, the n value of the T6 treated CG 7075 Al alloy is 0.095. In Ref. [37], the engineering strain corresponding to the ultimate tensile strength of the T6 treated 7075 Al alloy of average grain size 1 m was reported to be 5%. Given that the melt spun + aged 7075 Al alloy exhibits similar precipitates and average grain size to those in the T6 treated 7075 Al alloy in Ref. [37], it is considered appropriate to assume that the work-hardening ability of the melt spun + aged 7075 Al alloy is similar to that of the T6 treated 7075 Al alloy in Ref. [37]. In other words, we can presumably take 5% as the engineering strain corresponding to the ultimate tensile strength in the melt spun + aged 7075 Al alloy, i.e., n=0.048. Using Eq. 8, the calculated improvements in microhardness imparted by various strengthening mechanisms are 1826 - 2101 MPa and 1374 - 1680 MPa for the melt spun + aged 7075 Al alloy and the solid-solution treated + aged CG counterpart, respectively. Here the upper and lower bounds of the improvements in microhardness in the two materials correspond to the upper and lower bounds of the contributions of solid-solution strengthening as discussed above. Thus, the calculated increase in microhardness of the melt spun + aged 7075 Al alloy over that of the solid-solution treated + aged CG counterpart is 452 - 421 MPa, which is higher than the measured one (307 MPa as presented in section 3.1). We argue that the lower value of the measured increase in microhardness is likely attributed to the presence of pores in the melt-spun + aged 7075 Al alloy.

5. Conclusions In the present study, the microhardness and the corresponding microstructure of the 7075 Al alloy ribbons melt spun plus aged at 120 C for 24 h (the aging recipe of T6 temper), as well as the microstructural evolution during aging of the as-melt spun 7075 Al alloy, were investigated. Our results indicate that the melt spun + aged 7075 Al alloy ribbons exhibit 17% (307 MPa) higher Vickers microhardness than that of the CG counterpart T6 treated, i.e., 22

solid-solution treated at 480 C for 5 h plus aged at 120C for 24 h. The increased microhardness of the melt spun + aged 7075 Al alloy over that of the solid-solution treated + aged CG counterpart was attributed to the former’s microstructural characteristics: FG structure (average grain size 1.8 m) consisting of sub-micrometric and micrometric sized grains, intragranular subgrains, and nanoscale precipitates with 30 nm in inter-precipitate distance in the former, relative to CG structure, the absence of intragranular subgrains and nanoscale precipitates with 33 nm in inter-precipitate distance in the latter. The microstructure in the melt spun + aged 7075 Al alloy was attained via recovery of the as-melt spun FG 7075 Al alloy containing a medium density of dislocations, and almost no recrystallization and grain growth occurred, during aging. A slightly higher concentration of vacancies in the as-melt spun 7075 Al alloy than that in the solid-solution treated CG 7075 Al alloy promoted the formation of the slightly higher volumetric density of precipitates and thus of the slightly smaller inter-precipitate distance in the melt spun + aged 7075 Al alloy than those in the solid-solution treated + aged CG 7075 Al alloy. Based on the microstructural analysis, the contributions of precipitation, grain boundary, dislocation, solid solution and subgrain boundary strengthening to the improvements in the yield strengths of the melt-spun + aged 7075 Al alloy and of the solid-solution treated + aged CG 7075 Al alloy were calculated. By accurately correlating the calculated improvements in yield strengths with the improvements in microhardness using the modified Tabor relation, the increased microhardness of the melt-spun + aged 7075 Al alloy over that of the solid-solution treated + aged CG 7075 Al alloy was rationalized. Based on the calculations, we suggest that the higher microhardness of the melt spun + aged 7075 Al alloy can primarily be attributed to refined grain size, the presence of subgrain boundaries and smaller inter-precipitate distance in the alloy.

23

Acknowledgments Y.J. Lin was supported by the One-Hundred Talents Project (Grant No. 2010100005), Hebei Province, China and the Fundamental Research Funds for the Central Universities (WUT: 2016IVA003). L.M. Wang was funded by National Natural Science Foundation of China (Grant No. 51131002). H.L. Geng’s contribution to OM analysis of CG 7075 Al alloy is gratefully acknowledged.

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J.R.

Cahoon,

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2(1971)1979-1983.

28

Kutzak,

Metall.

Mater.

Trans.

B,

Fig. 1The photograph of the melt spinning apparatus used in the present study.

29

a

b

30

c

d

Fig. 2The Vickers microhardness and the fitted normal distribution curves of (a) the melt spun + aged 7075 Al alloy, and of (b), (c) and (d) the solid-solution treated + aged CG 7075 Al alloy on the planes determined by RD and ND, RD and TD, and ND and TD, respectively. 31

а

Fig. 3 The typical planar-view OM microstructure of the melt spun + aged 7075 Al alloy in the nearly half thickness of the ribbons: (a) the OM micrograph showing sub-micrometric and micrometric sized equiaxed grains, and (b) the statistical distribution of grain sizes and the fitted lognormal distribution curve.

32

33

34

Fig. 4The typical planar-view TEM microstructure of the melt spun + aged 7075 Al alloy in the nearly half thickness of the ribbons: (a) the TEM BF image and (b) the TEM DF image corresponding to (a) showing subgrain structure at grain interiors and subgrain boundaries comprising dense dislocation walls, (c) the SAED pattern from the region marked by the white circle in (a) showing misorientations between subgrains, and (d) the higher magnification TEM BF image from the region marked by the white rectangle in (a) showing nanoscale precipitates.

35

Fig. 5The typical OM microstructure of the solid-solution treated + aged CG 7075 Al alloy: (a) the OM micrograph showing CG structure, and (b) the statistical distribution of grain sizes and the fitted lognormal distribution curve. 36

Fig. 6The typical TEM microstructure of the solid-solution treated + aged 7075 Al alloy on the plane determined by RD and ND: (a) the TEM BF image showing a low density of dislocations, and (b) the higher magnification TEM BF image from the region marked by the white rectangle in (a) showing nanoscale precipitates.

37

a

b

38

Fig. 7XRD results of both the melt spun + aged 7075 Al alloy and the solid-solution treated + aged CG 7075 Al alloy: (a) the XRD patterns, (b)  / tan2  as a function of 2

 / tan sin  , together with the straight line obtained using a least-square fit for

the melt spun + aged 7075 Al alloy and the straight lines corresponding to diffraction peak pairs 111-222 and 200-400 for the solid-solution treated + aged CG 7075 Al alloy, and (c) the apparent lattice constant a () against the value of the Nelson-Riley function, (cos 2  / sin   cos 2  /  ) / 2 , as well as the straight line acquired by fitting the data points using the least-square method.

39