3.1. Microstructural Morphologies
The macrographs of 2014-O and 2014-T6 welds, at different welding parameters, are shown in Fig. 3 and Fig. 4, respectively. All the macrographs are taken at 20x magnification. The advancing and retreating sides of the joint are indicated by ‘AS’ and ‘RS’ respectively. The typical four regions of the friction stir welded joint are also labeled as NZ, TMAZ, HAZ and BM. From the figures it is clearly observed that the joints in both temper conditions have very different microstructural morphologies. Figure 3 shows 2014-O joints welded at different welding speeds. It is seen that the weld nuggets have partial ‘onion rings’ features in all the three welds. The fraction of these features reduces as the welding speed is increased. On the other hand, in 2014-T6 joints, Fig. 4, the ‘onion rings’ features are present in greater amount over the entire width of the NZ. The presence of ‘onion rings’ like features shows that the joints experienced sufficient plastic flow during welding process and the plasticized material was deposited in multiple layers [29]. It was observed by Sutton et al that ‘onion ring’ like features are due to the banded microstructure. These bands are formed due to uneven distribution of hard particles [30]. In 2014-O base material, majority of the precipitates exist in stable state and a relatively less fraction of metastable precipitates is present in the material in the form of Al2Cu, AlCuMn and AlCuMg [31]. At low welding speed, the peak temperature during FSW is high enough to dissolve and solutionize some fraction of the metastable particles into the aluminum matrix. However, the large amount of stable precipitates remains unaffected and are accumulated in high-strain regions giving rise to formation of bands of low and high particle density [30]. These bands give rise to partial ‘onion rings’ morphology in the NZ of 2014-O on advancing side. It is worth noting that, as the welding speed in increased, the peak temperature is reduced and hence the proportion of ‘onion rings’ features also reduces, Fig. 3(c).
For 2014-T6 base material, the majority of the precipitates exist in the metastable state and only a small fraction of stable precipitates is present in the material. During FSW process an appreciable amount of these metastable precipitates is dissolved in the aluminum matrix. As discussed above, the stable precipitates remain unaffected and are segregated in high strain regions, resulting in the formation of high and low particle density bands in the NZ. Such bands appear as ‘onion rings’ features in the NZ of 2014-T6 joints.
From Figs. 1 and 2 it can also be observed that in 2014-O joints, the size of weld bead remains almost the same at different welding speeds. However, in case of 2014-T6 joints, there is a marked effect of welding speed on the dimensions of weld bead as it significantly reduces with the increase of welding speed. This reduction in weld size is due to lesser amount of material flow around the rotating pin at high welding speed. Same observations were made by Venkateswarlu et. al. in their study on 2219-T62 and 2219-T87 alloys [32].
The optical micrographs of base metals (BM) in ‘O’ and ‘T6’ conditions are shown in Fig. 5. The microstructure represents typical features of a rolled material consisting of elongated grains oriented along the rolling direction consisting of fine Al2Cu precipitates, distributed in the α-Al matrix and dark colored particles. These particles are generally Al2CuMg and (Cu,Fe,Mn)Al6 or Al7Cu2Fe [33]. The average grain size of 2014-O and 2014-T6 BM was 19.7µm and 27µm, respectively.
The NZ microstructures of 2014-O and 2014-T6 joints are presented in Fig. 6. It was observed that the grain size in the NZ of 2014-O is coarser than 2014-T6. This may be attributed to the difference in their temper conditions which can significantly affect the mechanism of recrystallization. However, it was in contradiction to the observations made by Venkateswarlu et al in 2219-O and 2219-T6 aluminum alloys, where grain size of NZ was large in T6 alloys than the O alloys [34]. This may also be ascribed to the difference in hardness of both the materials. As 2014-O has much lower hardness than 2014-T6 samples, the mechanical agitation and subsequent material flow and temperature rise will be more severe in 2014-O than in 2014-T6 leading to coarsening of recrystallized grains during cooling from weld temperature to room temperature in the former than in the latter. The microstructure in both temper states consists of recrystallized fine-equiaxed grains in the NZ. This can be explained in terms of rotation and linear motion of tool during FSW process which causes severe plastic deformation by tool pin and high thermal effects by tool shoulder, resulting in the semi-solid material in this region. This semi-solid material forcefully flows along the tool pin surface of the rotating tool causing severe shear strains. Due to severe plastic strains at elevated temperature, dynamically recrystallized grains are nucleated at the grain boundaries of initially large and elongated grains in the NZ [35, 36]. The same trend was found in 2014-O and 2014-T6 samples welded at different welding speeds.
The microstructural variation from NZ to TMAZ are presented in Figs. 7 and 8 for 2014-O and 2014-T6 joints, respectively. In 2014-O joints, the boundary between NZ and TMAZ is not clear on the RS (Fig. 7a) whereas a narrow band of transition of grains from NZ to TMAZ was observed on AS (Fig. 7b). The same trend was also observed in other 2014-O samples which were welded at travel speeds of 200 and 280 mm/min. However, 2014-T6 joints showed a very distinct transition interface between the NZ and TMAZ on RS as well as AS irrespective of the travel speed. The same phenomena have been observed by other workers [9, 37, 38] and is explained in terms of the difference in plastic flow of material on advancing and retreating sides of a weld. The interface is more distinct on the side having higher shear forces and plastic strains. Since on AS, the directions of tool travel and plastic deformation are in the same direction, the shear forces and plastic strains are higher on this side, resulting in clear interfaces between NZ and TMAZ.
The region of TMAZ experiences both high temperature and plastic deformation to lesser extent than the NZ. The combined effect of deformation and temperature results in a microstructure which consists of elongated and rotated grains. The recrystallization mainly begins at NZ/TMAZ interface. The grain size of TMAZ is much larger than the NZ. Due to the thermal effects, the strengthening precipitates grow and get overaged thus causing drop in hardness in this region [35].
Figure 9 shows the HAZ in 2014-O and 2014-T6 samples. The size of HAZ in 2014-T6 is relatively larger than 2014-O. In both cases, there is no discernable boundary between HAZ and the base metal because the microstructural transition is not significant. HAZ experiences only thermal fluctuations and no plastic deformation takes place in this region. As a result, the grains are not recrystallized and the grain size is similar to or slightly larger than the base metal. Due to the thermal effects, hardening precipitates can dissolve or coarsen resulting in reduction of hardness. The extent of precipitate dissolution and coarsening depends on the base material condition.
3.2. Microhardness of the joints
Microhardness test were conducted on the transverse cross-sections of the weld joints. The microhardness profiles of 2014-O and 2014-T6 joints, welded at different welding parameters, are presented in Figs. 10 and 11 respectively. In 2014-O temper condition, the hardness is increased in the welded region with respect to BM whereas, and in 2014-T6 temper condition, it is decreased as compared to the BM. This may be ascribed to the difference in their BM temper conditions. The BM in 2014-O condition is mainly composed of α-Al grains in which precipitates are not dispersed homogeneously. Whereas, BM in 2014-T6 condition consists of α-Al grains with thoroughly distributed fine strengthening precipitates [39]. From the hardness profiles in Fig. 8, it can be observed that hardness increases in all 2014-O joints, irrespective of the welding parameters used. Also the increase in hardness is almost the same in all cases. Same observations were also made by [40–42] in different friction stir welded aluminum alloys. This increase in hardness in ‘O’ temper condition is due to the fact that the grain refinement took place in the weld region especially in the NZ, thus giving rise to hardness in this region. Since the temperature is high enough there may be some precipitation of strengthening particles during cooling from high temperature to room temperature [43, 44]. It was also observed that the rise in hardness did not vary significantly with change in welding speed.
A decrease in hardness occurred in weld zone of all the samples in 2014-T6 condition as opposed to 2014-O samples. It was due to the fact that high temperature and severe plastic strains dissolved the hardening precipitates back into the α-Al matrix of NZ. The samples in 2014-T6 condition showed a typical ‘W’ shaped profile across the welded region, Fig. 11. The partial recovery in hardness of NZ is due to grain refinement and re-precipitation of hardening phases during cooling from welding temperature to room temperature [43, 44]. The loss in hardness was maximum in the TMAZ and HAZ which was due to dissolution and coarsening of hardening precipitates and grain growth [45, 46]. The hardness loss was not same for all the samples of 2014-T6 condition at different welding parameters. The samples welded at low travel speed (120mm/min) showed a maximum loss in hardness by giving a value of minimum hardness of 85 HV, whereas, the samples welded at high travel speed (280 mm/min) showed a minimum hardness upto 92 HV. This may be caused by high thermal fluxes at low travel speed resulting in greater volume of precipitate dissolution than at high travel speed where low thermal effects dissolved lesser amount of the hardening precipitates. Same interpretations were made by other researchers in FSW of different aluminum alloys [34, 35].
3.3. Mechanical properties of the joints
The mechanical properties of the 2014-O and 2014-T6 joints, welded at different welding parameters, are presented graphically in Figs. 12 and 13, respectively, alongwith the BM properties. The yield strength, ultimate tensile strength, percentage elongation and percentage of joint efficiencies of and the BM properties are also summarized in Table 4 for comparison.
From Fig. 12, it is evident that the yield strength and ultimate tensile strength of 2014-O joints were almost same as that of BM, however, the joints showed about 15% drop in elongation. Thus the joint efficiency was 100% as shown in Table 4. The decrease in elongation is due to the high hardness in the NZ, Fig. 10. During tensile test, the load concentrates more in the low hardness region of the joint and final fracture takes place in this region [45]. For 2014-O joints the low hardness region is base metal, see Fig. 10. Therefore, NZ offered greater resistance to plastic deformation during tensile test, thus reducing the overall elongation. Similar observations have been made by other researchers in friction stir welded 6061-O and 7075-O alloys [45, 46].
The tensile samples of 2014-T6 condition showed different mechanical behavior as compared to 2014-O samples. The mechanical properties of the 2014-T6 joints were lower than the base metal as illustrated in Fig. 13. The mechanical properties of the friction stir welded joints are determined by welding defects and the hardness distribution across the weld joint [35]. In a defect free joint, hardness profile across the weld mainly controls the mechanical properties. It is also known that the friction stir welded joint consists of a non-uniform microstructure with different zones alongwith their interfaces, all of which have different mechanical properties [9, 47]. From Fig. 11, it can be noted that 2014-T6 joints have low hardness regions, such as NZ, TMAZs and HAZs as compared to BM. The mechanical properties of these regions are lower than the base metal. As discussed earlier, the stress is concentrated in these regions during the tensile test and the joints show lower mechanical properties than the BM. The 2014-T6 joints also showed a remarkable loss in ductility, Table 4. This may be due to the excessive loss of hardness in the weld region. During tensile test the load is confined to this region and major portion of plastic deformation takes place within weld region while BM remains almost intact due to higher hardness. Similar observations were also made in other friction stir welded aluminum alloys [35, 45, 46].
The yield strength efficiency (YSE), ultimate tensile strength efficiency (UTSE) and elongation efficiency (EE) for 2014-O and 2014-T6 joints are summarized in Table 4. The YSE and UTSE and EE are the ratios of the average yield strength, ultimate tensile strength and elongation of the joints, respectively, to those of the BMs [29]. It can be observed that efficiencies of different mechanical properties strongly depend on BM condition. The 2014-O BM has a complete stable condition and the YS and UTS efficiencies of 2014-O joints are almost same as that of BM irrespective of the welding parameter used. However, the EE is lower than BM, being 90% at low welding speed and 85% at high welding speed. In 2014-T6, the joints showed much lower weld efficiencies than the BM. The YSE remained around 50% of BM, whereas UTSE was around 70% of the BM. The EE in heat treated condition reduced significantly, being 50% at low welding speed and 32% at high welding speed. In general, it was observed that the strength and elongation efficiencies are decreased when the stability for precipitation of the BM decreases [35].
Table 4
Tensile properties of the base metals and the welded joints
Joint ID
|
2014-O
|
2014-T6
|
Value*
|
Std. Dev.
|
Efficiency, %
|
Value
|
Std. Dev.
|
Efficiency, %
|
BM
|
|
|
|
|
|
|
• YS [MPa]
|
92
|
1
|
-
|
441
|
3
|
-
|
• UTS [MPa]
|
189
|
1
|
-
|
495
|
5
|
-
|
• Elongation [%]
|
17
|
2
|
-
|
15
|
1
|
-
|
120
|
|
|
|
|
|
|
• YS [MPa]
|
90
|
3
|
98
|
220
|
2
|
50
|
• UTS [MPa]
|
194
|
3
|
103
|
345
|
4
|
70
|
• Elongation [%]
|
15
|
2
|
86
|
8
|
1
|
50
|
200
|
|
|
|
|
|
|
• YS [MPa]
|
86
|
3
|
93
|
237
|
2
|
54
|
• UTS [MPa]
|
179
|
2
|
95
|
359
|
5
|
72
|
• Elongation [%]
|
15
|
3
|
90
|
6
|
2
|
42
|
280
|
|
|
|
|
|
|
• YS [MPa]
|
85
|
7
|
92
|
242
|
7
|
55
|
• UTS [MPa]
|
186
|
4
|
98
|
359
|
3
|
73
|
• Elongation [%]
|
14
|
1
|
85
|
5
|
0.4
|
32
|
* average value of three sample
|
3.4. Fracture locations and fracture surfaces of the joints
The analysis of fracture locations and fracture surfaces was carried out to understand the mechanical behavior of 2014-O and 2014-T6 welded joints. The fracture location represents the weakest region of the joint. Analysis of fractured tensile samples provides useful information about structure-property relationship [33]. From the Table 5, which shows the representative fracture locations of 200 samples, it can be seen that 2014-O and 2014-T6 have different fracture locations. Same fracture locations were observed in other samples of a group whatever the welding parameters were used. All the tensile samples in ‘O’ condition fractured in the BM at RS of the weld joint, Table 4. Since BM had lowest hardness for 2014-O joints, Fig. 10, all the stress concentration occurred in BM.
In 2014-T6 samples, the fracture occurred on RS at NZ/TMAZ interface because it was the weakest region of the weld with minimum hardness, Fig. 11. At the same time, there was drastic variation of microstructure at NZ/TMAZ interface. The TMAZ in these samples has twisted coarse grained microstructure as compared to NZ and HAZ with partially dissolved and coarsened strengthening particles. This significant difference in the internal structure gave rise to the weakest region in the weld joint. Due to low hardness and distinct microstructure, the stress was concentrated in this region during the tensile test.
SEM fractographic analysis has been carried out to understand the mechanism of fracture in the samples. The SEM fractograph of the fracture surface of the 2014-O tensile sample is shown in Fig. 14. Since all the samples, welded at different welding parameters, fractured from BM, only a representative fractograph of 200 sample is shown here. The fracture surface of 2014-O joints revealed a typical ductile fracture consisting of deep dimples, of varying size and shape, which result from microvoid nucleation and coalescence. The size of the dimples shows that appreciable plastic deformation took place in the BM which led to deep dimples and high ductility. The notable necking in the vicinity of fracture also supports this observation. Within the dimples, different particles were also observed. Hence it can be stated that samples in ‘O’ condition fractured by the microvoid coalescence at coarse constituent particles [48].
Figure 15 shows the SEM fractographs of fractured samples in 2014-T6 condition at different welding parameters. Since fracture in ‘T6’ condition always occur in weld zone, all the fractographs show a different morphology than the BM. The fracture surfaces show a fine grain size that become finer at high welding speed, Fig. 15 (c). Although the fracture surfaces represent a ductile fracture morphology, yet the dimple size was much finer and shallower than the 2014-O samples which shows less plastic deformation during loading and a low ductility [35]. This observation is in agreement with the elongation data given in Table 4.