Characterization of electron beam welded AA2024

Vacuum 85 (2010) 268e282

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Characterization of electron beam welded AA2024 P. Wanjara a, *, M. Brochu b a b

National Research Council Canada, Institute for Aerospace Research, Aerospace Manufacturing Technology Centre, 5145 Decelles Avenue, Montreal, Quebec H3S 2S4, Canada McGill University, Department of Mining and Materials Engineering, 3610 University Street, Wong Building, Montreal, Quebec H3A 2B2, Canada

a r t i c l e i n f o

a b s t r a c t

Article history: Received 4 February 2010 Received in revised form 22 June 2010 Accepted 22 June 2010

For aerospace manufacturing, the perseverance for improving performance (high strength to density ratio) and reducing weight and costs has motivated consideration of welding techniques applicable to aluminum alloys. During fusion welding of aluminum alloy (AA) 2024, the avoidance of defects (e.g., porosity, oxides, solidification cracking, undercutting) and the optimization of the microstructureproperty characteristics are of critical concern. In this work, AA2024 was electron beam (EB) welded as part of a study to determine the influence of parametric conditions on the characteristics of the weldment to optimize the joining process. Specifically, the evolution in the weld geometry, microstructure and mechanical properties was examined as a function of the process conditions, including beam current, beam focus, beam oscillation, and welding speed. For optimized parametric conditions, microstructural examination of the joints revealed narrow fusion and heat-affected zones comprising of dendritic structures without the occurrence of defects that enabled a maximized joint efficiency. Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved.

Keywords: Aluminum alloy 2024 Electron beam welding Mechanical properties Microstructure Thermal analysis

1. Introduction In recent years, there is renewed interest for altering the fabrication of aerospace components by replacing conventionally manufactured assemblies, produced either by machining forgings and castings or mechanically fastening [1], with welded joints. In particular, the aerospace industry is continually striving for improvements in performance with maximized strength-to-weight ratios as well as reductions in the fabrication costs through shorter lead-times and lower buy-to-fly ratios. Hence, the utilization of welding without concomitant increases in the section thickness of the joint, would present considerable advantages in the manufacturing process for aerospace components, both in terms of weight savings and cost reductions (shorter production and leadtimes, lower buy-to-fly ratio). At present, the weldability of current aerospace aluminum alloys, including AA2024, suffers from various difficulties, including the occurrence of a large degraded fusion zone [2,3], solidification cracking during welding [4] and porosity in the weld [5,6]. Also, there is a need to develop optimized post-weld microstructures and properties [7] and control the residual stresses and distortion in the weld assembly [8]. Although tungsten inert gas (TIG) welding is usually selected as the fusion joining technique for 2000 and 5000 series aerospace alloys, and for structural 6000 series

* Corresponding author. Tel.: þ1 514 283 9380; fax: þ1 514 283 9445. E-mail address: [email protected] (P. Wanjara).

extrusions, previous work on electron beam (EB) welding of heat treatable aluminum alloys has indicated superior mechanical properties due to the lower heat input conditions and narrow profile of the fusion region [9e11]. Specifically, a higher energy density process, such as EB welding, has a greater potential for joining aluminum alloys due to the keyhole process that permits faster heating and cooling rates, which, in turn, enables a narrower weld bead with a minimal heat-affected zone (HAZ). Hence, the different penetration mechanism (keyhole), environment conditions (vacuum), and thermal history of heating, melting and cooling for EB welding as compared to TIG welding may provide the reduction in thermal strain needed to minimize metallurgical damage during joining of aluminum alloys. Although previous work on EB welding has been performed on various aluminum alloys [10e15] over a range of thickness values, a systematic investigation of the role of processing conditions on the weld bead and penetration profiles, microstructural characteristics and mechanical properties remains lacking. In this study, such an investigation was undertaken to develop processing conditions for optimized and controlled EB welding of AA2024. 2. Experimental methods 2.1. Material condition AA2024 in the T3 condition (solution heat treated, cold worked and naturally aged) was received in sheet form with a nominal thickness of 5 mm and a chemical composition as given in Table 1.

0042-207X/$ e see front matter Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.vacuum.2010.06.007

P. Wanjara, M. Brochu / Vacuum 85 (2010) 268e282

welding system and oriented to permit welding perpendicular to the rolling (or L) direction.

Table 1 Nominal chemical composition of AA2024 (wt.%). Element

269

Cu

Mg

Mn

Si

Fe

Zn

Ti

Cr

Other

Al

4.5

1.5

0.6

0.11

0.25

0.09

0.02

0.01

0.05

Bal.

The as-received microstructure reveals the typical pancake-shaped grain structure (Fig. 1) with the grains elongated in both the longitudinal and long transverse directions. The average grain size of the wrought as-received sheet was roughly 34 mm in the short transverse (ST) direction, 77 mm in the longitudinal (L) or rolling direction, and 36 mm in the long transverse (LT) direction. The detection of specific elements using energy dispersive spectrometry (EDS) suggests that the following intermetallic compounds were present in the as-received material, including Al2CuMg (S phase), Al6(Fe, Mn, Cu) and Al7Cu2Fe, which are the most common phases in AA2024-T3 [16,17]. In AleCueMg alloys, the main hardening phase depends on the ratio of Cu to Mg (in wt.%). For the AA2024-T3 material used in the present study, since the ratio is roughly 3:1, the equilibrium precipitating phase is reported to be Al2CuMg, the S phase [17], which is consistent with the microstructural results obtained for the present work by SEM and EDS of the parent material. The average Vickers microhardness of the AA2024-T3 sheet in the as-received condition was 145  6 HRV.

2.3. Welding condition Autogenous welds, roughly 10 cm in seam length, were joined along the center of the joint using a Sciaky W2000 EB welding machine (60 kV, 700 mA, 42 kW) operating with a vacuum atmosphere better than 6  103 Pa. For EB welding of AA2024-T3 sheets, a single pass welding procedure was applied with systematic variation in the process parameters to examine the influence of beam current, beam focus, beam oscillation and welding speed. Specifically, at a constant accelerating voltage (AV), beam focus (BF) and welding speed (WS), the beam current (BC) was systematically increased to define the operating window permitting full penetration welding whilst minimizing weld reinforcement, undercutting, concavity, porosity and cracking. The effect of beam focus location was examined by varying the z-axis position of the beam from a distance of 2 mm above the top surface to 15 mm below the bottom surface. For optimized beam current and focus conditions, the electron beam was oscillated to have a circular path with a diameter (f) between 2.5 and 3.0 mm at a constant frequency of 10 kHz. Finally to obtain the optimum processing window for joining 5 mm thick AA2024, the influence of the welding speed on the characteristics of weld was examined between 27.5 and 38 mm s1.

2.2. Joining condition 2.4. Sample preparation After sectioning the AA2024-T3 sheet into 10 cm  15 cm coupons, the edges to be welded were milled and prepared for butt welding by degreasing the joint surfaces with acetone, followed by grinding with a scouring pad to remove the surface oxides, and final cleaning in ethanol to remove any surface debris according to ISO standard TR 17671-7. It is noteworthy that careful preparation of the joint surface has been reported to be instrumental for reducing the level of porosity in aluminum alloy weldments [15]. The coupons were clamped with a torque of 14 kg m by using bolts in a restraint fixture designed for ‘no gap or clearance’ (tight fitting square butt) at the groove and a backing plate of AA2024-T3 material. The fixture was placed on the rotary table of the EB

After EB welding, each weld was examined visually and up to 20 using a stereomicroscope (Olympus SZX-12) to characterize the penetration, weld bead features and surface quality. The welded sheets were then sectioned transverse to the EB welds (L direction) to extract specimens for microscopic examinations. Each sample was prepared for metallographic analysis by mounting in a cold setting epoxy resin followed by applying automated grinding and polishing techniques. To reveal the grain boundary structure in the as-received AA2024-T3 sheet, the specimens were immersed in a modified Graff and Sargent reagent (15.5 ml HNO3, 2 ml HF, and 3 g CrO3 in 84 ml H2O) for 1e2 min, while the general

Fig. 1. Optical micrographs of the initial structural appearance of the as-received AA2024-T3 sheet.

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microstructure of the fusion zone and HAZ was revealed using Keller’s reagent (3 ml HCl, 5 ml HNO3, 2 ml HF in 190 ml H2O). 2.5. Metallographic characterization

0.4

a

4.9

0.0

4.7

-0.4

4.5

-0.8

4.3

-1.2 20

40 60 Beam Current (mA)

Penetration (mm)

5.1

0.4

4.9

0.0

4.7

-0.4

4.5

-0.8

4.3

-1.2 40 60 Beam Current (mA)

2.7. Mechanical property testing 2.7.1. Microhardness testing The microhardness profile across the welded sheets was measured using a Vickers microhardness machine (Struers Duramin A300) with a 200 g test load. An average of at least three hardness values was performed for each location to obtain a nominal microhardness profile for each EB welding process condition. 2.7.2. Shear punch testing The room temperature mechanical properties were determined by the shear punch test method, which is a small-specimen testing technique that enables the characterization of the flow behavior of the material, including elasticity followed by yielding, work hardening to a maximum stress and fracture [18,19]. The shear punch test (SPT), initially developed for determining the strength and ductility from a small test volume of irradiated material [20], is based on blanking a circular disk from a fixed sheet specimen. In

80

b

20

To evaluate the precipitation evolution in the weld region, differential scanning calorimetry (DSC) measurements were performed in a TA Instrument Q1000 equipped with auto-correction for baseline drift. The DSC samples were roughly 1 mm thick discs with a diameter of 3 mm and a weight of 20 mg. For the welded condition, at least three samples were extracted from within the fusion zone in a direction parallel to the weld joint. For each sample, the thermal cycle consisted of heating at 10  C min1 up to 500  C in a cell containing a nitrogen shielding gas atmosphere to minimize oxidation.

Concavity/Reinforcement (mm)

Penetration (mm)

5.1

Concavity/Reinforcement (mm)

An optical microscope (Olympus GX-71) coupled with image analysis software (AnalySIS Five) was used to correlate the structural characteristics of the weldments (fusion zone and HAZ) to the welding parameters. The grain size of the wrought microstructure of the as-received material was measured by using the circular intercept method with the image analysis software according to standard ASTM E112. For statistically representative results, 300e400 grains were sampled for the as-received metal in each direction (L, LT and ST). To determine the dendrite arm spacing, the mean value was calculated from the measurement (DAS module in AnalySIS Five) of at least 30e50 linear sections, each comprising of 5e10 dendrite arms. For statistically representative results, an average of at least 5 weld cross-sections were used for quantification of the weld characteristics (fusion zone size, fusion zone area, concavity, HAZ area), which were determined using linear and area functions in AnalySIS Five. For analysis of the precipitates in the asreceived material and after EB welding in the fusion and HAZ regions, a JEOL-840 scanning electron microscope was used at an accelerating voltage of 15 keV and a working distance of 15 mm, as well as with EDS for elemental analysis.

2.6. Differential scanning calorimetry

80

Fig. 2. Effect of beam current on weld penetration depth, weld reinforcement and concavity for a GTWD of (a) 17.5 cm and (b) 20 cm. Welding conditions of 50 keV at 30 mm s1 with the beam focus point located at the center of the sheet thickness (BF ¼ 2.5 mm). Open and closed markers represent incomplete penetration and complete penetration, respectively.

Fig. 3. Definitions for measurements performed on the EB welding of AA2024-T3 (SL is seam length, SW is seam width, RR is root reinforcement, WR is weld reinforcement and WP is weld penetration).

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this work, rectangular slices from the parent material and fusion zone were cut and ground (600 grit SiC finish) to roughly 0.33  0.02 mm in thickness and placed between a die and a washer assembly fitted on a lower housing. A flat tip cylindrical punch (1.5 mm in diameter), fitted in an upper housing, was forced through the specimen, the action of which punched a circular disk. Since the upper and lower housings were attached to a servohydraulic materials testing system (MTS Model 810), load-displacement data could be obtained from shear punch testing. As correlations between the shear punch data and tensile test data have been identified, the tensile flow properties, including yield strength, ultimate strength and elongation can be determined for the weld microstructure using the following empirical equation for the strength [19e22]:

ss ¼

L ¼ Cs st 2 prt

271

with a GTWD of 17.5 cm and 20 cm requires a beam current of 45 mA and 50 mA, respectively. For values below the beam current threshold, incomplete penetration was observed and weld reinforcement between 2 and 4% of the sheet thickness was measured, based on the nomenclature given in Fig. 3. Alternatively, the onset of complete weld penetration at the beam current threshold was observed to result in a concavity of roughly 2e4% of the sheet thickness. Regardless of the GTWD, beyond the beam current threshold value, the extent of

(1)

where ss (MPa) is the yield or ultimate stress in SPT, L (N) is either the yield (Ly) or maximum (Lm) load from the shear loaddisplacement curve, r (mm) is the punch radius, t (mm) is the specimen thickness, st (MPa) is the corresponding yield or maximum stress in tension and CS is an empirical correlation coefficient that was determined in this work to be 0.50 and 0.55 for the yield and ultimate tensile strengths, respectively. On this basis, these empirical coefficient values can be used to calculate the tensile yield strength and ultimate tensile strength, hereafter referred to as YS and UTS, from the corresponding measured values of ssYS or ssUTS obtained from SPT. In this work, the 95% confidence interval on the predicted stress value for a given load value was determined to be 2% for the YS and 1.5% for the UTS. It is noteworthy that the load-displacement behavior of the AA2024 microstructures examined in the present work exhibited gradual yielding, and, to minimize errors in identifying the point of deviation from linearity, the value for Ly was determined at 0.1% strain for each measured SPT curve. The empirical equation used for determining the ductility was [19,21]:

3s ¼

Df  Dy ¼ C d 3t t

(2)

where 3s (%) is the elongation in SPT, Df (mm) is the displacement to fracture, Dy (mm) is the displacement to the yield, t (mm) is the specimen thickness, 3t (%) is the tensile elongation and Cd, the correlation coefficient for the ductility, was determined to be 1. Intuitively, this linear equation is satisfying in that Df  Dy should be zero for a completely brittle material (i.e., no plasticity) and Df  Dy should approach the specimen thickness for a completely ductile material. The 95% confidence interval on the percent elongation for a given Df/t was determined to be 1.5% of the elongation. 3. Results and discussion 3.1. Macroscopic examination 3.1.1. Effect of beam current The integrity of the EB welds in AA2024-T3 was examined for defects (e.g., inadequate penetration, concavity, cracking, etc.) for the various welding conditions. Specifically with the other parameters constant, the effect of beam current, shown in Fig. 2 indicates that for a particular gun-to-work distance (GTWD) value, a threshold value is needed to obtain complete weld penetration in the 5 mm AA2024-T3 sheet during welding. At 50 keV and with the beam focus located at the center of the sheet thickness (BF of 2.5 mm), it was determined that complete weld penetration

Fig. 4. EB welds in 5 mm thick AA2024-T3 at 2.5 kW, 30 mm s1, 2.5 mm BF and 20 cm GTWD: (a) butt pass and (b) butt pass and cosmetic pass (at 25 mA and 3 mm f beam).

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the concavity was observed to rapidly increase beyond the maximum permissible value of 10%, as specified by EB welding standards SAE AMS 2681B, SAE AMS 2680C and AWSD17. Also, the effect of the GTWD over the limited range investigated was negligible, albeit considerable deviation from the optimum value of 20 cm has been conveyed to influence the welding characteristics [22]. Hence, the application of EB welding for autogenous joining of AA2024-T3 material must be performed at the beam current threshold to minimize weld concavity, as illustrated in Fig. 4a. Further reductions in the depth of the concave surface can be achieved through the application of a cosmetic pass after welding,

which, as illustrated in Fig. 4b, can decrease the concavity to less than 1%. Previous work on EB welding of aluminum alloys has indicated that parametric conditions giving uniform weld width and complete penetration without internal weld defects, tend to promote undercutting and/or concavity at the weld surface and the application of a second pass with or without a filler is frequently considered for improving the upper bead contour [9]. 3.1.2. Effect of beam focus position Variation in the location of the beam focus from the top (BF 0 mm) to below the bottom surface of the weld (BF 15 mm) was

Fig. 5. EB welds in 5 mm thick AA2024-T3 at 2.5 kW, 30 mm s1 and 20 cm GTWD: (a) BF of 0 mm, (b) BF of 2.5 mm, (c) BF of 5 mm and (d) BF of 10 mm.

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observed to influence the morphology of the fusion zone and mostly the appearance of the upper bead, as indicated in Figs. 5 and 6, respectively. In particular, location of the beam focus on the top surface (Fig. 5a) and in the center of the sheet thickness (Fig. 5b) was observed to result in a relatively uniform weld width as compared to focusing at, or below, the bottom surface (Fig. 5c and d). This, of course, is due to the significance of the beam focus position for the beam spot size and energy density distribution in the material. The location of the beam focus is where the spot size is smallest or minimal for the welding parameters selected. Hence for welding with a focused beam located on the top surface or at the center of the sheet thickness, the spot size increases from this minimal value within the sheet: (1) from top to bottom of the sheet or (2) from middle to top and middle to bottom (hourglass). Under such conditions, the material experiences a high energy density beam that localizes the heat input (steep temperature gradient) to the weld interface, which is typical of EB welding. Alternatively, a defocused beam (beam focus located at or below the bottom surface) results in a lower energy density in the material. Under such conditions, the heat inputted extends appreciably beyond the weld interface and is manifested by the widening of the weld region, as shown in Fig. 5. In terms of the appearance of the bead, location of the beam focus on the top surface (Fig. 6a and e) resulted in an

273

irregular upper bead as compared to focusing within the sheet thickness (Fig. 6b and f). The location of the focal point at the bottom surface was also observed to increase the tendency for solidification cracking on the weld surface (arrow in Fig. 6g) and, with subsequent lowering of the focus, the irregularity of the upper bead was observed to increase (Fig. 6d and h). The average size of the fusion zone width (Fig. 7a) was the smallest for beam focusing on the top surface (1.4 mm) as compared to focusing in the center of the sheet thickness (1.6 mm), at the bottom surface (1.85) and at a distance of 5 mm from the bottom surface (2.1 mm), as illustrated in Fig. 7a. Similarly, the fusion zone area continually increased from roughly 7e10 mm2 when the beam focus changed progressively from the top of the sheet to below the bottom surface (Fig. 7b). Alternatively, the size of the concavity and the HAZ area appeared to be minimized (0.23e0.27 mm and 1.1e1.2 mm2) when the beam was focused within the sheet thickness (top, middle or bottom), as illustrated in Fig. 7ced, which may be related to the effect of beam focus position on the energy density distribution and weld shape. Hence, it appears that a BF of 2.5 mm (i.e., focus midway within the sheet thickness) results in a balance between the upper bead appearance (regular bead pattern, no spattering, no cracking) and the weld zone characteristics (small and uniform fusion zone

Fig. 6. Effect of beam focus on the characteristics of the upper bead for single pass EB welds produced at 2.50 kW, 30 mm s1 and 20 cm GTWD: (aed) without oscillation and (eeh) with a 2.5 mm f beam.

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Fig. 7. Effect of bam focus on the (a) fusion zone size, (b) concavity, (c) fusion zone area and (d) HAZ area for single pass EB welds produced at 2.50 kW, 30 mm s1 and 20 cm GTWD; open and closed markers represent welding with and without oscillation (2.5 mm f beam).

width, small fusion zone and HAZ size and small concavity). Using a power input of 2.5 kW with the beam focused in the center of the sheet thickness (GTWD 20 cm), sound welds in AA2024 can be manufactured at a welding speed of 30 mm s1 without porosity and cracking in the 5 mm thick sheets. The results on the beam focus position are in agreement with previous work on EB welding of stainless steel, which has indicated that for beam powers of up to 4 kW, the focus must be below the top surface, while above 4 kW (for deeper welds) the focus position is not critical [23]. 3.1.3. Effect of beam oscillation The effect of applying beam oscillation during welding on the characteristics of the weld morphology is shown in Fig. 8 for a single pass and with a cosmetic pass. As compared to welding conditions without beam oscillation, the application of an oscillated beam 2.4 mm in diameter (f) was observed to increase the fusion zone size by w25% (Fig. 7a), reduce the concavity slightly (Fig. 7b) from 0.37 mm (7% of thickness) to 0.23 (5% of thickness) and increase the fusion zone area by w30% (Fig. 7c). For the HAZ area, depending on the beam focus position, there is an increase of roughly 0e13% for the oscillated beam welding conditions as compared to that without oscillation. When applying beam oscillation, the influence of the beam focus position between the top surface and the center of the sheet was observed to be relatively small on the fusion zone size, the size of the concavity, and the areas of both the fusion zone and HAZ. Also, when the beam focus

was positioned beyond the bottom surface, there was practically no change in the weld characteristics. This is attributed to the size of the oscillated beam being independent of the beam focus position. For the application of beam oscillation during welding, an increase in the fusion zone size and area represents a sizeable region over which the microstructural characteristics of the parent material have been altered. Previous work on laser and EB welding of AA2024 has indicated that during melting and solidification of the fusion region, the altering of the precipitates in AA2024-T3 results in a reduction in the strength of the weldments [5,24]. Hence the need to minimize the size of the fusion zone in AA2024-T3 weldments for improved mechanical properties is likely to render the application of beam oscillation during welding counter-productive, with the exception of the cosmetic pass where the circular path can promote smoothing of the weld bead surface to minimize the concavity (Figs. 4b and 8b). 3.1.4. Effect of welding speed With all other parameters held constant, analysis of the upper weld bead for different welding speeds (27.5e38 mm s1) indicated that beyond 34 mm s1 considerable spattering of the molten weld pool occurred on the surface, as illustrated in Fig. 9. Also an increase in the welding speed from 27.5 mm s1 to 38 mm s1 was observed to decrease the concavity from 0.31 to 0.12 (Fig. 10a) and the fusion zone area from roughly 7.6 mm2e7.3 mm2 (Fig. 10b), which is most likely related to the decrease in the linear heat input (at constant

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275

conditions for EB welding of AA2024 at a GTWD of 20 cm were determined to be 2.5 keV, 2.5 mm BF and 34 mm s1, and these were observed to give a relatively uniform weld width of 1.55 mm and a small concavity of 0.15 mm, as illustrated in Fig. 12. It is also noteworthy that despite the similar average weld width of the fusion zone at 30 mm s1 (1.6 mm) and 34 mm s1 (1.55 mm), the standard deviation of these values was determined to be 0.45 and 0.12 respectively. This, of course, is related to the increase in the weld width uniformity with increasing welding speed and further supports 34 mm s1 as the optimum welding speed for the conditions investigated. As compared to reported EB welding data [24] on 5 mm thick AA2024-T351 (6.5 kW, 22 mm s1, 300 J mm1), the optimum linear energy input determined in this work (2.5 kW, 34 mm s1, 73.5 J mm1) was considerably lower, and suggests that the resulting metallurgical damage to the weld assembly would be smaller. 3.2. Microscopic examination 3.2.1. Microstructure of fusion region The fusion region (w1.4e1.7 mm in size) in the full penetration welds consisted of thin, elongated and fine dendrites (Fig. 13aec) indicative of the high heating and cooling rates involved in EB welding of AA2024, especially under the high welding speed conditions, as in the present study (27e38 mm s1). In particular, the elongated dendritic network in the fusion zone was decorated with copper-rich (Al2Cu) precipitates (bright phase) along the interdendritic boundaries on account of copper segregation during solidification (Fig. 13c). Adjacent to the fusion zone, a narrow overaged HAZ (up to 0.2 mm thick) was observed with some grain boundary liquation and coarsened Al2Cu precipitates (Fig. 13def). As compared to previous work on EB welding of 5 mm thick AA2024-T351 (300 J mm1) that has indicated a fusion zone size of roughly 2.5 mm [24], the fusion zone size in the present work is smaller (1.5 mm) and is due to the lower linear energy input applied (73.5 J mm1).

Fig. 8. EB weld in 5 mm thick AA2024-T3 at 2.5 kW, 30 mm s1, 2.5 mm BF, 2.4 mm f beam and 20 cm GTWD: (a) butt pass and (b) butt pass and cosmetic weld (at 25 mA and 3 mm f beam).

power) from 90.8 J mm1 to 65.6 J mm1. The morphology of the fusion zone was also observed to change and an increase in the weld width uniformity occurred with increasing welding speed, as shown in Fig. 11 for 27.5 mm s1 and 38 mm s1. Of interest, also, is the undercut and underfill on the bottom surface at a welding speed of 38 mm s1, as shown in Fig. 11b. Hence beyond 34 mm s1, the lack of weld penetration as well as the occurrence of an irregular bead at the top and bottom surfaces precludes further increases in the welding speed, despite the progressive reduction in the concavity and fusion zone area. Thus, the optimum process

Fig. 9. EB welds in 5 mm thick AA2024-T3 at 2.5 kW, 2.5 mm BF and 20 cm GTWD: (a) 30 mm s1, (b) 34 mm s1 and (c) 38 mm s1.

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7.8

a Fusion Zone Area (mm 2 )

Concavity (mm)

0.4

0.2

0

b

7.6

7.4

7.2

25

30 35 Welding Speed (mm·s-1)

40

25

30 35 Welding Speed (mm·s-1)

40

Fig. 10. Effect of welding speed on (a) concavity and (b) fusion area for single pass EB welds produced at 2.50 kW, BF of 2.5 mm and 20 cm GTWD.

3.2.2. Composition of fusion region EDS of the weld produced at the optimum process conditions (2.5 kW, 30 mm s1, 73.5 J mm1) indicated that the magnesium content in the fusion zone (1.1e1.3 wt.%) was lower than the parent material (1.5 wt.%), as given in Table 2. This decrease (w20%) in the magnesium content is likely due to evaporation losses from the weld pool as opposed to a redistribution effect as the magnesium content in the interdendritic particles in the fusion zone (1.2 wt.%) and Al2Cu particles in the HAZ (1.2 wt.%) are similar to that of the adjacent regions (in the fusion zone dendrites at 1.1 wt.% and in the HAZ matrix at 1.4 wt.%). This finding is consistent with previous work on predicating the tendency for vaporization from molten weld pools of aluminum alloys, which established both Mg and Zn as the dominant metal vapors in the absence of compound formers such as oxygen and nitrogen [25]. Also, work on EB welding of AA2024-T351 has reported a lower magnesium content in the fusion zone (2 wt.%) as compared to parent material (2.7 wt.%) and attributed the losses (w26%) to vaporization during melting of the weldment [24]. For the optimum EB welding conditions (2.5 kW, 30 mm s1, 73.5 J mm1), EDS of the weld indicated that the copper content, determined by spot analyses, was lower for the dendritic regions in the center of the fusion zone (1.8 wt.%) as compared to measurements near the fusion boundary (4.1 wt.%), HAZ (4.2 wt.%) and parent material (4.5 wt.%). However an area analysis using EDS showed a similar copper level for the fusion zone (4.4 wt.%) and parent material (4.5 wt.%), suggesting that copper losses did not occur during EB welding. Hence, this decrease in the copper content may be related to the formation of Al2Cu precipitates in the interdendritic region of the fusion zone (55.6 wt.% Cu) and matrix of the HAZ (57.4 wt.% Cu), which through a redistribution effect renders the adjacent regions depleted in copper. This finding is supported by previous work on vaporization from molten weld pools of aluminum alloys, which established that copper, with a vapour pressure less than aluminum, does not vaporize during welding [25]. Also, work on autogenous TIG welding of AA2024 (48e90 J mm1) has revealed that the gradient in the copper content across the fusion zone is related to the composition of the first solid to form, 2 Cok to 3 Cok [1], where Co is 4.46 wt% Cu for the initial alloy concentration of a binary AleCu alloys of the same Cu composition as AA2024 and k is 0.17 for the equilibrium partition coefficient for binary AleCu alloys. Hence, the calculated composition of the first solid to form in the fusion zone would be between 1.5 and 2.28 wt.% Cu at the weld centerline, which supports the measured value of 1.8 wt.% obtained in the present work. Finally for EB welding of AA2024-T351, previous work has also reported a depletion in Cu in both the fusion zone (1.1 wt.%)

and HAZ (2.3 wt.%), as compared to the parent material (3.6 wt.%), which was attributed to the redistribution of Cu within the fusion zone and the concomitant formation of Al2Cu precipitates [24]. 3.2.3. Dendrite arm spacing During fusion welding of aluminum alloys, the resulting microstructure in the weldment is dependent on the solidification time of the weld metal, i.e., the average cooling rate between the liquidus and non-equilibrium solidus temperatures [26]. Knowing that the solidification time is dependent on the heat input conditions [27], the secondary dendrite arm spacing was examined for the different welding conditions studied in this work. Specifically, analysis of the secondary dendrite arm spacing (l) was performed as a function of the solidification parameter (Q/L)1\2 [27e29], where Q is the heat input and L is the unit length (i.e., AV$BC$WS1), as illustrated in Fig. 14 for the data obtained from EB welded AA2024-T3. Over the range of heat input per unit length values examined in this work (42e120 J mm1), the dendrite arm spacing was observed to increase from roughly 1.5 to 2.5 mm. In addition, the plot of dendrite arm spacing as a function of the solidification parameter was observed to be linear. This finding is consistent with previous work on weld metal substructure characterization that has indicated linear increases in the dendrite arm spacing with solidification time or solidification parameter (3.4e33 kJ cm1) for autogenous arc welds in AA2014 [26]. Previous work on double pass MIG welding of AA2024 (13 mm thick) has reported dendrite arm spacing values between 6 and 12 mm for heat input conditions of 870 J mm1/pass [30]. Using a lower heat input (60e120 J mm1) during TIG welding of AA2024 (1.6 mm) resulted in a finer fusion zone microstructure with dendrite arm spacing values ranging from 3 to 6 mm for relatively slower solidification conditions (as compared to EB welding) of 500e1800  C s1 [1]. Considering that work on aluminum alloys has also related the solidification rate (R) to the dendrite arm spacing [31,32], the empirical relationships given for casting [33,34] and laser melting [35] of Al-4.5 wt.% Cu, as well as laser welding of AA7020 [36] were examined to determine the solidification rate for EB welding conditions in the present work as follows:

l ¼ ac R1=3

(3)

where ac has been reported to be 45 [35] or 50 [33,34]

ln l ¼ 3:553  0:287 ln R ½33

(4)

Using the average dendrite arm spacing value of 2 mm calculated from the measurements (1.5e2.5 mm) obtained in this work, the average solidification rates were determined to be 1.1e1.6  104  C s1 and 2.2  104  C s1 using equations (3) and (4), respectively. Previous

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work on laser and EB melting of Al-4.5 wt.% Cu surfaces has reported that for solidification rates between 106$and 105  C s1 the dendrite arm spacing in the solidified zone ranges from 0.3 to 1 mm, respectively [35]. Reported data for slightly slower solidification conditions of 104  C s1 in Al-4.5 wt.% Cu has indicated a dendrite arm spacing value between 1 and 2 mm [36]. These findings support strongly, the dendrite arm spacing values obtained for the fusion region of the EB welds in AA2024 that were produced over the heat input conditions of 42e120 J cm1. Also, using the average dendrite spacing value of 2 mm for a heat input of 73.5 J mm1 and the equations given in literature for

277

Al-4.5 wt.% Cu, the solidification rate determined for EB welding of AA2024 with optimized parametric conditions was calculated to be about 104  C s1, which is in good agreement with data on the processing of aluminum alloys, as compared in Table 3. For each EB weld condition, it was also consistently observed that the dendritic structure was finer towards the weld centerline as compared to the fusion boundary region, as illustrated in Fig. 13d. Specifically, a progressive increase in the dendrite arm spacing was measured between the weld centerline and fusion boundary, such that the values near the fusion boundary region were determined to be about 1.5 times greater than those reported in Fig. 14. Previous work has indicated that this difference in dendrite arm spacing value between the fusion boundary and weld centerline may be attributed to the following: (1) Changes in the local solidification rate at different positions on the advancing solid/liquid interface which would generate faster weld-pool solidification at the center of the weldment as compared to the fusion boundary region [28], and/or (2) Increased solute segregation at the weld centerline that reduces the dendrite coarsening rate during solidification and results in a finer dendritic structure [26]. Ignoring solute segregation effects, for the optimum EB welding condition (73.5 J mm1) in the present work, the changes in the microstructural condition from the weld centerline (lavg 2 mm) to the fusion boundary (lavg 3.5 mm) suggests nearly on order of magnitude decrease in the solidification rate (i.e., from 104 to 103  C s1).

3.3. Precipitation evolution The DSC thermograms of the AA2024-T3 condition of the parent material and the fusion zone are given in Fig. 15. For the parent material microstructure, one exothermic peak (w262  C) and three endothermic peaks (177  C, 220  C and 478  C) were observed to occur. Although the characteristics of the precipitation sequence for

Fig. 11. EB welds at 2.50 kW, 2.5 mm BF and 20 cm GTWD: (a) 27.5 mm s1 and (b) 38 mm s1.

Fig. 12. EB weld at 2.50 kW, 2.5 mm BF, 20 cm GTWD and 34 mm s1.

278

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Fig. 13. Microstructure of EB welded AA2024: (aec) weld and (def) HAZ.

the AleCueMg alloy system have been contentious [37e43], recently a progression for the formation of Guinier Preston Bagaryatsky (GPB) zones has been reported to occur during continuous heating [44]:

Supersaturated Solid Solution/GPB zones/S0 /SðAl2 CuMgÞ (5)

Table 2 EDXS analysis for EB welded AA2024 (wt.%). Region

Al

Mg

Comments

PM (area) FZ (area)

92.9 93.8

Cu 4.5 4.4

1.5 1.2

e Loss of Mg

Dendrites in center of FZ (spot) Dendrites in FZ close to FB (spot) FZ interdendritic particles (spot) HAZ matrix (spot) HAZ particles (spot)

96.5 94.0 44.8 94.0 41.2

1.8 4.1 55.6 4.2 57.4

1.1 1.3 1.2 1.4 1.2

Loss of Mg and Cu Loss of Mg and Cu Al2Cu precipitates Loss of Mg and Cu Al2Cu precipitates

PM ¼ Parent Metal; FZ ¼ Fusion Zone; FB ¼ Fusion Boundary; HAZ ¼ Heat-Affected Zone.

Also, the microstructure for AA2024-T3 has been acknowledged to consist of GPB zones with Al2CuMg (S phase) as the principal precipitating phase, as determined by the EDS analysis of the parent material (Fig. 4 and Table 2). Previous work on the AleCueMg system has indicated that depending on the heat treatment condition of the material prior to differential thermal analysis, the DSC thermograms can exhibit up to five peaks, two exothermic precipitation peaks and three endothermic peaks between room temperature and 500  C [42]. Specifically, an exothermic peak below w150  C occurs due to the formation of GPB zones during the natural aging stage and this peak has been found to be absent during DSC testing of the aged material condition (such T3), presumably due to the prior existence of these precipitation phases in the structure. For the parent material, the appearance of the two endothermic peaks, labeled A1 and A2 in Fig. 15, has been associated with the dissolution of the GPB zones. In particular, peak A1 has been observed to become increasingly less prominent with increasing aging time and the presence of this peak indicates that some of the GPB zones dissolve at a slightly lower temperature range. Thus, the dissolution of the remaining GPB zones at a higher temperature range (A2), at

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0.15

3

B

0.1 Heat Flow (W/g)

Dendrite Arm Spacing (um)

279

2

Parent

0.05 FZ

0 A1 A2

-0.05

C

1

-0.1

20

25

30

50

35

Solidification Parameter SQRT {Q/V} (SQRT{J/cm})

150

250

350

450

Temperature (°C)

Fig. 14. Variation of dendrite arm spacing with heat input for EB welded AA2024-T3.

around 215e225  C, indicates the greater inherent stability of these latter precipitates. According to previous work on AleCueMg alloys, the appearance of the second endothermic peak (A2) may be related to the dissolution of GPB zone-dislocation complexes [42]. Between 260 and 265  C, an exothermic peak (labeled B in Fig. 15) was observed for the parent material microstructure. Previous work has reported that during heating of AleMgeCu alloys, a metastable S0 phase, which is described as a slightly strained version of the S phase [41], forms. However, since practical differentiation between these two phases on the DSC curve is not possible, it is usual to note the phase as S0 (S) [41e43]. Thus, the exothermic peak B is associated with the precipitation of S0 (S) precipitates from the solid solution resulting from the dissolution of the GPB zones. Subsequently, the occurrence of an endothermic peak (labeled C in Fig. 15) for the parent material microstructure at about 475e480  C can be related to the dissolution of these S0 (S) precipitates. The DSC thermogram of the fusion zone microstructure was observed to be relatively similar to that of the parent material. Specifically, a similar absence of an initial exothermic peak below w150  C was observed and this may be attributed to the presence of Al2Cu particles in the interdendritic regions, as determined from the EDS results. However, instead of two endothermic peaks associated with the dissolution of the GPB zones in the parent material, only a single peak at roughly 225  C was observed for the fusion zone structure. This appears to suggest that the precipitates associated with the GPB zones in the microstructure of the fusion region are of higher stability than those found in the parent material. This finding agrees with previous work on AleCueMg alloys that have indicated the absence of the lower temperature

Fig. 15. DSC thermograms of the parent material (AA2024-T3) and a place in the fusion zone of an EB weld performed under optimum processing conditions (73.5 J mm1). Peak A corresponds to the dissolution of GPB zones, peak B corresponds to the precipitation of S0 (S) and peak C to the dissolution of the S0 (S) precipitates.

range peak after prolonged (24 h) aging at 130  C, which resulted in the stabilization of the GPB zones [42]. Using this endothermic peak associated with the dissolution of the GPB zones, previous work has indicated that quantification of the GPB zones is possible [30,45]. Specifically, it has been reported that as long as the overlap with the ensuing exothermic reaction (peak labeled B) is sufficiently small, measurement of the area under this peak can be used to determine the volume fraction of GPB zones present in the material [46]. Hence in the present work it was assumed that the AA2024 parent material in the T3 condition contained only GPB zones and that their volume fraction in this state was the maximum possible for the alloy. With these assumptions, calculation of the area under the endothermic peaks (A1 and A2) thus provides a reference value Aref, which is characteristic of the volume fraction of GPB zones present in the AA2024 material prior to welding. Hence, the relative fraction of GPB zones in the fusion zone after EB welding (with any processing condition) can be calculated as the ratio between the GPB zone dissolution peak area A of the sample and that of the AA2024-T3 parent material (Aref), or A/Aref (Fig. 15), as given in Table 4. The ensuing exothermic peak for the fusion zone structure was observed to be shifted slightly to higher temperatures (270e275  C) and the peak height was diminished as compared to that observed for the parent material. Previous work has indicated that area associated with this peak can be used to quantify the

Table 4 DSC results for fusion zone of EB welded AA2024. Weld condition

Table 3 Comparison of the solidification rates various processes. Process

Material

Conditions

Cooling rate

Reference

TIG welding Laser welding Laser welding Laser/EB melting

AA2024 AA7020 AA2024 Al-4.5Cu

1.6 mm thick butt 3 mm thick butt Bead on plate Rapidly solidified surface

102e103  C s1 103  C s1 105  C s1 105e106  C s1

[1] [36] [6] [35]

EB welding

AA2024

5 mm thick butt

104  C s1

Current work

2.5 kW, focus in center, 34 mm s1 without oscillation 2.5 kW, focus in center, 34 mm s1 3 mm f beam

Relative volume fraction of precipitates (%)

Peak temperature ( C)

GPB zones (A/Aref)

S0 (S) phase (1  B/Bref)

Endothermic Peak

Exothermic Peak

90

30

225

271

69

35

225

278

280

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160

700

Weld Center

P WO 500

140

W

Load (N)

Vickers Microhardness

600

120

400 300 200

100 -2.5

a -1.5 -0.5 0.5 1.5 Distance Across Weld

100

2.5 0 0

Vickers Microhardness

160

Weld Center

0.05 0.1 0.15 Displacement (mm)

0.2

Fig. 17. Load-displacement curves in a shear punch test for the various microstructures including (P) parent metal, (W) weld metal without beam oscillation and (WO) weld metal with beam oscillation.

140

120

100 -2.5

b -1.5 -0.5 0.5 1.5 Distance Across Weld

2.5

Fig. 16. Microhardness profiles of EB welded joints in 5 mm AA2024: (a) 2.5 kW, 30 mm s1, focus at center and 20 cm GTWD and (b) same as (a) butt with beam oscillation (3.0 mm f).

content of S0 (S) particles [30,45]. Essentially it is assumed that in the T3 condition S0 (S) particles are not present and the exothermic peak of AA2024 parent material in this condition relates again the maximum volume fraction of S0 (S) particles, which provide a reference peak Bo [46]. Hence, in the case of a mixed initial microstructure (such as in the fusion zone) containing both GPB and S0 (S) precipitates, the peak (B) only reveals the complementary of the S0 (S) initially present. Consequently, the relative fraction of S0 (S) precipitates initially present in the EB welded samples (Table 3) can be calculated as 1  B/Bo. From the results obtained for the relative volume fraction of the GPB zones and S0 (S) phase in the fusion zone of EB welded AA2024, it appears that welding without beam oscillation using the optimum processing conditions (2.5 keV, 34 mm s1, 20 cm GTWD,

beam focus in center) results in a precipitation behavior having a closer resemblance to the parent material. Alternatively, with the application of beam oscillation (3.0 mm f circular path) considerable reduction in the quantity of GPB zones was observed with an enhanced precipitation of the S0 (S) phase. With regard to the precipitation rate, previously it has been observed that shifting of the exothermic peak to lower temperatures is a result of accelerated kinetics, while shifting of the endothermic peak to higher temperatures is caused by the deceleration of the dissolution rate [42]. In the present work, it was observed that the DSC thermogram for the fusion zone exhibited a shifting of the endothermic and exothermic peaks to higher temperatures relative to the parent material (211  C and 262  C, respectively). Hence the dissolution of the GPB zones in the fusion region of EB welded occurs earlier than for the parent material. Alternatively, the S0 (S) precipitates in the EB welds produced without beam oscillation appear later compared to the welded structure with beam oscillation. Since the strengthening condition of AA2024 is related to the precipitation characteristics, these differences in the fraction of GPB zones and S0 (S) phase between the EB welded and parent materials may be related to the mechanical properties, as discussed below. 3.4. Mechanical properties 3.4.1. Microhardness results Representative microhardness profiles for EB welded AA2024 are illustrated in Fig. 16. For the as-welded condition, the average hardness in the fusion region was measured to be the lowest, at around 105 HV, as compared to the parent material (145  6 HV). The microhardness in the HAZ was observed to be between that measured for the fusion zone and material. Also, as indicated in

Table 5 Mechanical property results from shear punch testing. Microstructural condition

Strength (MPa)

Ductility (%)

Yield

Joint efficiency

Maximum

Joint efficiency

Percent elongation

Joint efficiency

Parent Metal Weld Metal, 73.5 J mm1 without oscillation Weld Metal, 73.5 J mm1 3 mm f beam

347 318 308

e 91.6% 88.8%

488 358 341

e 73.3% 69.8%

19.9 9.2 8.6

e 46.2% 43.2%

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Fig. 16, the size of the fusion region (w1 mm) and the presence of a narrow HAZ (<0.2 mm) differentiated from the microhardness profiles corroborate with the microstructural results illustrated in Fig. 7. In comparison, the hardness profile for the as-welded condition with beam oscillation revealed a larger fusion zone (w2 mm), a wider HAZ (w0.5 mm) and a lower average hardness, at around 102 HV. For AA2024-T3, the reported hardness value in the solution treated condition (137 HV) [47] corroborate the measurements obtained in this work for the as-received wrought microstructure of the 5 mm thick sheet of AA2024-T3. For the as-welded condition, previous work on TIG [30,38], laser [5], electron beam [24] and friction stir [46] welding of AA2024-T3 have reported microhardness values of 90e100, 90e110, 108 and 120 HV, respectively for the weld microstructure, which supports the values obtained in this work. Previous results also support the microhardness in the HAZ; the values obtained were between those of the weld metal microstructure and the parent material for the as-welded condition in AA2024-T3 produced by TIG, laser, or EB welding [5,24,30]. In addition, considering that the microhardness values measured for the fusion zone of the EB AA2024-T3 material were lower than the parent material, the lower deformation resistance of the weld metal solidification microstructure can be related to the loss of strengthening precipitates associated with copper, as determined from the DSC results. Of great significance, nonetheless, is small size of the weld zone (1 mm) produced by EB welding as compared to previous work on AA2024 using laser (3 mm for 3 mm thickness), friction stir (20 mm for 6 mm thickness) and TIG (20 mm for 8 mm thickness) welding [5,24,30], which would render overall better properties during manufacturing. 3.4.2. Shear punch test results The load-displacement curves given in Fig. 17 demonstrate the flow behavior during shear punch testing of the parent material as compared to the as-welded microstructures in the fusion zone for an oscillated beam (3 mm f) and non-oscillated beam using the optimum processing condition (73.5 J mm1). Examination of these load-displacement data of the shear punch evaluations on the different microstructures was then used to determine the values of the yield load, maximum load and displacement to failure to obtain, through empirical formulations [21], the tensile mechanical property results (Table 5). Overall, the microstructure in the fusion zone of the EB weldments showed lower strength and ductility values as compared to the wrought structure of the as-received parent material. In particular, it was determined that the properties determined by shear punch testing of the parent material corroborate well with the reported tensile yield (345e350 MPa) and maximum (483e493 MPa) strength values as well as the ductility (19e19.2% elongation) for AA2024-T3 [24,47]. In addition, the reduction in the average yield and maximum strength values (Table 5) for the fusion zone microstructure (with and without beam oscillation) is in agreement with the lower hardness measured that was attributed to the loss of precipitation strengthening. Moreover, the use of the shear punch test has the capability of assessing the tensile ductility, which shows that the weld metal microstructure has a considerably lower elongation than the wrought as-received AA2024. These shear punch test results suggest that the fusion zone of AA2024 is particularly prone to strain concentration stemming from the combination of lower strength and ductility. Previous work on EB welded AA2024 has indicated that the joint efficiencies are w90 and w70% in terms of the yield and ultimate tensile strengths [24], which were attributed to strain concentration in the fusion zone and significant losses in the ductility. Nonetheless, as compared to the room temperature strength of TIG welded AA2024 that has

281

been reported to have a joint efficiency of about 50% [3], it appears that an improvement in mechanical properties is possible by EB welding, most likely due to lower heat input conditions of the latter, which achieves overall a smaller thermally affected region and a finer microstructure in the weldment. 4. Conclusions Autogenous electron beam welding of AA2024 was examined under varying joining conditions to determine the effect of process parameters on the weld structural and mechanical property characteristics. It was determined that for welding 5 mm thick sheet material with a butt joint configuration, a threshold value for the beam current exists, beyond which the surface concavity exceeds the specification requirements. For the beam focus, variation in the focal point from above the top surface to a position below the bottom of the sheet, indicated that the regularity in the weld shape and quality was optimum when the beam was centered in the vicinity of the mid-section of the work-piece thickness. Alternatively, since the application of beam oscillation was observed to considerably increase the size of the weld area (w30%) with only a small reduction (from 7 to 5%) in the extent of the concavity, to improve the bead shape for the autogenously welded AA2024, a cosmetic pass was determined to be preferable for improving the surface characteristics of a full penetration pass performed using a focused beam. Hence, optimum welding of the AA2024 material was determined for beam focus conditions at the center of the sheet thickness, without beam oscillation, with a power input of 2.5 kW and a speed of 34 mm s1. Under such conditions, the weld microstructure was observed to consist of a columnar dendritic structure with copper-rich precipitates in the interdendritic regions, while the small HAZ contained some grain boundary liquation and coarsened Al2Cu particles. Overall the EB welded joints exhibited a minimum hardness in the weld zone similar to AA2024 joined using TIG or laser welding due to the dissolution of the strengthening phases. However the smaller relative size of the weld zone (1 mm) of the EB welded AA2024 renders lower overall metallurgical damage as compared to other welding processes.

Acknowledgments This work was conducted under project 46M3-J006 in the Metallic Products Group of the Aerospace Manufacturing Technology Center (AMTC) of the NRC Institute for Aerospace Research (IAR). The authors would like to thank Mr. Eric Poirer, Mr. Maxime Harvey and S. Gravereaux of NRC-IAR-AMTC for their technical assistance during the interim of this project.

References [1] Norman AF, Drazhner V, Prangnell PB. Effect of welding parameters on the solidification microstructure of autogenous TIG welds in an AleCueMgeMn alloy. Mater Sci Eng A 2000;259:53e64. [2] Lefebvre F, Wang S, Starink MJ, Sinclair I. Microstructural features of fusion welded 2024-T351. Mater Sci Forum 2002;396e402:1555e60. [3] Preston RV, Shercliff HR, Withers PJ, Smith SD. Finite element modelling of tungsten inert gas welding of aluminium alloy 2024. Sci Technol Weld Joining 2003;8:10e8. [4] Kou S, Le Y. Welding parameters and the grain-structure of weld metal e a thermodynamic consideration. Metall Mater Trans A 1988;19:1075e82. [5] Hua B, Richardson IM. Autogenous laser keyhole welding of aluminum alloy 2024. J Laser Appl 2005;17:70e80. [6] Milewski JO, Lewis GK, Wittig JE. Microstructural evaluation of low and high duty cycle Nd:YAG laser beam welds in 2024-T3 aluminum. Weld J 1993; 72:341se6s. [7] Grong Ø. Metallurgical modelling of welding. In: Bhadeshia HKDH, editor. 2nd ed. London, U.K.: The Institute of Materials; 1997. p. 228.

P. Wanjara, M. Brochu / Vacuum 85 (2010) 268e282

[8] Owen RA, Preston RV, Withers PJ, Shercliff HR, Webster PJ. Neutron and synchrotron measurements of residual strain in TIG welded aluminium alloy 2024. Mater Sci Eng A 2003;346:159e67. [9] Saunders HL. Welding aluminium: theory and practice. 3rd ed. Washington, DC: The Aluminum Association; 1997. p. 12.4. [10] Ohkubo M, Hara M, Sugamata M, Kaneko J. Electron beam welding of AleCueMg and AleLieCu alloys, aluminium alloys: their physical and mechanical properties. In: Sato T, Kumai S, Kobayashi T, Murakami Y, editors. Proceedings of the 6th international conference on aluminium alloys (ICAA-6), vol. 3. Tokyo, Japan: Japan Institute of Light Metals; 1998. p. 1495e500. [11] Koçak M, Dos Santos JF, Riekehr S, Cam G, Jennequin G, Gonthier-Maurin P. “Electron beam weldability and fracture toughness testing of Al-Alloys”, welding and melting by electron and laser beams (CISFFEL 6). In: Proceedings of the 6th international conference, vol. 2. Cedex, France: Institut de Soudure; 1998. p. 921e3. [12] Sunwoo AJ, Morris JW. Aging characteristics of electron-beam welding and gas tungsten arc fusion zones of AleCueLi alloy 2090. Metall Mater Trans A 1991;22:903e13. [13] Arata Y, Ohsumi M, Hayakawa Y. Electron beam welding of high-strength Al alloy. Trans JWRI 1976;6:19e25. [14] DeBony Y, Griffaton B, Savornin J. Electron beam welding of thick aluminum alloy sheets. Mater Tech 1974;7:305e11. [15] Murphy JL, Huber RA, Lever WE. Joint preparation for electron-beam welding thin aluminum alloy 5083. Weld J 1990;69:125se32s. [16] Buchheit RG, Grant RP, Hiava PF, Mckenzie B, Zende GL. Local dissolution phenomena associated with S phase (Al2CuMg) particles in aluminum alloy 2024-T3. J Electrochem Soc 1997;144:2621e8. [17] Modolfo LF. Aluminum alloys: structure and properties. London, U.K: Butterworth; 1976. 491e683. [18] Lucas GE, Sheckherd JW, Odette GR. In: Corwin WR, Lucas GE, editors. The use of small-scale specimens for testing irradiated material, 1983. Philadelphia, PA, U.S.A: American Society for Testing and Materials; 1986. p. 112e39. Albuquerque, NM, U.S.A., ASTM-STP-888. [19] Wanjara P, Drew RAL, Yue S. Application of small specimen testing technique for mechanical property assessment of discontinuously reinforced composites. J Mater Sci Technol 2006;22:61e71. [20] Guduru RK, Darling KA, Kishore R, Scattergood RO, Koch CC, Murty KL. Evaluation of mechanical properties using shear-punch testing. Mater Sci Eng A 2005;395:307e14. [21] Wanjara P, Jahazi M. Application of shear punch testing to study microstructure-property relationships in electron beam welded 17-4 PH stainless steel. CMQ 2009;48(2):317e26. [22] Private communications Sciaky. [23] Mladenov G, Vutova K. Electron beam welding of metals in vacuum. Bulgarian Acad Sci News 2005;6(22):1e2, http://www.bas.bg/basnews/novosti/ novosti200506Us.pdf. [24] Çam G, Ventzke V, Dos Santos JF, Koçak M, Jennequin G, Gonthier-Maurin P. Characterisation of electron beam welded aluminium alloys. Sci Technol Weld Joining 1999;4:317e23. [25] Block-Bolten A, Eagar TW. Metal vaporisation from weld pools. Metall Mater Trans B 1984;15:461e9.

[26] Lanzafame JN, Kattamis TZ. Solidification structure and tensile properties of 2014 aluminum alloy welds. Weld J 1973;52:226se32s. [27] Brown PE, Adams Jr CM. Fusion zone structures and properties in aluminum alloys. Weld J 1960;39:520se4s. [28] Davies GJ, Garland JG. Solidification structures and properties of fusion welds. Int Met Rev 1975;2:83e106. [29] Jordan MF, Coleman MC. Fused-zone structures in arc-welded deposits of a commercial aluminiumemagnesiumemanganese alloy. Br Weld J 1968; 39:552e8. [30] Lefebvre F, Ganguly S, Sinclair I. Micromechanical aspects of fatigue in a MIG welded aluminium airframe alloy: Part 1. Microstructural characterization. Mater Sci Eng A 2005;397:338e45. [31] Flemings MC. Solidification processing. NY, USA: McGraw-Hill; 1974. [32] Eskin D, Du Q, Ruvalcaba D, Katgerman L. Experimental study of structure formation in binary AleCu alloys at different cooling rates. Mater Sci Eng A 2005;405:1e10. [33] Pratt RC, Jones H, Davies GJ. Effect of process variables on microstructure and solidification kinetics in static and centrifugal casting of AleCu alloys. In: Proceedings of the solidification processing (Sheffield, UK, 21e24 Sept. 1987); 1988. p. 349e52. [34] Melo MLNM, Rizzo EMS, Santos RG. Predicting dendrite arm spacing and their effect on microporosity formation in directionally solidified AleCu alloy. J Mater Sci 2005;40:1599e609. [35] Munitz A. Microstructure of rapidly solidified laser molten Al-4.5 Wt Pct Cu surfaces. Metall Mater Trans B 1985;16:149e61. [36] Nascente PAP, Bolfarini C, Benassi CL, Alcantara NG, Santos JF. Surface and microstructural characterization of laser beam welds in an aluminum alloy. J Vac Sci Technol A 2002;20:1416e9. [37] Bagaryatsky YA. Zh Tekhn Fiz 1950;20:424e7 (Crys. Structure, Experimental, 7). [38] Bagaryatsky YA. Dokl Akad SSSR 1952;87:397. [39] Charai A, Walther T, Alfonso C, Zahra AM, Zahra CY. Coexistence of clusters, GPB zones, S00 -, S0 - and S-phases in an Al-0.9% Cu-1.4% Mg alloy. Acta Mater 2000;48:2751e64. [40] Zahra AM, Zahra CY. Positron lifetime study of an Al-1.7 at.% Mg-1.1 at.% Cu alloy. Philos Mag Lett 2002;82:9e12. [41] Gupta AK, Gaunt P, Chaturvedi MC. Crystallography and morphology of the S prime-phase precipitate in an Al(CuMg) alloy. Philos Mag 1987;A55:375e87. [42] Jena AK, Gupta AK, Chaturvedi MC. A differential scanning calorimetric investigation of precipitation kinetics in the Al-1.53 wt% Cu-0.79 wt% Mg alloy. Acta Metall 1989;37:885e95. [43] Shih HC, Ho NJ, Huang JC. Kinetic study of precipitation behavior in AleCueMg and 2024 Aluminum Alloys. Metall Mater Trans A 1996;27:2479e94. [44] Wang SC, Starink MJ. The assessment of GPB2/S00 structures in AleCueMg alloys. Mater Sci Eng A 2004;386:156e63. [45] Deschamps A, Genevois C, Nicolas M, Perrard F, Bley F. Study of precipitation kinetics: towards non-isothermal and coupled phenomena. Philos Mag 2005;85:3091e112. [46] Genevois C, Deschamps A, Denquin A, Doisneau-Cottignies B. Quantitative investigation of precipitation and mechanical behaviour for AA2024 friction stir welds. Acta Mater 2005;53:2447e58. [47] MatWeb, http://www.matweb.com; 2006. ˇ

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