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Effects of forge axis force and backing plate thermal diffusivity on FSW of AA6056
Materials Science & Engineering A 558 (2012) 394–402
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Effects of forge axis force and backing plate thermal diffusivity on FSW of AA6056 P. Upadhyay n, A.P. Reynolds Department of Mechanical Engineering, University of South Carolina, 300 Main Street, Columbia, SC 29208, United States
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 May 2012 Received in revised form 3 August 2012 Accepted 6 August 2012 Available online 10 August 2012
Using three sets of tool rotation rate and welding speed, series of friction stir welds were made at various forge axis forces on different backing plates with widely varying thermal diffusivity on 4.2 mm thick AA6056. Temperature during the process was measured at the probe core using a thermocouple. Because of the use of different backing plates and forge forces it was possible to obtain a relatively wide range of peak temperature in the nugget. This provided a unique opportunity to investigate changes in joint microstructure and property (1) over a wide range of peak temperature while keeping welding and rotational speed constant (2) at similar peak temperature with different welding and rotational speeds. The results show that for the studied gauge thickness, metallurgically significant temperature variations can be achieved without changing the rotation and welding speed while producing defect free welds. The backing plate conductivity and forge force both independently affect the weld process parameters such as nugget temperature and tool torque significantly. A significant increase in the tool torque and hence power is observed with the increase in the forge force alone. & 2012 Elsevier B.V. All rights reserved.
Keywords: Friction stir welding Forge force Backing plate Thermal boundary conditions AA6056 Temperature measurements Torque
1. Introduction and background It has been established that the tool rotation rate primarily controls the peak temperature achieved at the deformation zone during friction stir welding [1–3]. For a series of welds the measured average torque and power has also been inversely correlated to the peak stir zone temperature and hence to the rotation rate. Fig. 1 shows the correlations of peak temperature measured at the tool probe (Probe T) and measured tool torque for welds made with 6.4 mm thick AA7050 in laboratory air and under-water [4]. Similar trends of peak temperature with change in rpm have been reported elsewhere in the literature [1,2]. One concern with this reasoning is that, in such series of welds it is not only the tool rpm that is varied. Assuming a force control scheme, during a series of welds with varying tool rotation rate, the FSW practitioner will invariably have to adjust the applied forge force as well in order to produce good quality welds with minimal flash. A ‘‘good’’ forge force for a specific welding condition will thus be established in due course of trial and error. ‘‘Good’’ here can be loosely defined as weld with minimal flash and no internal defects. Note that there will always be a small window of forge force that will produce ‘‘good welds’’. The size of the window
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Corresponding author. Tel.: þ1 803 777 3118; fax: þ1 803 777 0106. E-mail addresses: [email protected], [email protected] (P. Upadhyay). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.018
depends on factors including the alloy welded and the tool geometry. For an empirically established ‘‘good’’ forge force an increase in rpm will lead to increase in the average deformation zone temperature rendering the material more plastic. Without any forge force adjustments some of the plasticized material will expel away from under the shoulder instead of being consolidated behind the tool, resulting in excessive flash and, possibly, volumetric defects. For a given set of welding parameters (rpm and welding speed), plate geometry, alloy type and tool design the ‘‘good’’ forge force may change with changes in thermal boundary conditions (BC) that are present at the work piece. The rate of heat flux through the bottom of the work piece mostly depends on the backing plate (BP) thermal diffusivity and hence, as has been pointed out in literature [5–8], the most important thermal BC for FSW is the thermal condition at the backing plate. It was hence thought necessary to understand effects of forge force in conjunction with varying levels of thermal diffusivity of the backing plate. In the published literature the effects of forge force have not received as much attention as tool rotation and welding speed effects. Some qualitative study of general effects of either forge force or plunge depth can be found scattered in the literature. During a parametric study using position control Johnson [9] showed weld forge force increased with the increase in the plunge depth. Nishihara and Nagasaka [10] reported that tool and backing plate temperature increased when the plunge depth
P. Upadhyay, A.P. Reynolds / Materials Science & Engineering A 558 (2012) 394–402
Fig. 1. Measured peak probe temperature and torque plotted against the rotational speed for in air (IA) and underwater (UW) welds made in 6.35 mm thick AA7050. Adapted from [4].
was increased. Elagovan et al. [11] studied the effect of tool pin profiles and forge force on FSW of AA6061. Out of the three forge forces considered the intermediate one was reported to show superior mechanical properties. Since no temperature measurements were reported the effect of thermal exposure on mechanical property could not be elucidated. Excessive forge force probably led to higher heat input resulting in high flash and thinning of the weld zone. On the other hand a less than sufficient forge force might have led to insufficient heat input and hence inadequate material flow. At about same time Record et al. [12] reported a design of experiment based approach where nine welding parameters were used. They reported that welding speed and location of weld relative to the edge of the plate also had some minor effect on forge force in a position control mode. Effects of thermal boundary condition changes at the bottom of the work piece have also not been thoroughly examined in the open literature. As a part of quench sensitivity study Nelson et al. [3] reported that the use of a heated backing plate resulted in higher peak temperature and lower cooling rate, resulting in inferior mechanical properties in AA7075 welds. Rosales et al. [13] reported welds made with AA2024 and AA6013 where steel, copper and ceramic coated backing plates were used at three different combinations of rotational and welding speeds while the forge force was kept constant. Stir zone temperatures were not measured while the in-plate far field temperature measurements varied significantly when using different backing plates indicating the significance of backing plate conductivity. Macrographs of the welds indicate thinning of the cross-section in ceramic coated backing plate, probably because of high flash content. As the application of friction stir welding widens it will be important to understand the behavior of process variables like stir zone temperature, torque and in plane forces with the change in thermal boundary conditions. This understanding will be helpful in a variety of scenarios. In modeling for instance, contact conductance between backing plate and work piece is still elusive but has been deemed central in quantifying the heat flux out from the bottom of the work piece. Value of this contact conductance may depend on applied forge force, surface temperature and backing plate thermal conductivity. A look at the literature shows that a wide range of contact conductance values have been used by different authors ranging from 0.2 to 105 W/m2 K [4,6–8]. Whether or not the contact conductance could vary such a great deal can be debated. Although not addressed in this paper, it is fitting to call for a need to establish some physically valid upper and lower bound at least in the order of magnitude sense. The experimental results discussed in this paper including peak
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temperature and power correlations with welds performed using various backing plates in conjunction with parametric finite element modeling may provide useful insight into the problem. A number of studies have been conducted to understand the structure-property-processing relationship in 6XXX series alloys [14–19]. A fine distribution of the metastable needle shaped b00 (Mg5 Si6) precipitate is considered to be responsible for peak aged condition. In addition to the well-known b00 phase, a lath shaped hexagonal precursor phase to Q’ phase has also been attributed to strengthening. During the welding process the optimized microstructure of the peak aged base material is altered at various levels because of the high temperature cycle. This cycle will produce some level of dissolution of precipitates, diffusion of solute, and increase in vacancy concentration. Thus some non-strengthening phases form and some strengthening precipitates dissolve or coarsen, resulting in varying degrees of loss of strength. The purpose of most welding research work is to minimize this effect while keeping an acceptable rate of production. To this end, qualitative assessment of weld cross-section normally begins with micro hardness tests that provide a first glimpse into changes that the material has gone through during the welding. Welds made on peak or overaged precipitation hardening alloys, such that the nugget approaches a temperature close to Solution Heat Treatment (SHT) temperature, typically result in a characteristic ‘‘W’’ shaped hardness profile after some artificial or natural aging [2,20,21]. The nugget which reaches such a peak temperature would be in a condition similar to that obtained after SHT and quenching; therefore, re-precipitation of strengthening phases may occur during post weld cooling and subsequent natural and artificial aging. It has been found when the nugget temperature is beyond the overaging regime (discussed subsequently), the nugget hardness increases with the increase in the stir zone temperature (achieved through increase in rpm) until reaching a plateau close to SHT [1,2]. This has been attributed to greater availability of solute in the matrix for reprecipitation into strengthening phases during subsequent aging at higher peak temperature than at lower peak temperature. On either side of the nugget there is a quick drop in hardness continuously falling across the thermo-mechanically affected zone (TMAZ), reaching global minima in the heat affected zone. This arises due to over aging of the heat affected zone (HAZ) [1,22,23]. If the weld is performed at relatively low power, such that the stir zone peak temperature is around 350 1C, the characteristic W shape in the hardness profile is replaced by a U shaped one. With a peak weld temperature near 350 1C, the hardness in the HAZ and the nugget will be similar to each other and less than peak aged base metal [2,4]. For precipitation hardening alloy it has been found that the peak temperature exposure of 350 1C is most damaging to strength [24]. Upon tensile testing of FSW samples of AA7075, Mahoney et al. [25] found that failure occurred along the heat affected zone where the peak T recorded was between 300 1C–350 1C. Hwang and Chou [26] and Sato et al. [1] simulated thermal cycles at heat affected zone using several peak temperatures in AA7075 and AA6063. Results from both works indicated that minimum hardening effects occurred on samples that were exposed to a peak T of 350 1C. Subsequent TEM/SAD analysis of the overaged nugget and/or HAZ region from several works [1,21,25,26] show coarsening and dissolution of strengthening precipitates and presence of non-strengthening b0 precipitates in 6XXX series alloys. Although the sequence of precipitation is more complicated than what has been presented here, for the purpose of this work the following simple statements can be made. The kinetics of precipitate coarsening are maximum near 350 1C for 6XXX alloys. In other words the rate of formation of non-strengthening b0 in 6XXX series is at peak near this temperature. During its formation, b0 particles take away a significant amount of solute
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from the matrix which otherwise would have been available for reprecipitation of strengthening b00 phase during the post weld aging. In the regions of the weld where the peak T is around 350 1C (mostly in the HAZ and when weld power is relatively small, in the nugget as well), the situation will be aggravated when the welding speed is lower. Lower welding speed will cause the material to remain in the b0 phase formation temperature range for a relatively longer time causing the solute depletion to increase, resulting in lesser strength. At peak temperature level farther away from 350 1C, approaching the SHT temperature, more and more precipitate will dissolve into the matrix avoiding precipitate coarsening such that during subsequent heat treatment fine strengthening precipitates are formed resulting in a stronger nugget zone.
the nugget center using the Mean Linear Intercept (MLI) method. Three views at a magnification of 200 were examined. Grain boundary intersections were counted on a test line of length 0.5 mm (100 mm at 200 ). Five test lines per view were used. Micro-hardness tests were performed using a Vickers hardness indenter on naturally aged and post weld aged samples. Post weld heat treatment of samples was carried out at 190 1C for 4 h. This time and temperature combination is consistent with the industrial practice leading to the T6 or peak aged temper. A Vickers hardness indenter, with a load of 500 g and a load application time of 10 s was used for measurement of hardness as a function of distance from the weld centerline on transverse cross sections along the plate. For all the hardness traverses plots reported in this work the advancing side is on the left.
2. Experimental procedure
3. Results and discussions
Bead on plate welds were made on 4.2 mm thick rolled plates of the aluminum alloy, 6056-T451. The temper designation T451 indicates that the alloy has gone through solution heat treatment, quenching and subsequent stretching to relieve internal stresses [27]. The solution treatment temperature (SHT) for the alloy is 529 1C. Welds were produced on a hydraulically powered MTS FSW Process Development System (PDS) using Z-axis (forge) force control. During the welding process the PDS records a number of relevant process variables as a function of time. The process data recorded includes variables such as traverse and rotational speeds of the tool; longitudinal, transverse and vertical force (X,Y and Z respectively) experienced by the tool and tool plunge depth. Real time torque required by the spindle motor is also measured using a torque transducer. The tool used for production of all welds was of a two piece design with a 17.8 mm diameter, single scroll, H13 tool steel shoulder and a probe fabricated out of MP-159 (a high temperature cobalt based super alloy) in the shape of a truncated cone (81 taper) with threads and three flats. The probe was 4.1 mm long, with a diameter of 6.5 mm at the intersection with the shoulder. Temperature during welding was monitored and recorded using a K-type thermocouple spot welded into the probe on the axis of rotation approximately at the probe mid-plane height. Series of welds were made with three sets of rotational and translational speeds using various forge forces and backing plates with widely varying thermal diffusivity. Three sets of welding and rotational speeds from 320 rpm–3.4 mm s 1, 640 rpm– 6.8 mm s 1 and 960 rpm–10.2 mm s 1 were used. Welding speed and rotational speed were chosen such that the advance per revolution was 0.635 mm in all cases. Thirteen millimeter thick backing plates (BP) made of material with widely varying thermal diffusivity viz. AA2099, tool steel, Al6XN (an Allegheny Ludlum super austenitic stainless steel), Ti-6-4, and ceramic floor tile were used. The corresponding forge forces used can be inferred from Figs. 3 and 4. A forge force of 12.8 kN deemed as ‘‘good’’ for the steel BP with welds performed at 640 rpm and 6.8 mm s 1 was used as a baseline. Levels of forge force changes of 710%, 20%, 30% were then used to perform the set of experiments for different backing plates and the three sets of rotational and translation speed. Standard metallographic preparation technique was employed to grind and polish samples to satisfactory level before performing characterization by optical microscopy and hardness testing. Samples were ground using automatic and/or manual grinding machines with 180, 240, 320, 400, 600, 800 grit silicon carbide paper. They were then polished using Aluminum oxide powder of 5 mm and 3 mm followed by colloidal silica ( o0.05 mm). Samples were etched using Keller’s reagent. Grain size was measured at
3.1. Peak temperature correlations Representative temperature transients recorded by thermocouples located at the probe core are shown in Fig. 2 for welds made at 640 rpm–6.8 mm s 1 with aluminum backing plate. For all the forge forces used the probe T reaches a fairly steady value after half the weld length. Welds made with other backing plates also exhibit qualitatively similar temperature trends. The temperature measured inside the probe can be thought of as some average of the temperature of material in contact with the probe and may not be reflective of the maximum temperature at the nugget. However this measured temperature has been found to be repeatable and provides reasonable representation of nugget temperature with changes in welding parameters [4,28]. Fig. 3 shows the measured steady state peak probe temperature achieved at several levels of forge forces using four different backing plate materials for the 640 rpm–6.8 mm s 1 welds. These data were extracted from graphs like Fig. 2. Excluding ‘‘bad’’ welds indicated by cross marks viz: three welds made at relatively low forge forces that resulted in surface defects and three welds at relatively high forge force that resulted in high level of flash and thinning; the rest of the runs produced good quality welds with little to no flash and no defects. This figure illustrates substantial changes in process response that can be brought about by changes in either forge force, backing plate diffusivity or both while keeping rpm and welding speed
Fig. 2. Temperature transients measured by thermocouple at the probe core for welds made at 640 rpm–6.8 mm s 1 with aluminum backing plate at different forge forces as indicated in the legend.
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Fig. 3. Peak probe T for welds made at 6.8 mm s 1 using various indicated backing plate materials. Crosses indicate welds with defects or excessive flash.
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speed than lower rpm/lower welding speed. This indicates that increase in temperature caused by increase in the tool rpm (and higher weld power) overwhelms the corresponding temperature reduction effect of increased welding speed. If the lines of data from different backing plates and same parameter set are approximated as parallel this shows that the effect of backing plate diffusivity on achieved peak probe T is greater at lower welding and/or rotational speed and diminishes as they increase. Demonstratively at 3.4 mm s 1 the change from Al BP to Ti BP at a forge force level of 14.2 kN can cause a 90 1C change while for 10.2 mm s 1 the change in temperature is no greater than 40 1C. This suggests that the ability to alter nugget zone temperature by changing the backing plate is decreased with increased rotational and welding speed. The individual effect of welding speed and rotational speed are difficult to gather from above data set because the welds were performed at the same advance per revolution. With higher rotation rate the nugget temperature approaches the solidus T causing the material flow stress to decrease. This decrease will in turn limit power generation by plastic deformation and hence temperature increase. Thus, rate of temperature increase due to addition of power will be higher at lower temperature. On the other hand at higher welding speed the viscous diffusion rate due to higher rpm and welding speed dominates the conduction heat loss (thermal diffusion). The only way to settle this is to conduct some additional experiments with either one of the two parameters constant. Since welding and rotational speeds were constant the shear and shear rate around the probe can be assumed to be nearly constant, these results provides a unique opportunity to investigate microstructure where shear and shear rate are constant while the Peak T is significantly varying. This is predicated on the assumption that, for a given welding speed and rpm, the kinematics of the flow does not vary greatly. 3.2. Power/torque correlations
Fig. 4. Peak probe T for welds made at 3.4 mm s 1 (dashed lines) and 10.2 mm s 1 (dotted lines) using indicated backing plate materials.
constant. For instance considering only good quality welds with no defect and minimal flash, combined variation of forge force and backing plate diffusivity can result in as much as 70 1C difference in probe temperature. Change in forge force or backing plate alone can result in 50 1C difference in probe temperature (As shown by dotted vertical and horizontal lines). At forge axis force of 11.3 kN, for instance, a ceramic tile BP weld can achieve a peak probe T of 460 1C, while under equivalent rotational and translation rate steel BP achieves only 431 1C. Aluminum BP likewise attains 423 1C. Expectedly the peak probe temperatures line up in accordance with the backing plate diffusivity and the applied forge force. Similar results are observed from other parameter set combinations (see Fig. 4). This trend of increase in stir zone temperature by changing only the forge force or thermal boundary at the back plate is qualitatively similar to the trend seen in Fig. 1 obtained from series of experiments where the tool rotation rates were varied. The peak probe T extracted from weld parameter set from 320 rpm–3.4 mm s 1 (dashed lines) and 960 rpm–10.2 mm s 1 (dotted lines) are plotted in Fig. 4 indicated by arrows are the differences between peak probe temperatures achieved between two extreme cases of backing plates for two extreme cases of welding speeds. Note that the peak temperature obtained for equivalent forge force is higher with higher rpm/higher welding
Similar substantial change in required torque and hence power resulted from forge force and backing plate variation albeit the relationships are more complicated than temperature relationship described above. Weld power calculated from the measured torque is plotted against the applied forge force in Fig. 5 for the three parameter sets. Consider first the 320 rpm–3.4 mm s 1 set where a reasonably linear relationship exists between weld power and forge force. Data sets with higher diffusivity backing
Fig. 5. Weld power calculated from measured torque plotted against the applied forge force for various backing plates shown in different symbols. Different welding speeds are shown in different line styles as indicated.
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plate exhibit a greater change in power with the change in forge force as indicated by difference in slope of the lines. For instance a forge force increase from 14.2 kN to 18.5 kN produced an increase in power from 1.5 kW to 2.5 kW in Al BP. An equivalent forge force change in Ti-6-4 BP from 11.4 kN to 15.6 kN only changes the power from 1.5 kW to 1.75 kW. With higher welding and rotational speed sets the power response with change in forge force is reduced further. In the case of Ti backing plate at 960 rpm–10.2 mm s 1, there is virtually no change in power despite a 50% increase in forge force. The same is true for steel BP at 960 rpm–10.2 mm s 1 except at the lowest forge force. The aluminum backing plate on the other hand shows substantial increase in the torque similar to that seen in the lower rpm welding. Note that welds made with the aluminum backing plate do not reach to a torque plateau for tested forge force levels; however, it can be anticipated that torque would level off if higher forge forces were used shown by continuation of lines beyond data points. Further increase in forge force on Al BP was not considered to avoid production of a lap weld between the work piece and the backing plate. Similar relationship is shown for 640 rpm–6.4 mm s 1 set. In this case gradual increase in the torque with the increase in forge force is seen until power plateaus for all but Al BP. Note that for a given parameter set the value of power at the plateau increases with the increasing diffusivity of backing plate used. This can be explained by observing the probe T vs. forge force graph in Fig. 3. Owing to greater heat retention, relatively high probe temperature is consistently observed for lower diffusivity backing plate indicative of higher shoulder interface temperature. This high temperature in turn leads to greater softening of the material hence material deforms at lower shear stress leading to lower required torque. 3.3. Temperature vs power trends In Fig. 6a–c measured probe T for three sets of welding parameters are plotted against the weld power for different backing plates. Again note that ceramic tile backing plate was used rather than Ti-6-4 for 6.8 mm s 1 case as shown in the legend in Fig. 6c. For the Al BP, the relationship between power and probe temperature is decidedly linear regardless of the tool rotation and translation rate. For all of the Al BP welds, the temperature increases with increasing power and by inference from this data, the power increases with increasing forge force regardless of the welding speed and rotational speed used. Considering Fig. 6b and c the relationship between probe T and power is not so straightforward with lower diffusivity backing plates. In the case of steel BP, the 320 rpm–3.4 mm s 1 data and some of the 640 rpm–6.8 mm s 1 data exhibit relatively linear relationships with power, although the relationship is not unique:
for a given power level, the 320 rpm–3.4 mm s 1 welds exhibit higher peak T than the 640 rpm–6.8 mm s 1 welds. This may be because the slower welds take longer to complete and the probe temperature increases with weld time; therefore, the peak T, measured at the end of the welds, varies with weld time and actual weld zone T. Even after saturation in the torque/power required is reached at higher forge force in the case of low conductivity backing plate viz. Steel and Ti-6-4 as discussed in previous section, the measured probe T continues to rise. For example the maximum power for the 640 rpm–6.8 mm s 1 welds is approximately 3 kW and four welds (made with the four highest values of forge force ) all have the same power level but different probe T. All of the 960 rpm–10.2 mm s 1 welds have approximately same power levels but the measured temperatures vary by as much as 40 1C. Similar behavior is observed for the lowest diffusivity anvils (Ti-64/Tile). 3.4. Nugget grain size and microhardness results and correlations The previous sections sufficiently demonstrate our ability to affect process response variables such as required torque and weld zone temperature as well as, potentially, enabling production of good quality welds with lower forge force by modifying of thermal boundaries however, of equal importance are the ability to use the thermal boundary condition to influence weld microstructure and properties. In the following section the effects of forge force and backing plate diffusivity on hardness distributions and nugget grain size are discussed and correlated. In Fig. 7 the grain sizes measured at the nugget center using MLI method are plotted against the measured peak probe temperature. Data sets for two different welding speeds are shown by closed (6.8 mm s 1) and open (10.2 mm s 1) symbols. Welds made with different backing plates are shown using differently shaped symbols as indicated in the legend. There is almost a linear trend of increase in the grain size with the increase in the measured probe temperature. This ‘‘linear’’ relationship is true regardless of the boundary condition at the bottom of the work piece and the welding speed. The apparent linear relationship across two welding speeds suggests that for the given case grain size is relatively insensitive to the time at temperature during the welding and is governed mainly by the peak temperature. This is consistent with the fact that metallurgical processes such as grain growth are exponentially dependent on temperature and exhibit much weaker dependence on time. Shown in the Fig. 8a are the transverse micro hardness profiles for naturally aged samples taken from welds made with a ceramic tile backing plate at three different forge forces for parameter set of 640 rpm–6.8 mm s 1. The respective forge force used and measured midplane peak T achieved during the welding is
Fig. 6. Peak probe T plotted against the weld power for (a) Aluminum, (b) Steel, and (c) Ti backing plates. Different symbols indicate different welding speeds.
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indicated in the legend. This set of graphs illustrates changes in mechanical property that occur in weld cross section with the change in peak temperature due to the change in forge force alone. There is a significant change in the shape as well as the absolute value of the hardness distribution in the nugget with changes in the applied forge force. This change can be explained on the basis of the measured peak probe T during the welding. As discussed in the introduction section, the average nugget
Fig. 7. Midplane nugget grain size plotted against measured peak probe temperature for two different welding speeds.
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hardness increases with the increase in the probe T. For instance with a low measured peak T of 398 1C obtained at a low forge force of 5.7 kN, the weld nugget is in over-aged condition exhibiting a U shaped hardness distribution. As the forge force is increased the nugget hardness increases while the heat affected zones on both the sides remain over-aged. Notice the nugget hardness asymmetry in welds made at intermediate forge force. Several such cases have been observed in the current work and they will be dealt with in a separate section. With sufficiently high forge force of 12.8 kN, where the measured peak T was recorded to be 490 1C nugget hardness approaches base metal value and asymmetry vanishes. Notice also that minimum hardness values for all the considered samples located at HAZ for W shaped distributions and all over the nugget for U shaped are in the narrow range of 70–74 HV. This is reasonable and expected since thermal boundary conditions and welding speeds for all the cases were identical such that time at elevated temperature for precipitate coarsening and dissolution has not changed substantially with the change in forge force. Fig. 8b, c and d show similar hardness results for welds made at 3.4, 6.8 and 10.2 mm s 1 respectively using high conductivity aluminum as backing plate material in contrast to low conductivity ceramic tile in Fig. 8a. Corresponding forge forces used and peak T attained at the tool midplane are indicated in the legend. In the case of welds made at 3.4 mm s 1 (Fig. 8b) despite a temperature increase from 392 1C to 431 1C there is very small increase in the nugget hardness. All the three profiles appear to be in an overaged condition as will be discussed subsequently. In the case of welds made at higher welding speed of 10.2 mm s 1 and higher rpm, (Fig. 8d), all the three forge forces resulted in W
Fig. 8. Vickers hardness profiles on weld transverse cross-section at midplane for naturally aged samples (a) ceramic floor tile BP at 6.8 mm s 1 (b) aluminum BP at 3.2 mm s 1 (c) aluminum BP at 6.8 mm s 1 (d) alum BP at 10.2 mm s 1. Note that advancing side is on the left for all the hardness traverses.
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shaped hardness distributions such that with higher applied forge force the nugget hardness is higher and approaches the base metal hardness. The effect of achieved peak temperature on the nugget hardness, all other things being equal, can be appreciated by comparing the hardness distribution among Al (Fig. 8c) and ceramic tile BP (Fig. 8a) at parameter set of 640 rpm–6.8 mm s 1. With same forge force of 8.5 kN for instance weld made with Al BP consists of weld cross-section that is overaged hardness while with Tile BP the weld nugget has achieved some degree of hardness recovery. The only difference between these two cases is that the Tile BP attained peak temperature that is 50 1C higher than Al BP. This relationship between measured peak temperature and nugget hardness is further explored in Fig. 9 where the arithmetic average of edge to edge midplane nugget hardness are plotted against the measured peak mid-plane probe T. Note that welds made at same welding speed are represented by same symbol regardless of the backing plate used. Whereas the closed symbols indicate data from naturally aged samples, the open symbols represent data obtained from samples that were heat treated at 190 1C for 4 h. Note that not all data shown in naturally aged (NA) condition were tested after heat treatment (HT) so some NA samples do not have corresponding HT data points. Note also that relationship between peak probe temperature and nugget hardness are different amount heat treated and as welded samples. For post weld heat treated samples, the hardness value plateaus at higher peak probe temperature range, whereas for naturally aged case no such plateau is observed. Presumably at a high temperature range, in post weld heat treated samples the additional solute obtained because of further increase in the probe T is not significant to noticeably increase the hardness. In Fig. 10 the difference in average nugget hardness between heat treated and naturally aged samples are plotted against the measured peak T. First consider welds made at 960 rpm–10.2 mm s 1and 640 rpm–6.8 mm s 1. Excluding boxed data points at two lowest peak T cases for 6.8 mm s 1(they will be dealt with shortly), the nugget hardness generally increases with the increase in peak probe temperature for both as welded and PWHT cases. A hardness recovery of 15–25 HV is observed for all these cases (Fig. 10). Now consider boxed data points in Figs. 9 and 10 (most at 3.4 mm s 1and two at 6.8 mm s 1) where there is no apparent change in hardness despite a change in peak probe T from 360 to 430 1C. The hardness distribution for all these data points
Fig. 10. Change in average nugget hardness after heat treatment plotted against the corresponding peak T. Boxed points belong to overaged nugget that are unresponsive to heat treatment.
resembles the ones shown in Fig. 8b (U shaped). Note that this also happens to be the regime where no significant hardness recovery after PWHT takes place as evident from o5 HV recovery after heat treatment (Fig. 10). In other words they are all overaged. These are welds exposed to a relatively lower peak temperature and are at temperature regime where precipitation kinetics for coarsening of strengthening precipitates is at the peak as noted in the background section. Note that beyond certain temperature levels, welds made at both the welding speeds (3.4 mm s 1 and 6.8 mm s 1) start showing significant hardness recovery as well as sensitivity to peak probe T. It is interesting to note that this temperature beyond which significant hardness recovery after PWHT ( 415 HV) is observed appears to be correlated with the welding speed (thus time at temperature). At 3.4 mm s 1 for instance this ‘‘transition’’ temperature is 440 1C while for 6.8 mm s 1 the transition occurs at significantly lower temperature of 390 1C as indicated by the corresponding jump in the data point from the box shown by arrows. Alternatively this means that transition from U shaped (overaged throughout the nugget) to W shaped (significant hardness recovery at the nugget with overaged at HAZ) hardness distribution occurred at significantly lower temperature at higher welding speed. Note that for welds made at the highest welding speed of 10.2 mm s 1no overaged nugget hardness was observed at equivalent low temperature and all the hardness measured were of W shaped. The authors posit that significant hardness recovery at this welding speed might be anticipated at still lower temperature. This shows that time at temperature has a significant effect on precipitation kinetics in addition to the peak T. Alternatively this suggests that temperature considered most damaging for the strength of weld nugget depends on the welding speed and should not be thought as a fixed temperature regime. There may also be an effect of strain which is not accounted for in a purely temperature based analyses. Over a range of welding speed of 0.8–3.8 mm s 1 Reynolds et al. [2] presented similar trends of PWHT hardness recovery with increase in the welding speed in AA7050. The peak T obtained from an input torque model was correlated with PWHT hardness recovery showing that a temperature range of 330– 370 1C rendered the welds overaged and unresponsive to PWHT. 3.5. On hardness asymmetry
Fig. 9. Average nugget hardness plotted against peak T for 3 welding speeds (shown in different symbols). Closed symbols are data from naturally aged samples. Open symbols are for heat treated samples.
In few cases the measured hardness value did not stay constant throughout the nugget. As seen in some hardness profiles with
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intermediate peak temperature range, in Fig. 8a, c and d the hardness is highest at the advancing side edge of the nugget and gradually decreases from advancing to retreating side eventually reaching minimum at heat affected zone on the retreating side. To investigate the effect of peak deformation zone temperature on the degree of asymmetry, the difference between hardness measured at advancing and retreating side of the nugget (DVHNasymmetry) is plotted against the corresponding measured peak probe T for two welding speeds of 6.8 and 10.2 mm s 1 in Fig. 11. The data set from welding speed of 3.4 mm s 1 is not shown in Fig. 11 since no significant nugget asymmetry (the highest is 4.5 HV) is observed. For 6.8 mm s 1 after crossing the overaged regime (all data set with Peak T4400 1C), the nugget asymmetry increases with the increase in the probe T reaching a maximum of DHV¼18.5 at peak T of 450 1C.With further increase in the probe T the asymmetry vanishes, indicating flat nugget hardness distribution. In the case of 10.2 mm s 1 maximum asymmetry is observed at the peak T of 430 1C (lowest peak T for 10.2 mm s 1 case) after which the asymmetry decreases significantly indicating that the nugget hardness becomes flat beyond the peak T of 470 1C. Similar asymmetries in nugget hardness distributions have been reported by Wade and Reynolds [29] on welds made at 720 rpm and 10.2 mm s 1 in AA6019. Equivalent weld made with identical welding speed at a higher rpm of 1680 on the other hand resulted in a reasonably flat nugget hardness without asymmetry. This was observed for both as welded and PWHT samples. It has been proposed and argued in the literature [30–32] using simulation results, some analytical treatment and inherent asymmetrical flow pattern in the process, that advancing side experiences higher temperature and strain rate in comparison to retreating side. Using grain size as a ‘‘proxy’’ for peak T attained Wade and Reynolds [29] showed indirect experimental evidence of temperature asymmetry in the nugget for the first time. Having a wider window of peak T range over three rotational speeds Fig. 11 provides ample opportunity to expand the argument. As explained in the introduction section, away from the overaging regime ( 350 1C), incrementally higher peak T results in higher solute being available for reprecipitation during subsequent aging resulting in incrementally greater strength. This strength/peak T relationship is linear only up to a certain peak T beyond which the incremental amount of additional solute obtained by further peak T increase is not significant to noticeably increase the hardness. In other words saturation in the precipitation of strengthening phases is reached beyond a critical peak T in the vicinity of SHT. Until that saturation is reached the region towards the advancing side
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Fig. 12. HAZ minimum hardness of naturally aged samples vs. welding speed. Different symbols indicate different backing plates used.
being exposed to greater peak T will be stronger than the retreating side resulting in asymmetric hardness distribution. This explanation is consistent with the observation in Fig. 11 where asymmetry disappears when the hardness reaches peak strength.
3.6. HAZ minimum hardness While hardness measurements in relation with peak probe T is interesting, it is the minimum hardness that typically occurs at the HAZ which determines the strength of the joint. The HAZ minimum hardness measured from the naturally aged samples are plotted against the welding speed in Fig. 12. Different symbols indicate welds made under different backing plates. There is a familiar trend of increase in minimum hardness with the increase in the welding speed which has been observed in the past in several precipitation hardening alloys [4]. Unfortunately the minimum hardness does not show any significant sensitivity to the backing plate material. Spread within hardness data among same welding speed is too narrow to distinguish between the various cases. Hence it can be asserted that for this gauge and temper, welding speed has greater effect on minimum HAZ hardness than thermal BC at the bottom of the plate (backing plate material). Accordingly this suggests that the welding speed is primary determinant of effective time at elevated temperature.
4. Conclusions
Fig. 11. Difference between hardness measured at advancing and retreating edge of the nugget (DVHNasymmetry) plotted against the corresponding measured peak probe T for welding speeds of 6.8 and 10.2 mm s 1.
(1) Forge force has a significant effect on the stir zone temperature of otherwise similar welds. Defect free welds with metallurgically significant temperature variations can be achieved by changing the forge force alone within controlled limits. (2) The backing plate diffusivity has a very important role in the peak stir zone temperature achieved. For the considered plate gauge, thermal and microstructural through thickness homogeneity can be achieved by the use of low diffusivity backing plate material. (3) Similar peak temperature at nugget can be achieved by using a wide range of forge force with the use of different BPs. Same applied forge force can yield a wide range of peak temperature if different backing plates are used. (4) The tool torque increases with the increase in forge force until reaching a plateau at a high forge force level. With the use of
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low diffusivity backing plate, tool torque required for otherwise equivalent welding conditions is reduced. (5) For welds made at two different welding speeds, nugget grain size has a nearly linear relationship with the measured peak probe temperature regardless of backing plates used. (6) The measured peak temperature beyond which significant recovery in the nugget hardness after PWHT occurs depends on the welding speed. In the case studied, while for 6.8 mm s 1, the overaged nugget insensitive to PWHT was avoided beyond the peak T of 390 1C, for 3.4 mm s 1 a peak temperature of 440 1C was necessary to get similar effect. (7) While the use of different backing plates made a significant effect on peak temperature at the deformation zone, the rate of cooling was not affected in significant ways as evident from the minimum hardness correlations.
Acknowledgments The authors acknowledge the financial support of the Center for Friction Stir Processing which is a National Science Foundation I/UCRC supported by Grant No. EEC-0437341. The authors thank Dr. Wei Tang and Daniel Wilhelm, Department of Mechanical Engineering, University of South Carolina, Columbia, SC, USA for their help in preparing the weld joints. References [1] Y.S. Sato, H. Kokawa, M. Enomoto, S. Jogan, Metall. Mat. Trans. A 30 (1999) 2429–2437. [2] A.P. Reynolds, W. Tang, Z. Khandkar, J.A. Khan, K. Lindner, Sci. Technol. Weld. Join. 10 (2005) 190–199. [3] T.W. Nelson, R.J. Steel, W.J. Arbegast, Sci. Technol. Weld. Join. 8 (2003) 283–288. [4] P. Upadhyay, A.P. Reynolds, Mater. Sci. Eng. A 527 (2010) 1537–1543. [5] X.Guo Tang, J.C. McClure, L.E. Murr, J. Mater. Process. Manuf. Science 7 (1999) 163–172.
[6] M.Z.H. Khandkar, J.A. Khan, A.P. Reynolds, Sci. Technol. Weld. Join. 8 (n.d.) 165–174. [7] P.A. Colegrove, H.R. Shercliff, R. Zettler, Sci. Technol. Weld. Join. 12 (n.d.) 284–297. [8] Henrick Schmidt,Modeling of Thermomechancial Condition in Friction stir welding,Ph.D. Thesis, Denmark, Department of Manufacturing Engineering and Management, DTU, 2004. [9] R. Johnson, others, in: Third International Symposium on Friction Stir Welding, Japan, 2001. [10] T. Nishihara, Y. Nagasaka, in: Proceedings of the 4th International Symposium on Friction Stir Welding, 2003, pp. 14–16. [11] K. Elangovan, V. Balasubramanian, M. Valliappan, Int. J. Adv. Manuf. Technol. 38 (2007) 285–295. [12] J. Record, J. Covington, T. Nelson, C. Sorensen, B. Webb, Weld. J.N.Y. 86 (2007) 97. [13] M.J.C. Rosales, N.G. Alcantara, J. Santos, R. Zettler, MSF 636–637 (2010) 459–464. [14] D. Bakavos, P.B. Pragnell, R. Dif, in: n.d. [15] D. Chakrabarti, D.E. Laughlin, Prog. Mater. Sci. 49 (2004) 389–410. [16] I. Dutta, S.M. Allen, J. Mater. Sci. Lett. 10 (1991) 323–326. [17] G.A. Edwards, K. Stiller, G.L. Dunlop, M.J. Couper, Acta Mater. 46 (1998) 3893–3904. [18] C. Gallais, A. Denquin, Y. Bre´chet, G. Lapasset, Mater. Sci. Eng. A 496 (2008) 77–89. [19] A. Simar, Y. Bre´chet, B. de Meester, A. Denquin, C. Gallais, T. Pardoen, Prog. Mater. Sci. 57 (2012) 95–183. [20] Y.S. Sato, M. Urata, H. Kokawa, Metall. Mater Trans A 33 (2002) 625–635. [21] L.-E. Svensson, L. Karlsson, H. Larsson, B. Karlsson, M. Fazzini, J. Karlsson, Sci. Technol. Weld. Join. 5 (2000) 285–296. [22] K.A.A. Hassan, P.B. Prangnell, A.F. Norman, D.A. Price, S.W. Williams, Sci. Tech. Weld. Join. 8 (2003) 257–268. [23] K.V. Jata, K.K. Sankaran, J.J. Ruschau, Metall. Mat. Trans. A 31 (2000) 2181–2192. [24] P.L. Threadgill, A.J. Leonard, H.R. Shercliff, P.J. Withers, Int. Mater. Rev. 54 (2009) 49–93. [25] M.W. Mahoney, C.G. Rhodes, J.G. Flintoff, W.H. Bingel, R.A. Spurling, Metall. Mat. Trans. A 29 (1998) 1955–1964. [26] R.Y. Hwang, C.P. Chou, Scr. Mater. 38 (1997) 215–221. [27] H. Chandler, Heat Treater’s Guide: Practices and Procedures for Nonferrous Alloys, ASM International, 1996. [28] R. Brown, W. Tang, A.P. Reynolds, Mater. Sci. Eng. A 513–514 (2009) 115–121. [29] M. Wade, A.P. Reynolds, Sci. Technol. Weld. Join. 15 (2010) 64–69. [30] J.-H. Cho, D.E. Boyce, P.R. Dawson, Mater. Sci. Eng. A 398 (2005) 146–163. [31] H. Schmidt, J. Hattel, Model Simul. Mater. Sci. Eng. 13 (2005) 77–93. [32] A.P. Reynolds, Scr. Mater. 58 (2008) 338–342.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effects of forge axis force and backing plate thermal diffusivity on FSW of AA6056 P. Upadhyay n, A.P. Reynolds Department of Mechanical Engineering, University of South Carolina, 300 Main Street, Columbia, SC 29208, United States
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 May 2012 Received in revised form 3 August 2012 Accepted 6 August 2012 Available online 10 August 2012
Using three sets of tool rotation rate and welding speed, series of friction stir welds were made at various forge axis forces on different backing plates with widely varying thermal diffusivity on 4.2 mm thick AA6056. Temperature during the process was measured at the probe core using a thermocouple. Because of the use of different backing plates and forge forces it was possible to obtain a relatively wide range of peak temperature in the nugget. This provided a unique opportunity to investigate changes in joint microstructure and property (1) over a wide range of peak temperature while keeping welding and rotational speed constant (2) at similar peak temperature with different welding and rotational speeds. The results show that for the studied gauge thickness, metallurgically significant temperature variations can be achieved without changing the rotation and welding speed while producing defect free welds. The backing plate conductivity and forge force both independently affect the weld process parameters such as nugget temperature and tool torque significantly. A significant increase in the tool torque and hence power is observed with the increase in the forge force alone. & 2012 Elsevier B.V. All rights reserved.
Keywords: Friction stir welding Forge force Backing plate Thermal boundary conditions AA6056 Temperature measurements Torque
1. Introduction and background It has been established that the tool rotation rate primarily controls the peak temperature achieved at the deformation zone during friction stir welding [1–3]. For a series of welds the measured average torque and power has also been inversely correlated to the peak stir zone temperature and hence to the rotation rate. Fig. 1 shows the correlations of peak temperature measured at the tool probe (Probe T) and measured tool torque for welds made with 6.4 mm thick AA7050 in laboratory air and under-water [4]. Similar trends of peak temperature with change in rpm have been reported elsewhere in the literature [1,2]. One concern with this reasoning is that, in such series of welds it is not only the tool rpm that is varied. Assuming a force control scheme, during a series of welds with varying tool rotation rate, the FSW practitioner will invariably have to adjust the applied forge force as well in order to produce good quality welds with minimal flash. A ‘‘good’’ forge force for a specific welding condition will thus be established in due course of trial and error. ‘‘Good’’ here can be loosely defined as weld with minimal flash and no internal defects. Note that there will always be a small window of forge force that will produce ‘‘good welds’’. The size of the window
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Corresponding author. Tel.: þ1 803 777 3118; fax: þ1 803 777 0106. E-mail addresses: [email protected], [email protected] (P. Upadhyay). 0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.018
depends on factors including the alloy welded and the tool geometry. For an empirically established ‘‘good’’ forge force an increase in rpm will lead to increase in the average deformation zone temperature rendering the material more plastic. Without any forge force adjustments some of the plasticized material will expel away from under the shoulder instead of being consolidated behind the tool, resulting in excessive flash and, possibly, volumetric defects. For a given set of welding parameters (rpm and welding speed), plate geometry, alloy type and tool design the ‘‘good’’ forge force may change with changes in thermal boundary conditions (BC) that are present at the work piece. The rate of heat flux through the bottom of the work piece mostly depends on the backing plate (BP) thermal diffusivity and hence, as has been pointed out in literature [5–8], the most important thermal BC for FSW is the thermal condition at the backing plate. It was hence thought necessary to understand effects of forge force in conjunction with varying levels of thermal diffusivity of the backing plate. In the published literature the effects of forge force have not received as much attention as tool rotation and welding speed effects. Some qualitative study of general effects of either forge force or plunge depth can be found scattered in the literature. During a parametric study using position control Johnson [9] showed weld forge force increased with the increase in the plunge depth. Nishihara and Nagasaka [10] reported that tool and backing plate temperature increased when the plunge depth
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Fig. 1. Measured peak probe temperature and torque plotted against the rotational speed for in air (IA) and underwater (UW) welds made in 6.35 mm thick AA7050. Adapted from [4].
was increased. Elagovan et al. [11] studied the effect of tool pin profiles and forge force on FSW of AA6061. Out of the three forge forces considered the intermediate one was reported to show superior mechanical properties. Since no temperature measurements were reported the effect of thermal exposure on mechanical property could not be elucidated. Excessive forge force probably led to higher heat input resulting in high flash and thinning of the weld zone. On the other hand a less than sufficient forge force might have led to insufficient heat input and hence inadequate material flow. At about same time Record et al. [12] reported a design of experiment based approach where nine welding parameters were used. They reported that welding speed and location of weld relative to the edge of the plate also had some minor effect on forge force in a position control mode. Effects of thermal boundary condition changes at the bottom of the work piece have also not been thoroughly examined in the open literature. As a part of quench sensitivity study Nelson et al. [3] reported that the use of a heated backing plate resulted in higher peak temperature and lower cooling rate, resulting in inferior mechanical properties in AA7075 welds. Rosales et al. [13] reported welds made with AA2024 and AA6013 where steel, copper and ceramic coated backing plates were used at three different combinations of rotational and welding speeds while the forge force was kept constant. Stir zone temperatures were not measured while the in-plate far field temperature measurements varied significantly when using different backing plates indicating the significance of backing plate conductivity. Macrographs of the welds indicate thinning of the cross-section in ceramic coated backing plate, probably because of high flash content. As the application of friction stir welding widens it will be important to understand the behavior of process variables like stir zone temperature, torque and in plane forces with the change in thermal boundary conditions. This understanding will be helpful in a variety of scenarios. In modeling for instance, contact conductance between backing plate and work piece is still elusive but has been deemed central in quantifying the heat flux out from the bottom of the work piece. Value of this contact conductance may depend on applied forge force, surface temperature and backing plate thermal conductivity. A look at the literature shows that a wide range of contact conductance values have been used by different authors ranging from 0.2 to 105 W/m2 K [4,6–8]. Whether or not the contact conductance could vary such a great deal can be debated. Although not addressed in this paper, it is fitting to call for a need to establish some physically valid upper and lower bound at least in the order of magnitude sense. The experimental results discussed in this paper including peak
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temperature and power correlations with welds performed using various backing plates in conjunction with parametric finite element modeling may provide useful insight into the problem. A number of studies have been conducted to understand the structure-property-processing relationship in 6XXX series alloys [14–19]. A fine distribution of the metastable needle shaped b00 (Mg5 Si6) precipitate is considered to be responsible for peak aged condition. In addition to the well-known b00 phase, a lath shaped hexagonal precursor phase to Q’ phase has also been attributed to strengthening. During the welding process the optimized microstructure of the peak aged base material is altered at various levels because of the high temperature cycle. This cycle will produce some level of dissolution of precipitates, diffusion of solute, and increase in vacancy concentration. Thus some non-strengthening phases form and some strengthening precipitates dissolve or coarsen, resulting in varying degrees of loss of strength. The purpose of most welding research work is to minimize this effect while keeping an acceptable rate of production. To this end, qualitative assessment of weld cross-section normally begins with micro hardness tests that provide a first glimpse into changes that the material has gone through during the welding. Welds made on peak or overaged precipitation hardening alloys, such that the nugget approaches a temperature close to Solution Heat Treatment (SHT) temperature, typically result in a characteristic ‘‘W’’ shaped hardness profile after some artificial or natural aging [2,20,21]. The nugget which reaches such a peak temperature would be in a condition similar to that obtained after SHT and quenching; therefore, re-precipitation of strengthening phases may occur during post weld cooling and subsequent natural and artificial aging. It has been found when the nugget temperature is beyond the overaging regime (discussed subsequently), the nugget hardness increases with the increase in the stir zone temperature (achieved through increase in rpm) until reaching a plateau close to SHT [1,2]. This has been attributed to greater availability of solute in the matrix for reprecipitation into strengthening phases during subsequent aging at higher peak temperature than at lower peak temperature. On either side of the nugget there is a quick drop in hardness continuously falling across the thermo-mechanically affected zone (TMAZ), reaching global minima in the heat affected zone. This arises due to over aging of the heat affected zone (HAZ) [1,22,23]. If the weld is performed at relatively low power, such that the stir zone peak temperature is around 350 1C, the characteristic W shape in the hardness profile is replaced by a U shaped one. With a peak weld temperature near 350 1C, the hardness in the HAZ and the nugget will be similar to each other and less than peak aged base metal [2,4]. For precipitation hardening alloy it has been found that the peak temperature exposure of 350 1C is most damaging to strength [24]. Upon tensile testing of FSW samples of AA7075, Mahoney et al. [25] found that failure occurred along the heat affected zone where the peak T recorded was between 300 1C–350 1C. Hwang and Chou [26] and Sato et al. [1] simulated thermal cycles at heat affected zone using several peak temperatures in AA7075 and AA6063. Results from both works indicated that minimum hardening effects occurred on samples that were exposed to a peak T of 350 1C. Subsequent TEM/SAD analysis of the overaged nugget and/or HAZ region from several works [1,21,25,26] show coarsening and dissolution of strengthening precipitates and presence of non-strengthening b0 precipitates in 6XXX series alloys. Although the sequence of precipitation is more complicated than what has been presented here, for the purpose of this work the following simple statements can be made. The kinetics of precipitate coarsening are maximum near 350 1C for 6XXX alloys. In other words the rate of formation of non-strengthening b0 in 6XXX series is at peak near this temperature. During its formation, b0 particles take away a significant amount of solute
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from the matrix which otherwise would have been available for reprecipitation of strengthening b00 phase during the post weld aging. In the regions of the weld where the peak T is around 350 1C (mostly in the HAZ and when weld power is relatively small, in the nugget as well), the situation will be aggravated when the welding speed is lower. Lower welding speed will cause the material to remain in the b0 phase formation temperature range for a relatively longer time causing the solute depletion to increase, resulting in lesser strength. At peak temperature level farther away from 350 1C, approaching the SHT temperature, more and more precipitate will dissolve into the matrix avoiding precipitate coarsening such that during subsequent heat treatment fine strengthening precipitates are formed resulting in a stronger nugget zone.
the nugget center using the Mean Linear Intercept (MLI) method. Three views at a magnification of 200 were examined. Grain boundary intersections were counted on a test line of length 0.5 mm (100 mm at 200 ). Five test lines per view were used. Micro-hardness tests were performed using a Vickers hardness indenter on naturally aged and post weld aged samples. Post weld heat treatment of samples was carried out at 190 1C for 4 h. This time and temperature combination is consistent with the industrial practice leading to the T6 or peak aged temper. A Vickers hardness indenter, with a load of 500 g and a load application time of 10 s was used for measurement of hardness as a function of distance from the weld centerline on transverse cross sections along the plate. For all the hardness traverses plots reported in this work the advancing side is on the left.
2. Experimental procedure
3. Results and discussions
Bead on plate welds were made on 4.2 mm thick rolled plates of the aluminum alloy, 6056-T451. The temper designation T451 indicates that the alloy has gone through solution heat treatment, quenching and subsequent stretching to relieve internal stresses [27]. The solution treatment temperature (SHT) for the alloy is 529 1C. Welds were produced on a hydraulically powered MTS FSW Process Development System (PDS) using Z-axis (forge) force control. During the welding process the PDS records a number of relevant process variables as a function of time. The process data recorded includes variables such as traverse and rotational speeds of the tool; longitudinal, transverse and vertical force (X,Y and Z respectively) experienced by the tool and tool plunge depth. Real time torque required by the spindle motor is also measured using a torque transducer. The tool used for production of all welds was of a two piece design with a 17.8 mm diameter, single scroll, H13 tool steel shoulder and a probe fabricated out of MP-159 (a high temperature cobalt based super alloy) in the shape of a truncated cone (81 taper) with threads and three flats. The probe was 4.1 mm long, with a diameter of 6.5 mm at the intersection with the shoulder. Temperature during welding was monitored and recorded using a K-type thermocouple spot welded into the probe on the axis of rotation approximately at the probe mid-plane height. Series of welds were made with three sets of rotational and translational speeds using various forge forces and backing plates with widely varying thermal diffusivity. Three sets of welding and rotational speeds from 320 rpm–3.4 mm s 1, 640 rpm– 6.8 mm s 1 and 960 rpm–10.2 mm s 1 were used. Welding speed and rotational speed were chosen such that the advance per revolution was 0.635 mm in all cases. Thirteen millimeter thick backing plates (BP) made of material with widely varying thermal diffusivity viz. AA2099, tool steel, Al6XN (an Allegheny Ludlum super austenitic stainless steel), Ti-6-4, and ceramic floor tile were used. The corresponding forge forces used can be inferred from Figs. 3 and 4. A forge force of 12.8 kN deemed as ‘‘good’’ for the steel BP with welds performed at 640 rpm and 6.8 mm s 1 was used as a baseline. Levels of forge force changes of 710%, 20%, 30% were then used to perform the set of experiments for different backing plates and the three sets of rotational and translation speed. Standard metallographic preparation technique was employed to grind and polish samples to satisfactory level before performing characterization by optical microscopy and hardness testing. Samples were ground using automatic and/or manual grinding machines with 180, 240, 320, 400, 600, 800 grit silicon carbide paper. They were then polished using Aluminum oxide powder of 5 mm and 3 mm followed by colloidal silica ( o0.05 mm). Samples were etched using Keller’s reagent. Grain size was measured at
3.1. Peak temperature correlations Representative temperature transients recorded by thermocouples located at the probe core are shown in Fig. 2 for welds made at 640 rpm–6.8 mm s 1 with aluminum backing plate. For all the forge forces used the probe T reaches a fairly steady value after half the weld length. Welds made with other backing plates also exhibit qualitatively similar temperature trends. The temperature measured inside the probe can be thought of as some average of the temperature of material in contact with the probe and may not be reflective of the maximum temperature at the nugget. However this measured temperature has been found to be repeatable and provides reasonable representation of nugget temperature with changes in welding parameters [4,28]. Fig. 3 shows the measured steady state peak probe temperature achieved at several levels of forge forces using four different backing plate materials for the 640 rpm–6.8 mm s 1 welds. These data were extracted from graphs like Fig. 2. Excluding ‘‘bad’’ welds indicated by cross marks viz: three welds made at relatively low forge forces that resulted in surface defects and three welds at relatively high forge force that resulted in high level of flash and thinning; the rest of the runs produced good quality welds with little to no flash and no defects. This figure illustrates substantial changes in process response that can be brought about by changes in either forge force, backing plate diffusivity or both while keeping rpm and welding speed
Fig. 2. Temperature transients measured by thermocouple at the probe core for welds made at 640 rpm–6.8 mm s 1 with aluminum backing plate at different forge forces as indicated in the legend.
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Fig. 3. Peak probe T for welds made at 6.8 mm s 1 using various indicated backing plate materials. Crosses indicate welds with defects or excessive flash.
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speed than lower rpm/lower welding speed. This indicates that increase in temperature caused by increase in the tool rpm (and higher weld power) overwhelms the corresponding temperature reduction effect of increased welding speed. If the lines of data from different backing plates and same parameter set are approximated as parallel this shows that the effect of backing plate diffusivity on achieved peak probe T is greater at lower welding and/or rotational speed and diminishes as they increase. Demonstratively at 3.4 mm s 1 the change from Al BP to Ti BP at a forge force level of 14.2 kN can cause a 90 1C change while for 10.2 mm s 1 the change in temperature is no greater than 40 1C. This suggests that the ability to alter nugget zone temperature by changing the backing plate is decreased with increased rotational and welding speed. The individual effect of welding speed and rotational speed are difficult to gather from above data set because the welds were performed at the same advance per revolution. With higher rotation rate the nugget temperature approaches the solidus T causing the material flow stress to decrease. This decrease will in turn limit power generation by plastic deformation and hence temperature increase. Thus, rate of temperature increase due to addition of power will be higher at lower temperature. On the other hand at higher welding speed the viscous diffusion rate due to higher rpm and welding speed dominates the conduction heat loss (thermal diffusion). The only way to settle this is to conduct some additional experiments with either one of the two parameters constant. Since welding and rotational speeds were constant the shear and shear rate around the probe can be assumed to be nearly constant, these results provides a unique opportunity to investigate microstructure where shear and shear rate are constant while the Peak T is significantly varying. This is predicated on the assumption that, for a given welding speed and rpm, the kinematics of the flow does not vary greatly. 3.2. Power/torque correlations
Fig. 4. Peak probe T for welds made at 3.4 mm s 1 (dashed lines) and 10.2 mm s 1 (dotted lines) using indicated backing plate materials.
constant. For instance considering only good quality welds with no defect and minimal flash, combined variation of forge force and backing plate diffusivity can result in as much as 70 1C difference in probe temperature. Change in forge force or backing plate alone can result in 50 1C difference in probe temperature (As shown by dotted vertical and horizontal lines). At forge axis force of 11.3 kN, for instance, a ceramic tile BP weld can achieve a peak probe T of 460 1C, while under equivalent rotational and translation rate steel BP achieves only 431 1C. Aluminum BP likewise attains 423 1C. Expectedly the peak probe temperatures line up in accordance with the backing plate diffusivity and the applied forge force. Similar results are observed from other parameter set combinations (see Fig. 4). This trend of increase in stir zone temperature by changing only the forge force or thermal boundary at the back plate is qualitatively similar to the trend seen in Fig. 1 obtained from series of experiments where the tool rotation rates were varied. The peak probe T extracted from weld parameter set from 320 rpm–3.4 mm s 1 (dashed lines) and 960 rpm–10.2 mm s 1 (dotted lines) are plotted in Fig. 4 indicated by arrows are the differences between peak probe temperatures achieved between two extreme cases of backing plates for two extreme cases of welding speeds. Note that the peak temperature obtained for equivalent forge force is higher with higher rpm/higher welding
Similar substantial change in required torque and hence power resulted from forge force and backing plate variation albeit the relationships are more complicated than temperature relationship described above. Weld power calculated from the measured torque is plotted against the applied forge force in Fig. 5 for the three parameter sets. Consider first the 320 rpm–3.4 mm s 1 set where a reasonably linear relationship exists between weld power and forge force. Data sets with higher diffusivity backing
Fig. 5. Weld power calculated from measured torque plotted against the applied forge force for various backing plates shown in different symbols. Different welding speeds are shown in different line styles as indicated.
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plate exhibit a greater change in power with the change in forge force as indicated by difference in slope of the lines. For instance a forge force increase from 14.2 kN to 18.5 kN produced an increase in power from 1.5 kW to 2.5 kW in Al BP. An equivalent forge force change in Ti-6-4 BP from 11.4 kN to 15.6 kN only changes the power from 1.5 kW to 1.75 kW. With higher welding and rotational speed sets the power response with change in forge force is reduced further. In the case of Ti backing plate at 960 rpm–10.2 mm s 1, there is virtually no change in power despite a 50% increase in forge force. The same is true for steel BP at 960 rpm–10.2 mm s 1 except at the lowest forge force. The aluminum backing plate on the other hand shows substantial increase in the torque similar to that seen in the lower rpm welding. Note that welds made with the aluminum backing plate do not reach to a torque plateau for tested forge force levels; however, it can be anticipated that torque would level off if higher forge forces were used shown by continuation of lines beyond data points. Further increase in forge force on Al BP was not considered to avoid production of a lap weld between the work piece and the backing plate. Similar relationship is shown for 640 rpm–6.4 mm s 1 set. In this case gradual increase in the torque with the increase in forge force is seen until power plateaus for all but Al BP. Note that for a given parameter set the value of power at the plateau increases with the increasing diffusivity of backing plate used. This can be explained by observing the probe T vs. forge force graph in Fig. 3. Owing to greater heat retention, relatively high probe temperature is consistently observed for lower diffusivity backing plate indicative of higher shoulder interface temperature. This high temperature in turn leads to greater softening of the material hence material deforms at lower shear stress leading to lower required torque. 3.3. Temperature vs power trends In Fig. 6a–c measured probe T for three sets of welding parameters are plotted against the weld power for different backing plates. Again note that ceramic tile backing plate was used rather than Ti-6-4 for 6.8 mm s 1 case as shown in the legend in Fig. 6c. For the Al BP, the relationship between power and probe temperature is decidedly linear regardless of the tool rotation and translation rate. For all of the Al BP welds, the temperature increases with increasing power and by inference from this data, the power increases with increasing forge force regardless of the welding speed and rotational speed used. Considering Fig. 6b and c the relationship between probe T and power is not so straightforward with lower diffusivity backing plates. In the case of steel BP, the 320 rpm–3.4 mm s 1 data and some of the 640 rpm–6.8 mm s 1 data exhibit relatively linear relationships with power, although the relationship is not unique:
for a given power level, the 320 rpm–3.4 mm s 1 welds exhibit higher peak T than the 640 rpm–6.8 mm s 1 welds. This may be because the slower welds take longer to complete and the probe temperature increases with weld time; therefore, the peak T, measured at the end of the welds, varies with weld time and actual weld zone T. Even after saturation in the torque/power required is reached at higher forge force in the case of low conductivity backing plate viz. Steel and Ti-6-4 as discussed in previous section, the measured probe T continues to rise. For example the maximum power for the 640 rpm–6.8 mm s 1 welds is approximately 3 kW and four welds (made with the four highest values of forge force ) all have the same power level but different probe T. All of the 960 rpm–10.2 mm s 1 welds have approximately same power levels but the measured temperatures vary by as much as 40 1C. Similar behavior is observed for the lowest diffusivity anvils (Ti-64/Tile). 3.4. Nugget grain size and microhardness results and correlations The previous sections sufficiently demonstrate our ability to affect process response variables such as required torque and weld zone temperature as well as, potentially, enabling production of good quality welds with lower forge force by modifying of thermal boundaries however, of equal importance are the ability to use the thermal boundary condition to influence weld microstructure and properties. In the following section the effects of forge force and backing plate diffusivity on hardness distributions and nugget grain size are discussed and correlated. In Fig. 7 the grain sizes measured at the nugget center using MLI method are plotted against the measured peak probe temperature. Data sets for two different welding speeds are shown by closed (6.8 mm s 1) and open (10.2 mm s 1) symbols. Welds made with different backing plates are shown using differently shaped symbols as indicated in the legend. There is almost a linear trend of increase in the grain size with the increase in the measured probe temperature. This ‘‘linear’’ relationship is true regardless of the boundary condition at the bottom of the work piece and the welding speed. The apparent linear relationship across two welding speeds suggests that for the given case grain size is relatively insensitive to the time at temperature during the welding and is governed mainly by the peak temperature. This is consistent with the fact that metallurgical processes such as grain growth are exponentially dependent on temperature and exhibit much weaker dependence on time. Shown in the Fig. 8a are the transverse micro hardness profiles for naturally aged samples taken from welds made with a ceramic tile backing plate at three different forge forces for parameter set of 640 rpm–6.8 mm s 1. The respective forge force used and measured midplane peak T achieved during the welding is
Fig. 6. Peak probe T plotted against the weld power for (a) Aluminum, (b) Steel, and (c) Ti backing plates. Different symbols indicate different welding speeds.
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indicated in the legend. This set of graphs illustrates changes in mechanical property that occur in weld cross section with the change in peak temperature due to the change in forge force alone. There is a significant change in the shape as well as the absolute value of the hardness distribution in the nugget with changes in the applied forge force. This change can be explained on the basis of the measured peak probe T during the welding. As discussed in the introduction section, the average nugget
Fig. 7. Midplane nugget grain size plotted against measured peak probe temperature for two different welding speeds.
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hardness increases with the increase in the probe T. For instance with a low measured peak T of 398 1C obtained at a low forge force of 5.7 kN, the weld nugget is in over-aged condition exhibiting a U shaped hardness distribution. As the forge force is increased the nugget hardness increases while the heat affected zones on both the sides remain over-aged. Notice the nugget hardness asymmetry in welds made at intermediate forge force. Several such cases have been observed in the current work and they will be dealt with in a separate section. With sufficiently high forge force of 12.8 kN, where the measured peak T was recorded to be 490 1C nugget hardness approaches base metal value and asymmetry vanishes. Notice also that minimum hardness values for all the considered samples located at HAZ for W shaped distributions and all over the nugget for U shaped are in the narrow range of 70–74 HV. This is reasonable and expected since thermal boundary conditions and welding speeds for all the cases were identical such that time at elevated temperature for precipitate coarsening and dissolution has not changed substantially with the change in forge force. Fig. 8b, c and d show similar hardness results for welds made at 3.4, 6.8 and 10.2 mm s 1 respectively using high conductivity aluminum as backing plate material in contrast to low conductivity ceramic tile in Fig. 8a. Corresponding forge forces used and peak T attained at the tool midplane are indicated in the legend. In the case of welds made at 3.4 mm s 1 (Fig. 8b) despite a temperature increase from 392 1C to 431 1C there is very small increase in the nugget hardness. All the three profiles appear to be in an overaged condition as will be discussed subsequently. In the case of welds made at higher welding speed of 10.2 mm s 1 and higher rpm, (Fig. 8d), all the three forge forces resulted in W
Fig. 8. Vickers hardness profiles on weld transverse cross-section at midplane for naturally aged samples (a) ceramic floor tile BP at 6.8 mm s 1 (b) aluminum BP at 3.2 mm s 1 (c) aluminum BP at 6.8 mm s 1 (d) alum BP at 10.2 mm s 1. Note that advancing side is on the left for all the hardness traverses.
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shaped hardness distributions such that with higher applied forge force the nugget hardness is higher and approaches the base metal hardness. The effect of achieved peak temperature on the nugget hardness, all other things being equal, can be appreciated by comparing the hardness distribution among Al (Fig. 8c) and ceramic tile BP (Fig. 8a) at parameter set of 640 rpm–6.8 mm s 1. With same forge force of 8.5 kN for instance weld made with Al BP consists of weld cross-section that is overaged hardness while with Tile BP the weld nugget has achieved some degree of hardness recovery. The only difference between these two cases is that the Tile BP attained peak temperature that is 50 1C higher than Al BP. This relationship between measured peak temperature and nugget hardness is further explored in Fig. 9 where the arithmetic average of edge to edge midplane nugget hardness are plotted against the measured peak mid-plane probe T. Note that welds made at same welding speed are represented by same symbol regardless of the backing plate used. Whereas the closed symbols indicate data from naturally aged samples, the open symbols represent data obtained from samples that were heat treated at 190 1C for 4 h. Note that not all data shown in naturally aged (NA) condition were tested after heat treatment (HT) so some NA samples do not have corresponding HT data points. Note also that relationship between peak probe temperature and nugget hardness are different amount heat treated and as welded samples. For post weld heat treated samples, the hardness value plateaus at higher peak probe temperature range, whereas for naturally aged case no such plateau is observed. Presumably at a high temperature range, in post weld heat treated samples the additional solute obtained because of further increase in the probe T is not significant to noticeably increase the hardness. In Fig. 10 the difference in average nugget hardness between heat treated and naturally aged samples are plotted against the measured peak T. First consider welds made at 960 rpm–10.2 mm s 1and 640 rpm–6.8 mm s 1. Excluding boxed data points at two lowest peak T cases for 6.8 mm s 1(they will be dealt with shortly), the nugget hardness generally increases with the increase in peak probe temperature for both as welded and PWHT cases. A hardness recovery of 15–25 HV is observed for all these cases (Fig. 10). Now consider boxed data points in Figs. 9 and 10 (most at 3.4 mm s 1and two at 6.8 mm s 1) where there is no apparent change in hardness despite a change in peak probe T from 360 to 430 1C. The hardness distribution for all these data points
Fig. 10. Change in average nugget hardness after heat treatment plotted against the corresponding peak T. Boxed points belong to overaged nugget that are unresponsive to heat treatment.
resembles the ones shown in Fig. 8b (U shaped). Note that this also happens to be the regime where no significant hardness recovery after PWHT takes place as evident from o5 HV recovery after heat treatment (Fig. 10). In other words they are all overaged. These are welds exposed to a relatively lower peak temperature and are at temperature regime where precipitation kinetics for coarsening of strengthening precipitates is at the peak as noted in the background section. Note that beyond certain temperature levels, welds made at both the welding speeds (3.4 mm s 1 and 6.8 mm s 1) start showing significant hardness recovery as well as sensitivity to peak probe T. It is interesting to note that this temperature beyond which significant hardness recovery after PWHT ( 415 HV) is observed appears to be correlated with the welding speed (thus time at temperature). At 3.4 mm s 1 for instance this ‘‘transition’’ temperature is 440 1C while for 6.8 mm s 1 the transition occurs at significantly lower temperature of 390 1C as indicated by the corresponding jump in the data point from the box shown by arrows. Alternatively this means that transition from U shaped (overaged throughout the nugget) to W shaped (significant hardness recovery at the nugget with overaged at HAZ) hardness distribution occurred at significantly lower temperature at higher welding speed. Note that for welds made at the highest welding speed of 10.2 mm s 1no overaged nugget hardness was observed at equivalent low temperature and all the hardness measured were of W shaped. The authors posit that significant hardness recovery at this welding speed might be anticipated at still lower temperature. This shows that time at temperature has a significant effect on precipitation kinetics in addition to the peak T. Alternatively this suggests that temperature considered most damaging for the strength of weld nugget depends on the welding speed and should not be thought as a fixed temperature regime. There may also be an effect of strain which is not accounted for in a purely temperature based analyses. Over a range of welding speed of 0.8–3.8 mm s 1 Reynolds et al. [2] presented similar trends of PWHT hardness recovery with increase in the welding speed in AA7050. The peak T obtained from an input torque model was correlated with PWHT hardness recovery showing that a temperature range of 330– 370 1C rendered the welds overaged and unresponsive to PWHT. 3.5. On hardness asymmetry
Fig. 9. Average nugget hardness plotted against peak T for 3 welding speeds (shown in different symbols). Closed symbols are data from naturally aged samples. Open symbols are for heat treated samples.
In few cases the measured hardness value did not stay constant throughout the nugget. As seen in some hardness profiles with
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intermediate peak temperature range, in Fig. 8a, c and d the hardness is highest at the advancing side edge of the nugget and gradually decreases from advancing to retreating side eventually reaching minimum at heat affected zone on the retreating side. To investigate the effect of peak deformation zone temperature on the degree of asymmetry, the difference between hardness measured at advancing and retreating side of the nugget (DVHNasymmetry) is plotted against the corresponding measured peak probe T for two welding speeds of 6.8 and 10.2 mm s 1 in Fig. 11. The data set from welding speed of 3.4 mm s 1 is not shown in Fig. 11 since no significant nugget asymmetry (the highest is 4.5 HV) is observed. For 6.8 mm s 1 after crossing the overaged regime (all data set with Peak T4400 1C), the nugget asymmetry increases with the increase in the probe T reaching a maximum of DHV¼18.5 at peak T of 450 1C.With further increase in the probe T the asymmetry vanishes, indicating flat nugget hardness distribution. In the case of 10.2 mm s 1 maximum asymmetry is observed at the peak T of 430 1C (lowest peak T for 10.2 mm s 1 case) after which the asymmetry decreases significantly indicating that the nugget hardness becomes flat beyond the peak T of 470 1C. Similar asymmetries in nugget hardness distributions have been reported by Wade and Reynolds [29] on welds made at 720 rpm and 10.2 mm s 1 in AA6019. Equivalent weld made with identical welding speed at a higher rpm of 1680 on the other hand resulted in a reasonably flat nugget hardness without asymmetry. This was observed for both as welded and PWHT samples. It has been proposed and argued in the literature [30–32] using simulation results, some analytical treatment and inherent asymmetrical flow pattern in the process, that advancing side experiences higher temperature and strain rate in comparison to retreating side. Using grain size as a ‘‘proxy’’ for peak T attained Wade and Reynolds [29] showed indirect experimental evidence of temperature asymmetry in the nugget for the first time. Having a wider window of peak T range over three rotational speeds Fig. 11 provides ample opportunity to expand the argument. As explained in the introduction section, away from the overaging regime ( 350 1C), incrementally higher peak T results in higher solute being available for reprecipitation during subsequent aging resulting in incrementally greater strength. This strength/peak T relationship is linear only up to a certain peak T beyond which the incremental amount of additional solute obtained by further peak T increase is not significant to noticeably increase the hardness. In other words saturation in the precipitation of strengthening phases is reached beyond a critical peak T in the vicinity of SHT. Until that saturation is reached the region towards the advancing side
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Fig. 12. HAZ minimum hardness of naturally aged samples vs. welding speed. Different symbols indicate different backing plates used.
being exposed to greater peak T will be stronger than the retreating side resulting in asymmetric hardness distribution. This explanation is consistent with the observation in Fig. 11 where asymmetry disappears when the hardness reaches peak strength.
3.6. HAZ minimum hardness While hardness measurements in relation with peak probe T is interesting, it is the minimum hardness that typically occurs at the HAZ which determines the strength of the joint. The HAZ minimum hardness measured from the naturally aged samples are plotted against the welding speed in Fig. 12. Different symbols indicate welds made under different backing plates. There is a familiar trend of increase in minimum hardness with the increase in the welding speed which has been observed in the past in several precipitation hardening alloys [4]. Unfortunately the minimum hardness does not show any significant sensitivity to the backing plate material. Spread within hardness data among same welding speed is too narrow to distinguish between the various cases. Hence it can be asserted that for this gauge and temper, welding speed has greater effect on minimum HAZ hardness than thermal BC at the bottom of the plate (backing plate material). Accordingly this suggests that the welding speed is primary determinant of effective time at elevated temperature.
4. Conclusions
Fig. 11. Difference between hardness measured at advancing and retreating edge of the nugget (DVHNasymmetry) plotted against the corresponding measured peak probe T for welding speeds of 6.8 and 10.2 mm s 1.
(1) Forge force has a significant effect on the stir zone temperature of otherwise similar welds. Defect free welds with metallurgically significant temperature variations can be achieved by changing the forge force alone within controlled limits. (2) The backing plate diffusivity has a very important role in the peak stir zone temperature achieved. For the considered plate gauge, thermal and microstructural through thickness homogeneity can be achieved by the use of low diffusivity backing plate material. (3) Similar peak temperature at nugget can be achieved by using a wide range of forge force with the use of different BPs. Same applied forge force can yield a wide range of peak temperature if different backing plates are used. (4) The tool torque increases with the increase in forge force until reaching a plateau at a high forge force level. With the use of
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low diffusivity backing plate, tool torque required for otherwise equivalent welding conditions is reduced. (5) For welds made at two different welding speeds, nugget grain size has a nearly linear relationship with the measured peak probe temperature regardless of backing plates used. (6) The measured peak temperature beyond which significant recovery in the nugget hardness after PWHT occurs depends on the welding speed. In the case studied, while for 6.8 mm s 1, the overaged nugget insensitive to PWHT was avoided beyond the peak T of 390 1C, for 3.4 mm s 1 a peak temperature of 440 1C was necessary to get similar effect. (7) While the use of different backing plates made a significant effect on peak temperature at the deformation zone, the rate of cooling was not affected in significant ways as evident from the minimum hardness correlations.
Acknowledgments The authors acknowledge the financial support of the Center for Friction Stir Processing which is a National Science Foundation I/UCRC supported by Grant No. EEC-0437341. The authors thank Dr. Wei Tang and Daniel Wilhelm, Department of Mechanical Engineering, University of South Carolina, Columbia, SC, USA for their help in preparing the weld joints. References [1] Y.S. Sato, H. Kokawa, M. Enomoto, S. Jogan, Metall. Mat. Trans. A 30 (1999) 2429–2437. [2] A.P. Reynolds, W. Tang, Z. Khandkar, J.A. Khan, K. Lindner, Sci. Technol. Weld. Join. 10 (2005) 190–199. [3] T.W. Nelson, R.J. Steel, W.J. Arbegast, Sci. Technol. Weld. Join. 8 (2003) 283–288. [4] P. Upadhyay, A.P. Reynolds, Mater. Sci. Eng. A 527 (2010) 1537–1543. [5] X.Guo Tang, J.C. McClure, L.E. Murr, J. Mater. Process. Manuf. Science 7 (1999) 163–172.
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