1

Effect of filler metal feed rate and composition on microstructure and mechanical

properties of fibre laser welded AA 2024-T3

J. Ahn1*, L. Chen2, E. He2, C. M. Davies1 and J. P. Dear1

1 Department of Mechanical Engineering, Imperial College London, South Kensington Campus,

London, UK SW7 2AZ

2 Science and Technology on Power Beam Lab, Beijing Aeronautical Manufacturing Technology

Research Institute, China

ABSTRACT

The influence of aluminium alloy 4043 filler wire feed rate on the weld quality and mechanical properties

of high power 5 kW fibre laser welded aluminium alloy 2024-T3 was investigated. Loss of volatile

alloying elements such as magnesium and other elements including copper and silicon which all

contributed to the hot crack sensitivity was measured using energy dispersive X-ray spectroscopy at

different filler wire feed rates. High feed rates of above 4.0 m/min produced instabilities, whereas, low

feed rates below 2.0 m/min did not sufficiently modify the chemical composition of the weld pool. The

optimum feed rate was found to be in the range between 2 and 3 m/min, where the corresponding

dilution ratio of around 9-12% in the weld pool with less than 0.6% silicon content reduced the

percentage of Mg2Si and also decreased the solidification temperature and total shrinkage during

freezing. The addition of filler metal reduced the risk of welding defects and improved ductility to over

3.5% and a fairly higher tensile strength of around 380 MPa than without. Microstructural examination

showed that the addition of filler wire increased the number of finer dimples within the weld, resulting in

a purely ductile fracture behaviour, as well as reduced micro hot cracks and porosities.

KEY WORDS

AA 2024-T3; Fibre laser; Welding; Digital Image Correlation; Mechanical properties; Microstructure

*corresponding author: joseph.ahn08@imperial.ac.uk; Tel: +44 7948 532 667

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1 INTRODUCTION

While pure aluminium is not susceptible to hot cracking due to the absence of low melting point

eutectic at the grain boundary, the addition of alloying elements such as copper, magnesium and

silicon to the 2024 alloy makes its composition crack susceptible. As a result, it is considered to have

poor weldability because of its high solute content, resulting in a wide solidification temperature range

and increased tendency to form low melting constituents along grain boundary, known as the eutectic

composition with lower freezing points than the base metal; and also due to high coefficient of thermal

expansion and large solidification shrinkage. Hot cracking is a high temperature cracking mechanism

and it can refer to either solidification cracking which occurs intergranularly along the grain boundary

of the weld metal during solidification due to the stresses acting on the segregated low melting point

constituents [1], or liquation cracking which occurs in the coarse heat affected zone along the fusion

boundary as a result of grain boundary liquation of the low melting point constituents.

The weld metal composition in autogenous welding of AA 2024-T3 is identical to the base metal

composition, so it is solidification cracking susceptible. Due to the higher equilibrium vapour pressure

and lower boiling point of volatile alloying elements such as zinc (917°C) and magnesium (1091°C)

compared to aluminium (2519°C), they are excessively and selectively vaporised during laser keyhole

welding, and as a consequence, segregation of these elements and dissolution of strengthening

precipitates [2] result in degradation of mechanical properties of the weld [3,4]. The vaporisation

process involves transport of the alloying element from bulk to surface of the weld pool and then

vaporisation at the liquid/vapour interface into the surround gas phases [5]. Moon and Metzbower [6]

observed depletion of magnesium in the fusion zone of laser beam welded AA 5456, where a

reduction from 5% in the base metal to 4% in the fusion zone occurred. The 20% reduction was

considered to be the reason for the reduced mechanical properties of the weld. Alferi et al. [7,8] also

observed loss of magnesium of 53% when butt welding 3.2 mm thick AA 2024 at a laser power of 1.6

kW, a welding speed of 10 mm/s and -0.5 mm defocus using disk laser. The reduced Mg content was

found to be responsible for keyhole instability which induced macro porosity formation. A part of the

reason is because the addition of elements such as zinc, silicon, copper and iron reduce the solubility

of hydrogen in liquid aluminium due to strong bonding of Al atoms to these elements, whereas, strong

attractive interactions between hydrogen and magnesium, lithium and titanium increase the hydrogen

solubility [9]. The 2024 alloy also contains small quantities of silicon which induces the formation of

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non-strengthening coarse Mg2Si precipitates, and iron which induces the formation of Cu2FeAl7

phase, both of which reduce the fracture toughness of the weld significantly and deplete the solid

solution strengthening Cu solutes in the matrix needed for aged hardening [10].

In order to reduce the cracking susceptibility of AA 2024-T3, it is necessary to adjust the chemical

composition and the solidification microstructure of the weld metal to a less susceptible level. Hot

cracking sensitivity of the 2024 alloy can be reduced by controlling the solidification process during

welding by optimization of welding parameters and avoiding the crack sensitive composition in the

weld metal through the use of a suitable filler metal with different chemistry and a lower solidification

point than that of the base metal [11]. It is possible to reduce the crack sensitivity by introducing minor

eutectic alloy such as Al-Si via the addition of a crack resistant filler metal of the right composition,

usually from the 4xxx series alloys. More eutectics are provided by the filler wire to heal the cracks

formed during the solidification of the weld pool. In addition, using filler wire also improves the process

stability, reduces the tendency to porosity formation and leads to wider volume of the weld pool by

compensating for the loss of material due to vaporisation [12]. Improved welding process stability

means that macro pores caused by keyhole collapse and instability can be avoided. In addition, the

increased silicon content in the molten pool by filler metal addition effectively lowers its hydrogen

solubility and therefore, reduces the tendency to micro porosity formation as less hydrogen is

dissolved in the molten pool. When filler metal is used, the weld metal composition is determined by

the ratio between the amount of the melted base metal and the filler metal in the weld seam. The

dilution ratio is usually around 20% for laser beam welding and due to the low fraction of filler metal in

the weld, the eutectic alloy 4043 with 5% silicon was used as filler metal in this investigation. The

major alloying element of the 4xxx series aluminium-silicon alloys is silicon in amounts up to a

maximum of 12% in 4047, with a low melting point suitable as a filler material for welding AA 2024-T3

to eliminate solidification cracking and undercut, without causing brittleness in the resulting welded

joint [13]. The 4047 wire however, has lower ductility and is more expensive than the 4043 wire so

4043 was chosen over 4047 for its better ductility. The high percentages of silicon in 4043 filler metal

was important for the prevention of solidification cracking and it helped reduce the total shrinkage

during freezing. The low solidification temperature and narrow freezing range of around 5°C of the

4043 filler metal ensured that the base metal completely solidifies prior to the weld and reach its

maximum strength before become subjected by shrinkage stresses so that the duration in which liquid

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metal is affected by this contraction is minimised during cooling. The filler metal was the last part of

the weld to solidify. The 4043 filler wire also combined with the base metal to lower the percentage of

MgSi2 in the weld and reduce its crack sensitivity. However, the addition of filler metal to the weld also

led to a mismatching mechanical properties and reduced the yield and tensile strengths of the weld

compared to the base metal. As well as that, in the case where the wire feed rate was too fast and not

optimised, it resulted in weld pool instability and formation of welding defects. Therefore, the optimum

values for wire feed rate was investigated to produce the best weld quality with reduced crack

sensitivity and mechanical performance.

2 EXPERIMENTAL PROCEDURES

Heat treatable aluminium alloy (Al-Cu-Mg) alloy 2024 sheets in the T3 temper condition (i.e. solution

heat-treated, cold worked and naturally aged) with 3 mm was used. A 0.6 mm diameter consumable

4043 aluminium filler alloy of a nominal composition Al-5%Si was used when welding 2024 with a filler

wire. The chemical compositions of AA 2024-T3 and Aa 4043 are listed in Table 1.

Table 1 Chemical composition of AA 2024-T3 and AA 4043 (Wt. %)

Material Al Cu Mg Mn Cr Si

AA 2024-T3 92.1 5.9 1.0 0.6 0.1 0.3 AA 4043 94.8 - - - - 5.2

A 5 kW continuous wave (CW) ytterbium fibre laser system YLS-5000 from IPG Photonics was used

in TEM01* mode for laser welding. The beam output of 200 W single-mode active fibres was combined

to deliver the beam output to the workpiece via a feed fibre with a diameter of 200 µm. To protect the

feed fibre from damage during welding, the feed fibre was coupled into a processing fibre with a

diameter of 300 µm which was connected directly to the laser processing head [14], and the beam

diameter at focus was 630 μm. The wavelength of fibre laser was 1070 nm, the beam quality factor,

M2 was around 7.3, the divergence half angle of the focused beam was 12.5 mrad and the Rayleigh

length was about 3.1 mm. The focal length of focusing lens was 300 mm, and the diameter of the

focusing lens was 50 mm. The focal length of the collimator lens was 100 mm and the diameter of the

collimator lens was 50 mm. A beam parameter product (BPP) of less than 2.5 mm mrad was formed.

The filler wire feed rate was the only parameter which was changed while all the other parameters

were kept constant. An independent investigation was conducted to optimise the remaining

parameters such as laser power, welding speed and focal position. All specimens were welded using

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a laser power of 4.9 kW, a welding speed of 3.0 m/min, +4 mm defocus and helium gas shielding as

shown in Figure 1.

Figure 1 Full penetration bead on plate welding of AA 2024-T3 using AA 4043 filler metal at different feed rates

The effect of varying the filler wire feed rate was studied as shown in Table 2. A BINZEL Master-

Feeder system was used to supply filler metal into the leading edge of the weld pool, ahead of the

laser beam impingement point at an angle of 45 or 60° with the workpiece. Industrial grade helium

with 99.999% purity was used at a flow rate of 15-20 l/min. The shielding gas was supplied to protect

both top and underside of the weld. The coaxial shielding gas was delivered via the weld nozzle to

protect molten pool and the back protecting shield gas was supplied via the shielding gas path in the

copper insert to protect back weld. Both top and bottom surfaces of each specimen were brushed with

a stainless steel wire brush and then cleaned using an industrial grade unbuffered 99.9% pure

acetone before welding.

Table 2 Fibre laser welding parameters used as a function of filler metal feed rate and dilution ratio

Laser power Speed Focal position Wire feed rate Dilution ratio (kW) (m/min) (mm) (m/min) (%)

4.9 3.0 +4.0 0.0 0.0 4.9 3.0 +4.0 1.5 6.3 4.9 3.0 +4.0 2.0 8.5 4.9 3.0 +4.0 2.5 10.3 4.9 3.0 +4.0 3.0 12.3 4.9 3.0 +4.0 4.0 16.2 4.9 3.0 +4.0 5.0 20.1 4.9 3.0 +4.0 7.0 27.4

The microstructural constituents of the weld are revealed by using suitable chemical etchants. Keller’s

reagent which is a mixture of 95% distilled water, 2.5% HNO3, 1.5% HCl and 1.0% HF was used to

etch AA 2024-T3 weld specimens by immersing for 10-30 seconds fresh. Macroscopic and

microscopic inspection of transverse sections of metallographic welded specimens followed the test

procedures specified in BS EN 1321 [15]. Microscopic examination with a magnification within 50 to

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500 times also with etching to reveal features of welded joints. An optical microscope (OM), Zeiss

Axio Scope A1 was used for microstructural examination. Energy dispersive X-ray spectroscopy

(EDX) in an environmental scanning electron microscope (SEM), Hitachi S-3400N VPSEM was used

to determine chemical compositions of the specimens at an accelerating voltage of 15 kV, an

emission current of 76 μA, a working distance of 6.8 mm, an elevation of 35° and a live time of 50

seconds. Scanning electron microscope (SEM) was also used to examine the fracture surface of

selected tensile specimens after tensile testing in the base metal and the welded joints at various

magnifications to determine the fracture behaviour and the presence of welding defects.

The quality of welds produced in this experiment were evaluated against a set of welding acceptance

criteria from several international standards on welding as shown in Table 3:

Table 3 Weld quality assessment criteria applied to 3 mm thick AA2024-T3 [16–18]

Standard Level Face width (mm)

Root width (mm)

Porosity (mm)

Undercut (mm)

Underfill (mm)

Reinforcement (mm)

AWS D17.1 Class A N/A N/A ≤ 0.99 ≤ 0.05 ≤ 0.13 ≤ 0.99 BS EN ISO 13919-2 stringent B N/A N/A ≤ 0.90 ≤ 0.15 ≤ 0.15 ≤ 0.65

BS EN 4678 AA ≤ 4.00 ≤ 2.50 ≤ 0.90 ≤ 0.15 ≤ 0.30 ≤ 0.55

The main criteria assessed were the top and bottom weld widths, the ratio of root to face width, Rw,

the depth of undercut and underfill, the size of weld porosity and the height of reinforcement or excess

weld penetration. Any weld showing a lack of penetration or crack were rejected according to AWS

D17.1, BS EN ISO 13919-2 and BS EN 4678. Other welding imperfections such as overlap and

spatter were also identified. The Rw was used to evaluate the processing stability of full penetration

welding and a value of 0.6 determined by Chen et al. [19].

The influence of filler wire feed rate on microstructural transformations and variations of local

hardness profiles of AA 2024-T3 welds was evaluated by measuring hardness in the direction

perpendicular to the transverse weld cross-sections along three lines located 0.5 mm away from the

top and the bottom surfaces, with a 1 mm gap between them. Micro-hardness measurements were

performed on weld cross-sections perpendicular to the weld line using a Zwick Roell Z2.5 (ZHU 0.2)

hardness testing machine at a load of 100 g and a dwell period of 15 seconds at a speed of 60 μms-1

to characterise the whole hardness profile across the weld seams up to the base metal. To avoid work

hardening contributions from the adjacent indents, indentations were separated by 200 μm and the

lines of indentation were separated by 500 μm.

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Uniaxial tensile tests were conducted at ambient temperature on a 200 kN Instron electromechanical

universal testing machine in ram displacement control at a constant crosshead speed of 1 mm/min.

Tensile properties were determined from two butt-welded 3 mm thick AA 2024-T3 tensile specimens

one welded with and the other without filler metal. Digital image correlation (DIC) technique was used

to capture digital images of the deforming surface of the specimens during uniaxial tensile

deformation every second to compute the corresponding displacement and strain. Images were

captured using the 2D DIC technique at the minimum focus distance of 0.45 m and at a resolution of

3456 × 2304 pixels. The DIC area of analysis was selected manually to process the images only

within the specimen, with a resolution of 310 × 1800 pixels. Additionally, a facet size of 20 pixels and

a facet step of 16 pixels, giving an overlap of 4 pixels, were set up according to the configuration with

the average speckle diameter and the resolution of the image in order to ensure that the results are

accurate.

Table 4 lists the welding parameters of the specimens tested. In order to evaluate quantitatively the

differences in mechanical properties between the autogenously welded specimen and the specimen

welded using an optimum feed rate of 2.6 m/min as found in this investigation, the development of strain

distribution during loading was determined.

Table 4 Differernt sets of welding parameters used for tensile testing macro-tensile welded specimens

Welding mode Laser power

(kW) Welding speed

(m/min) Focal position

(mm) Wire feed rate

(m/min) Dilution ratio

(%)

Autogenous 2.9 1.5 +4 - 0.0 Filler 2.9 1.5 +4 2.6 10.5

3 RESULTS AND DISCUSSIONS

Microstructure of the weld with and without filler metal was examined at 50x and 500x magnifications

as shown in Figure 2. The specimen which was welded autogenously contained large solidification

cracks on the weld top surface, whereas, the one welded in the presence of filler metal displayed no

visible crack within the weld. It was found that in both cases the centre of the FZ showed the

formation of characteristic equiaxed dendrites, and columnar dendrites near the FZ boundary.

Epitaxial continuous growth of columnar grains was observed in the direction of thermal gradients in

the FZ, with the same crystallographic orientation to that at the FZ line. Dendritic growth was reported

by Watkins et al. [20] for alloys containing less than 5% weight copper, in this case AA 2024, with the

dendrites being α-Al and with either CuAl2 precipitates or CuAl2-Al eutectic as the inter-dendritic

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phase. The FZ of AA 2024 mainly consisted of α-Al phase with a surrounding eutectic CuMgAl2 phase

as it contained magnesium.

a) b)

c) d)

Figure 2 Microstructure of the AA 2024-T3 weld a) autogenous and b) with filler metal, and dendrites (equiaxed and columnar) at the weld centre line c) autogenous and d) with filler metal

During welding of AA 2024, the low melting point eutectic with a wide range of freezing temperatures

segregated in the grain boundaries and formed the low melting point constituents, which were

rejected by the solidifying columnar grains. The amount of eutectic liquid between grains were large

enough to form a thin, continuous grain boundary film during solidification at a depressed liquidus and

solidus temperatures because of magnesium, compared to the bulk solidus temperature. The solidus

temperature was further suppressed due to a lack of diffusion resulting from rapid non-equilibrium

solidification during welding. The shrinkage strains was proportional to the coherence range between

the first formation of mushy stage by dendrite interlocking and the solidus, so a wider coherence

range increased the tendency for solidification cracking [21]. When the amount of liquid available

during the freezing process was insufficient to fill in the spaces between the solidifying grains at the

centre, then micro-cracks such as those observed in Figure 2 were formed due to the lack of material

and high shrinkage strains in the weld pool [22]. Equiaxed dendritic structure on the other hand,

reduced the crack sensitivity due to the abundance of liquid metal between grains which were able to

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deform more easily under stresses [23], and the lower coherent temperature range resulting from the

formation of equiaxed dendrites at a later stage in freezing [24]. In addition, the fine isotropic grain

structure of equiaxed grains unlike coarse anisotropic columnar grains, increased the resistance to

crack formation and propagation [25] by distributing the low melting point segregates over a larger

grain boundary area and also relieved local shrinkage strains developed during freezing more

efficiently [24]. Equiaxed grain formation is important for the grain refinement of welds but due to the

high solidification rate and thermal gradient, it is often considered difficult to obtain. Instead, columnar

grain growth is favoured and there is a small chance of equiaxed grain formation, resulting in

predominantly coarse, low ductility columnar grain structure in the FZ [26]. The use of 4043 filler metal

which has a freezing range of around 5°C enabled rapid solidification of welds and reduced the time

for shrinkage during solidification, and therefore, micro-cracks were not observed in specimens

welded with filler wire.

The weld face and root dimensions were initially measured as shown in Figure 3. It was found that the

top and the bottom weld width of the specimens increased with increasing filler wire feed rate. The

change in the bottom width with feed rate was relatively small but quite large for the top width. The top

weld width of specimens measured at all feed rates passed the criterion in BS EN 4678, of 4.0 mm,

whereas the bottom weld width was above the maximum of 2.5 mm at all feed rates and therefore,

failed the criterion. However, as the weld quality in general was good in terms of welding defects and

Rw, and also because there were no weld quality acceptance criteria related to weld width specified in

AWS D17.1 and BS EN ISO 13919-2, it was not possible to judge the specimens using the criteria on

face and root weld widths in BS EN 4678 alone but also had to consider other factors as well. It was

found that the magnesium content in the weld measured by EDX decreased with increasing feed rate

from around 0.80% at 0 m/min to 0.65% at 7.0 m/min, which obviously increased the silicon content

as well, from around 0.10% at 0 m/min to 0.88% at 7.0 m/min, illustrated in more details in Figure 5.

The Rw was above 0.6 at all feed rates so the processing stability for full penetration welding was

high. A trend was observed where the Rw decreased with increasing feed rate, meaning that as

mentioned above, the rate of change in the top width was greater than the bottom width with the feed

rate. Underfill was observed when autogenous welding, above the maximum limit of 0.13 mm in AWS

D17.1 and 0.15 in BS EN ISO 13919-2 but below 0.30 mm in BS EN 4678. The depth of underfill was

reduced by more than half at 1.5 m/min which passed all criteria so welding with filler wire reduced

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the formation of underfill defects and even eliminated at higher feed rates. As expected the height of

reinforcement or excess weld metal increased with increasing feed rate and at 7.0 m/min, it was

above the most stringent limit of 0.55 mm in BS EN 4678 but below that in BS EN ISO 13919-2 and

AWS D17.1, of 0.65 and 0.99 mm respectively. For the rest of the specimens, it was less than 0.55

mm. The depth of undercut was below the maximum limit of 0.15 mm in BS EN ISO 13919-2 and BS

EN 4678 at all feed rates, but was above the 0.05 mm in AWS D17.1 at the two highest feed rates of

5.0 and 7.0 m/min. The problems with welding defects including surface porosity, reinforcement and

undercut were found to be the most significant at 5.0 m/min and 7.0 m/min. It was possible that the

feed rate was too high at these feed rates which supplied too much filler metal to the weld pool for the

given laser power and welding speed. High feed rates produced instabilities, whereas, low feed rates

did not sufficiently modify the chemical composition of the weld pool [11].

a) b)

Figure 3 a) Relationship between weld width, magnesium and silicon content and filler metal feed rate at a laser power of 4.9 kW, a welding speed of 3.0 m/min, +4 mm defocus and with helium shielding gas, and b) the resultant weld width ratio, undercut, underfill and reinforcement

The weld shape was similar at all feed rates with an hourglass shape but with larger weld widths with

increasing feed rate. The small underfill observed at 0 m/min was reduced at 1.5 m/min and

completely removed at 2.0 m/min. The weld quality was good between 2.0 and 4.0 m/min but at 5.0

m/min, clustered pores with a maximum diameter of 0.21 mm. The criterion in AWS D17.1 for surface

porosity specifies at least 8 times the size of larger adjacent imperfection and a smaller maximum

pore size of 0.75 mm compared to 0.99 mm for subsurface pores. BS EN 4678 and BS EN ISO

13919-2 on the other hand, specify the same diameter of 0.99 mm for surface pores but with the

distance between the individual pores in clustered porosity greater than ¼ of the material thickness.

Although the size of surface pores was small, they were too close to each other so were

unacceptable. These surface pores were not observed at 7.0 m/min but instead large undercut

0.0

0.5

1.0

1.5

2.0

0.0

1.0

2.0

3.0

4.0

5.0

0 2 4 6 8

We

ight (%

)

We

ld w

idth

(m

m)

Filler wire feed rate (m/min)

TopBottomSiMg

0.00

0.10

0.20

0.30

0.40

0.50

0.0

0.2

0.4

0.6

0.8

1.0

1.2

0.0 2.0 4.0 6.0 8.0

Imperf

ections (

mm

)

Rw,

Rein

forc

em

ent (m

m)

Filler wire feed rate (m/min)

RwReinforcementUndercutUnderfill

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defects on the top surface and excessive penetration on the bottom surface was found. Therefore, it

was concluded that the best weld quality in terms of morphology was produced when the filler metal

feed rate was in the optimum range of 1.5 to 4.0 m/min.

0 m/min 1.5 m/min 2.0 m/min 2.5 m/min

3.0 m/min 4.0 m/min 5.0 m/min 7.0 m/min

Figure 4 Transverse sections of welds and weld top bead profiles produced with different filler metal feed rate at a laser power of 4.9 kW, a welding speed of 3.0 m/min, +4 mm defocus and with helium shielding gas

The effect of silicon addition on the crack sensitivity of AA 2024-T3 is shown in Figure 5. The weight

content of magnesium and silicon in the weld was measured using energy dispersive X-ray

spectroscopy on specimens welded with different filler metal feed rates. It was found that the

magnesium level dropped whereas, the silicon level increased with increasing feed rate. As it can be

seen from the crack sensitivity curves for aluminium in Figure 5, the crack sensitivity is the maximum

when the Cu content is approximately 3%, Si is 1%, and Mg is 1.5%. AA 2024-T3 contained

approximately 4.5% Cu which may initially have indicated that it has relatively low crack sensitivity.

However, it also contained a small amount of Mg close to the critical level of 1.5% in the base metal,

which increased the crack sensitivity by widening the coherence range, and depressing the solidus

temperature but not the highest temperature of coherence [27]. Segregation of the low boiling point

alloying elements such as Zn and Mg caused hot cracking at the grain boundaries due to the

shrinkage strains during the solidification process. The presence of Si as well as Mg in the AA 2024-

T3 base metal increased the risk of inducing coarse Mg2Si precipitates so the maximum content of Si

in the 2024 alloy was required to be less than 0.7% [28]. According to Davis [29], the sensitivity

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decreases rapidly if the Si content exceeds 1.5%. The dilution of the weld pool with excess silicon by

welding with the 4043 filler metal effectively reduced the percentage of Mg2Si in the weld by

combining with the base metal. Also, the addition of silicon to the weld lowered the solidification

temperature and decreased the total shrinkage during freezing as mentioned previously to prevent

cracking. As a result, the peak of the solidification crack sensitivity curve for Al-Mg and Al-Mg2Si

shifted away from the crack sensitive ranges. The silicon content in the welded specimens were

detected to be less than the recommended 0.6% to avoid the crack sensitive range of Al-Si up to the

feed rate of 3.0 m/min but above 0.6% at higher feed rates of 4.0, 5.0 and 7.0 m/min. Therefore, the

solidification crack sensitivity was minimised by welding at a filler wire feed rate of 2.0 to 3.0 m/min

with a dilution ratio of 8.5 to 12.3%.

a) b)

Figure 5 a) Weight percentage (%) of main alloying elements in the weld as a function of filler metal feed rate obtained using energy dispersive X-ray spectroscopy (EDX) and b) aluminium crack sensitivity curves showing the effects of different alloy additions (Figure 5 b) modified from [21,30])

3.1 Micro-hardness

Micro-hardness testing in the transverse weld bead cross-sections as illustrated in Figure 6 showed

that all AA 2024-T3 welds were under-matched with the lowest hardness in the FZ. The hardness in

the BM was the highest as expected. It was also found that the hardness in the HAZ was greater than

in the FZ but lower than in the BM. Micro-hardness increased as a function of the distance from the

weld centre in which the FZ had a hardness of around 90-100 HV, the HAZ hardness in the range of

100-120 HV and the BM hardness in the range of 130-140 HV. The HAZ adjacent to the FZ showed

hardness values close to that in the FZ whereas, the HAZ adjacent to the BM showed a hardness

close that in the BM. Since the extent of the FZ and the HAZ was very small, the resulting hardness

gradient was very steep. On the other hand, the hardness distribution was relatively uniform across

the FZ in most specimens.

0

1

2

3

4

5

Cu Mg Si

We

ight

(%)

Vf = 0 m/minVf = 1.5 m/minVf = 2.0 m/minVf = 2.5 m/minVf = 3.0 m/minVf = 4.0 m/minVf = 5.0 m/minVf = 7.0 m/min

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The dissolution or loss of strengthening precipitates and alloying elements, softening in the FZ, and

over-aging in the HAZ were the main causes of hardness degradation during welding process. The

effect of grain growth with respect to strength was of minor importance but instead mainly influenced

by modification of precipitates [31]. Softening in the FZ was caused by microstructural changes as a

result of very high temperatures experienced in the FZ and the associated rapid heating and cooling

rates during welding. The heating action of the laser led to segregation of the strengthening elements,

magnesium and copper, and their hybrids (intermetallic compounds), formation and growth of non-

strengthening coarse precipitates, dissolution of strengthening precipitates and uniform re-distribution

of precipitating elements during heating which then froze due to fast cooling rates [32]. In addition,

softening can also be attributed to violent vaporization of low boiling point magnesium and element

variation resulting from the filler dilution [33] as observed in Figure 5 a) showing lower Cu and Mg

contents with increasing filler metal feed rate. The hardening effect was therefore, removed and the

mechanical properties of the weld degraded. The hardness in the FZ was similar to the hardness

measured in a fully solution treated and quenched AA 2024 of around 80 HV [34]. Even though the FZ

partially recovered its hardness by natural ageing at room temperature for several days after welding,

the effect was small due to inhomogeneous distribution of solute atoms. Loss of volatile elements

such as magnesium and zinc for strengthening also contributed to lowering the hardness in the FZ by

affecting the weld pool chemistry. The welding thermal cycle also affected the precipitation behaviour

in the HAZ such as dissolution, precipitation and coarsening so the HAZ was divided into two different

microstructural regions of partially melted zone and over-aged zone. The hardness in the partially

melted zone decreased due to dissolution of strengthening precipitates during melting and

segregation of alloys during solidification. In the over-aged zone, coarsening of the strengthening

semi-coherent S ́phase as well as transformation to the non-strengthening incoherent stable S phase

reduced the hardness [35]. As it can be seen from Figure 6, changing the filler metal feed rate does

not significantly affect the heat input and therefore, its effect on micro-hardness was relatively small

but rather influenced the weld width, where increasing the feed rate increased both the face and the

root weld widths.

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a)

b)

c) d)

e) f)

g) h)

Figure 6 Micro-indentation hardness distributions of fibre laser welded AA 2024-T3 welds as a function of filler wire feed rate

50

75

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-3 -2 -1 0 1 2 3

Vic

ker's

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ness (

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Distance from weld centre (mm)

Top

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W15P= 4.9 kW, V = 3.0 m/min,f = +4 mm

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Distance from weld centre (mm)

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W22P= 4.9 kW, V= 3.0f = +4 mm, w = 1.5 m/min

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W23P= 4.9 kW, V= 3.0f = +4 mm, w = 2.0 m/min

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W24P= 4.9 kW, V= 3.0f = +4 mm, w = 2.5 m/min

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W25P= 4.9 kW, V= 3.0f = +4 mm, w = 3.0 m/min

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W26P= 4.9 kW, V= 3.0f = +4 mm, w = 4.0 m/min

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W27P= 4.9 kW, V= 3.0f = +4 mm, w = 5.0 m/min

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W28P= 4.9 kW, V= 3.0f = +4 mm, w = 7.0 m/min

15

3.2 Global and local tensile properties

Figure 7 shows the full field longitudinal strain distribution at different percentage of the fracture time

at 25, 50, 75 and 99%. The strain profiles across the weld showed a sharp strain gradient from the

weld centreline to the BM at the onset of final fracture. The maximum strain localisation was located in

the FZ for both specimens, whereas, the minimum occurred in the BM. The strain distribution was

uniform and symmetrical about the weld centreline throughout the deformation with relatively larger

local strains of around 11% in the FZ for both specimens welded with or without filler metal.

Figure 7 Full field longitudinal strain distributions in the loading direction for fibre laser welded AA 2024-T3 showing the development of strain localisation relative to the time to fracture

The strain measured in the BM was of the order of only 0.5-1.0% at failure which indicated that it was

still under elastic loading. Although the boundaries of different characteristic regions corresponding to

the BM, the HAZ and the FZ cannot be easily identified from the strain maps, their locations were

measured and marked outside the processed regions prior to testing and also it was possible to

determine their extent by examining the highly non-uniform strain distribution across the weld at different

load levels. It was obvious from these strain maps that the stiffness of the welded specimen was a result

of the stiffness of the three different microstructural regions

The global tensile behaviour of the welded specimens and the BM and the corresponding mechanical

properties are shown in Figure 8. It was possible to create any size and number of gauge lengths on

the processed DIC images and so the global stress and strain curves were determined for a 25 mm

gauge length equal to that of the extensometer used. For mechanical characterisation of the welded

11

9

7.5

6

4.5

3

1.5

0

εy (%)11

9

7.5

6

4.5

3

1.5

0

εy (%)Autogenous Welded with filler metal

Percentage of failure timePercentage of failure time

25% 50% 75% 99% 25% 50% 75% 99%

16

joints, the elastic modulus, yield strength, ultimate tensile strength and elongation to failure were

determined as listed in Table 5.

Table 5 Results of tensile properties determined from tensile testing welded joints

Welding mode UTS (MPa) YS (MPa) Elongation (%) E (GPa)

Filler 370 277 3.7 63.6 Autogenous 364 283 2.4 66.6 Unwelded 463 358 14.9 69.1

Global tensile test results showed significant losses in ductility and tensile strength in the welded

specimens compared to the unwelded BM due to plastic strain localisation and increased constraint

within the lower strength weld region of the welded joint for the composite gauge length. Only moderate

variations in the yield strength and elastic modulus were observed, while considerable differences in

elongation to failure and ultimate tensile strength were measured. The addition of filler metal reduced

the risk of welding defects and resulted in a significantly greater ductility over 3.5% and a fairly higher

tensile strength of around 380 MPa than the other specimen.

Figure 8 Global stress and strain curves obtained from tensile testing welded specimens

The influence of local material behaviours in the various weld zones on the overall weld response was

determined using the DIC by assuming an iso-stress condition for all specimens, where the global stress

was considered as the corresponding local stress at any point within the analysed displacement data

field. The local strain data were derived from the DIC by measuring the local strain over each individual

region, which were then plotted against the global stress data to obtain the local tensile properties in

the FZ and the HAZ on either side of the FZ (HAZ1, HAZ2) as shown in Figure 9, As it was only possible

to obtain the full tensile response in the weakest region where strain localisation occurred, the stress

and strain curves in the stronger regions such as the BM were not obtained. The high hardening rate

0

100

200

300

400

500

0.00 0.05 0.10 0.15

Str

ess (M

Pa)

Strain

P=2.9 kW, V=1.5 m/min, f=+4 mm, filler

P=2.9 kW, V=1.5 m/min, f=+4 mm

BM

17

and low ductility observed in the global stress and strain curves of the welded specimens proved that

strain localisation occurred in the weaker regions of the weld. In fact, as all the AA 2024-T3 specimens

were under-matched, the local strain evolution in the FZ or the HAZ was successfully calculated up to

complete fracture, which was not possible to determine from standard tensile specimens without the

DIC because of the narrow size of the FZ. It was found that the longitudinal strains measured in the

weld was much higher than the measured global fracture strains which indicated that the strain

distribution within the 25 mm gauge length was not uniform but highly localised in the weaker FZ or

HAZ so the overall behaviour was dominated by that of the weakest component of the specimen and

minimal plastic deformation occurred outside the weld [36]. The maximum strain was reached in the FZ

for both specimens.

P=2.9 kW, V=1.5 m/min, f=+4 mm, filler metal feed rate=2.6 m/min (10.5% dilution ratio)

P=2.9 kW, V=1.5 m/min, f=+4 mm

Figure 9 Local mechanical responses in FZ and HAZ constructed from full field DIC tensile tests compared to overall responses

Figure 10 shows the relationship between longitudinal strain distributions across the weld at the onset

of final fracture obtained using the DIC and micro-hardness distributions in and around the respective

welds from micro-indentation hardness testing. The strain values in the BM were small tending to zero

with increasing distance from centre of weld, whereas, the strain increased in the HAZ at the same

positions where the hardness decreased so the hardness measurements were used to confirmed the

extent of the HAZ as well as the FZ where the hardness was the minimum with only small variations.

The strain distribution was more uniform in the specimen welded with filler wire, with a wider weld width.

0

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100

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200

250

300

350

400

450

0.00 0.02 0.04 0.06 0.08 0.10 0.12

Tru

e s

tre

ss (

MP

a)

True strain

FZ

HAZ1

HAZ2

Global

0

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100

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250

300

350

400

450

0.00 0.02 0.04 0.06 0.08 0.10 0.12

Tru

e s

tre

ss (

MP

a)

True strain

FZ

HAZ1

HAZ2

Global

18

P=2.9 kW, V=1.5 m/min, f=+4 mm, filler metal feed rate=2.6 m/min (10.5% dilution ratio)

P=2.9 kW, V=1.5 m/min, f=+4 mm

Figure 10 Local strain distributions across the welded joint obtained using the DIC at the onset of final fracture superimposed on micro-indentation hardness distributions measured in the transverse weld cross-sections

3.3 Microstructural observations using SEM

Scanning electron microscopy (SEM) analysis was performed on the fracture surfaces of the AA 2024-

T3 DIC tensile specimens as shown in Figure 11 to examine the fracture behaviour and the existence

of welding defects such as hot cracks and porosities. It was observed that for both specimens inspected,

failure occurred within the weld, either in the FZ or in the HAZ/FZ boundary but not in the BM due to

weld under-match.

For the specimen which was welded without filler wire, a rough and irregular fracture surface, consisting

of microvoids and dimples that act as microvoid nucleation sites, was observed. This indicated that the

specimen failed in a more ductile manner where fracture initiated by microvoid coalescence and then

by dimple rupture. Rough fracture surfaces with fine equiaxed dimples are the characteristic features

of ductile transgranular fracture mode. Hydrogen induced spherical shaped pores of different sizes

ranging from around 100 to 300 μm with round tip dendrites, were identified on the fracture surface in

clusters which were responsible for crack initiation during deformation. Porosities reduced the effective

cross-sectional area of the welded joints and therefore, caused stress concentrations which

deteriorated the strength of the joints in proportion to the reduction of the cross-sectional area.

For the specimen which was welded with filler wire at an optimum rate of 2.6 m/min (10.5% dilution

ratio), even greater amount of finer dimples was found to dominate the fracture surfaces which is the

characteristic feature of purely ductile fracture. A significant amount of localised microscopic weld

plasticity was observed which improved the tensile strength and ductility of the respective welds as

previously determined from tensile testing. While intergranular inter-dendritic micro hot cracks and

80

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150

0.0

2.0

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ain

(%

)

Distance from weld centre (mm)

Top

Middle

BottomBM FZHAZ HAZ BM

80

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0.0

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Distance from weld centre (mm)

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BM FZHAZ HAZ BM

19

porosities were detected in the other specimen, they were almost or completely absent in this specimen

so a significant improvement in tensile strength and ductility were obtained. In fact, porosity level less

than 3% of the total volume does not affect much the yield and tensile strength of a material but may

affect ductility [37].

P=2.9 kW, V=1.5 m/min, f=+4 mm, filler metal feed rate=2.6 m/min (10.5% dilution ratio)

P=2.9 kW, V=1.5 m/min, f=+4 mm

Figure 11 SEM fracture morphology of DIC tensile test specimens at 500x magnification showing different modes of failure and presence of welding defects

4 CONCLUSIONS

Welding with filler metal reduced the crack sensitivity of AA 2024-T3 but it was also important to

optimise the filler metal feed rate to avoid the formation of welding defects and keyhole instability.

High feed rates produced instabilities, whereas, low feed rates did not sufficiently modify the chemical

composition of the weld pool. The optimum feed rate was found to be around 2-3 m/min (8.5-12.3%

dilution ratio) which minimised the crack sensitive range of Al-Mg, Al-Mg2Si and Al-Si according to the

measurements obtained by EDX.

Changing the filler metal feed rate does not significantly affect the heat input and therefore, its effect

on micro-hardness was relatively small but rather influenced the weld width, where increasing the

feed rate increased both the face and the root weld widths.

20

Only small variations in the yield strength and elastic modulus were observed, while considerable

differences in elongation to failure and ultimate tensile strength were measured. The addition of filler

metal reduced the risk of welding defects and improved ductility to over 3.5% and a fairly higher tensile

strength of around 380 MPa than without. The strain distribution was more uniform in the specimen

welded with filler wire, with a wider weld width.

Microstructural examination showed that the addition of filler wire increases the number of finer dimples

within the weld, resulting in a purely ductile fracture behaviour. A significant amount of localised

microscopic weld plasticity was observed which improved the tensile strength and ductility.

5 ACKNOWLEDGEMENT

The strong support from the Aviation Industry Corporation of China (AVIC) and Beijing Aeronautical

Manufacturing Technology Research Institute (BAMTRI) for this funded research is much appreciated.

The research was performed at the AVIC Centre for Structural Design and Manufacture at Imperial

College London. Dr C M Davies acknowledges the support of EPSRC under grant number

EP/I004351/1.

21

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