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St52 steel composite produced by friction stir lap welding
Materials Science & Engineering A 772 (2020) 138775
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Microstructure and mechanical properties of IF/St52 steel composite produced by friction stir lap welding Mohammad Reza Roodgari, Roohollah Jamaati *, Hamed Jamshidi Aval Department of Materials Engineering, Babol Noshirvani University of Technology, Shariati Ave., Babol 47148–71167, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Laminated steel composite Low carbon steels Friction stir lap welding Microstructure Mechanical properties
In the present research, a new composite including interstitial free (IF) steel as the outer layer and St52 as the inner layer was fabricated by friction stir lap welding (FSLW). The microstructures and mechanical properties of initial and composite samples were investigated. It was found that the quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min was excellent due to the higher strain rate. It was observed that the microstructure of composite samples consisted of fine equiaxed ferrite, Widmanstatten ferrite, martensite, grain boundary ferrite, and aggregates of ferrite þ cementite. The increase of tool traverse speed from 40 to 70 and 100 mm/min at a constant tool rotation speed of 1000 rpm decreased the cooling time leading to a higher cooling rate and the formation of martensite phase. The grain size of the composite produced by the traverse speed of 40 mm/min in the first and second HAZ regions was larger than that of the 70 mm/min due to the higher heat input and therefore lower cooling rate of the sample welded by 40 mm/min. The hardness of the center of the stirred zone (SZ) was higher than that of the advancing and retreating sides. The St52 layer of composites showed the microhardness value of ~190 HV whereas the initial St52 sheet had a microhardness around 160 HV. The composite produced by the traverse speed of 70 mm/min had the highest yield and ultimate tensile strength. This high strength achieved without a remarkable decrease in elongation due to the formation of a heterogeneous microstructure. The fracture surfaces indicated that the bonding between IF and St52 steels in the composites produced by traverse speeds of 70 and 100 mm/min was high enough strong to maintain integrity up to a high level of deformation during the tensile test. Finally, with increasing the tool traverse speed, the size and depth of dimples decreased.
1. Introduction Interstitial free (IF) steel and low carbon steels widely used in the automotive industry due to their properties [1–6]. On the one hand, the IF steel exhibits excellent formability, surface quality, and fatigue resistance but low strength. On the other hand, the low carbon steels have low cost, moderate strength, and good ductility, but the low quality of the surface [4–8]. Therefore, a new approach is essential to improve the surface quality of low carbon steels and increase the strength of the interstitial free steel without scarifying its excellent cold formability. It was reported that various techniques such as explosive welding [9–11], roll bonding [12–14], diffusion bonding [15–17], and friction stir lap welding (FSLW) [18–22] can be used to produce multi-layered composites. Park et al. [7] fabricated IF/TWIP/IF steel sheets by hot roll bonding. The interfaces between TWIP and IF steels were well bonded without pores or voids. The fabricated sheets could cover a wide
range of yield strength (YS) (320–498 MPa), ultimate tensile strength (UTS) (545–878 MPa), and total elongation (TE) (48–54%) via con trolling the thickness of TWIP and IF steels. Gladkovsky et al. [9] investigated the microstructure and mechanical properties of pure copper/AISI 1020 steel/pure copper composites fabricated by explosive welding. The welded multi-layered composite had a UTS (440 MPa) higher than the pure copper (240 MPa), but lower than the low carbon steel (495 MPa). On the other hand, the total elongation (TE) of explo sively welded join (42%) was lower than the pure copper (56%) but higher than the AISI 1020 low carbon steel (35%). Chen et al. [15] investigated the effect of diffusion bonding on the microstructure and mechanical behavior of low-carbon ferritic/martensitic steel. The tensile properties of the diffusion bonded joint (560 MPa) was better than that of the initial sample (535 MPa), since the martensite lath near the diffusion bonding joint was finer than the matrix. In addition, the impact toughness of the diffusion bonded joint at room temperature was slightly
* Corresponding author. E-mail address: [email protected] (R. Jamaati). https://doi.org/10.1016/j.msea.2019.138775 Received 30 September 2019; Received in revised form 30 November 2019; Accepted 3 December 2019 Available online 6 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Table 1 The chemical composition of the IF and St52 steels (in wt. %). IF steel St52
C
Si
Mn
Cr
Ni
Mo
P
S
Ti
Cu
Fe
0.002 0.16
0.01 0.13
0.14 0.44
– 0.04
0.018 0.03
– 0.01
0.01 0.01
0.02 0.01
0.055 0.002
0.01 0.04
Bal. Bal.
best knowledge, the effect of FSLW on the microstructure and me chanical properties of low carbon steel covered by IF steel has not been studied yet. This surface modification via fabrication of two-layered composite combines advantages of low carbon steel (moderate strength and low cost) and interstitial free steel (excellent surface quality and formability). The purpose is to produce an IF steel layer without defects and with metallurgical bonding to the St52 substrate. 2. Experimental procedure Interstitial free steel sheet of 0.7 mm thickness and AISI 1016 (St52) steel sheet of 2 mm thickness was used as the coating and the substrate materials, respectively. The chemical compositions of steels are pre sented in Table 1. The samples with the dimensions of 150 mm ⨯ 50 mm were cut from the as-received sheets. To achieve a high quality of bonding between IF and St52 steels, it is essential to clean the surfaces to be bonded. Therefore, the surfaces were ground by SiC paper with 120 grit and then cleaned by ethanol. Finally, the IF and St52 steels were stacked and bonded together with friction stir lap welding. A tungsten carbide tool with a pin diameter of 6 mm, pin height of 0.5 mm, shoulder diameter of 20 mm, and tool tilt angle of 3� was used. The FSLW was conducted with a constant tool rotation rate of 1000 rpm and three different tool traverse speeds of 40, 70, and 100 mm/min. The schematic of lap weld of IF and St52 steel sheets are shown in Fig. 1. The samples for optical microscopy (OM) were cross-sectioned on the two-layered composite perpendicular to the welding direction as indicated in Fig. 1. The samples were mounted, ground with SiC papers, polished with alumina suspension, and then etched in 2% Nital solution. The Vickers microhardness test was carried out (Fig. 1) on KOOPA MH3 device under a load of 100 g for 10 s dwell time. Tensile samples of 12 mm gauge length and 3 mm gauge width were machined parallel to
Fig. 1. The schematic illustration of the FSLWed two-layered composite and the position of the samples.
lower than that of the initial sample. Recently, Argade et al. [23] employed friction stir lap welding to coating AISI 316 L austenitic stainless steel on an AISI 1018 plain carbon steel substrate. After the FSLW, the average grain sizes of stainless steel and low carbon steel decreased to 3 μm and 2 μm, respectively. The hardness of FSLWed material on the AISI 316 L austenitic stainless steel side and AISI 1018 low carbon steel side was 250 HV and 200 HV, respectively. The welded joint had YS, UTS, and TE of 434 MPa, 592 MPa, and 25%, respectively, which was lower than austenitic stainless steel. However, the FSLWed joint exhibited higher UTS than low carbon steel (519 MPa). To overcome the shortcomings of IF steel and low carbon steels, production of IF steel coating on the low carbon steel substrate by friction stir lap welding considered in the present work. To the author’s
Fig. 2. The microstructures of the initial (a,b) IF and (c,d) St52 steel sheets at two different magnifications. 2
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Fig. 3. The interface microstructure for the weld produced by 40 mm/min.
Fig. 4. The interface microstructure for the weld produced by 70 mm/min.
the FSLW direction as shown in Fig. 1. Meanwhile, tensile samples with the same size were machined from the initial IF and St52 steel sheets. The tensile tests were carried out using a SANTAM STM-250 testing machine under a crosshead speed of 1 mm/min at ambient temperature. Total elongation measured as the difference in gauge length before and after the test. Finally, the fracture surfaces of initial and welded samples were observed by scanning electron microscopy (SEM) (FEI/Quanta model).
iron oxide, more intermixing, and increasing the elemental diffusion rate, which helps to eliminate the voids. It should be noted that the composites produced by tool traverse speeds of 40 and 100 mm/min had smooth interfaces. On the contrary, for the traverse speed of 70 mm/min, the interface between interstitial free steel and low carbon steel was diffuse. This result indicates that the composite fabricated via the tool traverse speed of 70 mm/min had a higher rate of elemental diffusion and more material intermixing during FSLW. This can be due to the generation of high enough strain rate and temperature during 70 mm/min traverse speed. However, the strain rate of 40 mm/min traverse speed and the temperature of 100 mm/min were not high. Figs. 4 and 5 show the material intermixing at the center of the stirred zone (SZ) is much more than advancing and retreating sides. This result indicates that vertical material flow occurred from the low carbon steel to the interstitial free steel, or vice-versa. The microstructures of interfaces for the tool traverse speeds of 40, 70, and 100 mm/min at higher magnification are shown in Figs. 6–8, respectively. As seen, the microstructures mainly consist of fine equi axed ferrite, Widmanstatten ferrite, grain boundary ferrite, aggregates of ferrite þ cementite, and martensite, which was consistent with the researches of low carbon steels [24–28]. The presence of Widmanstatten ferrite and martensite indicates that austenite formed during FSLW upon heating. The maximum temperature during friction stir lap welding is related to tool rotation and traverse speeds as given in the following equation [23]: � �α Tmaximum ω2 ¼K (1) Tmelting 2:362υ � 10000
3. Results and discussion 3.1. Microstructure The optical micrographs of as-received IF and St52 steels at two different magnifications are depicted in Fig. 2. As seen, the IF steel has a single-phase ferritic microstructure with an average grain size of 21 μm. On the other hand, the microstructure of St52 consists of ferrites with a mean grain size of 12 μm and about 19% fine pearlite grains. Figs. 3–5 show the weld interface for the tool traverse speeds of 40, 70, and 100 mm/min. From Fig. 3, there is a narrow dark line along with the interface between IF and St52 steels. The interface produced by the traverse speed of 40 mm/min indicates partial microstructural discon tinuity and the presence of a few microvoids or iron oxides inside the narrow dark line. The iron oxide may be created by surface oxidation of IF and St52 steels during the FSLW. On the other hand, a defect-free interface with intermixing between interstitial free steel and low car bon steel is clearly visible from Figs. 4 and 5. The quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min is excellent since no deterioration occurred around the interface. This can be attributed to the higher strain rate during FSLW with the traverse speed of 70 and 100 mm/min than that of 40 mm/min. The higher strain rate increases the material flow around the pin leading to easier breaking of
where Tmaximum is the estimated maximum temperature (� C) during the FSLW, Tmelting is the melting temperature (� C) of the base metals, ω is the tool rotation speed (rpm), υ is the tool traverse speed (mm/min), and K and α are constants (K ¼ 0.7 and α ¼ 0.05) [23]. Considering the tool 3
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Fig. 5. The interface microstructure for the weld produced by 100 mm/min.
rotation speed of 1000 rpm, the estimated maximum temperatures in the IF and St52 steels for three different tool traverse speeds are presented in Table 2. The temperature range for samples during FSLW was between 979 and 1074 � C. These temperatures lead to the transformation of ferrite and pearlite phases of plain carbon steel to the austenite phase. Consequently, during cooling, the Widmanstatten ferrite and martensite can be created depending on the cooling rate. On the other hand, the formation of fine equiaxed ferrite can be attributed to the occurrence of dynamic recrystallization (DRX) in the austenite phase during FSLW. This produces the small austenite grains and then during the cooling cycle, the fine equiaxed ferrite can be formed in the prior austenite grain boundary. It should be noted that there is no martensite in the microstructure of the composite produced by the traverse speed of 40 mm/min (see Fig. 6). However, the martensite is formed for the tool traverse speeds of 70 and 100 mm/min (see Figs. 7 and 8). The heat gradient is responsible for the differences in the formation of martensite. As shown in previous researches [29,30], with increasing the tool traverse speed, the heat gradient from the maximum temperature of the stirred zone increases. The increase of tool traverse speed from 40 to 70 and 100 mm/min at a constant tool rotation speed of 1000 rpm decreases the cooling time leading to higher cooling rate and formation of the martensite phase. The optical micrographs representing the microstructures in the top, middle, and bottom of the three distinct regions in the stirred zone (retreating side, the center of the stirred zone, and advancing side) for the tool traverse speeds of 40, 70, and 100 mm/min are shown in Figs. 9–11, respectively. Also, the Fe–Fe3C phase diagram with different microstructures and the approximately related temperatures for the composite samples are depicted in Fig. 12. These figures indicate the heterogeneous microstructure in the composite samples. According to the pin height of 0.5 mm, it can be said that the IF steel has only stirred zone during FSLW. As seen in Figs. 9–11, the microstructures of IF steel
Fig. 6. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 40 mm/min.
can be divided into two regions. The region close to the surface of composite has fine ferrite grains (first SZ). However, very large grains observed in the region far from the surface (second SZ). This difference can be attributed to the shear deformation of IF steel caused by the tool movement. In other words, the tool surface shear strain, and therefore, the number of new DRX grains in the first SZ is larger than the second SZ 4
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Fig. 7. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 70 mm/min.
Fig. 8. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 100 mm/min.
of IF steel. Consequently, the grain size of the first SZ is lower than that of the second SZ. As shown in Fig. 12, the produced composites have two different regions including stirred zone and heat-affected zone (HAZ) in the St52. It should be noted that there is no distinct thermomechanically-affected zone (TMAZ) due to allotropic transformation in carbon steel during the
cooling as also reported in previous researches [25,31]. Figs. 9–12 demonstrate two different types of microstructures in the stirred zone of the St52. The SZ for the traverse speed of 40 mm/min mainly consists of Widmanstatten ferrite and fine equiaxed ferrite. On the other hand, the microstructure of the stirred zone of composites produced by 70 and 100 mm/min traverse speeds mainly containing 5
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From Fig. 12, the grain size of the composite produced by the tra verse speed of 40 mm/min in the first and second HAZ regions is larger than that of the 70 mm/min. This result can be attributed to the higher heat input and therefore lower cooling rate of the sample welded via 40 mm/min as compared to the 70 mm/min weld. The higher cooling rate in the traverse speed of 70 mm/min provides a shorter time for the austenite grains to grow above the Ac3 for the first HAZ and between the Ac1 and Ac3 for the second HAZ. Thus, the size of new ferrite and pearlite grains in the first and second HAZ of the sample welded via 70 mm/min are lower than 40 mm/min.
Table 2 The estimated maximum temperatures (� C) for IF and St52 steels. Samples IF steel St52
Traverse speed (mm/min) 40
70
100
1074 1025
1044 997
1026 979
Widmanstatten ferrite and martensite. As seen in Figs. 9–12, there are two different HAZ regions in the composites produced by FSLW. Fig. 12 shows the first HAZ has a unimodal distribution of grain size including fine grains owing to grain refinement. The first HAZ subjected to a maximum temperature above the Ac3. Therefore, the ferrite þ pearlite structure completely transforms into the austenite phase upon heating to just above Ac3. However, since the composite is not heated to a high enough temperature, the mobility of austenite grain boundaries is limited and extensive grain growth cannot occur. Then, during cooling, the austenite grains decompose into fine ferrite and pearlite grains. Fig. 12 illustrates the second HAZ has a bimodal grain size distri bution consisting of coarse and fine grains due to partial grain refine ment. According to the Fe–Fe3C diagram, the second HAZ subjected to a maximum temperature between the Ac1 and Ac3 in the αþγ region. During FSLW, first, the pre-existed ferrite and pearlite partially trans formed to the austenite. During subsequent cooling, the fine ferrite and pearlite grains nucleated at the austenite grain boundaries. The coarse proeutectoid ferrite grains remained almost unchanged. Therefore, a bimodal grain size distribution observed in the partial grain refinement region.
3.2. Mechanical properties The Vickers microhardness profiles across the vertical line (see Fig. 1) in the retreating side, the center of the SZ, and the advancing side for the traverse speeds of 40, 70, and 100 mm/min are shown in Fig. 13. In the IF steel, the hardness value of the first SZ is higher than that of the second SZ because the first SZ had smaller grains as shown in Figs. 9–11. In fact, the second SZ of IF steel has the lowest microhardness in the microhardness profiles. On the other hand, in the St52 steel, the hard ness of the regions close to the interface is higher and with increasing the distance from the interface, the hardness decreases. In the regions close to the St52 surface (bimodal area), a microstructure with coarser grain size observed for all samples (see Figs. 9–11). This result can be attrib uted to the formation of martensite, Widmanstatten ferrite, and fine equiaxed ferrite in the regions close to the IF/St52 interface as explained in the earlier section. From Fig. 13, it should be noted that the hardness of the center of the SZ is higher than that of the advancing and retreating sides. This can be related to the effect of tool pin on the shear
Fig. 9. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 40 mm/min. 6
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Fig. 10. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 70 mm/min.
deformation during FSLW. This effect is more severe for the regions close to the center of the stirred zone. Therefore, the fraction of martensite and Widmanstatten ferrite, as well as the grain sizes of ferrite and pearlite, is larger in the center of the SZ. For comparison, the microhardness profiles of the center of the stirred zone for the composites produced by the three different traverse speeds as well as initial IF and St52 steels are depicted in Fig. 14. The St52 layer of composites shows the microhardness value of ~190 HV whereas the initial St52 sheet has the microhardness around 160 HV. This result is owing to the formation of Widmanstatten ferrite, martensite, and fine ferrite and pearlite grains. On the other hand, the IF steel layer of composites exhibits the microhardness between 130 to 200 HV (for the traverse speeds of 40 and 100 mm/min) whereas the initial IF steel sheet has the microhardness value of ~100 HV. This is due to the grain refinement in the first SZ of IF steel. It should be noted that the hardness of the IF steel layer of composite processed by 70 mm/min traverse speed is extraordinary because of the intermixing between IF and St52 steels and formation of martensite as shown in Figs. 7(b) and 10. A sharp increase in the microhardness (around 290 HV) observed at the interface. Beyond this, the Vickers microhardness value gradually dropped to 195 HV on both layers. From Fig. 14, the hardness of the composite processed by the traverse speed of 40 mm/min is lower than that of 70 and 100 mm/min. This can be attributed to the formation of coarser ferrite grains and no martensite in the sample produced via 40 mm/min traverse speed as explained in the previous section. The tensile engineering stress-strain curves of the initial and pro cessed samples are depicted in Fig. 15, and corresponding YS, UTS, and TE extracted from the curves are given in Fig. 16. From Fig. 15, the yield point phenomenon observed in the stress-strain curves of the composite
samples except for the traverse speed of 70 mm/min. This can be attributed to the formation of martensite phase during FSLW. In the composite samples, with an increase in tool traverse speed from 40 to 70 mm/min, the YS and UTS increased from 381.8 and 463.1 MPa to 433.3 and 543.9 MPa, respectively, while total elongation decreased from 40.0% to 31.7%. As mentioned before, this result is related to the for mation of martensite and finer ferrite grains in the composite processed by the traverse speed of 70 mm/min. With further increasing the tra verse speed to 100 mm/min, the YS, UTS, and TE decreased to 390.4 MPa, 495.0 MPa, and 30.6%, respectively. The yield strength of the welds is higher than that of the base materials, while the total elongation of the composites is less than that of the initial samples. This is due to the presence of Widmanstatten ferrite, martensite, and fine equiaxed ferrite in the microstructure of composites. The composite produced by the traverse speed of 70 mm/min has the highest yield and ultimate tensile strength, which is consistent with the microhardness measurements in the center of the stirred zone. This high strength achieved without a remarkable decrease in elongation due to the formation of a heteroge neous microstructure. The good combination of strength and ductility is the most important factor in the applications of the engineering materials. It is important to note that the stress-strain behavior of IF and St52 steels during the tensile test is different. A schematic illustration of the elastoplastic deformation mechanism at the IF/St52 interface is shown in Fig. 17. As seen, there are three distinct steps in the stress-strain curves of composite samples. In step one, both IF and St52 steels are in elastic deformation. In step two, the interstitial free steel starts plastic deformation while the low carbon steel is under elastic deformation. In step three, both IF and St52 steels are in plastic deformation. Therefore, 7
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Fig. 11. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 100 mm/min.
as indicated in Fig. 17, different elastoplastic properties of the intersti tial free and low carbon steels produce a dislocation density gradient in both layers during the tensile test. The relationship between the ultimate tensile strength and total elongation is usually expressed by the so-called banana curve for various steels as shown in Fig. 18. In the past few decades, numerous researches on grain refinement using severe plastic deformation (SPD) techniques have been performed to obtain high strength interstitial free steel [2,4,6, 32–36]. Although the yield and ultimate tensile strength of interstitial free steel are remarkably increased, the problem of decreased elongation limits the application of such high strength steel as forming and struc tural materials. For the more industrial applications of interstitial free steel, the low strength needs to be solved without sacrificing the elon gation. The present study indicates that a good combination of strength and ductility could be obtained through FSLW, which is mainly due to the formation of a heterogeneous microstructure in both layers. As seen in Fig. 18, the produced composite lies in the high strength steel cate gory, while the initial samples (IF and St52 steels) were low strength steels. Therefore, the FSLW exhibits to be a promising process to fabri cate a high strength steel composite.
IF and St52 steels. The composite produced by the traverse speed of 40 mm/min failed by separation of the layers, at the interface, as indicated in Fig. 20(a). This result shows that the bonding between IF and St52 steels processed by FSLW was not enough strong to maintain integrity up to a high level of deformation during the tensile test. On the other hand, the delamination between the interstitial free and low carbon steels is not observed in the FSLW with the traverse speeds of 70 and 100 mm/ min, and the subsequent tensile test. As shown in Figs. 3–5, unlike the composite welded by 40 mm/min traverse speed, there was no sign of microstructural discontinuity between the IF and St52 steels for the traverse speeds of 70 and 100 mm/min. Consequently, the layers of the interstitial free and low carbon steels remained well bonded during the tensile test. From Fig. 20, obvious necking observed on the fracture surface of the composite samples produced by FSLW indicating a substantial plastic deformation before the failure. The fracture surfaces of both layers in the produced composites exhibit many dimples indicative of ductile fracture mode without any cleavage or quasi cleavage fracture mode. By comparing the fracture surfaces of the initial St52 and IF steels (Fig. 19) and the St52 and IF layers in the composites (Fig. 20), two different features can be observed. Firstly, the size of dimples in the fracture surfaces of initial materials is more uniform than that of the St52 and IF layers in the FSLWed composites. This is due to the formation of het erogeneous microstructure in both layers after the FSLW. Secondly, the IF and St52 layers in the composites exhibit smaller dimples than the initial materials indicating lower ductility, which is consistent with the tensile test results. By increasing the tool traverse speed, the size of dimples decreased. The size of dimples for the traverse speed of 40 mm/min is greater than
3.3. Fractography Fig. 19 presents the fracture surfaces of the initial IF and St52 steels after the tensile test. Numerous and deep dimples can be observed on the fracture surfaces of the interstitial free and low carbon steels, which indicates a completely ductile failure. Fig. 20 shows the fracture surfaces of the two-layered steel com posites after the tensile test. The red arrows indicate the interfaces of the 8
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Fig. 12. The Fe–Fe3C phase diagram with different microstructures and the approximately related temperatures for the composite samples produced by the traverse speeds of 40 and 70 mm/min.
that of 70 and 100 mm/min. This result can be attributed to the larger grain size of the composite produced by the traverse speed of 40 mm/ min as indicated in Figs. 9–12. By increasing the grain size, the fraction of grain boundaries decreased, the number of dimple nucleation reduced, and the size of dimples increased. On the other hand, with increasing the tool traverse speed, the depth of dimples decreased. The size of dimples for the traverse speeds of 70 and 100 mm/min is smaller than that of 40 mm/min. This is due to the formation of a hard phase (i. e., martensite) in the microstructure of the composite produced by the traverse speeds of 70 and 100 mm/min as shown in Figs. 9–12. Friction stir lap welding of interstitial free and low carbon steels can prove an excellent method to produce well-bonded two-layered steel composite. The FSLWed IF/St52 steel composites showed good tensile strength and elongation due to the formation of heterogeneous micro structure. These results allow the extended use of interstitial free steel in the automotive industry.
3. The increase of tool traverse speed from 40 to 70 and 100 mm/ min at a constant tool rotation speed of 1000 rpm decreased the cooling time leading to a higher cooling rate and the formation of martensite phase. 4. The heterogeneous microstructure was observed in the composite samples produced by FSLW. 5. The grain size of the composite produced by the traverse speed of 40 mm/min in the first and second HAZ regions was larger than that of the 70 mm/min due to the higher heat input and therefore lower cooling rate of the sample welded via 40 mm/min. 6. The hardness of the center of the SZ was higher than that of the advancing and retreating sides. This could be attributed to the effect of tool pin on the shear deformation during FSLW. 7. The St52 layer of composites showed the microhardness value of ~190 HV whereas the initial St52 sheet had the microhardness around 160 HV. This was owing to the formation of Widman statten ferrite, martensite, and fine ferrite and pearlite grains. 8. The composite produced by the traverse speed of 70 mm/min had the highest YS and UTS. This high strength achieved without a remarkable decrease in elongation due to the formation of a heterogeneous microstructure. 9. The fracture surfaces indicated the bonding between IF and St52 steels in the composites produced by traverse speeds of 70 and 100 mm/min was high enough strong to maintain integrity up to a high level of deformation during the tensile test. 10. With increasing the tool traverse speed, the size of dimples decreased. The size of dimples for the traverse speed of 40 mm/ min was greater than that of 70 and 100 mm/min. This result could be due to the larger grain size of the composite produced by the traverse speed of 40 mm/min.
4. Conclusions In the present study, two-layered steel composite consisted of IF and St52 steels sheets was successfully fabricated by friction stir lap welding. The main conclusions were drawn as follows: 1. The quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min was excellent due to the higher strain rate during FSLW with the traverse speed of 70 and 100 mm/min than that of 40 mm/min. 2. The microstructures of composite samples mainly consisted of fine equiaxed ferrite, Widmanstatten ferrite, grain boundary ferrite, aggregates of ferrite þ cementite, and martensite.
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Fig. 13. The Vickers microhardness profiles across the vertical line in the retreating side, the center of the SZ, and the advancing side for the traverse speeds of (a) 40, (b) 70, and (c) 100 mm/min.
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Fig. 14. The microhardness profiles of the center of the stirred zone for the composites produced by the three different traverse speeds as well as initial IF and St52 steels.
Fig. 15. The tensile engineering stress-strain curves of the initial and processed samples.
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Fig. 16. (a) Yield strength, (b) ultimate tensile strength, and (c) total elongation of the initial and processed samples.
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Fig. 17. Schematic illustration of the elastoplastic deformation mechanism at the IF/St52 interface during the tensile test.
Fig. 18. The tensile strength-elongation relationship in steels.
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Fig. 19. The fracture surface of (a) IF and (b) St52 steels.
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Fig. 20. The fracture surfaces of the composites produced by the traverse speeds of (a) 40, (b) 70, and (c) 100 mm/min.
Data availability
References
The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.
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Author contribution statement Mohammad Reza Roodgari: Investigation, Resources, Writing Original Draft. Roohollah Jamaati: Conceptualization, Methodology, Writing Review & Editing, Supervision. Hamed Jamshidi Aval: Conceptualization, Methodology, Writing Review & Editing, Supervision. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements The authors acknowledge the funding support of Babol Noshirvani University of Technology through Grant program No. BNUT/393044/98 and BNUT/370167/98.
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructure and mechanical properties of IF/St52 steel composite produced by friction stir lap welding Mohammad Reza Roodgari, Roohollah Jamaati *, Hamed Jamshidi Aval Department of Materials Engineering, Babol Noshirvani University of Technology, Shariati Ave., Babol 47148–71167, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Laminated steel composite Low carbon steels Friction stir lap welding Microstructure Mechanical properties
In the present research, a new composite including interstitial free (IF) steel as the outer layer and St52 as the inner layer was fabricated by friction stir lap welding (FSLW). The microstructures and mechanical properties of initial and composite samples were investigated. It was found that the quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min was excellent due to the higher strain rate. It was observed that the microstructure of composite samples consisted of fine equiaxed ferrite, Widmanstatten ferrite, martensite, grain boundary ferrite, and aggregates of ferrite þ cementite. The increase of tool traverse speed from 40 to 70 and 100 mm/min at a constant tool rotation speed of 1000 rpm decreased the cooling time leading to a higher cooling rate and the formation of martensite phase. The grain size of the composite produced by the traverse speed of 40 mm/min in the first and second HAZ regions was larger than that of the 70 mm/min due to the higher heat input and therefore lower cooling rate of the sample welded by 40 mm/min. The hardness of the center of the stirred zone (SZ) was higher than that of the advancing and retreating sides. The St52 layer of composites showed the microhardness value of ~190 HV whereas the initial St52 sheet had a microhardness around 160 HV. The composite produced by the traverse speed of 70 mm/min had the highest yield and ultimate tensile strength. This high strength achieved without a remarkable decrease in elongation due to the formation of a heterogeneous microstructure. The fracture surfaces indicated that the bonding between IF and St52 steels in the composites produced by traverse speeds of 70 and 100 mm/min was high enough strong to maintain integrity up to a high level of deformation during the tensile test. Finally, with increasing the tool traverse speed, the size and depth of dimples decreased.
1. Introduction Interstitial free (IF) steel and low carbon steels widely used in the automotive industry due to their properties [1–6]. On the one hand, the IF steel exhibits excellent formability, surface quality, and fatigue resistance but low strength. On the other hand, the low carbon steels have low cost, moderate strength, and good ductility, but the low quality of the surface [4–8]. Therefore, a new approach is essential to improve the surface quality of low carbon steels and increase the strength of the interstitial free steel without scarifying its excellent cold formability. It was reported that various techniques such as explosive welding [9–11], roll bonding [12–14], diffusion bonding [15–17], and friction stir lap welding (FSLW) [18–22] can be used to produce multi-layered composites. Park et al. [7] fabricated IF/TWIP/IF steel sheets by hot roll bonding. The interfaces between TWIP and IF steels were well bonded without pores or voids. The fabricated sheets could cover a wide
range of yield strength (YS) (320–498 MPa), ultimate tensile strength (UTS) (545–878 MPa), and total elongation (TE) (48–54%) via con trolling the thickness of TWIP and IF steels. Gladkovsky et al. [9] investigated the microstructure and mechanical properties of pure copper/AISI 1020 steel/pure copper composites fabricated by explosive welding. The welded multi-layered composite had a UTS (440 MPa) higher than the pure copper (240 MPa), but lower than the low carbon steel (495 MPa). On the other hand, the total elongation (TE) of explo sively welded join (42%) was lower than the pure copper (56%) but higher than the AISI 1020 low carbon steel (35%). Chen et al. [15] investigated the effect of diffusion bonding on the microstructure and mechanical behavior of low-carbon ferritic/martensitic steel. The tensile properties of the diffusion bonded joint (560 MPa) was better than that of the initial sample (535 MPa), since the martensite lath near the diffusion bonding joint was finer than the matrix. In addition, the impact toughness of the diffusion bonded joint at room temperature was slightly
* Corresponding author. E-mail address: [email protected] (R. Jamaati). https://doi.org/10.1016/j.msea.2019.138775 Received 30 September 2019; Received in revised form 30 November 2019; Accepted 3 December 2019 Available online 6 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Table 1 The chemical composition of the IF and St52 steels (in wt. %). IF steel St52
C
Si
Mn
Cr
Ni
Mo
P
S
Ti
Cu
Fe
0.002 0.16
0.01 0.13
0.14 0.44
– 0.04
0.018 0.03
– 0.01
0.01 0.01
0.02 0.01
0.055 0.002
0.01 0.04
Bal. Bal.
best knowledge, the effect of FSLW on the microstructure and me chanical properties of low carbon steel covered by IF steel has not been studied yet. This surface modification via fabrication of two-layered composite combines advantages of low carbon steel (moderate strength and low cost) and interstitial free steel (excellent surface quality and formability). The purpose is to produce an IF steel layer without defects and with metallurgical bonding to the St52 substrate. 2. Experimental procedure Interstitial free steel sheet of 0.7 mm thickness and AISI 1016 (St52) steel sheet of 2 mm thickness was used as the coating and the substrate materials, respectively. The chemical compositions of steels are pre sented in Table 1. The samples with the dimensions of 150 mm ⨯ 50 mm were cut from the as-received sheets. To achieve a high quality of bonding between IF and St52 steels, it is essential to clean the surfaces to be bonded. Therefore, the surfaces were ground by SiC paper with 120 grit and then cleaned by ethanol. Finally, the IF and St52 steels were stacked and bonded together with friction stir lap welding. A tungsten carbide tool with a pin diameter of 6 mm, pin height of 0.5 mm, shoulder diameter of 20 mm, and tool tilt angle of 3� was used. The FSLW was conducted with a constant tool rotation rate of 1000 rpm and three different tool traverse speeds of 40, 70, and 100 mm/min. The schematic of lap weld of IF and St52 steel sheets are shown in Fig. 1. The samples for optical microscopy (OM) were cross-sectioned on the two-layered composite perpendicular to the welding direction as indicated in Fig. 1. The samples were mounted, ground with SiC papers, polished with alumina suspension, and then etched in 2% Nital solution. The Vickers microhardness test was carried out (Fig. 1) on KOOPA MH3 device under a load of 100 g for 10 s dwell time. Tensile samples of 12 mm gauge length and 3 mm gauge width were machined parallel to
Fig. 1. The schematic illustration of the FSLWed two-layered composite and the position of the samples.
lower than that of the initial sample. Recently, Argade et al. [23] employed friction stir lap welding to coating AISI 316 L austenitic stainless steel on an AISI 1018 plain carbon steel substrate. After the FSLW, the average grain sizes of stainless steel and low carbon steel decreased to 3 μm and 2 μm, respectively. The hardness of FSLWed material on the AISI 316 L austenitic stainless steel side and AISI 1018 low carbon steel side was 250 HV and 200 HV, respectively. The welded joint had YS, UTS, and TE of 434 MPa, 592 MPa, and 25%, respectively, which was lower than austenitic stainless steel. However, the FSLWed joint exhibited higher UTS than low carbon steel (519 MPa). To overcome the shortcomings of IF steel and low carbon steels, production of IF steel coating on the low carbon steel substrate by friction stir lap welding considered in the present work. To the author’s
Fig. 2. The microstructures of the initial (a,b) IF and (c,d) St52 steel sheets at two different magnifications. 2
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Fig. 3. The interface microstructure for the weld produced by 40 mm/min.
Fig. 4. The interface microstructure for the weld produced by 70 mm/min.
the FSLW direction as shown in Fig. 1. Meanwhile, tensile samples with the same size were machined from the initial IF and St52 steel sheets. The tensile tests were carried out using a SANTAM STM-250 testing machine under a crosshead speed of 1 mm/min at ambient temperature. Total elongation measured as the difference in gauge length before and after the test. Finally, the fracture surfaces of initial and welded samples were observed by scanning electron microscopy (SEM) (FEI/Quanta model).
iron oxide, more intermixing, and increasing the elemental diffusion rate, which helps to eliminate the voids. It should be noted that the composites produced by tool traverse speeds of 40 and 100 mm/min had smooth interfaces. On the contrary, for the traverse speed of 70 mm/min, the interface between interstitial free steel and low carbon steel was diffuse. This result indicates that the composite fabricated via the tool traverse speed of 70 mm/min had a higher rate of elemental diffusion and more material intermixing during FSLW. This can be due to the generation of high enough strain rate and temperature during 70 mm/min traverse speed. However, the strain rate of 40 mm/min traverse speed and the temperature of 100 mm/min were not high. Figs. 4 and 5 show the material intermixing at the center of the stirred zone (SZ) is much more than advancing and retreating sides. This result indicates that vertical material flow occurred from the low carbon steel to the interstitial free steel, or vice-versa. The microstructures of interfaces for the tool traverse speeds of 40, 70, and 100 mm/min at higher magnification are shown in Figs. 6–8, respectively. As seen, the microstructures mainly consist of fine equi axed ferrite, Widmanstatten ferrite, grain boundary ferrite, aggregates of ferrite þ cementite, and martensite, which was consistent with the researches of low carbon steels [24–28]. The presence of Widmanstatten ferrite and martensite indicates that austenite formed during FSLW upon heating. The maximum temperature during friction stir lap welding is related to tool rotation and traverse speeds as given in the following equation [23]: � �α Tmaximum ω2 ¼K (1) Tmelting 2:362υ � 10000
3. Results and discussion 3.1. Microstructure The optical micrographs of as-received IF and St52 steels at two different magnifications are depicted in Fig. 2. As seen, the IF steel has a single-phase ferritic microstructure with an average grain size of 21 μm. On the other hand, the microstructure of St52 consists of ferrites with a mean grain size of 12 μm and about 19% fine pearlite grains. Figs. 3–5 show the weld interface for the tool traverse speeds of 40, 70, and 100 mm/min. From Fig. 3, there is a narrow dark line along with the interface between IF and St52 steels. The interface produced by the traverse speed of 40 mm/min indicates partial microstructural discon tinuity and the presence of a few microvoids or iron oxides inside the narrow dark line. The iron oxide may be created by surface oxidation of IF and St52 steels during the FSLW. On the other hand, a defect-free interface with intermixing between interstitial free steel and low car bon steel is clearly visible from Figs. 4 and 5. The quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min is excellent since no deterioration occurred around the interface. This can be attributed to the higher strain rate during FSLW with the traverse speed of 70 and 100 mm/min than that of 40 mm/min. The higher strain rate increases the material flow around the pin leading to easier breaking of
where Tmaximum is the estimated maximum temperature (� C) during the FSLW, Tmelting is the melting temperature (� C) of the base metals, ω is the tool rotation speed (rpm), υ is the tool traverse speed (mm/min), and K and α are constants (K ¼ 0.7 and α ¼ 0.05) [23]. Considering the tool 3
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Fig. 5. The interface microstructure for the weld produced by 100 mm/min.
rotation speed of 1000 rpm, the estimated maximum temperatures in the IF and St52 steels for three different tool traverse speeds are presented in Table 2. The temperature range for samples during FSLW was between 979 and 1074 � C. These temperatures lead to the transformation of ferrite and pearlite phases of plain carbon steel to the austenite phase. Consequently, during cooling, the Widmanstatten ferrite and martensite can be created depending on the cooling rate. On the other hand, the formation of fine equiaxed ferrite can be attributed to the occurrence of dynamic recrystallization (DRX) in the austenite phase during FSLW. This produces the small austenite grains and then during the cooling cycle, the fine equiaxed ferrite can be formed in the prior austenite grain boundary. It should be noted that there is no martensite in the microstructure of the composite produced by the traverse speed of 40 mm/min (see Fig. 6). However, the martensite is formed for the tool traverse speeds of 70 and 100 mm/min (see Figs. 7 and 8). The heat gradient is responsible for the differences in the formation of martensite. As shown in previous researches [29,30], with increasing the tool traverse speed, the heat gradient from the maximum temperature of the stirred zone increases. The increase of tool traverse speed from 40 to 70 and 100 mm/min at a constant tool rotation speed of 1000 rpm decreases the cooling time leading to higher cooling rate and formation of the martensite phase. The optical micrographs representing the microstructures in the top, middle, and bottom of the three distinct regions in the stirred zone (retreating side, the center of the stirred zone, and advancing side) for the tool traverse speeds of 40, 70, and 100 mm/min are shown in Figs. 9–11, respectively. Also, the Fe–Fe3C phase diagram with different microstructures and the approximately related temperatures for the composite samples are depicted in Fig. 12. These figures indicate the heterogeneous microstructure in the composite samples. According to the pin height of 0.5 mm, it can be said that the IF steel has only stirred zone during FSLW. As seen in Figs. 9–11, the microstructures of IF steel
Fig. 6. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 40 mm/min.
can be divided into two regions. The region close to the surface of composite has fine ferrite grains (first SZ). However, very large grains observed in the region far from the surface (second SZ). This difference can be attributed to the shear deformation of IF steel caused by the tool movement. In other words, the tool surface shear strain, and therefore, the number of new DRX grains in the first SZ is larger than the second SZ 4
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Materials Science & Engineering A 772 (2020) 138775
Fig. 7. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 70 mm/min.
Fig. 8. The interface microstructure of (a) retreating side, (b) center of the stirred zone, and (c) advancing side for the traverse speed of 100 mm/min.
of IF steel. Consequently, the grain size of the first SZ is lower than that of the second SZ. As shown in Fig. 12, the produced composites have two different regions including stirred zone and heat-affected zone (HAZ) in the St52. It should be noted that there is no distinct thermomechanically-affected zone (TMAZ) due to allotropic transformation in carbon steel during the
cooling as also reported in previous researches [25,31]. Figs. 9–12 demonstrate two different types of microstructures in the stirred zone of the St52. The SZ for the traverse speed of 40 mm/min mainly consists of Widmanstatten ferrite and fine equiaxed ferrite. On the other hand, the microstructure of the stirred zone of composites produced by 70 and 100 mm/min traverse speeds mainly containing 5
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Materials Science & Engineering A 772 (2020) 138775
From Fig. 12, the grain size of the composite produced by the tra verse speed of 40 mm/min in the first and second HAZ regions is larger than that of the 70 mm/min. This result can be attributed to the higher heat input and therefore lower cooling rate of the sample welded via 40 mm/min as compared to the 70 mm/min weld. The higher cooling rate in the traverse speed of 70 mm/min provides a shorter time for the austenite grains to grow above the Ac3 for the first HAZ and between the Ac1 and Ac3 for the second HAZ. Thus, the size of new ferrite and pearlite grains in the first and second HAZ of the sample welded via 70 mm/min are lower than 40 mm/min.
Table 2 The estimated maximum temperatures (� C) for IF and St52 steels. Samples IF steel St52
Traverse speed (mm/min) 40
70
100
1074 1025
1044 997
1026 979
Widmanstatten ferrite and martensite. As seen in Figs. 9–12, there are two different HAZ regions in the composites produced by FSLW. Fig. 12 shows the first HAZ has a unimodal distribution of grain size including fine grains owing to grain refinement. The first HAZ subjected to a maximum temperature above the Ac3. Therefore, the ferrite þ pearlite structure completely transforms into the austenite phase upon heating to just above Ac3. However, since the composite is not heated to a high enough temperature, the mobility of austenite grain boundaries is limited and extensive grain growth cannot occur. Then, during cooling, the austenite grains decompose into fine ferrite and pearlite grains. Fig. 12 illustrates the second HAZ has a bimodal grain size distri bution consisting of coarse and fine grains due to partial grain refine ment. According to the Fe–Fe3C diagram, the second HAZ subjected to a maximum temperature between the Ac1 and Ac3 in the αþγ region. During FSLW, first, the pre-existed ferrite and pearlite partially trans formed to the austenite. During subsequent cooling, the fine ferrite and pearlite grains nucleated at the austenite grain boundaries. The coarse proeutectoid ferrite grains remained almost unchanged. Therefore, a bimodal grain size distribution observed in the partial grain refinement region.
3.2. Mechanical properties The Vickers microhardness profiles across the vertical line (see Fig. 1) in the retreating side, the center of the SZ, and the advancing side for the traverse speeds of 40, 70, and 100 mm/min are shown in Fig. 13. In the IF steel, the hardness value of the first SZ is higher than that of the second SZ because the first SZ had smaller grains as shown in Figs. 9–11. In fact, the second SZ of IF steel has the lowest microhardness in the microhardness profiles. On the other hand, in the St52 steel, the hard ness of the regions close to the interface is higher and with increasing the distance from the interface, the hardness decreases. In the regions close to the St52 surface (bimodal area), a microstructure with coarser grain size observed for all samples (see Figs. 9–11). This result can be attrib uted to the formation of martensite, Widmanstatten ferrite, and fine equiaxed ferrite in the regions close to the IF/St52 interface as explained in the earlier section. From Fig. 13, it should be noted that the hardness of the center of the SZ is higher than that of the advancing and retreating sides. This can be related to the effect of tool pin on the shear
Fig. 9. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 40 mm/min. 6
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Fig. 10. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 70 mm/min.
deformation during FSLW. This effect is more severe for the regions close to the center of the stirred zone. Therefore, the fraction of martensite and Widmanstatten ferrite, as well as the grain sizes of ferrite and pearlite, is larger in the center of the SZ. For comparison, the microhardness profiles of the center of the stirred zone for the composites produced by the three different traverse speeds as well as initial IF and St52 steels are depicted in Fig. 14. The St52 layer of composites shows the microhardness value of ~190 HV whereas the initial St52 sheet has the microhardness around 160 HV. This result is owing to the formation of Widmanstatten ferrite, martensite, and fine ferrite and pearlite grains. On the other hand, the IF steel layer of composites exhibits the microhardness between 130 to 200 HV (for the traverse speeds of 40 and 100 mm/min) whereas the initial IF steel sheet has the microhardness value of ~100 HV. This is due to the grain refinement in the first SZ of IF steel. It should be noted that the hardness of the IF steel layer of composite processed by 70 mm/min traverse speed is extraordinary because of the intermixing between IF and St52 steels and formation of martensite as shown in Figs. 7(b) and 10. A sharp increase in the microhardness (around 290 HV) observed at the interface. Beyond this, the Vickers microhardness value gradually dropped to 195 HV on both layers. From Fig. 14, the hardness of the composite processed by the traverse speed of 40 mm/min is lower than that of 70 and 100 mm/min. This can be attributed to the formation of coarser ferrite grains and no martensite in the sample produced via 40 mm/min traverse speed as explained in the previous section. The tensile engineering stress-strain curves of the initial and pro cessed samples are depicted in Fig. 15, and corresponding YS, UTS, and TE extracted from the curves are given in Fig. 16. From Fig. 15, the yield point phenomenon observed in the stress-strain curves of the composite
samples except for the traverse speed of 70 mm/min. This can be attributed to the formation of martensite phase during FSLW. In the composite samples, with an increase in tool traverse speed from 40 to 70 mm/min, the YS and UTS increased from 381.8 and 463.1 MPa to 433.3 and 543.9 MPa, respectively, while total elongation decreased from 40.0% to 31.7%. As mentioned before, this result is related to the for mation of martensite and finer ferrite grains in the composite processed by the traverse speed of 70 mm/min. With further increasing the tra verse speed to 100 mm/min, the YS, UTS, and TE decreased to 390.4 MPa, 495.0 MPa, and 30.6%, respectively. The yield strength of the welds is higher than that of the base materials, while the total elongation of the composites is less than that of the initial samples. This is due to the presence of Widmanstatten ferrite, martensite, and fine equiaxed ferrite in the microstructure of composites. The composite produced by the traverse speed of 70 mm/min has the highest yield and ultimate tensile strength, which is consistent with the microhardness measurements in the center of the stirred zone. This high strength achieved without a remarkable decrease in elongation due to the formation of a heteroge neous microstructure. The good combination of strength and ductility is the most important factor in the applications of the engineering materials. It is important to note that the stress-strain behavior of IF and St52 steels during the tensile test is different. A schematic illustration of the elastoplastic deformation mechanism at the IF/St52 interface is shown in Fig. 17. As seen, there are three distinct steps in the stress-strain curves of composite samples. In step one, both IF and St52 steels are in elastic deformation. In step two, the interstitial free steel starts plastic deformation while the low carbon steel is under elastic deformation. In step three, both IF and St52 steels are in plastic deformation. Therefore, 7
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Fig. 11. The microstructures in the top, middle, and bottom of the retreating side, the center of the stirred zone, and advancing side for the tool traverse speed of 100 mm/min.
as indicated in Fig. 17, different elastoplastic properties of the intersti tial free and low carbon steels produce a dislocation density gradient in both layers during the tensile test. The relationship between the ultimate tensile strength and total elongation is usually expressed by the so-called banana curve for various steels as shown in Fig. 18. In the past few decades, numerous researches on grain refinement using severe plastic deformation (SPD) techniques have been performed to obtain high strength interstitial free steel [2,4,6, 32–36]. Although the yield and ultimate tensile strength of interstitial free steel are remarkably increased, the problem of decreased elongation limits the application of such high strength steel as forming and struc tural materials. For the more industrial applications of interstitial free steel, the low strength needs to be solved without sacrificing the elon gation. The present study indicates that a good combination of strength and ductility could be obtained through FSLW, which is mainly due to the formation of a heterogeneous microstructure in both layers. As seen in Fig. 18, the produced composite lies in the high strength steel cate gory, while the initial samples (IF and St52 steels) were low strength steels. Therefore, the FSLW exhibits to be a promising process to fabri cate a high strength steel composite.
IF and St52 steels. The composite produced by the traverse speed of 40 mm/min failed by separation of the layers, at the interface, as indicated in Fig. 20(a). This result shows that the bonding between IF and St52 steels processed by FSLW was not enough strong to maintain integrity up to a high level of deformation during the tensile test. On the other hand, the delamination between the interstitial free and low carbon steels is not observed in the FSLW with the traverse speeds of 70 and 100 mm/ min, and the subsequent tensile test. As shown in Figs. 3–5, unlike the composite welded by 40 mm/min traverse speed, there was no sign of microstructural discontinuity between the IF and St52 steels for the traverse speeds of 70 and 100 mm/min. Consequently, the layers of the interstitial free and low carbon steels remained well bonded during the tensile test. From Fig. 20, obvious necking observed on the fracture surface of the composite samples produced by FSLW indicating a substantial plastic deformation before the failure. The fracture surfaces of both layers in the produced composites exhibit many dimples indicative of ductile fracture mode without any cleavage or quasi cleavage fracture mode. By comparing the fracture surfaces of the initial St52 and IF steels (Fig. 19) and the St52 and IF layers in the composites (Fig. 20), two different features can be observed. Firstly, the size of dimples in the fracture surfaces of initial materials is more uniform than that of the St52 and IF layers in the FSLWed composites. This is due to the formation of het erogeneous microstructure in both layers after the FSLW. Secondly, the IF and St52 layers in the composites exhibit smaller dimples than the initial materials indicating lower ductility, which is consistent with the tensile test results. By increasing the tool traverse speed, the size of dimples decreased. The size of dimples for the traverse speed of 40 mm/min is greater than
3.3. Fractography Fig. 19 presents the fracture surfaces of the initial IF and St52 steels after the tensile test. Numerous and deep dimples can be observed on the fracture surfaces of the interstitial free and low carbon steels, which indicates a completely ductile failure. Fig. 20 shows the fracture surfaces of the two-layered steel com posites after the tensile test. The red arrows indicate the interfaces of the 8
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Fig. 12. The Fe–Fe3C phase diagram with different microstructures and the approximately related temperatures for the composite samples produced by the traverse speeds of 40 and 70 mm/min.
that of 70 and 100 mm/min. This result can be attributed to the larger grain size of the composite produced by the traverse speed of 40 mm/ min as indicated in Figs. 9–12. By increasing the grain size, the fraction of grain boundaries decreased, the number of dimple nucleation reduced, and the size of dimples increased. On the other hand, with increasing the tool traverse speed, the depth of dimples decreased. The size of dimples for the traverse speeds of 70 and 100 mm/min is smaller than that of 40 mm/min. This is due to the formation of a hard phase (i. e., martensite) in the microstructure of the composite produced by the traverse speeds of 70 and 100 mm/min as shown in Figs. 9–12. Friction stir lap welding of interstitial free and low carbon steels can prove an excellent method to produce well-bonded two-layered steel composite. The FSLWed IF/St52 steel composites showed good tensile strength and elongation due to the formation of heterogeneous micro structure. These results allow the extended use of interstitial free steel in the automotive industry.
3. The increase of tool traverse speed from 40 to 70 and 100 mm/ min at a constant tool rotation speed of 1000 rpm decreased the cooling time leading to a higher cooling rate and the formation of martensite phase. 4. The heterogeneous microstructure was observed in the composite samples produced by FSLW. 5. The grain size of the composite produced by the traverse speed of 40 mm/min in the first and second HAZ regions was larger than that of the 70 mm/min due to the higher heat input and therefore lower cooling rate of the sample welded via 40 mm/min. 6. The hardness of the center of the SZ was higher than that of the advancing and retreating sides. This could be attributed to the effect of tool pin on the shear deformation during FSLW. 7. The St52 layer of composites showed the microhardness value of ~190 HV whereas the initial St52 sheet had the microhardness around 160 HV. This was owing to the formation of Widman statten ferrite, martensite, and fine ferrite and pearlite grains. 8. The composite produced by the traverse speed of 70 mm/min had the highest YS and UTS. This high strength achieved without a remarkable decrease in elongation due to the formation of a heterogeneous microstructure. 9. The fracture surfaces indicated the bonding between IF and St52 steels in the composites produced by traverse speeds of 70 and 100 mm/min was high enough strong to maintain integrity up to a high level of deformation during the tensile test. 10. With increasing the tool traverse speed, the size of dimples decreased. The size of dimples for the traverse speed of 40 mm/ min was greater than that of 70 and 100 mm/min. This result could be due to the larger grain size of the composite produced by the traverse speed of 40 mm/min.
4. Conclusions In the present study, two-layered steel composite consisted of IF and St52 steels sheets was successfully fabricated by friction stir lap welding. The main conclusions were drawn as follows: 1. The quality of interfacial bonding for the tool traverse speeds of 70 and 100 mm/min was excellent due to the higher strain rate during FSLW with the traverse speed of 70 and 100 mm/min than that of 40 mm/min. 2. The microstructures of composite samples mainly consisted of fine equiaxed ferrite, Widmanstatten ferrite, grain boundary ferrite, aggregates of ferrite þ cementite, and martensite.
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Fig. 13. The Vickers microhardness profiles across the vertical line in the retreating side, the center of the SZ, and the advancing side for the traverse speeds of (a) 40, (b) 70, and (c) 100 mm/min.
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Fig. 14. The microhardness profiles of the center of the stirred zone for the composites produced by the three different traverse speeds as well as initial IF and St52 steels.
Fig. 15. The tensile engineering stress-strain curves of the initial and processed samples.
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Fig. 16. (a) Yield strength, (b) ultimate tensile strength, and (c) total elongation of the initial and processed samples.
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Fig. 17. Schematic illustration of the elastoplastic deformation mechanism at the IF/St52 interface during the tensile test.
Fig. 18. The tensile strength-elongation relationship in steels.
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Fig. 19. The fracture surface of (a) IF and (b) St52 steels.
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Fig. 20. The fracture surfaces of the composites produced by the traverse speeds of (a) 40, (b) 70, and (c) 100 mm/min.
Data availability
References
The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.
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Author contribution statement Mohammad Reza Roodgari: Investigation, Resources, Writing Original Draft. Roohollah Jamaati: Conceptualization, Methodology, Writing Review & Editing, Supervision. Hamed Jamshidi Aval: Conceptualization, Methodology, Writing Review & Editing, Supervision. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements The authors acknowledge the funding support of Babol Noshirvani University of Technology through Grant program No. BNUT/393044/98 and BNUT/370167/98.
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