Aluminium alloys AA 5083 and AA 6082 have been friction stir welded and the mechanical properties and microstructures of the welds have been evaluated.
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Friction stir welding (FSW) is a relatively new technique, developed by TWI1 in the early 1990s, for the ef®cient joining of aluminium and other low melting point alloys. The plates to be joined are clamped with a sturdy ®xture to a backing plate, counteracting the lifting force arising during welding. A rotating tool with a cylindrical shoulder is pushed into the material at the joint. The tool is then moved along the joint, frictionally heating the material to high temperatures (but still below the melting temperature). The heat decreases the resistance to plastic deformation of the material, which then easily moves behind the tool and forms a solid state weld as the stirred material is consolidated. Figure 1 shows a workpiece and the tool used to carry out the process. The process has been further developed by ESAB AB, especially in terms of the handling of the plates to be welded and the design of the equipment. These developments have led to commercially available `friction super stir welding' ISSN 1362 ± 1718
equipment.2 The ®rst installation, which has been running since 1996 with encouraging results, is used for welding plate and panels of aluminium (up to 6616 m in size) for the marine sector. In this application, two main base materials are used: AA 5083 and AA 6082. Before practical application, the mechanical properties of FSW joints in these two materials were assessed by a programme stipulated by Det Norske Veritas and found to be adequate. A general understanding of the microstructure of friction stir welds and how this is related to hardness and tensile properties is beginning to emerge. Some studies of the process itself, the temperature distribution around the weld during the FSW process, and the microstructural developments during FSW in some aluminium alloys have been reported recently.3 ± 13 The ®ndings of these reports are outlined in more detail below. However, there are also many aspects that are not yet understood and a detailed understanding of the microstructure of FSW joints and how it is related to, for example, process variables has yet to be developed. The aim of the present paper is to make a detailed assessment of the microstructure of friction stir welds in two aluminium alloys, AA 5083 (a non-hardenable alloy) and AA 6082 (a precipitation hardening alloy), using both optical and electron microscopy. These techniques allow changes in both the size and structure of grains, and effects on coarser particles as well as very ®ne scale precipitation, to be observed, thus covering the whole range of metallurgical processes occurring during FSW. By combining this information with mechanical data, a comprehensive description of friction stir welds in these two types of aluminium alloys is presented. Furthermore, by comparison with literature data, similarities and dissimilarities with other aluminium alloys have been identi®ed.
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To provide useful background information, some of the important ®ndings in earlier investigations are reported in more detail ®rst. This discussion is divided into three sections, covering microstructure, thermal cycles, and mechanical properties. First, however, some characteristics of the alloys involved are presented brie¯y. In the main, three alloy types are discussed. Two of these (AA 5083 and AA 6082) were used in the present experimental programme and the third (AA 7075) was studied in Refs. 5 and 6. Alloy AA 5083 is an Al ± 5Mg nonhardenable alloy, whereas AA 6082 is an age hardenable alloy containing y1%Si and 0.9%Mg, further strengthened with y0.7%Mn. Alloy AA 7075 is an age hardenable alloy containing y5.5%Zn, 2.5%Mg, and 1.5%Cu. In addition, brief discussion of AA 6061 and AA 6063, which are also age hardenable, is given. Alloy AA 6061 has similar silicon and magnesium contents to AA 6082, but higher copper content and lower manganese content. Alloy AA 6063 has Science and Technology of Welding and Joining
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Properties of friction stir welded Al alloys
A parent material unaffected by weld; B heat affected zone (HAZ); C,D thermomechanically affected zone (TMAZ); C unrecrystallised area found in Al alloys; D recrystallised nugget found in aluminium alloys
For aluminium alloys, the TMAZ can be distinguished from the nugget zone. In the TMAZ, the combination of high temperature and large strains causes deformation of the grain structure, but no recrystallisation takes place. Beyond is the heat affected zone (HAZ), which is affected by the heat but not by deformation.
1: workpiece; 2: probe; 3: shoulder of tool
2
% 3
silicon and magnesium contents of y0.5% and no other strengthening elements. Aluminium alloys always contain some impurities, mainly iron and silicon. Owing to the low solid solution solubility of many elements, formation of intermetallic particles is very common. Iron, silicon, manganese, copper, and chromium, for example, form intermetallic compounds. Thus, the possible intermetallics that can form are numerous.14 They are usually fairly coarse and easily resolved by optical microscopy. However, in manganese bearing alloys, a relatively ®ne dispersion of particles, of the Al ± Mn ± Si type and having a bcc crystal lattice, forms during homogenisation. These dispersoids can have a profound effect as nucleation sites for ®ner precipitates, but in themselves are too coarse to have any signi®cant effect on tensile properties. The non-hardenable alloys receive their strength from solid solution hardening and from deformation (i.e. dislocation strengthening). In the age hardenable alloys, hardening occurs through precipitation of very ®ne scale precipitates. The precipitation sequence is complex and contains many steps, usually starting from precipitation of Guinier ± Preston (GP) zones. For example, in Mg ± Si containing alloys, the GP zones can transform to hardening b0-Mg2Si or non-hardening b9-Mg2Si.15 In AA 7075 the hardening precipitate is of the type MgZn2.
Published results on the microstructure of friction stir welds concern mainly two aspects, namely the general microstructure of the welded zone, including the formation mechanisms of the central, heavily stirred part of the weld, and the effect of welding on the ®ne scale precipitation in hardenable aluminium alloys. In a few cases the effect of welding on the inclusions in the materials has been touched upon. Following Threadgill,3 the microstructure in a crosssection of a FSW joint can be divided into several zones (Fig. 2). In the centre lies the nugget, which usually has a ®ne grain size and contains a characteristic `onion ring' structure. It is generally assumed that the ®ne grain size is a result of a recrystallisation process.3 However, it should be noted that for aluminium alloys recrystallisation is con®ned to the nugget zone. In other friction stir welded alloys, recrystallisation occurs much more easily and over a much wider zone, the thermomechanically affected zone (TMAZ). Thus, in general, the nugget is considered as a part of the TMAZ. Science and Technology of Welding and Joining
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Liu 4 reported a TEM investigation of AA 6061. They noted the large difference in grain size between the base material and the nugget zone, similar to the ®ndings of Threadgill. However, the grain size in the nugget zone was not extremely small, as, for example is observed in alloys undergoing superplastic deformation. They also examined the dislocation content in the different regions and found that the nugget zone had a much lower dislocation density than the base material. Based on their observations, Liu argued that dynamic continuous recrystallisation was the formation mechanism for the microstructure of the nugget zone, a process which occurs by formation of subgrains. The subgrain boundaries subsequently increase in misorientation by accommodating dislocations generated by intragranular slip. Precipitation was also examined by TEM and it was noted that there were fewer large particles in the nugget zone than in the unaffected base material. The larger particles were rich in iron, but large AlSi particles were also found. However, the medium sized precipitates (0.1 ± 0.5 mm) were generally smaller in the nugget zone, while the number density variation was not mentioned. These particles had a stoichiometry of approximately AlxMgSi, where &~8 ± 10. In the TMAZ, a range of precipitates, both coherent and non-coherent, were found. Some of these particles were lathlike and contained aluminium, magnesium, silicon, and copper, and some were GP type precipitates on {100} cube planes.
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A detailed assessment of friction stir welds in AA 7075, concerning both microstructure and properties, was reported in Refs. 5 and 6. Of particular interest here is the effect of the welding process on the ®ne scale precipitates and how this affects the mechanical properties. However, before referring to the precipitates, another interesting observation should be noted. Rhodes 5,6 found that the onion ring structure in this alloy re¯ected differences in grain size: the rings contained alternating bands, having either a very ®ne (3 ± 5 mm) or a ®ne (5 ± 10 mm) grain size. The base metal contained three populations of small precipitates: one group in the grain boundaries (y30 ± 40 nm) and two within the grains. Intragranular precipitates, both assumed to contribute to the strength of the material, were either `coarse' (50 ± 70 nm) or `very ®ne' (10 ± 15 nm). The exact crystallography of the precipitates could not be determined, but it seemed that both the coarse
,
and the ®ne precipitates belonged to the same family. The coarser precipitates were elongated in the rolling direction. In the nugget zone, the coarse particles were somewhat coarser (60 ± 80 nm) and randomly distributed. The interpretation drawn was that they had dissolved and reprecipitated. The ®ne size precipitates population, however, had vanished. In transverse tensile testing of the welded joints of this alloy, the fracture appeared outside the nugget zone. The microstructure of the area close to the fracture surface was found to contain Mg ± Zn particles in only two size classes. The larger precipitates were 30 ± 40 nm in size and the smaller ones were 15 ± 30 nm in size. Precipitates were disc shaped with a thickness of 10 nm. It was assumed that both sizes contributed to the strengthening of the material. It was also noted that the density of coarser particles was higher in the fracture zone than in the base metal, while the density of the smaller particles was lower, with accompanying lower strength. In the hotter parts of the HAZ (closer to the nugget zone), the smallest particles (10 ± 15 nm in size) had vanished but the density of the larger ones (50 ± 100 nm in size) had increased. The larger particles consisted both of Mg ± Zn precipitates and chromium bearing dispersoids. It would be expected that this region would have the lowest yield strength, since there were no small precipitates, which are mainly held responsible for the strengthening effect. However, the strength of this area was found to be higher than the area further away from the nugget zone, where fracture had taken place. Rhodes 5,6 argued that instead the larger particles contributed to strength.
Sato 7 investigated the ®ne scale precipitates in AA 6063, both in FSW joints and in plates that had gone through simulated thermal cycles with different peak temperatures. In both types of specimen they observed the same precipitation reaction. The ®ne, needle shaped b0Mg2Si precipitates, which dominated in the base metal structure, were dissolved. Closer to the weld the amount of the rod shaped, incoherent b9-Mg2Si precipitates increased. However, as the distance to the weld centre decreased, the amount of this coarser precipitate also decreased and ®nally it was totally absent. By comparing the changes in precipitation between the welded and simulated specimens, Sato concluded that all precipitates were dissolved when the temperature was above 400³C and that the minimum in hardness occurred at a position where the peak temperature had been y350³C.
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The larger inclusions often present in aluminium alloys are rarely mentioned. Larsson 8 compared fusion welds and friction stir welds. In friction stir welds, there was a higher proportion of coarser inclusions. It was suggested that these were fragments from the very coarse inclusions in the base material. In the fusion welds a high density of ®ner inclusions was found.
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To estimate the maximum temperatures and cooling conditions occurring during FSW, measurements provided by Mahoney 6 and Backlund 9 were used. Mahoney presented a detailed map of peak temperatures at various positions relative to the nugget zone boundary. In the paper of Backlund fewer measurements were presented, although both peak temperatures and cooling curves were given. Figure 3 shows a plot of peak temperature as a function of distance from the nugget zone boundary, with data from both references. Backlund gave measurements as a
Properties of friction stir welded Al alloys
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6 7 v. 33 & 8 4 et al. $ 7 et al.5!9 3 3 7 3 33 8
function of the distance to the weld centreline and not to the nugget zone boundary. In the present work it was assumed that the nugget zone extends 5 mm from the weld centreline, so that results from the two investigations could be compared. It can be seen that there is a large difference in peak temperatures between the two investigations. This is probably a result of the differences in energy input to the weld. The plate thickness was similar in both cases, but there was a large difference in welding speed. Mahoney used a travel speed of 127 mm min21 while Backlund used 500 mm min21. There might also have been other variations in welding data, but details regarding these were not given. Frigaard 10 developed a two-dimensional numerical heat ¯ow model to calculate the thermal cycle during welding. This was used to analyse the temperature variations in front of the rotational tool and the steady state condition. The authors showed that a process parameter $0/, where $0 is the net heat input power, the welding travel speed, and the plate thickness, controls the thermal cycle.
Threadgill3 reported hardness variations across welded joints for different friction stir welded aluminium alloys, which can be considered to re¯ect the ultimate tensile strength of these joints. The hardness curves had a fairly common appearance, typical of welded aluminium alloys. For non-hardenable alloys, such as AA 5083 ± O (i.e. in an annealed state), hardness is more or less constant across the weld. However, the hardness will vary across a weld for non-hardenable, deformation hardened alloys. Typically, hardness decreases to a minimum value just outside the nugget zone. The hardness of the nugget zone itself is often slightly higher than the minimum value. The same pattern is also observed for age hardenable alloys, aged to higher hardness. Mahoney 6 presented a detailed evaluation of the tensile properties of friction stir welded AA 7075, both longitudinally in the nugget zone and transverse to the welded joint. They showed that for tests in the longitudinal direction the fracture surface appeared to be related to the onion ring structure. In the transverse tensile tests, fracture occurred in the HAZ outside the weld nugget. Science and Technology of Welding and Joining
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Fatigue properties of the present friction stir welds were reported in Ref. 8 and these data are presented below to complete the discussion of this subject. Further reports on fatigue properties of friction stir welds can be found in Refs. 16 ± 18.
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Friction stir welds were produced with AA 5083 ± O (4.6Mg ± 0.6Mn ± 0.3Si) and AA 6082 ± T6 (0.7Mg ± 0.5Mn ± 0.9Si) using different combinations of plate thickness and welding speed, according to Table 1. These welds were produced as part of a much larger programme aimed at de®ning the tolerance box for friction stir welds in these alloys. The welds in Table 1 were all within the tolerance box and are thought to represent typical welding parameters that could be used industrially. They were therefore selected for further investigations.
Specimens for optical microscopy were prepared by grinding down to 4000 mesh paper, followed by polishing down to 1 mm diamond. The ®nal preparation step was chemical polishing in a solution of OP-S suspension, KOH, and water. To reveal the microstructure, etching was carried out in a solution of hydrated FeCl3 and HF. One specimen of AA 6082 was prepared for observation in polarised light, by electropolishing in 10% perchloric acid in ethanol, followed by anodisation, using Bakers reagent (5%HBF4 in pure water) at 30 V for y1 min. Microstructures were studied using a light optical microscope and a scanning electron microscope. The latter was used both for higher resolution studies of the microstructure and also, using the backscatter mode, to reveal the grain size of the nugget zone. In addition the number density and size distribution of particles were estimated using a combination of light optical microscopy and SEM. Typical chemical compositions of the particles were determined qualitatively by EDS in the scanning electron microscope. Fracture surfaces from transverse tensile specimens and fatigue specimens were examined in the scanning electron microscope. Transmission electron microscopy was used for detailed examination of the ®ne scale precipitates. Specimens for the TEM investigation were prepared by electropolishing thin discs in a solution of nitric acid in methanol at 30 V and 220³C in a Struers Tenupol.
TMAZ, and nugget zone, out through the TMAZ and HAZ of the other base plate, and ending in the unaffected base material. In order to compare the hardness of various constituents in the nugget zone, detailed microhardness measurements (HV0.025) were made. To determine the location of the hardness minimum in AA 6082, hardness (HV0.025) was determined along a number of lines extending from the nugget zone, through the HAZ. The accuracy of the hardness determinations was estimated to 2 ± 3 units for HV1 and 4 ± 6 units for HV0.025. The ultimate tensile strength of the welded joints was determined using transverse tensile testing (with specimen thickness equal to plate thickness). Bend testing was performed over a mandrel with a diameter 6 times that of the specimen thickness, through an angle of 180³. Fatigue testing was conducted at room temperature and with constant amplitude, according to ASME Standard E466. The frequency was 140 Hz. The load ratio 0 (equal to minimum stress/maximum stress) was 0.1 and the stress range was 90 ± 150 MPa. Fatigue specimens were in the as welded condition, without machining of the weld top and root faces.
+ .* ; 3
All of the welds referred to in Table 1 were each made in one pass. For plate thicknesses greater than 15 mm, multipass welding is necessary A sturdy ®xture was used to hold the workpieces in the correct position during welding. Special attention was paid to avoiding root defects. Tool position must be carefully adjusted to ensure that the weld extends completely to the root side of the plate. If the reverse side of the plates is not accessible, 100% penetration cannot be guaranteed. The hole, which inevitably is created when welding is completed, was cut from the plates. In practical applications, it may be necessary to ®ll this hole using some other welding process. It can be seen that there is a large difference in the welding speed attainable with the two alloys. Alloy AA 6082 could be welded much more rapidly than AA 5083 for the same plate thickness, still achieving a sound weld.
2
3
Hardness testing (HV1) was conducted on cross-sections, along horizontal lines. The horizontal line was situated 1.65 mm (AA 5083) or 2.5 mm (AA 6082) above the root side, extending from one base plate, through the HAZ,
Alloy AA 5083 ± O had a relatively uniform grain size, slightly elongated in the rolling direction (Fig. 4). The grain sizes were on average 30 mm along and 20 mm across the rolling direction. Alloy AA 6082 ± T6 contained two sizes of grains, owing to partial recrystallisation. The recrystallised grains were y10 mm in size, while the non-recrystallised grains varied in size, typically between 50 and 150 mm (Fig. 4). Both base materials contained coarse inclusions, elongated in the rolling direction.
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Plate thickness, mm
Welding speed, cm min21
AA 5083
15 10 10 6 6
4.6 6.6 9.2 9.2 13.2
AA 6082
10 10 5 5
26.4 37.4 53 75
Alloy
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Figure 5 shows macrographs of three different sections of welds in the two alloys. Figure 5 shows the complex shape and the ring pattern in the nugget zone of friction stir welded AA 5083. It can be seen that the ring pattern is relatively well developed in the centre of the nugget zone, while the boundary between the nugget zone and the TMAZ is more clearly de®ned at one side of the weld (to the right in Fig. 5). This is because the tool travel and rotation directions coincide on this side. It can also be seen how the nugget zone becomes wider towards the top side of the plate. This is caused by the effect of the shoulder of the tool, which is much wider than the tool itself. When it touches the plate, the shoulder creates a stirring effect, which spreads down into the plate. On the top side, a small ridge can be
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Properties of friction stir welded Al alloys
289
areas inside rings. The rings were y20 ± 30 mm wide. No difference in grain size between ring areas and outside ring areas was found. Nor was any concentration or depletion of particles inside the rings observed. Using polarised light optical microscopy on anodised specimens, however, it was found that the grains inside the rings seemed to have less relative misorientation, compared to grains outside rings (Fig. 6). Experiments using backscattered electrons in the scanning electron microscope did not ®nd any signi®cant difference between areas inside and areas outside rings. The grains in the nugget zone had the typical equiaxed shape reported in other investigations. Thus, all traces of the base metal grain structure were effectively wiped out. Figure 7 shows the grain structure of the nugget zone in AA 6082; this structure is also representative of AA 5083. The grain size was y10 mm. The main microstructural effects in the TMAZ are shown in Fig. 8. Figure 8 illustrates how the grains were bent around the nugget zone as a result of the large deformation imposed by the stirring during welding. In Fig. 8, an SEM image shows how the grains in the TMAZ have been rotated, but not recrystallised. Also in Fig. 8 the ®ne grains in the nugget zone can be seen, illustrating the difference between recrystallised and non-recrystallised areas.
a
9
b AA 5083-O; AA 6082-T6, note large variation in grain size in this alloy and that dark patches in microstructure are caused by specimen preparation technique, not as result of any two phase structure of material
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a
b9-Mg1.7Si particles are quite small, with a length of y50 nm and a width of 2 ± 3 nm. In the nugget zone (Fig. 10) the dispersoids were clean (and no other small precipitates were seen either). To obtain full hardening of AA 6xxx type alloys, precipitation of GP zones or b0-Mg5Si6 should take place. Despite extensive examination, neither GP zones nor b0-Mg5Si6 could be seen in the TEM images. This was a common observation both for the nugget zone and the HAZ. The presence of these hardening precipitates would give hardness of over 100 HV1, as was found in the unaffected base material. The small b9-Mg1.7Si precipitates and the dispersoids give a much smaller contribution to strength and hardness.
%
The hardness of AA 5083 was y75 HV1 and was practically constant across the welds (Fig. 11). Horizontal hardness pro®les across welds in AA 6082 (Fig. 11) had a signi®cantly different appearance compared with AA 5083. The unaffected base material was harder (y110 HV1) and a decrease towards the weld was found. The minimum hardness (60 ± 65 HV1) was detected y8 mm from the weld centre, on each side of the weld, when the measurements were made along a line 2.5 mm from the bottom side of the weld. The hardness of the nugget zone itself was typically 70 ± 75 HV1. Microhardness measurements in rings and in between rings (Fig. 12) did not reveal any differences in hardness.
291
b optical micrograph showing major ¯ow around the nugget zone; scanning electron micrograph showing large scale rotation of grains in TMAZ
* 8 33 & *D4 E 3 33 &
To identify the position of lowest hardness, a number of low load (HV0.025) hardness measurements were made along lines at various distances from the plate surface. It was found that the minimum hardness for these measurements was 85 HV0.025. Figure 13 shows the isohardness curves, corresponding to 85 HV0.025, plotted on a sketch of the weld. It can be seen that the positions of the curves correspond well with the width of the tool shoulder at the top side of the weld. The isohardness curves then move closer to the nugget zone towards the bottom of the plates.
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33 & 3 C > 3
Tensile properties, measured transverse to the welding direction, are given in Table 3. The properties varied considerably between the two alloy types. For AA 5083 welds, the tensile strength varied between 303 and 344 MPa, depending on plate thickness and welding speed. Strength was lower in AA 6082, in the range 226 ± 254 MPa. It can be seen that there was considerable variation in strength within the same alloy type. However, there are no clear trends for the in¯uence of either plate thickness or welding speed. When the location of the fracture relative to the weld centre was also examined, an interesting pattern was found. For welds made in AA 5083, almost exclusively, fracture was found close to the centre of the weld. The fracture surface in general was inclined at about 45³. At the root side the fracture surface was very close to the original joint line. In a few bars, fracture started remote from the original joint Science and Technology of Welding and Joining
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Properties of friction stir welded Al alloys
a
a
b
b
nugget zone; zone of lowest hardness
9
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line. In these specimens the fracture surface was closer to the original joint line on the top side of the plate, but displaced between 3 and 7 mm from it. In AA 6082, fracture never occurred close to the original joint line. Instead, it occurred mostly close to the line where the shoulder of the tool had touched the top side of the weld. These fracture surfaces were also inclined, so at the bottom side of the weld the fracture surface was closer to original joint line, but still displaced y7 mm from it. In this case a few specimens also had a different appearance, with the fracture surface lying y5 ± 7 mm from the original joint
2
* 3 3 a & 7 b9 32 B ! b 33 &
line at the top side. The fracture surface was then inclined away from the centre of the weld, lying y15 mm away from the original joint line at the bottom of the plate.
;
Figure 14 shows the results of fatigue testing of the friction stir weldments and compares these with design curves from the European Convention for Structural Steelwork.19 In most specimens, fracture was initiated in the base material or in the centre of the weld. Only in a few specimens was
* 8 * Alloy
Plate thickness, mm
Welding speed, cm min21
Number of tests
Average tensile strength, MPa
AA 5083
15 10 10 6 6
4.6 6.6 9.2 9.2 13.2
6 6 6 6 6
318 344 331 312 303
AA 6082
10 10 5 5
26.4 37.4 53 75
4 6 4 4
226 236 254 254
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(a)
(b)
2
across welded joint in AA 5083 along line 1.65 mm from bottom side of plate, note almost constant hardness was found across whole welded joint; horizontally across friction stir weld in AA 6082, note hardness minima on either side of nugget
22
# F
fracture initiated in the weld metal ± base metal transition region. The fatigue stresses measured were in general well above the design curve requirement. There are some indications that lower speed resulted in a higher fatigue resistance of the weld, but this must be con®rmed by further testing. Welds in AA 5083 exhibited higher fatigue properties, compared with welds in AA 6082. Scatter was also larger for welds in AA 6082 than for welds in AA 5083.
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Results in the present investigation have shown that sound welds giving acceptable mechanical properties can be produced using FSW of the aluminium alloys AA 5083 and AA 6082. Discussion of the results will focus on three areas: (i) evolution of the microstructure in the nugget zone (ii) the relationship between microstructure and hardness/tensile properties of the welded joints (iii) the difference in weldability of the two alloys.
The division of the microstructure of a friction stir weld into a nugget zone, TMAZ, and HAZ is well established and the different zones can relatively easily be distinguished. The pattern found in the present investigation is well in accordance with previous observations. Development of the microstructure in the nugget zone has attracted most interest, particularly the ring structure which is so typical of the nugget zone.
5
It is generally agreed that the re®ned grain size of the nugget zone is a result of recrystallisation. As noted in the literature review, Liu 4 argued that in AA 6061 the grain size is determined by dynamic continuous recrystallisation. It is assumed that the same mechanism also operates in
! 3 8 3
AA 6082, since the compositional difference between these two alloys is relatively small. Studies of hot deformation of AA 5083 have shown that this alloy does not recrystallise by dynamic continuous recrystallisation but by geometric dynamic recrystallisation.20,21 According to Threadgill,3 the grain size in the nugget zone of many alloys is approximately equal to or less than 10 mm. Both alloys in the present investigation were found to conform approximately to this generalisation. Thus, even if the mechanisms of recrystallisation differ between the alloys, the end results are strikingly similar. The investigation showing a slight deviation from this pattern is the one by Mahoney ,6 in which two families of grain sizes appeared: one very ®ne (3 ± 5 mm) and one a little coarser (5 ± 10 mm). The differences between the results of Mahoney , the results summarised by Threadgill,3 and the results of the present investigation are certainly small and it is striking how similar the grain sizes of the nugget zones are, despite differences both in alloys and in process parameters. The general structure of the nugget zone is explained by dynamic recrystallisation, but the origin of the ring structure is still not explained. Observation of the ring structure in the three different sections (Fig. 5) shows that the rings in fact are three-dimensional surfaces curved behind the welding tool. The ring structure is closely related to the ®ns of the tool, the distance between them, and the weld travel and rotational speeds of the tool. The grain size variations noted by Mahoney 6 were not found in the present investigation. Observations made using polarised light optical microscopy (Fig. 6) indicate that the material inside the rings recrystallises in a different manner, forming grains with a smaller orientation difference and also, possibly, a more aligned crystal orientation.
8
Changes in particle size distribution produced by FSW were relatively small in both alloys (Table 2) for particles larger than 0.1 mm. The main difference between the alloys was a decreased number density of 1 ± 10 mm sized particles in AA 5083 and an increased number density in AA 6082. In AA 5083 fracturing is the most likely explanation for the decrease of 1 ± 10 mm sized particles and the increased number of smaller particles. Growth of very small precipitates might have contributed to the increase of 0.1 ± 1 mm particles. Coarsening of precipitates appears to be a more signi®cant effect in AA 6082, resulting in an increased number of both 0.1 ± 1 and 1 ± 10 mm sized particles. Science and Technology of Welding and Joining
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Properties of friction stir welded Al alloys
(a)
2
( 7 104
Growth of hardening precipitates is probable in the TMAZ, where the decrease in hardness and strength cannot be fully recovered by a post-weld aging treatment.9 It has been argued, however, that a signi®cant dissolution rather than growth of Mg2Si takes place in the nugget zone, owing to the higher temperatures in this zone.9 The increased number of the 0.1 ± 1 mm sized particles is, therefore, more likely to be explained in terms of fracturing of larger inclusions and growth of very small, nonhardening inclusions, similar to the behaviour of AA 5083. Behaviour of the very small particles studied by TEM is discussed below.
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107
108
109
105
106
107
108
109
(b)
G 3
Hardly any variation in hardness occurred across the weld in AA 5083. This is not surprising, since the base material was in an annealed state from the beginning. The main hardening mechanism for this alloy is deformation hardening, through formation of a high dislocation density. In the annealed state the dislocation content is low, as is also the case in the weld nugget. The effect of the variation in grain size on hardness is very small. Values observed are in good agreement with the results of Threadgill.3 The fracture position in the transverse tensile tests often coincided with the nugget area, indicating some weakness of this zone. There was a greater number density of small (0.1 ± 1 mm) particles in the nugget zone, although the number density of larger particles was lower. Both these particle sizes can affect the ®brous fracture, but since the smaller ones are more numerous, it is likely that they have the larger effect, which could explain the fact that fracture predominantly was located in this zone. In AA 6082 a marked minimum in hardness was noted 7 ± 8 mm from the weld centreline when hardness was measured along a line 2.5 mm from the root face. The isohardness curve for minimum hardness is then inclined at y45³, as shown in Fig. 13. The fracture position in most of the tensile specimens coincided well with the minimum hardness position (which is as expected). It should be noted that the strength of friction stir welded AA 6082 is actually lower than the strength of welded AA 5083, despite the fact that AA 6082 is much stronger in the unaffected base material. Hardness and tensile properties could be increased, however, although not fully recovered, by an aging treatment. It is of interest to understand the microstructural effect giving rise to this low hardness and tensile strength. However, to be able to interpret the results of the TEM investigation, approximate values, at least, of the peak temperatures at various places in the HAZ should be obtained. From Fig. 3 it can be deduced that the weld travel speed has a large in¯uence on the peak temperature distances from the nugget boundary. In the present investigation, the weld travel speed was fairly close to the travel speed used by Backlund 9 It is therefore estimated that the peak temperature at the position of the
105
104
in AA 5083; in AA 6082
2
