Friction stir welding of multilayered steel

Friction stir welding of multilayered steel J. Taendl*1, S. Nambu2, J. Inoue2, N. Enzinger1 and T. Koseki2 This study investigates the mechanical properties and microstructure of friction stir butt welded high strength/ductility multilayered steel consisting of 15 alternating layers of SUS 301 austenitic stainless steel (eight layers) and SUS 420J2 martensitic stainless steel (seven layers) with a total thickness of 1?2 mm. With optimised welding parameters, defect free welds with an ultimate tensile strength (UTS) of 1240 MPa and a fracture elongation of 13% were accomplished. This corresponds to a joint efficiency of 90%. In this case, fracture occurred in the heat affected zone as a result of a very pronounced hardness drop in the martensitic layers resulting from the formation of a large amount of grain boundary precipitates, which were formed at temperatures ,750uC slightly below Ac1. By applying post-weld heat treatment, the hardness drop in the martensitic layers was removed and the tensile properties were enhanced to UTS of 1310 MPa (95% joint efficiency) and a fracture elongation of 22%. Keywords: Friction stir welding, Multilayered material, High strength steel, Stainless steel, Butt configuration

Laminated composite materials consisting of alternating metal or metal containing layers dramatically improve many mechanical properties of single constituents, including fracture toughness, fatigue behaviour, impact behaviour, wear, corrosion, damping capacity and even deformability.1 Multilayered steel with enhanced deformation behaviour and tensile properties consisting of alternating layers of ductile austenitic steel and brittle martensitic steel has been subject to various investigations.2–6 Inoue et al.2 showed a transition of the fracture behaviour of the martensitic layer from brittle cleavage to ductile shear just by reducing the layer thickness in a hot rolled laminate. More than 50% elongation of as quenched martensite was achieved under tensile loading using a laminated composite consisting of conventional martensitic and austenitic steel.3 Nambu et al.4 demonstrated tremendous improvement in tensile ductility by increasing the bonding strength in a brittle/ductile multilayered structure. Detailed investigations concerning strain field computation5 and microtexture development for different cold rolling reductions6 were carried out for multilayered steel composed of martensitic and austenitic layers. However, no successful welding of that material has been reported so far. Friction stir welding (FSW) is a solid state welding process invented at The Welding Institute in Cambridge, UK in 1991.7 The basic concept provides a rotating tool with a specially designed pin and shoulder that is inserted in the abutting edges of the plates to be joined and traversed

along the joint line in order to heat up the material and generate a material flow to produce the joint. The heat is generated during the process by friction between the tool and the workpiece and by plastic deformation in the workpiece.8 Friction stir welding is very successfully used for joining relatively soft workpiece materials such as a wide range of aluminium alloys9 and is also considered a very suitable method for welding dissimilar materials.10 However, FSW still faces problems when it comes to welding of steel, even though the weld microstructure and the mechanical properties are in general very good.11 Owing to the higher strength of steel compared to aluminium and the resulting higher process forces and temperatures, expensive tools with outstanding mechanical properties need to be used. Thus, the application of FSW steel is currently limited to joining problems that cannot be solved by conventional techniques.12 Since the excellent tensile properties of multilayered steel are attributed to the laminated composition, the deterioration of the structure during a welding process is preferably low. Conventional fusion welding methods would certainly destroy the laminated structure due to melting by creating a rather undefined alloy.13 The FSW process, in contrast, provides solid state welding and thus less impact on the joining material. In this study, sheets of a laminated composite material consisting of alternating layers of SUS 301 austenitic stainless steel and SUS 420J2 martensitic stainless steel were welded in butt configuration using FSW. The weldability and the influence of different welding parameters on the mechanical properties and on the microstructure were investigated. Moreover, the effects of post-weld heat treatment (PWHT) were evaluated.

1

Experimental

Introduction

Institute for Materials Science and Welding, Graz University of Technology, Kopernikusgasse 24, Graz 8010, Austria 2 Department of Materials Engineering, the University of Tokyo, 7-3-1 Hongo, Bunkyo-ku Tokyo 113-8656, Japan *Corresponding author, email [email protected]

Material The material used in the present study was a multilayered steel composite consisting of 15 alternating

ß 2012 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 9 December 2011; accepted 11 January 2012 DOI 10.1179/1362171812Y.0000000003

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Friction stir welding of multilayered steel

a optical micrograph; b electron backscatter diffraction (EBSD) image of laminated structure; c, d EBSD images of layer interface 1 Microstructure of multilayered material consisting of 15 alternating layers of SUS 301 and SUS 420J2

layers of SUS 301 austenitic stainless steel (eight layers) and SUS 420J2 martensitic stainless steel (seven layers). The chemical compositions (mass-%) of the single constituent materials are shown in Table 1. In order to fabricate the composite, the 1 mm thick steel sheets were stacked and sealed at the edge by tungsten arc welding before hot rolling at 1100uC to a thickness of 1?6 mm. After a subsequent annealing heat treatment at 680uC for 10 min, the multilayered composite was cold rolled to a final thickness of 1?2 mm and finally heat treated at 1000uC for 2 min followed by cooling in N2 gas with a cooling rate of 5uC s21. Consequentially, the SUS 301 layers showed a microstructure composed of mainly austenite, some deformation induced martensite and chromium rich precipitates at the grain boundaries, whereas the SUS 420J2 layers showed a martensitic microstructure with some ferrite (Fig. 1). A comparison of the mechanical properties of the multilayered steel and that of the constituent materials is illustrated in Table 2.

Welding The multilayered steel sheets were cut to a size of 5006120 mm and friction stir butt welded in rolling direction. The welding distance for each parameter set was 100 mm. Temperature measurements using K-type thermocouples were accomplished on top of the steel sheets at a distance of 10 and 14 mm from the weld centre both on the advancing side (AS) and on the retreating side (RS) in order to catch expected temperature asymmetries.14 The tool was nitride coated and made of a sintered matrix consisting of 92 wt-% tungsten carbide and 8 wt-% cobalt. The hardness of the tool material equalled 1275 HV30. The tool consisted of a concave shoulder with a diameter of 17 mm and a conical probe with a tip diameter of 5?5 mm and a length of 0?85 mm.15 A special water cooled welding head was used to protect the tool from overheating. Based on previous studies,16,17 all welds were made in force control mode with a steady down force of 19 kN and a constant tilt angle of 1?5u between the tool rotation axis and the vertical axis. The tool rotation speed ranged from 300 to

1200 rev min21, and the welding speed ranged from 50 to 500 mm min21 in order to cover a broad parameter area. A summary of all the used welding parameters is given in Table 3.

Testing The mechanical properties of the welds were evaluated using tensile testing transverse to welding direction on the one hand and hardness measurement on the other hand. Individual measurements instead of statistical evaluations were performed, since for each parameter combination, only one weld could be produced. The dimensions of the tensile specimen were 12?5 mm in Table 2 Mechanical properties of SUS 301, SUS 420J2 and laminated composite consisting of SUS 301 and SUS 420J2 Material

Rm/MPa

Elongation/%

SUS 301 SUS 420J2 Multilayered steel

1145 1750 1370

44 3?5 26

Table 3 Summary of all used welding parameters Tilt angle/u

Vertical force Fz/kN

Spindle speed/ rev min21

100 150 100 150 200 50 80 100 150 200 300 200 300 200 300 400 500

300

400

1?5

Welding speed/ mm min21

19

550 700 800 1200

Table 1 Chemical compositions of component steels/mass-% Material

C

Si

Mn

P

S

Cr

Ni

Balance

Thickness/mm

SUS 301 SUS 420J2

0?1 0?29

0?51 0?61

0?78 0?44

0?028 0?023

0?001 0?003

16?80 13?13

6?54 0?22

Fe Fe

1 1


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2 Welded sheets

width, ,1?2 mm in thickness and 50 mm in gauge length according to the standard EN 10002-1:2001.18 Hardness tests across the weld seam in both martensitic and austenitic layers were carried out. The used hardness measurement method was HV0?1 Vickers hardness with a load of 0?981 N. The influence of a PWHT on the mechanical properties and on the microstructure was assessed by applying a similar heat treatment after welding as in the last step of the production process of the multilayered steel. More precisely, the PWHT included holding at a temperature of 1000uC for 2 min followed by air cooling. Investigations on the fracture behaviour and on the microstructure were performed using light optical microscopy (LOM), scanning electron microscopy (SEM), EBSD and transmission electron microscopy (TEM).


from the weld centre of 10 and 14 mm on both AS and RS. Apparently, the temperature on the AS was, in general, higher than on the RS due to an addition of tool rotation and welding speeds, resulting in higher friction and plastic deformation. This fact showed good correlation to previous investigations,16,22 even though the temperature difference between AS and RS was more pronounced in the present study. As expected, the maximum temperatures increased with decreasing revolution pitch, showing maxima at 10 mm distance on the AS equalling 390uC for a revolution pitch of 0?55 mm rev21 (Fig. 4a) and 470uC for a revolution pitch of 0?25 mm rev21 (Fig. 4b), compared to 765uC for a revolution pitch of 0?09 mm rev21 (Fig. 4c).

Results and discussion Temperature measurement, tensile testing and LOM were performed for every single parameter set. Hardness testing and detailed metallographic investigations were only carried out for selected samples.

Weldability Figure 2 illustrates the typical outcome of a welding experiment showing welded sheets with four weld seams accomplished using different welding parameters. Figure 3a shows an optical micrograph across a weld seam etched with V2A etchant.19 The weld was accomplished at a tool rotation speed of 800 rev min21 and a welding speed of 200 mm min21. The right side of the figure refers to the AS, and the left side refers to the RS. The position of the tool and the classification of the weld regions20 in parent material (PM), heat affected zone (HAZ), thermomechanically affected zone (TMAZ) and stir zone (SZ) are illustrated in Fig. 3b. No major problems concerning process stability were noticed; complete welding was feasible with appropriate parameter set-up.

Temperature measurement It has been reported21 that the heat input is related to the revolution pitch (mm rev21) and thereby on the ratio of welding speed (mm min21)/tool rotation speed (rev min21), as demonstrated in equation (1) Revolution pitch~

welding speed tool rotation speed

(1)

Figure 4 shows temperature cycles for different heat inputs measured on top of the welded steel sheets at a distance

3 a optical micrograph across weld seam and b micrograph including position of FSW tool and weld zones

a revpitch50?55 mm rev21; b revpitch50?25 mm rev21; c revpitch50?09 mm rev21 4 Temperature cycles at different distances from weld centre


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5 Influence of different welding speeds on tensile properties at constant tool rotation speed of 550 rev min21 and constant down force of 19 kN

Mechanical properties Figure 5 shows the results of tensile tests for different welding speeds at a constant tool rotation speed of 550 rev min21. For the lowest welding speed of 50 mm min21 (0?09 mm rev21), a tensile strength of 1070 MPa and a fracture elongation of at least 4?7% were reached. By increasing the welding speed, the tensile properties were improved, having their peak at a welding speed of 200 mm min21 resulting in a tensile strength of 1220 MPa and a fracture elongation of 8?5%. However, by further increasing the welding speed to 300 mm min21 (0?55 mm rev21), the strength and ductility of the joint dropped to 1100 MPa in ultimate tensile strength (UTS) and 4% maximum elongation. In short, it was demonstrated that the welding parameters played a dominant role for the weld quality by directly influencing the heat input. In Fig. 6a, a summary of the results of tensile tests for all investigated parameter combinations is given. The


tensile properties ranged from a low level of 690 MPa tensile strength and 1?5% fracture elongation to a maximum level of 1240 MPa tensile strength and 13% elongation. The matrix of welding parameters in Fig. 6b illustrates the range of revolution pitches that created the most favourable results. Tensile strengths exceeding 1175 MPa and maximum elongations higher than 8% were achieved at the same time for samples with a revolution pitch in a range from 0?25 to 0?4 mm rev21. This corresponded to a joint efficiency of more than 85%. In terms of productivity, it is worth mentioning that this criterion was also reached for a high welding speed of 500 mm min21. The optimal tensile properties were found for a tool rotation speed of 800 rev min21 and a welding speed of 200 mm min21 (0?25 mm rev21), resulting in a UTS of 1240 MPa and a maximum elongation of 13%. That equalled a joint efficiency of 90% according to strength and ,50% of the fracture elongation of the parent metal. For the optimised parameters, fracture occurred in HAZ on RS, as for all samples exceeding 85% joint efficiency. All the other samples fractured in the region of the SZ with or without previous necking in HAZ-RS or SZ. Fracture or necking on the AS has not been observed in any case. The hardness of the PM equalled about 315 HV0?1 in layers of SUS 301 and 550 HV0?1 in layers of SUS 420J2. Figure 7 shows the hardness profile across the weld in middle layers of both austenitic SUS 301 and martensitic SUS 420J2 for the optimised parameters with a revolution pitch of 0?25 mm rev21, as well as the hardness profile in the SUS 420J2 layers of a low heat input weld (0?55 mm rev21) and a high heat input weld (0?09 mm rev21). For the optimised parameters, the following observations could be deduced. For the austenitic layer, no significant deviation from the hardness level of the PM could be detected. A slight decrease in hardness towards the weld centre with a minimum of 259 HV0?1 was observed, no differences between the AS and RS were noticed. The maximum

6 a summary of results of tensile tests for all used parameter sets and b matrix of welding parameters and revolution pitches: filled diamonds refer to joint efficiency higher than 85% and fracture elongation higher 8%, and empty diamonds indicate tensile properties below this level


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7 Hardness profiles across weld for different heat inputs

measured hardness in austenitic zones equalled 321 HV0?1. A similar hardness pattern in austenitic layers was observed for all other revolution pitches. In contrast, the hardness profile in the martensitic layer showed strong deviations from the hardness of the PM. Directly beside the weld seam, the hardness of the PM was reached, whereas in the HAZ, a well pronounced hardness drop could be detected on both AS and RS. The minimum measured hardness was similar on both sides, equalling 339 HV0?1 on AS and 354 HV0?1 on RS, and also the width of HAZ did not differ significantly from one side to the other, whereas HAZ slightly enlarged with decreasing revolution pitch. However, asymmetry was observed, that is, HAZ on RS was closer to the weld centre than HAZ on AS. In the TMAZ, the hardness rose steadily from HAZ towards SZ on both sides, whereas the increase was more steady and quick on AS. The


maximum hardness was measured within the SZ equalling a value of 687 HV0?1. By comparing the hardness profiles of the SUS 420J2 layers for different heat inputs in Fig. 7, a rather systematic behaviour could be observed. Apparently, the HAZ was shifted further away from the weld centre with increasing heat input, whereas in any case, the HAZ on RS was closer to the centre than on AS. This asymmetry could be explained by comparing the hardness profiles in Fig. 7 to the temperature cycles in Fig. 4. The hardness profile for the highest heat input together with the correlating temperature cycles pointed out that the critical temperatures that led to the formation of the soft HAZ must have been ,750uC, since these temperatures were measured in HAZ on AS at 10 mm distance from the weld centre. In any other case, those high temperatures could not be observed. Thus, all the other HAZs were shifted towards the weld centre, to locations where the critical temperatures could be reached. Consequentially, the HAZ on RS for the lowest heat input was the one closest to the weld centre, and the HAZ on AS for the highest heat input was the one farthest away from the weld centre. Furthermore, it could be noticed that the average hardness in SZ for the highest heat input was only 575 HV0?1 compared to 650 HV0?1 for the lower heat input. It was considered that a lower cooling rate for high heat input welds was responsible for this circumstance.

Metallography Low heat input

Figure 8 shows an optical macrograph across the weld seam of a low heat input weld with a revolution pitch of 0?55 mm rev21 as well as detailed SEM images and an

a cross-section; b SEM image of pores at SZ-RS; c SEM image incomplete root; d pores at SZ-AS; e fractured sample; f SEM image fracture surface; g LOM of fracture path 8 Microstructure of lowest heat input weld


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a cross-section; b, c EBSD image of pores; d EBSD image of root; e fractured sample; f SEM image of fracture surface; g LOM of fracture path 9 Microstructure of highest heat input weld

illustration of the fracture surface. It was observed that the deterioration of the layered structure was limited to a narrow region within SZ, while the laminated composition was sustained in all other weld regions. For low heat input welds, a large number of pores could be detected in SZ, among which the largest ones had a diameter of more than 10 mm. The most pronounced pores in SZ on RS are represented in Fig. 8b, while Fig. 8d shows the largest pore on AS in SZ. Moreover, incomplete consolidation at the root was observed, as demonstrated in Fig. 8c. However, the observed defects were not considered as related specifically to the FSW of multilayered steel; rather, their presence was attributed to the low heat input, since similar observations have been reported for different materials.23 Figure 8e–g illustrates the fracture behaviour of low heat input welds, whereas Fig. 8e shows a macrophoto of the fractured weld, Fig. 8f shows the SEM image of the fracture surface, while Fig. 8g demonstrates the fracture path in transverse direction by means of optical microscopy. From the SEM image of the fracture surface, it could be deduced that weld defects occurred all along the weld seam within SZ. However, since fracture occurred directly at the centre of SZ without necking, as illustrated in Fig. 8e and g, the phenomenon of brittle fracture for low heat input welds was attributed to the presence of welding defects in the weld centre, whereas it was assumed that fracture initiated from the incomplete root. High heat input

Figure 9 shows the microstructure of a high heat input weld with a revolution pitch of 0?09 mm rev21. Apparently, the cross-section of the high heat input weld in Fig. 9a showed

a considerably wider SZ and significantly higher deterioration of the laminated structure than the cross-section of the low heat input weld shown in Fig. 8a. As shown in Fig. 9d, even the layers at the root showed severe deformation. However, even for high degrees of deterioration in SZ, the layers were basically identifiable, and no random distribution of fcc and bcc grains was noticed, which means that the structure was basically maintained, even in the weld nugget. Small pores were detected in regions of highly distorted layers, as illustrated in Fig. 9b and c. In addition, as the revolution pitch decreased, unfavourable thinning effects within SZ emerging from softened material that was pushed away from the weld centre were observed. For high heat input welds, fracture occurred within SZ, whereas slight necking was noticed on the RS of the weld seam, as illustrated in Fig. 9e. In the SEM image of the fracture surface (Fig. 9f), it was observed that even for high degrees of deterioration, the martensitic layers showed a rather ductile fracture behaviour. Figure 9g illustrates the fracture path in transverse direction, which proceeded directly through a region of highly distorted layers. As a consequence, the fracture at modest strength levels for high heat input welds was attributed to the coincidence of two phenomena, namely, structural impairment resulting from the deterioration of the layers and thinning effects within SZ, whereas the latter one was considered most influential. Optimised set of parameters

Figure 10a shows a macrograph of the cross-section of an optimised parameter weld with a revolution pitch of 0?25 mm rev21 resulting from a tool rotation speed of 800 rev min21 and a welding speed of 200 mm min21. While for low heat input welds a large number of weld


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a cross-section; b, c EBSD images of SZ; d, e EBSD images of TMAZ; f fractured sample; g SEM image of fracture surface; h LOM of fracture path 10 Microstructure of optimised parameter weld

defects was detected, for high heat input welds, structural impairment in combination with extensive thinning was observed; neither of those unfavourable effects occurred for the optimised parameters set. In SZ, no pores were detected, and also the root could be considered fully consolidated. Images (EBSD) of the interface between SUS 301 and SUS 420J2 layers in Fig. 10b–e illustrate the microstructural changes in different weld zones in detail. In SZ (Fig. 10b and c), full recrystallisation was observed within SUS 301 layers, resulting in a fully austenitic microstructure without any deformation induced martensite or chromium rich precipitates. Consequentially, the hardness level for SUS 301 was slightly lower in SZ than in PM. The SUS 420J2 layers in SZ

consisted of 100% fresh martensite as a result of full austenitisation followed by subsequent cooling, leading to a considerably higher hardness in SZ compared to PM. Figure 10d and e shows the microstructure in the layer interface in TMAZ, which was modified from microstructure similar to PM (Fig. 10e) to that similar to SZ (Fig. 10d) when approaching the weld centre. This behaviour showed good correlation to the hardness profile in Fig. 7. The microstructural changes in the HAZ of martensitic layers by means of LOM and TEM are demonstrated in Fig. 11. While no significant differences to the microstructure in PM could be deduced from EBSD images, the optical micrograph in Fig. 11a indicated grain boundary precipitation within HAZ,

a LOM of HAZ; b LOM of PM; c TEM image of precipitates in HAZ 11 Microstructure in HAZ of optimised weld


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12 Hardness profile across weld in both austenitic and martensitic layers in as welded condition and PWHT condition for revolution pitch of 0?27

whereas no such indications were found in PM (Fig. 11b). This appraisal was evidenced from TEM investigations, where a high amount of grain boundary precipitates could be detected in HAZ, as illustrated in Fig. 11c. Owing to that, the hardness drop in HAZ apparently originated from the formation of grain boundary precipitates and presumably from recovery effects within the martensitic microstructure due to exposure to elevated temperatures in a region of 750uC, slightly below Ac1. In HAZ, no microstructural modifications were observed in SUS 301 layers. For optimised parameters, fracture occurred at high strength levels in HAZ on RS. Figure 10f shows a macrophoto of the fractured sample, while Fig. 10g demonstrates an SEM image of the fracture surface, and Fig. 10h illustrates the fracture path through HAZ. From microstructural investigations and hardness measurements, it could be clearly deduced that fracture occurred in HAZ due to the formation of a soft, tempered martensitic microstructure in SUS 420J2 layers. As explained, HAZ on RS was closer to the weld centre than on AS in any case. Since process related thinning always occurs and increases as approaching the weld centre, the local reduction of the weld cross-section in HAZ on RS was more pronounced than in HAZ on AS, resulting in higher local stresses and leading to preferred fracture in HAZ-RS.

Post-weld heat treatment The purpose of the PWHT was the removal of the local weakness of the martensitic layers in HAZ by applying a heat treatment consisting of holding at 1000uC for 2 min followed by air cooling. Figure 12 shows a comparison of the hardness across the weld seam for a revolution pitch of 0?27 in both as welded and PWHT conditions. Apparently, the hardness drop in HAZ of the martensitic layer was removed by


PWHT due to the dissolution of the soft tempered martensite and precipitates while holding at 1000uC followed by a subsequent martensitic transformation during air cooling. Consequentially, the minimum hardness in HAZ was remarkably higher in PWHT condition, reaching a value of 508 HV0?1 compared to 322 HV0?1 in as welded condition. Furthermore, the increase in hardness towards the weld centre was reduced by PWHT, resulting in a lower maximum hardness within SZ of only 607 HV0?1 compared to 668 HV0?1 in the as welded condition. This observation was considered a consequence of the lower cooling rate during air cooling compared to the cooling rate during the FSW process. In austenitic layers, the hardness was, in general, slightly higher in the as welded condition, which might be attributed to the reduction of deformation induced martensite during PWHT. Table 4 gives an overview of the welding parameters and the resulting tensile properties with and without PWHT. Apparently, by applying PWHT, tremendous improvement of the tensile properties was achieved. The UTS could be enhanced to 1310 MPa, which corresponds to a joint efficiency of .95%. Moreover, the maximum elongation was improved significantly to 22%, equalling almost 85% of the PM. While fracture occurred in HAZ on RS for samples in as welded condition, PWHT samples fractured either in PM well away from the weld or in SZ. It is considered that after removing the soft HAZ, the potentially weakest part of the weld, therewith the strength determining factor is SZ. However, it was demonstrated that with appropriate welding parameters, the strength level of SZ was similar to that of the parent metal, resulting in almost identical mechanical properties. Figure 13 illustrates the fracture behaviour of PWHT samples. In case of fracture in PM (Fig. 13a–c), multiple necking occurred around the fracture zone, as evidenced from Fig. 13a. Furthermore, similar to previous studies,4 partial delamination was observed in several layer interfaces, as illustrated in Fig. 13b and c. However, in case of fracture in SZ (Fig. 13d–f), no significant delamination could be detected. Instead, Fig. 13c–f shows a strongly mixed area located directly at the centre of the fracture path.

Conclusions High strength multilayered steel sheets consisting of 15 alternating layers of SUS 301 austenitic stainless steel (eight layers) and SUS 420J2 martensitic stainless steel (seven layers) with an initial thickness of 1?2 mm were welded in butt configuration using FSW. Temperature measurement was carried out, the mechanical properties were assessed and the microstructure was analysed using LOM, SEM, EBSD and TEM. Furthermore, the influence of a PWHT on the weld properties was evaluated. The conclusions can be summarised as follows.

Table 4 Overview of PWHT welding parameters and tensile properties As welded condition

PWHT

Tool rotation speed/rev min21

Welding speed/ mm min21

Revolution pitch/ mm rev21

UTS/MPa

Maximum elongation/%

UTS/MPa

Maximum elongation/%

550

150

0?27

1210

8

1310

22


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a–c fracture in PM; d–f fracture in SZ 13 Fracture behaviour of PWHT samples

1. Friction stir welding of high strength multilayered steel sheets was accomplished successfully with a tool consisting of a sintered matrix of 92% tungsten carbide and 8% cobalt. With appropriate welding parameters, sound welds with a fully consolidated root could be produced. The optimised tensile properties were found for a tool rotation speed of 800 mm rev21 and a welding speed of 200 mm min 21 , resulting in a UTS of 1240 MPa and a maximum elongation of 13%. This corresponds to a joint efficiency of 90%. 2. The hardness across the weld seam did not change significantly for layers of SUS 301. For layers of SUS 420J2, the hardness profile showed a pronounced hardness drop in HAZ and a steady increase in hardness towards the weld centre, whereas the maximum hardness in SZ significantly exceeded that of PM. Asymmetry was observed, whereas HAZ on RS was systematically closer to the weld centre than HAZ on AS. The critical temperature for the formation of the weak HAZ was in the region of 750uC, slightly below Ac1. 3. Low heat input welds showed pores in SZ and an unconsolidated root, which led to fracture at low strength levels directly within SZ. High heat input welds showed maximum deterioration in the laminated structure combined with considerable thinning, which resulted in modest tensile properties and fracture in SZ. Welds with optimised parameters fractured in HAZ on RS due to the softening effects in HAZ. Fracture occurred on RS due to more pronounced thinning, leading to higher local stresses. 4. The SUS 301 layers showed a microstructure consisting of mainly austenite, some deformation induced martensite and chromium rich precipitates in PM and HAZ, which modified to a fully austenitic microstructure as approaching SZ due to recrystallisation. The SUS 420J2 layers showed a microstructure consisting of mainly martensite and some ferrite in PM, tempered martensite with grain boundary precipitates and some ferrite in HAZ and a fully martensitic microstructure within SZ.

5. By applying PWHT, the soft HAZ was removed. For welds with a strong SZ, a joint efficiency of 95% and a fracture elongation of 85% of the PM could be achieved. In those cases, fracture occurred either within SZ or in PM after considerable necking.

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