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Effect of coating on fiber laser welded joints of DP980 steels
Materials and Design 90 (2016) 516–523
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Effect of coating on fiber laser welded joints of DP980 steels Q.L. Cui a,b, D. Parkes b, D. Westerbaan c, S.S. Nayak c, Y. Zhou c, D. Liu d,⁎, F. Goodwin e, S. Bhole b, D.L. Chen b a
State Key Laboratory of Rolling and Automation, Northeastern University, 3-11 Wenhua Road, Shenyang, China Department of Mechanical and Industrial Engineering, Ryerson University,350 Victoria Street, Toronto, Ontario M5B 2K3, Canada c Department of Mechanical and Mechatronics Engineering, University of Waterloo, 200 University Avenue West, Waterloo, Ontario N2L 3G1, Canada d Teck Metals Ltd, 2380 Speakman Drive, Mississauga, Ontario L5K 1B4, Canada e International Zinc Association, Durham, NC 27713, USA b
a r t i c l e
i n f o
Article history: Received 4 May 2015 Received in revised form 16 October 2015 Accepted 17 October 2015 Available online 19 October 2015 Keywords: Coating Fiber laser welded DP980 Similar joints Fatigue
a b s t r a c t Advanced high strength steels (AHSS), specifically dual phase (DP) steels, are extensively used to reduce the weight of vehicles. Fiber laser welding (FLW) has been shown to provide welds with superior mechanical properties made at high welding speeds. Galvannealing (GA) and galvanizing (GI) coated sheet steels are extensively used for auto bodies to improve corrosion resistance. As such, the effect of coating type on the microstructure, mechanical properties and fatigue properties of FLW similar joints of DP980 was assessed. There existed spherical tempered martensite in the GA coated DP980 base metal (BM), most likely caused by the galvannealing process. Coating type did not present a noticeable effect on the microstructure of the welded zones. The GI coating, however, could evaporate violently during welding, resulting in high concavity and other welding defects which negatively affect the high-cycle fatigue property. Measures to mitigate the negative effects of the GI coating were discussed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Weight reduction of auto bodies plays an important role in improving the fuel economy of vehicles to meet regulations [1]. Advanced High Strength Steels (AHSS) are being developed to reduce weight while enhancing crashworthiness of auto bodies. Dual phase (DP) steel is popular favorable type of AHSS because of its good processability and balanced combination of strength and ductility [2,3]. Laser welding is a desirable joining technology because of its high energy density and fast speed result in a small heat affected zone (HAZ) with fine weld metal quality. In the automotive industry the demand for laser welded blanks (LWBs), also known as tailor welded blanks, has increased over the last decades owing to the need to reduce vehicle weight and CO2 emissions. LWBs consist of two or more sheets of similar or different materials, properties, thickness, and surface conditions which are welded together to form a blank. LWBs are generally fabricated in the butt joint configuration to form an assembly, which is subsequently formed to make the desired three-dimensional shapes. In general, LWBs are made with conventional mild (low carbon) steel or interstitial free steel; in recent years the weight reduction potential of LWBs has been realized via material down gauging by employing stronger steels ⁎ Corresponding author. E-mail address: [email protected] (D. Liu).
http://dx.doi.org/10.1016/j.matdes.2015.10.098 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
like AHSS. Thus, it is important to study the weldability of automotive steels like AHSS via laser welding [4,5]. Among the laser welding methods, fiber laser welding (FLW) produces the highest quality welds at a low cost to the manufacturer due to its high power with low beam divergence, flexible beam delivery, low maintenance costs, high efficiency and compact size [6,7]. Sheet steels with GI and GA coatings are commonly used to construct auto bodies for their excellent corrosion resistance. The GI coating essentially contains layer of elemental Zn with up to 0.5% Al; whereas, the GA coating includes up to 12% Fe because the Zn coating formed in the hot-dip process reacts with the steel substrate during the post-dip annealing process at temperatures over 500 °C to form Zn–Fe intermetallic coating. The effect of these coatings on the behavior of spot welded steels was studied by several researchers [8–10]. Different electrode life behaviors were reported in resistance spot welding of GA and GI coated steels. The aluminum content of the coating was the most significant factor affecting electrode life of GI coated sheet steels. Rich experience in spot welding of GA and GI coated steels affirmed that the GI coating tended to eject molten zinc with zinc vapor whereas GA coating enjoyed the reputation of good weldability. There are reports on laser weldability, microstructure, mechanical and fatigue properties of DP steels [5–7,11]. However, no study has been reported so far on the effect of coating types on the properties of FLW joints. Therefore, the present study reports a comparison of the GI and GA coating types on the
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performance in FLW of DP980 steel sheet. Weld concavity formed during FLW and its effect on the microstructure, microhardness, the tensile and fatigue properties has been reported. 2. Materials and experimental procedure The materials used in this work were DP980 sheet steel with either GI or GA coating. The steel thickness was 1.2 mm. The nominal coating weights per side were 60 g/m2 for the GI and 41 g/m2 for the GA coating. The GA coating weight is typically lower to control the powdering tendency. The chemical compositions of the steels are listed in Table 1. Normally, DP steel contains about 0.10 to 0.15 wt.% carbon and is generally alloyed with various amounts of Mn, Si, Cr and Mo to achieve the required tensile strength. The Carbon Equivalent, Ceq, of DP980 steel is calculated using the Ito–Bessyo equation [12], C eq ¼ C þ Si=30 þ ðMn þ Cu þ CrÞ=20 þ Mo=15 þ V=10 þ Ni=60 þ 5B:
ð1Þ
The two DP980 steels tested in the present study were low in C contents and relatively lean in alloying. They thereby had relatively low Ceq values of 0.236 wt.% (GA) and 0.245 wt.% (GI), as listed in Table 1. Steels with low Ceq typically have good weldability. As these two steels have low and similar Ceq values, it is expected that they both attain similar weldability. Fiber laser welding was carried out using an IPG Photonics YLS-6000 fiber laser system, with a focal length of 200 mm, attached to a Panasonic robotic arm. Welding were made along the rolling direction of the sheets. Butt welds of 100 mm × 200 mm sheets of each material were made to make 200 mm × 200 mm laser welded blanks as shown in Fig. 1a. The welding speed was 20.5 m/min and the laser power was 4 kW, which was chosen based on our earlier study to reduce weld concavity [5]. The fiber laser had a fiber core diameter of 0.3 mm with a laser beam spot size/diameter of 0.6 mm. Welding was performed with a 20° vertical head angle, i.e., the laser beam was at an angle of 70° to the surface of sheets (work pieces). No shielding gas was used during welding; instead, blown air was used to remove vapors from the weld pool and to protect the system optics. Metallographic samples were cut from the weld cross-section, then mounted, ground, polished, and etched with a 2% Nital solution. The etched samples were observed first using a light microscope attached with Clemex image analysis system software and then with a scanning electron microscope (JEOL JSM-6380) equipped with threedimensional fractographic analysis. Vickers microhardness was measured on the polished samples across the weld using a computerized microhardness tester with a load of 200 g and a dwell time of 15 s. Localized strain hardening caused by adjacent indentations was avoided by carefully spacing each indentation. To ensure the accuracy of each test result, two calibration tests were carried out using a standard reference test block before the microhardness tests on the welded joints were conducted. Tensile and fatigue test samples were machined to form transverse welds in accordance with ASTM: E8/E8M, the example of which is indicated by the dashed line in Fig. 1a with the geometry and dimensions of the test coupons shown in Fig. 1b. The specimens were machined in such a way that the weld line was positioned at the center of the gauge length (Fig. 1b). Tensile tests were conducted on a fully computerized united tensile testing machine at room temperature and with a strain rate of 1 × 10− 3 s− 1. An extensometer with a gauge length of
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50 mm and a strain limit of 20% was used to measure the strain during the tensile tests. Load control fatigue tests were performed following ASTM: E466 on a fully computerized Instron 8801 servo-hydraulic testing system. Tension-tension cyclic loading was employed to prevent potential buckling of the samples and a stress ratio of R = 0.1 was used. The test conditions were conducted at room temperature with a frequency of 50 Hz and the load was applied using a sinusoidal waveform. A minimum of two specimens were tested in the tensile tests and in the fatigue tests at each of the cyclic stress amplitude. The fatigue fracture surfaces were examined using scanning electron microscopy (SEM) to study the fatigue fracture mechanisms. 3. Results and discussion 3.1. Microstructure To understand the welding behavior, it is helpful to examine the microstructure of the base material and how it was obtained. Through the heat treatment steps in a continuous galvanizing line, the DP steel is intercritically annealed to form a mixture of ferrite and austenite. Subsequent rapid cooling avoided the formation of pearlite and bainite while promoting martensite formation [13]. The steel then went through a hold zone typically controlled at about 470 °C, before immersing in a molten zinc bath at 460 °C. The resulting microstructure of DP980 steel base metal (BM) is shown in Fig. 2a and b. Both the GA and GI coated DP980 had a microstructure consisting of a ferrite matrix with islands of martensite. It is worth noting that there existed spherical tempered martensite in the GA coated DP980 BM whereas there was none in the GI coated steel. The tempering was caused by the additional heating cycle in the galvannealing process. Microscopic examination of the welds showed the presence of martensite in both the fusion zone (FZ) and the HAZ, as shown in Fig. 3. In laser welding, the weld cooling was extremely rapid. It has been shown through modeling that the cooling rates in laser welding were on the order of 105 °C/s [14]. The critical cooling rate CR for steels to achieve martensitic transformation is given by the following equation, LogCR ¼ 7:42−3:13C−0:71Mn−0:37Ni−0:34Cr−0:45Mo:
ð2Þ
For the DP980 steels used in this study, the critical rate is about 400 °C/s. Therefore, a martensitic structure was obtained in both the FZ and the HAZ. There was limited amount of bainite in the FZ which was attributed to the high cooling rate during the FLW process [15]. The HAZ of the weld appeared to be similar in both coating conditions with an upper-critical HAZ, an inter-critical HAZ, and a subcritical HAZ. In the upper-critical HAZ, the temperature was above the Ac3 line, leading to a microstructure that consisted of mostly martensite with a very small amount of ferrite (Fig. 3c and d) [14]. In the intercritical HAZ (Fig. 3e and f), the temperature was between the Ac1 and Ac3 lines, resulting in the coexistence of ferrite and austenite, the latter of which subsequently transformed into martensite during cooling. In the sub-critical HAZ, the temperature was below the Ac1 line [7]. This caused the martensite that was normally present in the BM to temper and decompose into partially tempered martensite (PTM), as can be clearly seen in Fig. 3g and h. A “soft zone” was thus created where the steel was the weakest [16]. Tempering of BM martensite in the fusion welding of DP980 steel has been reported in several previous studies [17,18]. The amount of PTM in the sub-critical HAZ for both welded joints was very similar.
Table 1 Chemical composition of the DP980 in present study. Steel
C
Mn
Si
Al
Cr
Ni
Mo
Nb
N
Cu
Ti
Ceq
GI GA
0.15 0.1
1.45 2.19
0.33 0.33
0.05 0.04
0.02 0.21
0.01 0.01
0.002 0.238
0.001 0.016
0.01 0.01
0.02 0.02
0.002 0.022
0.245 0.236
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Fig. 1. (a) Geometry of the welded sheet, (b) transverse tensile joint machine from the dotted line in (a).
3.2. Microhardness Microhardness measurements of the GA and GI coated DP980 welds are shown in Fig. 4. Both specimens displayed similar trends. The hardness was the highest in the FZ, about 100 HV higher than that of the BM, due to its predominant martensitic microstructure (Fig. 3). Outside the FZ, the hardness decreased abruptly from the maximum value in the upper-critical HAZ to the minimum value in the sub-critical HAZ, corresponding to the change in microstructure from all martensite to a mixture of ferrite and partially tempered martensite. As the distance from the center of the weld further increased, the heat effect diminished and, hence, the hardness gradually recovered to the level of that of the BM. In this particular study, the hardness of the GA coated specimen was lower than that of the GI coated specimen with an average difference of about 36 HV. This was largely attributable to a) the difference in steel composition, in particular the carbon content and b) the effect of galvannealing. As shown in Table 1, the carbon content of GI coated DP980 was 0.15 wt.%, which was 50% higher than the 0.10% of its GA counterpart. Hence, the carbon content in martensite was higher for the GI coated steel, leading to higher hardness and strength. In addition, the galvannealing process caused tempering of the martensite (Fig. 2a), resulting in reduced hardness and strength.
3.3. Galvanized coatings and weld defects Cross-sectional views of both GI and GA coatings are shown in Fig. 5. The location of the weld was to the left of these micrographs. Fig. 5a shows a typical GI coating, taken at 0.74 mm from the edge of FZ. The coating consisted of one single layer of zinc solid solution, the composition of which matched the composition of a galvanizing bath. Between
the zinc layer and the steel substrate existed an Al–Fe intermetallic layer, which was too thin to observe at this magnification. This thin Al–Fe intermetallic layer prevented the reaction between Fe and Zn, and was therefore called the inhibition layer. The thickness of the GI coating, as shown in Fig. 5a, was about 12 μm. Fig. 5b shows a typical GA coating taken at 0.60 mm from the edge of FZ. The coating consisted of two layers of Zn–Fe intermetallic compounds. The upper layer (δ phase) had an average Fe content of 11.2 wt.%, which was typical for the so-called “all-δ” GA coating. The thickness of the GA coating was about 7 μm. Because thick GA coatings tend to have formability issues such as powdering and flaking, only thin GA coatings are normally produced. Fig. 5c was taken in the HAZ, 0.29 mm from the edge of FZ. At this location, the heat from welding was enough to melt the GI coating (melting point 420 °C). Silver lines of very fine Fe–Zn intermetallic precipitates can be clearly seen along the boundaries of large zinc grains. Fe–Zn intermetallic particles are also found at the coating/substrate interface due to the breakdown of the inhibition layer which typically occurs at temperatures above 500 °C. Considering the short exposure time to the peak temperature (fast heating and cooling in laser welding), to have this extent of Fe–Zn reaction, the peak temperature is estimated to be over 550 °C. This location was in the soft zone shown in Fig. 4, which also indicated an exposure to an elevated temperature. Fig. 5d shows the GA coating also at 0.29 mm from the edge of FZ. It is interesting to note that the coating in the left half of the micrograph had high porosity whereas the right half was largely dense. The coating porosity increased progressively to the left all the way to the edge of the FZ. To the right of the micrograph, the GA coating was completely dense like the one shown in Fig. 5b. Considering the diffusivity of zinc is two orders of magnitude higher than that of iron in Fe–Zn intermetallic compounds, it is possible that the porosity can form due to Kirkendall effect. However, the abrupt transition from high porosity to no porosity
Fig. 2. BM microstructure of DP980 (a) GA (b) GI (F: ferrite, M: martensite, PTM: partially tempered martensite).
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Fig. 3. Microstructure of the DP980 similar welding with GA and GI coating, FZ (a) GA and (b) GI, upper-critical HAZ (c) GA and (d) GI, inter-critical HAZ (e) GA and (f) GI, tempered subcritical HAZ (g) GA and (h) GI (F: ferrite, M: martensite, B: bainite, PTM: partially tempered martensite).
indicated that the liquid phase was involved in the rapid diffusion that led to the porosity. The δ phase coexisted with liquid zinc in the temperature range between 530 °C and 672 °C. The measured Fe content of 11.2% was near the upper limit of the δ phase composition range (7.0– 11.5%). Therefore, the temperature should be near the upper end of the temperature range and was calculated to be 658 °C using a thermodynamic database [19]. The thickness of the GI coating started to decrease when the distance to the FZ was 0.04 mm, partially due to the evaporation of zinc at temperatures exceeds its boiling point of 907 °C. This temperature was above the Ac3 temperature of DP980 and this location was therefore in the upper-critical HAZ, which was also reflected in the hardness measurement shown in Fig. 4. The GI coating
evaporated completely at the distance of 0.01 mm from the FZ. In comparison, a good portion of the GA coating survived direct contact with molten steel right to the edge of the FZ, exhibiting better stability at high temperatures. This agreed well with prior observations that GA coating displayed lower zinc activity [2,9,10]. The entire cross-sections of FZ of GI coated and GA coated samples are shown in Fig. 5e and f, respectively. In general, greater material loss in FZ was observed with GI coated samples, which led to larger concavity and more severe welding defects. It was confirmed that concavity formation was due to the metal ejection [5]. Relevant experience in spot welding indicated that the evaporation of zinc in the GI coating could cause ejection of molten metal whereas the GA coating exhibited good
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Fig. 6. Tensile curves of GI and GA coated DP980 similar welded joints. Fig. 4. Microhardness curves of GA and GI coated DP980–DP980 welded joints.
stability. Moreover, because the thickness of the GA coating is almost half that of the GI coating, it is reasonable to think that the GA coating in this case would offer lower resistance [8]. In the present study the thick GI coating used further deteriorated the performance of GI coated weld in comparison to the thin GA coating. The zinc content in the FZ of both GI and GA specimens was below the detection limit, indicating that essentially all zinc had evaporated in FZ. 3.4. Tensile properties Select curves of engineering stress versus engineering strain of GA or GI coated DP980 specimens with or without a weld (i.e., base metal) are displayed in Fig. 6, while the average properties of all specimens are listed in Table 2. The BM specimens exhibited good repeatability
whereas the welded specimens displayed large variation, especially in elongation. The average ultimate tensile strength (UTS) of GA BM was 936 MPa, which was notably lower than 1083 MPa of GI BM. The causes of lower ultimate tensile strength of GA BM, similar to those discussed earlier in microhardness, were (1) lower carbon content in steel chemistry and, (2) the softening of materials due to the GA process. The softening was also reflected in the slight increase in elongation, with 13.6% for GA BM versus 12.1% for GI BM. For the welded specimens, their UTS (946 MPa for GA and 999 MPa for GI) deviated marginally from their BM counterparts. The limited decrease in strength was attributable to the narrow and mild soft zones (Fig. 4) which were achieved by having a small laser beam spot size and high power intensity in conjunction with a high welding speed. The elongation of welded specimens was, however, significantly lower with 8.6% for GA and only 3.8% for GI. The elongation of the welded specimens was also much more scattered. For instance, the elongation of the four GI coated weld specimens
Fig. 5. (a) Normal GI coating outside of HAZ, 0.74 mm from the edge of FZ; (b) Normal GA coating outside of HAZ, 0.60 mm from the edge of FZ; (c) GI coating inside HAZ, 0.29 mm from the edge of FZ; (d) GA coating inside HAZ, 0.29 mm from the edge of FZ; (e) Weld of GI coated DP980; and (f) Weld of GA coated DP980.
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Table 2 Summary of fatigue and tensile properties. Steel
YS (MPa)
UTS (MPa)
Elongation %
Joint efficiency %
Fatigue limit (MPa)
Fatigue ratio
σ'f (MPa)
b
GA joints GA BM GI joints GI BM
667 644 717 707
946 934 999 1083
8.6 13.6 3.8 12.1
100 – 92.2 –
100 200 b100 250
0.11 0.21 b0.1 0.23
1370 733 4130 682
−0.153 −0.077 −0.260 −0.060
ranged from 1.3% to 6.2% with a standard deviation of 2.4%. It is, hence, not possible to select one single representative curve to reflect the scattering nature of this property. To better illustrate the complex situation and avoid overcrowdings of the figure, two curves, one with the lowest elongation and another with the greatest elongation, were displayed in Fig. 5 for each of the GA weld and GI weld specimens. Pictures of the fractured tensile specimens are shown in Fig. 7. The GA BM specimen (top specimen in Fig. 7a) fractured along a plane at an angle of about 60° with the pulling direction, indicating the ductile failure of slipping caused mainly by shear stress. Necking was also apparent, offering additional indication of ductile failure. The fracture process of the GI BM specimen (bottom specimen in Fig. 7a) appeared to initiate from two edges via slipping caused by shear stress, and then completed with fracture caused by normal stress at the center of the specimen. For the welded GA specimen with the smallest elongation (top specimen in Fig. 7b), fracture appeared to initiate from both edges in the sub-critical HAZ. Cracks tended to propagate within the narrow soft band of sub-critical HAZ (90° angle with the pulling direction) but at the same time were influenced by shear stress which tended to lead to a failure at a 45° angle. The reduced elongation was likely attributable to the more localized necking near the sub-critical HAZ. For the welded GA specimen with the greatest elongation (lower specimen in Fig. 7b), fracture occurred outside the HAZ, demonstrating that
welding did not cause significant weakening to this specimen. The nature of failure was no different from that of the GA BM specimen shown in Fig. 7a. For the welded GI specimen with the smallest elongation, its stress–strain curve shown in Fig. 6 indicated premature failure without apparent plastic deformation. Fracture appeared to occur along the HAZ/FZ interface on the left-hand side of the FZ (top specimen in Fig. 7c). Thus, the fracture surface was perpendicular to the pulling direction. For the welded GI specimen with the greatest elongation, fracture appeared to initiate along both HAZ/FZ interfaces on either side of the FZ, and then finished with a fracture crossing the FZ. It is worth noting that the microstructure in the FZ and upper-critical HAZ was mainly composed of martensite (Fig. 3a-d) and the hardness was the highest (Fig. 4). Failure initiated at the HAZ/FZ interface was, therefore, unlikely caused by low strength in these regions. Close examination of the weldment found that the welded GA specimen suffered only mild concavity and mild misalignment (Fig. 5f) while the welded GI specimen suffered severe concavity and notable misalignment (Fig. 5e). The crack (about 34 μm in length) in the upper-right corner of the FZ in Fig. 5e was found to extend another 135 μm into the FZ when examined carefully at a higher magnification. The difference in concavity resulted from the difference in the zinc vapor pressure which was higher for the pure zinc GI coating than for the GA coating where zinc was tied up with iron in more stable Fe–Zn intermetallic compounds. The volatility of pure zinc GI coating could cause a loss of molten metals in the FZ, resulting in a large concavity. The volatility could also generate turbulence and create weld defects, resulting in stress concentration which was particularly detrimental to hard and brittle martensitic steel. 3.5. Fatigue and fractography Fig. 8 displays the S-N curves of the GA and GI coated welded joints with their corresponding BMs. For stress amplitude of 350 MPa and above, there was no notable difference in fatigue life between welded GA specimens and their corresponding BM specimens. Put it in perspective, a stress amplitude of 350 MPa had a maximum stress of 778 MPa while a stress amplitude of 450 MPa had a maximum stress of
Fig. 7. Pictures of fractured tensile specimens corresponding to curves shown in Fig. 5. (a) BM of GA and GI (b) Lowest and greatest elongation of welded GA (c) Lowest and greatest elongation of welded GI
Fig. 8. Comparison of S–N curves of the GA and GI coated similar DP980 welded joints with its corresponding BM.
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1000 MPa which was higher than the UTS of GA BM (936 MPa as shown in Fig. 5) tested using a strain rate of 1 × 10−3 s−1. Granted, the strain rate in a fatigue test was over two orders of magnitude higher than the strain rate used in the tensile test. Nevertheless, the stress level was very high and both GA welded specimens tested using a stress amplitude of 450 MPa failed almost instantly in just over 1 s with the fracture occurred in the base metal at an angle of about 60° to the pulling direction, almost identical to the bottom specimen shown in Fig. 7b for the tensile test. At the stress amplitude of 350 MPa, the test started to resemble a typical fatigue test with specimens lasted for over 1 × 104 cycles (several minutes) and a steady crack growth stage (Stage II) being distinctly identified with flat and smooth fracture surfaces perpendicular to the loading direction. The Stage II fracture surfaces were close to the FZ/HAZ interface and their area was about 20% of the cross-sectional area of the GA specimens. The final ductile failure occurred with rapid crack growth along a plane about 45° in the thickness direction. This plane ran across the entire HAZ, which had a width of less than 0.8 mm as determined by the micro-hardness measurement shown in Fig. 4 and by the coating morphology analysis in Fig. 5. It is well known that low-cycle fatigue is closely related to plastic deformation. Hence, the softening effect of the GA process decreased the fatigue life at high stress amplitudes. That was the main reason why the GI BM specimens outperformed the GA BM specimens in Fig. 8 with stress amplitudes greater than or equal to 350 MPa. The softening in the HAZ had a less significant effect on the already softened GA welded specimens. Therefore, there was little difference between the GA welded and GA BM specimens. The softening effect in the HAZ was more significant for GI coated specimens (Fig. 4). Thus, more notable difference in fatigue life was observed between GI welded and GI BM specimens, as shown in Fig. 8. The severe concavity in GI welded specimens at the FZ/HAZ interface where cracking initiated further decreased their fatigue life. Unlike the low-cycle fatigue which is related to plastic deformation and sensitive to soft zones, the high-cycle fatigue is more sensitive to defects and stress concentration. Harders and Baeker [20] pointed out that the fatigue strength of a material under dynamic cyclic loading was much more sensitive to the manufacturing process and materials than the static strength, and in particular for the lower level cyclic load. The severe concavity formed a notch similar to a notched specimen studied by Chapetti et al. [21] where the stress concentration was generated, and consequently a decreased fatigue limit or fatigue life at low level stress amplitudes was observed. In addition to the concavity, severe defects formed in FZ of the welded GI specimens. Fig. 9 shows the top view of a weld of a GI specimen undergone a fatigue test with stress amplitude of 150 MPa. The upper portion of the picture shows severe turbulence of molten melt close to the FZ/HAZ interface. The lower portion of the picture shows another FZ/HAZ interface where Stage II fracture occurred. On the weld surface scattered a great number of
small pieces of surface oxide, presumably broken by the thermal stress during the fast cooling following the solidification of the weld. Dendritic grains are evidently visible along the center of the weld. The weld defects were markedly less severe for GA specimens. Note that the important Stage II crack growth took place at the FZ/HAZ interface (perpendicular to loading direction) and the Stage II area increased with decreasing stress amplitude. At the stress amplitude of 150 MPa, the Stage II area became over 60% of the cross-sectional area, thereby playing a very important role in determining the fatigue life of welded specimens. Consequently, the welded GA specimens performed significantly better than their GI counterparts at low stress magnitudes. For the lowest stress magnitude (100 MPa) tested in the present study, the welded GI specimens failed after 4 × 105 cycles on average, whereas the welded GA specimens survived the entire test of over 1 × 107 cycles without failure. For BM specimens, GA BM specimens had the disadvantage of a lower tensile strength and a brittle coating which was prone to microcracking. As a result, the GA BM specimens had shorter fatigue life than their GI BM counterparts, as shown in Fig. 8. The GI BM specimens survived the test at a stress amplitude of 250 MPa while the GA BM specimens survived the test at a lower stress amplitude of 200 MPa. Notable improvement in fatigue life can be expected, especially for the GI coated specimens, if the weld direction does not run perpendicular to the loading direction. As such, the critical Stage II crack propagation cannot take place along the defect-laden FZ/HAZ interface. According to the experience of the explosive welding [22–24], it is supposed that an angled, Chevronor curved weld can effectively force the Stage II crack to propagate away from the FZ/HAZ and into the BM, and hence improve the fatigue performance of the welded part. The fatigue limit and fatigue ratio (i.e., a ratio of fatigue limit to the UTS) are tabulated in Table 2. The decreases in fatigue limit for GA and GI welded specimens, with respect to their BM counterparts, were 50% and over 60%, respectively. The fatigue limit of the GA welded specimens was higher than that of the GI welded specimens. The fatigue ratios was 0.11 for the GA welded specimens and was below 0.1 (fatigue limit lower than the experimental range and was undetermined) for the GI welded specimens. These results suggested that while the tensile strength of the FLW DP980 joints with GI coating was higher than that of the joints with GA coating, the fatigue strength was actually better for GA welded specimens at a lower level of stress amplitudes (≤150 MPa). From Fig. 8, it can also be seen that GA welded joints exhibited a much higher fatigue strength in the low cycle fatigue regime below ~ 5 × 104 comparing with that of GI welded joints. The S-N plot in Fig. 8 is expressed using the Basquin equation [17], 0
σ a ¼ σ f 2N f
b
ð3Þ
where σa is the cyclic stress amplitude, σ′f is the fatigue strength coefficient defined by the stress intercept at 2Nf = 1, Nf is the number of cycles to failure (2Nf is the number of reversals to failure), and b is the fatigue strength exponent. The values of these parameters obtained by fitting the data points according to Eq. (3) are tabulated in Table 2. The fatigue life of the welded joints could be estimated based on the values of σ′f and b. Eq. (3) indicates that the higher the value of σ'f and the smaller the absolute value of b, which is always negative, the longer the fatigue life. From Table 2, it can be seen that the GA welded joints have a better b value and a higher fatigue limit which means that they would have a better fatigue life in comparison to the GI coated samples. 4. Conclusions
Fig. 9. The top surface of a welded GI specimen undergone a fatigue test with a stress amplitude of 150 MPa.
The aim of this study was to take a comparative look at the welding performance of GI and GA coatings. It should be noted that this study pertained only to the steel compositions and the coating weights selected. The following conclusions could be reached based on the study
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conducted to examine the differences between GA and GI coated DP980 steels: 1. Due to the annealing process of the GA coating process, there was more spherical tempered martensite in the GA coated DP980 BM than in the GI coated DP980 BM. Coating type had no effect on the microstructure of the welded zones. 2. The hardness of GA coating welded joint was generally lower than that of GI coating from the BM to the FZ. In the BM, this was partially attributable to the PTM that was seen in the BM of GA coated DP980. In FZ and HAZ, the hardness difference between the GA and GI welded joints stemmed mainly from the difference in chemical composition. The UTS and YS of the GA coated welded joints were lower than that of GI coated samples. This was affirmed by the hardness results. 3. The GA similar welded joints showed improved fatigue resistance compared to GI welded joints due to lower concavity resulting from the coating characteristics of GA and from its low coating thickness. Acknowledgments The authors would like to thank the National Natural Science Foundation of China (No. 51304045), the China Scholarship Council, Fundamental Research Funds for the Central Universities of China (No. N110407001), Natural Sciences and Engineering Research Council of Canada (NSERC) and AUTO21 Network of Centers of Excellence for providing financial support. The financial support from International Zinc Association (IZA) and Arcelor Mittal Dofasco is highly acknowledged. One of the authors (D.L. Chen) is grateful for the financial support by the Premier's Research Excellence Award (PREA), NSERC-Discovery Accelerator Supplement (DAS) Award, Canada Foundation for Innovation (CFI), and Ryerson Research Chair (RRC) program. The authors would like to thank Dr. J. Chen and Dr. Y.L. He (CANMET-Materials Technology Laboratory, Natural Resources Canada, Hamilton, Canada), Mr. E. Biro (Arcelor Mittal Global Research, Hamilton, Canada), and Dr. J. Villafuerte (Center Line (Windsor) Ltd., Windsor, Canada) for their support and helpful discussion. The assistance of Q. Li, A. Machin, J. Amankrah, and R. Churaman in performing the experiments is gratefully acknowledged. References [1] L.F. Mei, G.Y. Chen, X.Z. Jin, W.Q. ZhangY, Research on laser welding of high-strength galvanized automobile steel sheets, Opt. Lasers Eng. 47 (2009) 1117–1124.
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[2] X. Li, S. Lawson, ZhouY. Novel technique for laser lap welding of zinc coated sheet steels, Journal of Laser Applications 19 (4) (2007) 259–264. [3] Y.Y. Zhou, Y.S. Zhang, W. Hu, Effect of welding speed on microstructure, hardness and tensile properties in laser welding of advanced high strength steel, Sci. Technol. Weld. Join. 18 (7) (2013) 581–590. [4] World Auto Steel. Advanced high strength steel (AHSS) application guidelines. Ebook available at www. worldautosteel.org 2009; Version4.1,1–30. [5] D. Westerbaan, D. Parkes, S.S. Nayak, D.L. Chen, E. Biro, F. Goodwin, Y. Zhou, Effect of concavity on tensile and fatigue properties in fiber laser welding on automotive steels, Sci. Technol. Weld. Join. 1-9 (2013). [6] W. Xu, D. Westerbaan, S.S. Nayak, D.L. Chen, F. Goodwin, Y. Zhou, Tensile and fatigue properties of fiber laser welded high strength low alloy and DP980 dual-phase steel joints, Mater. Des. 43 (2013) 373–383. [7] D. Parkes, W. Xu, D. Westerbaan, S.S. Nayak, Y. Zhou, F. Goodwin, S. Bhole, D.L. Chen, Microstructure and fatigue properties of fiber laser welded dissimilar joints between high strength low alloy and dual-phase steels, Mater. Des. 51 (2013) 665–675. [8] M. Tumuluru, The effect of coatings on the resistance spot welding behavior of 780 MPa dual phase, Weld. J. 86 (2007) 161–169. [9] Y.G. Kim, I.J. Kim, J.S. Kim, Y.I. Chung, D.Y. Choi, Evaluation of surface crack in resistance spot welds of Zn-coated steel, Mater. Trans. 55 (1) (2014) 171–175. [10] I. Ibrahim, Y. Uematsu, T. Kakiuchi, Y. Tozaki, Fatigue behavior of dissimilar Al alloy/ galvanized steel friction stir spot welds fabricated by scroll grooved tool without probe, 3rd Japan–China Joint Symposium on Fatigue of Engineering Materials and Structures, 2014 Nov. 7–8, Takayama, Japan. [11] W. Xu, D. Westerbaan, S.S. Nayak, D.L. Chen, F. Goodwin, E. Biro, Y. Zhou, Microstructure and fatigue performance of single and multiple linear fiber laser welded DP980 dual-phase steel, Mater. Sci. Eng. 553 (2012) 51–58. [12] www.eng-tips.com/faqs. [13] K. Kuziakr, S. Waengler, Advanced high strength steels for automotive industry, Archives of Civil and Mechanical Engineering 2 (2008) 103–116. [14] J.E. Gould, S.P. Khurana, T. Li, Predictions of microstructures when welding automotive advanced high-strength steels, Weld. J. 5 (2006) 111–116. [15] E. Biro, J. McDermid, J. Embury, Y. Zhou, Softening kinectics in the subcritical heataffected zone of dual-phase steel welds, The Minerals, Metals & Materials Society 41A (2010) 2348–2356. [16] O. Kwon, S.C. Baik, Manufacture and Application of Advanced High Strength Steel Sheets for Auto Manufacture, POSCO, Pohang, 2005 785–790. [17] N. Farabi, D.L. Chen, Y. Zhou, Tensile properties and work hardening behavior of laser-welded dual-phase steel joints, J. Mater. Eng. Perform. 21 (2012) 222–230. [18] M. Hazratinezhad, N.B. Mostafa Arab, A.R. Sufizadeh, M.J. Torkamany, Mechanical and metallurgical properties of pulsed neodymium-doped yttrium aluminum garnet laser welding of dual phase steels, Mater. Des. 33 (2012) 83–87. [19] X.P. Su, N.Y. Tang, J.M. Toguri, Thermodynamic evaluation of the Fe–Zn system, J. Alloys Compd. 325 (2001) 129–136. [20] J. Harders, H. Baeker, Mechanical Behavior of Engineering Materials: Metals, Ceramics, Polymers, and Composites, 1st ed. Springer, New York, 2007. [21] M.D. Chapetti, N. Katsura, T. Tagawa, T. Miyata, Static strengthening and fatigue blunt-notch sensitivity in low-carbon steels, Int. J. Fatigue 23 (2001) 207–214. [22] M. Acarer, B. Gulenc, F. Findik, Investigation of explosive welding parameters and their effects on microhardness and shear strength, Mater. Des. 24 (8) (2003) 659–664. [23] M. Acarer, B. Gulenc, F. Findik, The influence of some factors on steel/steel bonding quality on their characteristics of explosive welding joints, J. Mater. Sci. 39 (2004) 6457–6466. [24] F. Findik, Recent developments in explosive welding, Mater. Des. 32 (2011) 1081–1093.
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Materials and Design journal homepage: www.elsevier.com/locate/jmad
Effect of coating on fiber laser welded joints of DP980 steels Q.L. Cui a,b, D. Parkes b, D. Westerbaan c, S.S. Nayak c, Y. Zhou c, D. Liu d,⁎, F. Goodwin e, S. Bhole b, D.L. Chen b a
State Key Laboratory of Rolling and Automation, Northeastern University, 3-11 Wenhua Road, Shenyang, China Department of Mechanical and Industrial Engineering, Ryerson University,350 Victoria Street, Toronto, Ontario M5B 2K3, Canada c Department of Mechanical and Mechatronics Engineering, University of Waterloo, 200 University Avenue West, Waterloo, Ontario N2L 3G1, Canada d Teck Metals Ltd, 2380 Speakman Drive, Mississauga, Ontario L5K 1B4, Canada e International Zinc Association, Durham, NC 27713, USA b
a r t i c l e
i n f o
Article history: Received 4 May 2015 Received in revised form 16 October 2015 Accepted 17 October 2015 Available online 19 October 2015 Keywords: Coating Fiber laser welded DP980 Similar joints Fatigue
a b s t r a c t Advanced high strength steels (AHSS), specifically dual phase (DP) steels, are extensively used to reduce the weight of vehicles. Fiber laser welding (FLW) has been shown to provide welds with superior mechanical properties made at high welding speeds. Galvannealing (GA) and galvanizing (GI) coated sheet steels are extensively used for auto bodies to improve corrosion resistance. As such, the effect of coating type on the microstructure, mechanical properties and fatigue properties of FLW similar joints of DP980 was assessed. There existed spherical tempered martensite in the GA coated DP980 base metal (BM), most likely caused by the galvannealing process. Coating type did not present a noticeable effect on the microstructure of the welded zones. The GI coating, however, could evaporate violently during welding, resulting in high concavity and other welding defects which negatively affect the high-cycle fatigue property. Measures to mitigate the negative effects of the GI coating were discussed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Weight reduction of auto bodies plays an important role in improving the fuel economy of vehicles to meet regulations [1]. Advanced High Strength Steels (AHSS) are being developed to reduce weight while enhancing crashworthiness of auto bodies. Dual phase (DP) steel is popular favorable type of AHSS because of its good processability and balanced combination of strength and ductility [2,3]. Laser welding is a desirable joining technology because of its high energy density and fast speed result in a small heat affected zone (HAZ) with fine weld metal quality. In the automotive industry the demand for laser welded blanks (LWBs), also known as tailor welded blanks, has increased over the last decades owing to the need to reduce vehicle weight and CO2 emissions. LWBs consist of two or more sheets of similar or different materials, properties, thickness, and surface conditions which are welded together to form a blank. LWBs are generally fabricated in the butt joint configuration to form an assembly, which is subsequently formed to make the desired three-dimensional shapes. In general, LWBs are made with conventional mild (low carbon) steel or interstitial free steel; in recent years the weight reduction potential of LWBs has been realized via material down gauging by employing stronger steels ⁎ Corresponding author. E-mail address: [email protected] (D. Liu).
http://dx.doi.org/10.1016/j.matdes.2015.10.098 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
like AHSS. Thus, it is important to study the weldability of automotive steels like AHSS via laser welding [4,5]. Among the laser welding methods, fiber laser welding (FLW) produces the highest quality welds at a low cost to the manufacturer due to its high power with low beam divergence, flexible beam delivery, low maintenance costs, high efficiency and compact size [6,7]. Sheet steels with GI and GA coatings are commonly used to construct auto bodies for their excellent corrosion resistance. The GI coating essentially contains layer of elemental Zn with up to 0.5% Al; whereas, the GA coating includes up to 12% Fe because the Zn coating formed in the hot-dip process reacts with the steel substrate during the post-dip annealing process at temperatures over 500 °C to form Zn–Fe intermetallic coating. The effect of these coatings on the behavior of spot welded steels was studied by several researchers [8–10]. Different electrode life behaviors were reported in resistance spot welding of GA and GI coated steels. The aluminum content of the coating was the most significant factor affecting electrode life of GI coated sheet steels. Rich experience in spot welding of GA and GI coated steels affirmed that the GI coating tended to eject molten zinc with zinc vapor whereas GA coating enjoyed the reputation of good weldability. There are reports on laser weldability, microstructure, mechanical and fatigue properties of DP steels [5–7,11]. However, no study has been reported so far on the effect of coating types on the properties of FLW joints. Therefore, the present study reports a comparison of the GI and GA coating types on the
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performance in FLW of DP980 steel sheet. Weld concavity formed during FLW and its effect on the microstructure, microhardness, the tensile and fatigue properties has been reported. 2. Materials and experimental procedure The materials used in this work were DP980 sheet steel with either GI or GA coating. The steel thickness was 1.2 mm. The nominal coating weights per side were 60 g/m2 for the GI and 41 g/m2 for the GA coating. The GA coating weight is typically lower to control the powdering tendency. The chemical compositions of the steels are listed in Table 1. Normally, DP steel contains about 0.10 to 0.15 wt.% carbon and is generally alloyed with various amounts of Mn, Si, Cr and Mo to achieve the required tensile strength. The Carbon Equivalent, Ceq, of DP980 steel is calculated using the Ito–Bessyo equation [12], C eq ¼ C þ Si=30 þ ðMn þ Cu þ CrÞ=20 þ Mo=15 þ V=10 þ Ni=60 þ 5B:
ð1Þ
The two DP980 steels tested in the present study were low in C contents and relatively lean in alloying. They thereby had relatively low Ceq values of 0.236 wt.% (GA) and 0.245 wt.% (GI), as listed in Table 1. Steels with low Ceq typically have good weldability. As these two steels have low and similar Ceq values, it is expected that they both attain similar weldability. Fiber laser welding was carried out using an IPG Photonics YLS-6000 fiber laser system, with a focal length of 200 mm, attached to a Panasonic robotic arm. Welding were made along the rolling direction of the sheets. Butt welds of 100 mm × 200 mm sheets of each material were made to make 200 mm × 200 mm laser welded blanks as shown in Fig. 1a. The welding speed was 20.5 m/min and the laser power was 4 kW, which was chosen based on our earlier study to reduce weld concavity [5]. The fiber laser had a fiber core diameter of 0.3 mm with a laser beam spot size/diameter of 0.6 mm. Welding was performed with a 20° vertical head angle, i.e., the laser beam was at an angle of 70° to the surface of sheets (work pieces). No shielding gas was used during welding; instead, blown air was used to remove vapors from the weld pool and to protect the system optics. Metallographic samples were cut from the weld cross-section, then mounted, ground, polished, and etched with a 2% Nital solution. The etched samples were observed first using a light microscope attached with Clemex image analysis system software and then with a scanning electron microscope (JEOL JSM-6380) equipped with threedimensional fractographic analysis. Vickers microhardness was measured on the polished samples across the weld using a computerized microhardness tester with a load of 200 g and a dwell time of 15 s. Localized strain hardening caused by adjacent indentations was avoided by carefully spacing each indentation. To ensure the accuracy of each test result, two calibration tests were carried out using a standard reference test block before the microhardness tests on the welded joints were conducted. Tensile and fatigue test samples were machined to form transverse welds in accordance with ASTM: E8/E8M, the example of which is indicated by the dashed line in Fig. 1a with the geometry and dimensions of the test coupons shown in Fig. 1b. The specimens were machined in such a way that the weld line was positioned at the center of the gauge length (Fig. 1b). Tensile tests were conducted on a fully computerized united tensile testing machine at room temperature and with a strain rate of 1 × 10− 3 s− 1. An extensometer with a gauge length of
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50 mm and a strain limit of 20% was used to measure the strain during the tensile tests. Load control fatigue tests were performed following ASTM: E466 on a fully computerized Instron 8801 servo-hydraulic testing system. Tension-tension cyclic loading was employed to prevent potential buckling of the samples and a stress ratio of R = 0.1 was used. The test conditions were conducted at room temperature with a frequency of 50 Hz and the load was applied using a sinusoidal waveform. A minimum of two specimens were tested in the tensile tests and in the fatigue tests at each of the cyclic stress amplitude. The fatigue fracture surfaces were examined using scanning electron microscopy (SEM) to study the fatigue fracture mechanisms. 3. Results and discussion 3.1. Microstructure To understand the welding behavior, it is helpful to examine the microstructure of the base material and how it was obtained. Through the heat treatment steps in a continuous galvanizing line, the DP steel is intercritically annealed to form a mixture of ferrite and austenite. Subsequent rapid cooling avoided the formation of pearlite and bainite while promoting martensite formation [13]. The steel then went through a hold zone typically controlled at about 470 °C, before immersing in a molten zinc bath at 460 °C. The resulting microstructure of DP980 steel base metal (BM) is shown in Fig. 2a and b. Both the GA and GI coated DP980 had a microstructure consisting of a ferrite matrix with islands of martensite. It is worth noting that there existed spherical tempered martensite in the GA coated DP980 BM whereas there was none in the GI coated steel. The tempering was caused by the additional heating cycle in the galvannealing process. Microscopic examination of the welds showed the presence of martensite in both the fusion zone (FZ) and the HAZ, as shown in Fig. 3. In laser welding, the weld cooling was extremely rapid. It has been shown through modeling that the cooling rates in laser welding were on the order of 105 °C/s [14]. The critical cooling rate CR for steels to achieve martensitic transformation is given by the following equation, LogCR ¼ 7:42−3:13C−0:71Mn−0:37Ni−0:34Cr−0:45Mo:
ð2Þ
For the DP980 steels used in this study, the critical rate is about 400 °C/s. Therefore, a martensitic structure was obtained in both the FZ and the HAZ. There was limited amount of bainite in the FZ which was attributed to the high cooling rate during the FLW process [15]. The HAZ of the weld appeared to be similar in both coating conditions with an upper-critical HAZ, an inter-critical HAZ, and a subcritical HAZ. In the upper-critical HAZ, the temperature was above the Ac3 line, leading to a microstructure that consisted of mostly martensite with a very small amount of ferrite (Fig. 3c and d) [14]. In the intercritical HAZ (Fig. 3e and f), the temperature was between the Ac1 and Ac3 lines, resulting in the coexistence of ferrite and austenite, the latter of which subsequently transformed into martensite during cooling. In the sub-critical HAZ, the temperature was below the Ac1 line [7]. This caused the martensite that was normally present in the BM to temper and decompose into partially tempered martensite (PTM), as can be clearly seen in Fig. 3g and h. A “soft zone” was thus created where the steel was the weakest [16]. Tempering of BM martensite in the fusion welding of DP980 steel has been reported in several previous studies [17,18]. The amount of PTM in the sub-critical HAZ for both welded joints was very similar.
Table 1 Chemical composition of the DP980 in present study. Steel
C
Mn
Si
Al
Cr
Ni
Mo
Nb
N
Cu
Ti
Ceq
GI GA
0.15 0.1
1.45 2.19
0.33 0.33
0.05 0.04
0.02 0.21
0.01 0.01
0.002 0.238
0.001 0.016
0.01 0.01
0.02 0.02
0.002 0.022
0.245 0.236
518
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Fig. 1. (a) Geometry of the welded sheet, (b) transverse tensile joint machine from the dotted line in (a).
3.2. Microhardness Microhardness measurements of the GA and GI coated DP980 welds are shown in Fig. 4. Both specimens displayed similar trends. The hardness was the highest in the FZ, about 100 HV higher than that of the BM, due to its predominant martensitic microstructure (Fig. 3). Outside the FZ, the hardness decreased abruptly from the maximum value in the upper-critical HAZ to the minimum value in the sub-critical HAZ, corresponding to the change in microstructure from all martensite to a mixture of ferrite and partially tempered martensite. As the distance from the center of the weld further increased, the heat effect diminished and, hence, the hardness gradually recovered to the level of that of the BM. In this particular study, the hardness of the GA coated specimen was lower than that of the GI coated specimen with an average difference of about 36 HV. This was largely attributable to a) the difference in steel composition, in particular the carbon content and b) the effect of galvannealing. As shown in Table 1, the carbon content of GI coated DP980 was 0.15 wt.%, which was 50% higher than the 0.10% of its GA counterpart. Hence, the carbon content in martensite was higher for the GI coated steel, leading to higher hardness and strength. In addition, the galvannealing process caused tempering of the martensite (Fig. 2a), resulting in reduced hardness and strength.
3.3. Galvanized coatings and weld defects Cross-sectional views of both GI and GA coatings are shown in Fig. 5. The location of the weld was to the left of these micrographs. Fig. 5a shows a typical GI coating, taken at 0.74 mm from the edge of FZ. The coating consisted of one single layer of zinc solid solution, the composition of which matched the composition of a galvanizing bath. Between
the zinc layer and the steel substrate existed an Al–Fe intermetallic layer, which was too thin to observe at this magnification. This thin Al–Fe intermetallic layer prevented the reaction between Fe and Zn, and was therefore called the inhibition layer. The thickness of the GI coating, as shown in Fig. 5a, was about 12 μm. Fig. 5b shows a typical GA coating taken at 0.60 mm from the edge of FZ. The coating consisted of two layers of Zn–Fe intermetallic compounds. The upper layer (δ phase) had an average Fe content of 11.2 wt.%, which was typical for the so-called “all-δ” GA coating. The thickness of the GA coating was about 7 μm. Because thick GA coatings tend to have formability issues such as powdering and flaking, only thin GA coatings are normally produced. Fig. 5c was taken in the HAZ, 0.29 mm from the edge of FZ. At this location, the heat from welding was enough to melt the GI coating (melting point 420 °C). Silver lines of very fine Fe–Zn intermetallic precipitates can be clearly seen along the boundaries of large zinc grains. Fe–Zn intermetallic particles are also found at the coating/substrate interface due to the breakdown of the inhibition layer which typically occurs at temperatures above 500 °C. Considering the short exposure time to the peak temperature (fast heating and cooling in laser welding), to have this extent of Fe–Zn reaction, the peak temperature is estimated to be over 550 °C. This location was in the soft zone shown in Fig. 4, which also indicated an exposure to an elevated temperature. Fig. 5d shows the GA coating also at 0.29 mm from the edge of FZ. It is interesting to note that the coating in the left half of the micrograph had high porosity whereas the right half was largely dense. The coating porosity increased progressively to the left all the way to the edge of the FZ. To the right of the micrograph, the GA coating was completely dense like the one shown in Fig. 5b. Considering the diffusivity of zinc is two orders of magnitude higher than that of iron in Fe–Zn intermetallic compounds, it is possible that the porosity can form due to Kirkendall effect. However, the abrupt transition from high porosity to no porosity
Fig. 2. BM microstructure of DP980 (a) GA (b) GI (F: ferrite, M: martensite, PTM: partially tempered martensite).
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Fig. 3. Microstructure of the DP980 similar welding with GA and GI coating, FZ (a) GA and (b) GI, upper-critical HAZ (c) GA and (d) GI, inter-critical HAZ (e) GA and (f) GI, tempered subcritical HAZ (g) GA and (h) GI (F: ferrite, M: martensite, B: bainite, PTM: partially tempered martensite).
indicated that the liquid phase was involved in the rapid diffusion that led to the porosity. The δ phase coexisted with liquid zinc in the temperature range between 530 °C and 672 °C. The measured Fe content of 11.2% was near the upper limit of the δ phase composition range (7.0– 11.5%). Therefore, the temperature should be near the upper end of the temperature range and was calculated to be 658 °C using a thermodynamic database [19]. The thickness of the GI coating started to decrease when the distance to the FZ was 0.04 mm, partially due to the evaporation of zinc at temperatures exceeds its boiling point of 907 °C. This temperature was above the Ac3 temperature of DP980 and this location was therefore in the upper-critical HAZ, which was also reflected in the hardness measurement shown in Fig. 4. The GI coating
evaporated completely at the distance of 0.01 mm from the FZ. In comparison, a good portion of the GA coating survived direct contact with molten steel right to the edge of the FZ, exhibiting better stability at high temperatures. This agreed well with prior observations that GA coating displayed lower zinc activity [2,9,10]. The entire cross-sections of FZ of GI coated and GA coated samples are shown in Fig. 5e and f, respectively. In general, greater material loss in FZ was observed with GI coated samples, which led to larger concavity and more severe welding defects. It was confirmed that concavity formation was due to the metal ejection [5]. Relevant experience in spot welding indicated that the evaporation of zinc in the GI coating could cause ejection of molten metal whereas the GA coating exhibited good
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Fig. 6. Tensile curves of GI and GA coated DP980 similar welded joints. Fig. 4. Microhardness curves of GA and GI coated DP980–DP980 welded joints.
stability. Moreover, because the thickness of the GA coating is almost half that of the GI coating, it is reasonable to think that the GA coating in this case would offer lower resistance [8]. In the present study the thick GI coating used further deteriorated the performance of GI coated weld in comparison to the thin GA coating. The zinc content in the FZ of both GI and GA specimens was below the detection limit, indicating that essentially all zinc had evaporated in FZ. 3.4. Tensile properties Select curves of engineering stress versus engineering strain of GA or GI coated DP980 specimens with or without a weld (i.e., base metal) are displayed in Fig. 6, while the average properties of all specimens are listed in Table 2. The BM specimens exhibited good repeatability
whereas the welded specimens displayed large variation, especially in elongation. The average ultimate tensile strength (UTS) of GA BM was 936 MPa, which was notably lower than 1083 MPa of GI BM. The causes of lower ultimate tensile strength of GA BM, similar to those discussed earlier in microhardness, were (1) lower carbon content in steel chemistry and, (2) the softening of materials due to the GA process. The softening was also reflected in the slight increase in elongation, with 13.6% for GA BM versus 12.1% for GI BM. For the welded specimens, their UTS (946 MPa for GA and 999 MPa for GI) deviated marginally from their BM counterparts. The limited decrease in strength was attributable to the narrow and mild soft zones (Fig. 4) which were achieved by having a small laser beam spot size and high power intensity in conjunction with a high welding speed. The elongation of welded specimens was, however, significantly lower with 8.6% for GA and only 3.8% for GI. The elongation of the welded specimens was also much more scattered. For instance, the elongation of the four GI coated weld specimens
Fig. 5. (a) Normal GI coating outside of HAZ, 0.74 mm from the edge of FZ; (b) Normal GA coating outside of HAZ, 0.60 mm from the edge of FZ; (c) GI coating inside HAZ, 0.29 mm from the edge of FZ; (d) GA coating inside HAZ, 0.29 mm from the edge of FZ; (e) Weld of GI coated DP980; and (f) Weld of GA coated DP980.
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Table 2 Summary of fatigue and tensile properties. Steel
YS (MPa)
UTS (MPa)
Elongation %
Joint efficiency %
Fatigue limit (MPa)
Fatigue ratio
σ'f (MPa)
b
GA joints GA BM GI joints GI BM
667 644 717 707
946 934 999 1083
8.6 13.6 3.8 12.1
100 – 92.2 –
100 200 b100 250
0.11 0.21 b0.1 0.23
1370 733 4130 682
−0.153 −0.077 −0.260 −0.060
ranged from 1.3% to 6.2% with a standard deviation of 2.4%. It is, hence, not possible to select one single representative curve to reflect the scattering nature of this property. To better illustrate the complex situation and avoid overcrowdings of the figure, two curves, one with the lowest elongation and another with the greatest elongation, were displayed in Fig. 5 for each of the GA weld and GI weld specimens. Pictures of the fractured tensile specimens are shown in Fig. 7. The GA BM specimen (top specimen in Fig. 7a) fractured along a plane at an angle of about 60° with the pulling direction, indicating the ductile failure of slipping caused mainly by shear stress. Necking was also apparent, offering additional indication of ductile failure. The fracture process of the GI BM specimen (bottom specimen in Fig. 7a) appeared to initiate from two edges via slipping caused by shear stress, and then completed with fracture caused by normal stress at the center of the specimen. For the welded GA specimen with the smallest elongation (top specimen in Fig. 7b), fracture appeared to initiate from both edges in the sub-critical HAZ. Cracks tended to propagate within the narrow soft band of sub-critical HAZ (90° angle with the pulling direction) but at the same time were influenced by shear stress which tended to lead to a failure at a 45° angle. The reduced elongation was likely attributable to the more localized necking near the sub-critical HAZ. For the welded GA specimen with the greatest elongation (lower specimen in Fig. 7b), fracture occurred outside the HAZ, demonstrating that
welding did not cause significant weakening to this specimen. The nature of failure was no different from that of the GA BM specimen shown in Fig. 7a. For the welded GI specimen with the smallest elongation, its stress–strain curve shown in Fig. 6 indicated premature failure without apparent plastic deformation. Fracture appeared to occur along the HAZ/FZ interface on the left-hand side of the FZ (top specimen in Fig. 7c). Thus, the fracture surface was perpendicular to the pulling direction. For the welded GI specimen with the greatest elongation, fracture appeared to initiate along both HAZ/FZ interfaces on either side of the FZ, and then finished with a fracture crossing the FZ. It is worth noting that the microstructure in the FZ and upper-critical HAZ was mainly composed of martensite (Fig. 3a-d) and the hardness was the highest (Fig. 4). Failure initiated at the HAZ/FZ interface was, therefore, unlikely caused by low strength in these regions. Close examination of the weldment found that the welded GA specimen suffered only mild concavity and mild misalignment (Fig. 5f) while the welded GI specimen suffered severe concavity and notable misalignment (Fig. 5e). The crack (about 34 μm in length) in the upper-right corner of the FZ in Fig. 5e was found to extend another 135 μm into the FZ when examined carefully at a higher magnification. The difference in concavity resulted from the difference in the zinc vapor pressure which was higher for the pure zinc GI coating than for the GA coating where zinc was tied up with iron in more stable Fe–Zn intermetallic compounds. The volatility of pure zinc GI coating could cause a loss of molten metals in the FZ, resulting in a large concavity. The volatility could also generate turbulence and create weld defects, resulting in stress concentration which was particularly detrimental to hard and brittle martensitic steel. 3.5. Fatigue and fractography Fig. 8 displays the S-N curves of the GA and GI coated welded joints with their corresponding BMs. For stress amplitude of 350 MPa and above, there was no notable difference in fatigue life between welded GA specimens and their corresponding BM specimens. Put it in perspective, a stress amplitude of 350 MPa had a maximum stress of 778 MPa while a stress amplitude of 450 MPa had a maximum stress of
Fig. 7. Pictures of fractured tensile specimens corresponding to curves shown in Fig. 5. (a) BM of GA and GI (b) Lowest and greatest elongation of welded GA (c) Lowest and greatest elongation of welded GI
Fig. 8. Comparison of S–N curves of the GA and GI coated similar DP980 welded joints with its corresponding BM.
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1000 MPa which was higher than the UTS of GA BM (936 MPa as shown in Fig. 5) tested using a strain rate of 1 × 10−3 s−1. Granted, the strain rate in a fatigue test was over two orders of magnitude higher than the strain rate used in the tensile test. Nevertheless, the stress level was very high and both GA welded specimens tested using a stress amplitude of 450 MPa failed almost instantly in just over 1 s with the fracture occurred in the base metal at an angle of about 60° to the pulling direction, almost identical to the bottom specimen shown in Fig. 7b for the tensile test. At the stress amplitude of 350 MPa, the test started to resemble a typical fatigue test with specimens lasted for over 1 × 104 cycles (several minutes) and a steady crack growth stage (Stage II) being distinctly identified with flat and smooth fracture surfaces perpendicular to the loading direction. The Stage II fracture surfaces were close to the FZ/HAZ interface and their area was about 20% of the cross-sectional area of the GA specimens. The final ductile failure occurred with rapid crack growth along a plane about 45° in the thickness direction. This plane ran across the entire HAZ, which had a width of less than 0.8 mm as determined by the micro-hardness measurement shown in Fig. 4 and by the coating morphology analysis in Fig. 5. It is well known that low-cycle fatigue is closely related to plastic deformation. Hence, the softening effect of the GA process decreased the fatigue life at high stress amplitudes. That was the main reason why the GI BM specimens outperformed the GA BM specimens in Fig. 8 with stress amplitudes greater than or equal to 350 MPa. The softening in the HAZ had a less significant effect on the already softened GA welded specimens. Therefore, there was little difference between the GA welded and GA BM specimens. The softening effect in the HAZ was more significant for GI coated specimens (Fig. 4). Thus, more notable difference in fatigue life was observed between GI welded and GI BM specimens, as shown in Fig. 8. The severe concavity in GI welded specimens at the FZ/HAZ interface where cracking initiated further decreased their fatigue life. Unlike the low-cycle fatigue which is related to plastic deformation and sensitive to soft zones, the high-cycle fatigue is more sensitive to defects and stress concentration. Harders and Baeker [20] pointed out that the fatigue strength of a material under dynamic cyclic loading was much more sensitive to the manufacturing process and materials than the static strength, and in particular for the lower level cyclic load. The severe concavity formed a notch similar to a notched specimen studied by Chapetti et al. [21] where the stress concentration was generated, and consequently a decreased fatigue limit or fatigue life at low level stress amplitudes was observed. In addition to the concavity, severe defects formed in FZ of the welded GI specimens. Fig. 9 shows the top view of a weld of a GI specimen undergone a fatigue test with stress amplitude of 150 MPa. The upper portion of the picture shows severe turbulence of molten melt close to the FZ/HAZ interface. The lower portion of the picture shows another FZ/HAZ interface where Stage II fracture occurred. On the weld surface scattered a great number of
small pieces of surface oxide, presumably broken by the thermal stress during the fast cooling following the solidification of the weld. Dendritic grains are evidently visible along the center of the weld. The weld defects were markedly less severe for GA specimens. Note that the important Stage II crack growth took place at the FZ/HAZ interface (perpendicular to loading direction) and the Stage II area increased with decreasing stress amplitude. At the stress amplitude of 150 MPa, the Stage II area became over 60% of the cross-sectional area, thereby playing a very important role in determining the fatigue life of welded specimens. Consequently, the welded GA specimens performed significantly better than their GI counterparts at low stress magnitudes. For the lowest stress magnitude (100 MPa) tested in the present study, the welded GI specimens failed after 4 × 105 cycles on average, whereas the welded GA specimens survived the entire test of over 1 × 107 cycles without failure. For BM specimens, GA BM specimens had the disadvantage of a lower tensile strength and a brittle coating which was prone to microcracking. As a result, the GA BM specimens had shorter fatigue life than their GI BM counterparts, as shown in Fig. 8. The GI BM specimens survived the test at a stress amplitude of 250 MPa while the GA BM specimens survived the test at a lower stress amplitude of 200 MPa. Notable improvement in fatigue life can be expected, especially for the GI coated specimens, if the weld direction does not run perpendicular to the loading direction. As such, the critical Stage II crack propagation cannot take place along the defect-laden FZ/HAZ interface. According to the experience of the explosive welding [22–24], it is supposed that an angled, Chevronor curved weld can effectively force the Stage II crack to propagate away from the FZ/HAZ and into the BM, and hence improve the fatigue performance of the welded part. The fatigue limit and fatigue ratio (i.e., a ratio of fatigue limit to the UTS) are tabulated in Table 2. The decreases in fatigue limit for GA and GI welded specimens, with respect to their BM counterparts, were 50% and over 60%, respectively. The fatigue limit of the GA welded specimens was higher than that of the GI welded specimens. The fatigue ratios was 0.11 for the GA welded specimens and was below 0.1 (fatigue limit lower than the experimental range and was undetermined) for the GI welded specimens. These results suggested that while the tensile strength of the FLW DP980 joints with GI coating was higher than that of the joints with GA coating, the fatigue strength was actually better for GA welded specimens at a lower level of stress amplitudes (≤150 MPa). From Fig. 8, it can also be seen that GA welded joints exhibited a much higher fatigue strength in the low cycle fatigue regime below ~ 5 × 104 comparing with that of GI welded joints. The S-N plot in Fig. 8 is expressed using the Basquin equation [17], 0
σ a ¼ σ f 2N f
b
ð3Þ
where σa is the cyclic stress amplitude, σ′f is the fatigue strength coefficient defined by the stress intercept at 2Nf = 1, Nf is the number of cycles to failure (2Nf is the number of reversals to failure), and b is the fatigue strength exponent. The values of these parameters obtained by fitting the data points according to Eq. (3) are tabulated in Table 2. The fatigue life of the welded joints could be estimated based on the values of σ′f and b. Eq. (3) indicates that the higher the value of σ'f and the smaller the absolute value of b, which is always negative, the longer the fatigue life. From Table 2, it can be seen that the GA welded joints have a better b value and a higher fatigue limit which means that they would have a better fatigue life in comparison to the GI coated samples. 4. Conclusions
Fig. 9. The top surface of a welded GI specimen undergone a fatigue test with a stress amplitude of 150 MPa.
The aim of this study was to take a comparative look at the welding performance of GI and GA coatings. It should be noted that this study pertained only to the steel compositions and the coating weights selected. The following conclusions could be reached based on the study
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conducted to examine the differences between GA and GI coated DP980 steels: 1. Due to the annealing process of the GA coating process, there was more spherical tempered martensite in the GA coated DP980 BM than in the GI coated DP980 BM. Coating type had no effect on the microstructure of the welded zones. 2. The hardness of GA coating welded joint was generally lower than that of GI coating from the BM to the FZ. In the BM, this was partially attributable to the PTM that was seen in the BM of GA coated DP980. In FZ and HAZ, the hardness difference between the GA and GI welded joints stemmed mainly from the difference in chemical composition. The UTS and YS of the GA coated welded joints were lower than that of GI coated samples. This was affirmed by the hardness results. 3. The GA similar welded joints showed improved fatigue resistance compared to GI welded joints due to lower concavity resulting from the coating characteristics of GA and from its low coating thickness. Acknowledgments The authors would like to thank the National Natural Science Foundation of China (No. 51304045), the China Scholarship Council, Fundamental Research Funds for the Central Universities of China (No. N110407001), Natural Sciences and Engineering Research Council of Canada (NSERC) and AUTO21 Network of Centers of Excellence for providing financial support. The financial support from International Zinc Association (IZA) and Arcelor Mittal Dofasco is highly acknowledged. One of the authors (D.L. Chen) is grateful for the financial support by the Premier's Research Excellence Award (PREA), NSERC-Discovery Accelerator Supplement (DAS) Award, Canada Foundation for Innovation (CFI), and Ryerson Research Chair (RRC) program. The authors would like to thank Dr. J. Chen and Dr. Y.L. He (CANMET-Materials Technology Laboratory, Natural Resources Canada, Hamilton, Canada), Mr. E. Biro (Arcelor Mittal Global Research, Hamilton, Canada), and Dr. J. Villafuerte (Center Line (Windsor) Ltd., Windsor, Canada) for their support and helpful discussion. The assistance of Q. Li, A. Machin, J. Amankrah, and R. Churaman in performing the experiments is gratefully acknowledged. References [1] L.F. Mei, G.Y. Chen, X.Z. Jin, W.Q. ZhangY, Research on laser welding of high-strength galvanized automobile steel sheets, Opt. Lasers Eng. 47 (2009) 1117–1124.
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