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Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input
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Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input Hakan Aydin, Tracy W. Nelson
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Received date: 24 April 2013 Revised date: 20 July 2013 Accepted date: 22 July 2013 Cite this article as: Hakan Aydin, Tracy W. Nelson, Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2013.07.090 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input Hakan AYDIN∗a, Tracy W. NELSONb
a)
Mechanical Engineering Department, Uludag University, 16059, Gorukle-Bursa, Turkey
b)
Mechanical Engineering Department, Brigham Young University, 435 CTB, Provo, UT 84602, USA.
Abstract The study was conducted to investigate the microstructure and mechanical properties of the hard zone in friction stir welded X80 pipeline steel at different heat inputs. Microstructural analysis of the welds was carried out using optical microscopy, transmission electron microscopy, and microhardness. Heat input during friction stir welding process had a significant influence on the microstructure and mechanical properties in the hard zone along the advancing side of the weld nugget. Based on the results, the linear relationships between heat input and post-weld microstructures and mechanical properties in the hard zone of friction stir welded X80 steels were established. It can be concluded that with decreasing heat input bainitic structure in the hard zone become finer and so hard zone strength increases. Keywords: X80 pipeline steel; Friction stir welding; Hard zone, Hardness; Tensile Properties.
∗
Corresponding author: Assistant Professor Dr. Hakan AYDIN Tel.: +90 224 294 06 52; fax.: +90 224 294 19 03; E-mail address: [email protected] Postal address: Uludag University, Engineering and Architecture Faculty, Mechanical Engineering Department, 16059, Gorukle-Bursa, Turkey.
1. Introduction The American Petroleum Institute (API) X80 steel has gained increasing attention from pipeline companies for use in oil and gas transmission pipelines [1]. The commercial X80 API-grade steels are produced by Thermo-mechanically Controlled Processing (TMCP), which produces uniform refined microstructures consisting of ferrite and bainite [2]. These refined microstructures provide superior combination of high-strength and excellent toughness. Welding is an essential process to make large-scale pipelines for long distance transportation of crude oil or natural gas under high pressure. During conventional fusion welding processes, linepipe steels often meet with some problems, such as significant loss of strength and toughness in the heat affected zone (HAZ). Moreover, in conventional fusion welding of linepipe steels, the HAZ is susceptible to hydrogen-assisted-cracking (HAC). The increasing relevance of X80 steel in oil and gas transportation operation requires research on more efficient and reliable joining processes. In recent years, friction stir welding (FSW) has shown considerable promise for joining aluminum alloys, as well as for magnesium, copper, titanium and steel [3-5]. The oil and gas industry has a number of potential applications for this technology. However, a more complete understanding of the technological process and the achieved mechanical behavior of FSW in steels is needed. FSW makes use of a nonconsumable welding tool to induce severe plastic deformation of the workpiece material, generating both heat and mixing across the joint [6,7]. FSW offers several advantages over conventional fusion welding processes, due to its low heat input and absence of melting and solidification process [5]. Advantages include better mechanical properties, low residual stress and distortion, and reduced occurrence of defects [3,4]. In addition, it is expected that hydrogen induced cracking in steels will be reduced or eliminated. These advantages are likely to make FSW an attractive process for joining linepipe steels.
While many studies have focused on FSW of the commercial Al alloys, such as 2XXX [8-11], 6XXX [12-15] and 7XXX [16-18], there is considerable interest in extending this technology to mild steel and other low to medium carbon ferrous alloys [2,19-22]. The mechanical properties of FSWed carbon steels are partially improved compared with the base metals, because FSW achieves grain refinement in the stir zone of the carbon steel, similar to Al alloys [19,20]. The microstructure and mechanical properties of friction stir welds are dependent on the chemical composition and the microstructure of the base metal (BM), the thermo-mechanical conditions during FSW process, and the cooling rate within the critical temperature range. The thermal cycles and high strain rates imposed during the FSW process of X80 steel lead to localized brittle hard zone (HZ) region in the weld. It has been shown that this localized HZ may play a significant role in the fracture toughness of FSW X80 steel [23]. Therefore, the microstructure and mechanical properties of this zone need to be characterized in more detail. Although there are some studies on microstructure and mechanical properties of FSW high strength low alloy (HSLA) steels in the literature [24-28], only a few focused on FSWed X80 steel [1,2,29,30]. Of these, there has been no work on the microstructural and mechanical properties of HZ region. Therefore, the objective of the present work was to evaluate the microstructure and mechanical properties of HZ region in FSWed X80 steel relative to different heat input (HI).
2. Experimental Details This study was performed on a commercial grade API X80 pipe-line steel with chemical composition (wt.%) listed in the Table 1. Before welding, oxide and surface scale were removed from both sides of the plate by grinding followed by degreasing with a methanol solvent. Test plates with dimensions of 1016 x 127 x 11 mm were friction stir welded parallel
to the rolling direction. Schematic illustration of the FSW process can be seen in Fig.1. A Polycrystalline cubic boron nitride (PCBN) CS4 tool was used for all welds (Fig.2). An argon gas atmosphere, at a flow rate of 1.1m3/h, was used to prevent surface oxidation of the weld and tool during the weld cycle. All welds were made under a depth-controlled process with a head tilt of 0.5º. Both single and double-sided welds were evaluated. When double sided, plates were flipped longitudinally to maintain advancing and retreating side symmetry. The FSW welding parameters with corresponding HI used in this investigation are listed in Table 2. Samples were removed transverse to the weld using a water-jet cutter from each weld for post-weld metallographic and microhardness characterization. After that, samples were mounted in bakelite then ground and polished successively through 1µm diamond paste, etched with 2% Nital reagent for 20 seconds, and analyzed optically using an Olympus GX51 microscope. Vickers micro-hardness measurements were performed on a LECO LM 100AT microhardness tester using a diamond pyramid indenter with a 500 gram load, 15 second dwell time and 400μm inter-indentation spacing. For micro-tensile test specimens, the FSW samples were cut parallel to the welding direction using a water-jet cutter. Then, longitudinal through-thickness tensile test specimens were removed from the weld zone of the samples by an electrical discharge cutting machine (EDM). The configuration and dimensions of the longitudinal tensile test specimens can be seen in Fig.3. Tensile tests were carried out on an Instron universal testing machine (Model No. 123), at room temperature and with a crosshead speed of 0.762 mm/min. Transverse samples for transmission electron microscopy (TEM) studies were removed as 3 mm diameter cylinders by cutting with the EDM from selected areas of the samples. 300 µm thick discs were sliced by a low speed diamond saw from these cylinders. Discs were
mechanically thinned from 300µm to 70µm by grinding with 1200 grit SiC water-proof abrasive paper. Finally, the electron transparent thin sections for TEM analysis were prepared by means of double jet electro-polishing, using a solution of 10% (volume fraction) perchloric acid and 90% glacial acetic acid with 20V below 0 ºC in the icy water. The thin foil specimens were examined in a Philips Tecnai F30 TEM FEG model transmission electron microscope at 300 kV to observe the microstructural details.
3. Results and Discussion
3.1. Microstructure Fully consolidated welds X80 steel were produced at all weld conditions. No macroscopic defects, such as voids, cracks or distortion, have been observed in the inspection of the welds. The macroscopic views of the welds at low magnification reveal distinct changes in microstructures at the different parameters investigated (Fig. 4). The various friction stir weld zones are indicated in Fig. 4a. The HZ is located at the top surface and in the weld nugget (WN) along the advancing side (AS). As can be seen in Fig. 4b-d, HI has a strong effect on as-welded microstructure, especially in the HZ region. The HZ and HAZ regions expand with increasing HI. In addition, higher HI produced a more swirl-like microstructure pattern that was asymmetric about the weld centerline. The optical microscopy observations performed on the cross-section of the single-sided welded sample revealed the distinct changes in microstructures of the weld zones (Fig.5). As can be seen in Fig.5a, the BM microstructure is mainly composed of elongated fine-grained polygonal ferrite, with the average grain size about 6 µm, and a small amount of refined upper bainite islands which are almost parallel to the rolling direction of the plate. The post-weld microstructures in the weld zone are a function of the thermo-mechanical cycles experienced
during the welding process. There is no evidence of BM microstructures in the WN and HZ regions (Fig. 5 b, c). The elongated grains in BM have completely transformed to the bainitic structure with well-defined laths in the HZ region. This bainitic microstructure consists of thin, relative straight and long, parallel ferrite laths with discontinuous carbide particles (second phases) at the lath boundaries. This microstructure indicates that the HZ region has reached a peak temperature in excess of A3, which is the temperature at which ferrite transforms to austenite on heating, and the cooling rate is sufficient to form lath bainite. In addition, Wei et. al. [26] reported that the HZ region experienced high strain and/or strain rate which aided in the transformation to austenite. Fairchild et al. [31] have also reported that the temperature and strain rates were very high at the top surface and on the AS compared to other regions in the weld. The WN microstructures are predominantly coarse bainitic structure with carbides along the lath boundaries (Fig.5c). Bainite lath boundaries in the WN region are more irregular that those observed in the HZ region. This is likely the result of deformation during the transformation in the high heat input welds. Additionally, the prior austenite grain (PAG) boundaries are discontinues in both the WN and HZ microstructure [26,27]. The elongated ferrite grains in the BM underwent recrystallization forming more refined equiaxed polygonal ferrite in the HAZ. It is likely that temperatures in the HAZ did not exceed the A3 temperature, but the temperature was sufficient to cause significant coarsening and spheroidization of the carbides to occur. The TEM images of the WN and HZ microstructures of the single-sided welded sample are shown in Fig.6. The microstructure in WN consists of nonequiaxed and interwoven nonparallel lathy ferrite with dense density dislocations, and coarse carbide phases distributed intra-lath and along the lath boundaries (Fig.6a,b). The dark phases in this matrix are
Martensite/Austenite (M/A) constituents or retained austenite [32]. The formation of M/A constituents may be attributed to the partitioning of carbon during the transformation to bainite and the post-transformation of carbon-enriched austenite. Blocky M/A constituents are generally beneficial to the improvement of toughness compared with M/A strip [33]. Thin M/A strips are observed in the well-defined lath boundary locations in the HZ microstructure (Fig.6c,d). Give this; the HZ may exhibit lower toughness than the WN region as it has more M/A strips as reported by Tribe [23]. The optical micrographs of the HZ region of the welded samples with different HI are shown in Fig.7. These figures show that the microstructures in the HZ are similar, consisting of predominantly classical lath-bainitic morphologies with more well defined PAG boundaries. It is clear from these micrographs that HI has a significant influence on microstructure in the HZ region. The HZ regions of all samples display relatively coarse PAGs which will have an effect on the resulting transformation product. The PAG size measurement was performed using a method described by Wei [35]. In this method, as seen in Fig.8, a circle is placed on the top of each PAG and the intercepted length on each diameter is measured. The grain size of a PAG is the average value of these measurements. The traditional ASME intercept or planimetric methods are not suitable to use due to inadequate numbers of PAGs in the micrographs. Wei [35] stated that this method was good to measure the large grain size in the image with limited number grains. Roughly 60% decrease was observed in the PAG size at the HZ in the lowest HI (1151 J/mm) compared to the highest HI (2848 J/mm). In addition, as can be seen in Fig.9a, PAG size increases linearly with increasing HI. Higher HI produces higher peak temperatures, which coarsens PAGs size. In addition, higher HI results in slower cooling rate [27,34]. It is well known that cooling rate plays a crucial role in determining weld microstructure. The polygonal and allotriomorphic grain boundary ferrite may occur at relatively slow cooling rates, while the intermediate transformation products, i.e., bainite, may
be observed at relatively high cooling rates. The bainitic lath structure becomes thinner with increasing cooling rate. Cooling rates in the HZ regions of the welds investigated were fast enough for the formation of lath bainite. Cota et al. [36] has reported that when the cooling rate increases, the reduction in the bainite start temperature led to an increase in the driving force for the nucleation rate of sub-units of ferrite and, consequently, a decreased in the width of bainitic laths. Higher HI and slower cooling rate provides favorable diffusion conditions promoting bainite coarsening in the HZ region. The lath width of the bainitic structure in the HZ region is finest as the lowest HI (fastest cooling rate) compared to the samples with the highest HI. This is shown in the TEM images in Fig.10. The bainite lath widths were measured using a mean intercept method on the TEM micrographs: the lath width at lowest HI is 0,64 µm compared to the 0,90 µm lath width at the highest HI. As expected, similar to the PAGs size, bainite lath width in the HZ regions also increases almost linearly with increasing HI (Fig.9b).
3.2. Microhardness In order to examine the property variations across the welds, microhardness was measured on transverse cross sections of each weld. Fig.11 shows the hardness maps from three different heat input FSW in X80 steel. These microhardness maps show several features in the weld zones: HAZ softening on both sides of the weld; slightly higher hardness of the WN compared to the BM; and the HZ that varies in size and distribution as a function of HI. The non-uniform distribution of hardness correlates to a non-uniform heat gradient generated in the stir zone. Higher strains on the AS of the weld result in an asymmetric heat generation across the stir zone. This also means that there is a larger difference in material properties between the AS and retreating side (RS) of the weld. In Fig.11, the hardness peaks commonly follow a line on the AS of the weld close to the location of the shoulder and edge of the weld
tool. This is likely the region of greatest shear strain which would result in greater heating [37]. If this were the case, higher peak temperatures would produce larger PAG size and subsequent transformation to completely lath-bainitic microstructure, as observed in the microstructure examination. The HAZ has the lowest hardness values due to the predominantly overtempered polygonal ferrite microstructure at elevated temperatures during FSW process. The minimum hardness value in the HAZ decreases due to the higher amount of heat transferred to the weld (Fig.11b,c). The influence of the HI on the microhardness of weld zone can be seen in Fig.11b,c. The hardness of the WN decreases with increasing HI. Although the area of HZ region expands almost to the RS of the weld with increasing HI (Fig.11b,c), the peak hardness in the HZ regions decreases with increasing HI similar to the rest of the WN (Fig.12). The increase in HZ area is likely the result of a large fraction of re-austenitization due to the increased heat input. The decrease in peak hardness in the HZ ad WN are the result of coarsening laths and PAG structure with increasing HI and/or slower cooling rate (Fig.13). Decrease in the HI leads to an increase of the peak hardness values in the HZ region from about 304 HV0.5 to about 333 HV0.5. Peak hardness values in the HZ regions of the samples decrease almost linearly with increasing HI, bainite lath width and PAG size (Figs.12 and 13). Additionally, it is well known that the strength/hardness of macroscopic material is proportional to the reciprocal of the square root of grain size, according to the Hall-Petch relation [38]. The linear behavior between hardness and PAG size follows a Hall- Petch type relationship too (Fig.13b). Both the expansion and decreasing peak hardness of the HZ region with increasing HI provides evidence of strong dependence of the HZ microstructure and properties on weld HI.
3.3. Tensile Properties The mechanical properties of each microstructural distinct region of the single-sided welded sample were evaluated. Table 3 shows the tensile properties and hardness values from the different weld zones of the single-sided welded sample. The HZ region exhibits the highest yield and ultimate tensile strength, but this is accompanied by the lowest elongation: 43% lower than the BM and 23% lower than the WN. Additionally, there is no evidence of grain coarsening in the HAZ. Given the lack of grain coarsening and slight tempering observed the FSW HAZ; it is likely that the HAZ in the FSW will exhibit much greater toughness than that observed in arc weldments of the same steel. The longitudinal tensile properties of the HZ regions for the various heat inputs investigated are summarized in Table 4. Fig.14 shows the effect of the HI on the strength values in the HZ regions of the samples, i.e. the yield strength and ultimate tensile strength. This figure shows that there is a fairly strong linear correlation between tensile properties and HI. Over the range of parameters investigated, the yield strength in the HZ decreased roughly 15% when the HI increased from 1151 J/mm to 2848 J/mm. Overall the tensile properties appear consistent with the microhardness values in the HZ regions (Fig.15). Variation in strength in the HZ regions of FSW X80 steels can be explained in terms of microstructure. FSW samples with lower HI had higher strength as a result of finer lath microstructures resulting from higher cooling rates. This tendency of increase in strength and hardness with decrease in HI is related with peak temperature and cooling rate. Cota et. al. [36] have reported that when the cooling rate increased, the reduction in the transformation start temperature led to finer bainitic structure and an increase dislocation density, producing an increase in the ultimate tensile and yield
strengths. In HSLA steels, the precipitation of fine microalloying carbides and carbonitrides are not the predominant factor for increasing strength. The predominant factor is fine bainitic microstructure. These results are in agreement with those from the literature [39-42] for HSLA low-carbon bainitic steels. The relationships between bainite lath width, PAG size and strength values in the HZ regions of the samples can be seen in Figs.16 and 17. The strength values in the HZ regions is highly dependent upon the bainite lath width and PAG size; yield and ultimate tensile strength in the HZ regions of the samples increase almost linearly with decreasing bainite lath width and PAG size. Similar to the linear Hall- Petch type relationship between hardness and PAG size, here again, the linear behavior between strength and PAG size follows strongly a Hall- Petch type relationship (Fig.16b). Elongation values in the HZ regions of the samples essentially tend to decrease with decreasing bainite lath width and PAG size, but a significant linear correlation has not been established between elongation and bainite lath width and PAG size. In short, lower the heat input produces smaller the PAG and finer grain/lath structure resulting in higher strength and lower ductility.
4. Conclusions The present study has focused on the microstructure and mechanical properties of HZ region in FSWed X80 steel relative to HI. Correlations between HI and post-weld microstructure/mechanical properties of HZ in FSW HSLA X80 line-pipe steels have been investigated. From this investigation, the following conclusions can be derived: • The HZ microstructure consists of upper bainitic structure with well-defined laths while BM presents typical rolled fine-grained polygonal ferrite microstructure with refined upper bainite islands.
• HI had a significant influence on the microstructure of the HZ. With increasing the HI, (1) the HZ width become larger and (2) the PAGs in the HZ become coarser. The coarser PAGs resulted in the coarser the bainitic structure, and the coarser the lath structure. Roughly 60% and 30% decrease were observed in the PAG size and lath width at the HZ in the lowest HI (1151 J/mm), respectively, compared to the highest HI (2848 J/mm). • The mechanical properties of HZ were higher than those of the other zones, but this zone had the lowest elongation. • Using welding parameters, which cause a higher HI, induced lower microhardness values in the weld zone. • The HI had an obvious effect on the mechanical properties of the HZ. With increasing the HI, (1) the strength values decreased and (2) elongation values had the tendency to increase. When the HI is 1151 J/mm, the yield and ultimate tensile strength were above 16.6% and 9.4% of the welded joint with 2848 J/mm HI, respectively, but the lowest elongation with 4.5%. • The linear relationships between HI and post-weld microstructures and mechanical properties in the HZ of FSW samples were established. With increasing the heat input, lath width and PAG size in the HZ increased linearly, and the peak hardness and strength values in the HZ decreased linearly. Elongation values in the HZ essentially tend to decrease with decreasing HI, lath width and PAG size, but a significant linear correlation has not been established between elongation and HI and post-weld microstructure.
Acknowledgments The authors want to acknowledge the support provided by Science Fellowships and Grant Programmes Department (BIDEB) of The Scientific and Technological Research Council of Turkey (TUBITAK). The authors are also grateful to Dr. J. Farrer for technical assistance, Brigham Young University.
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Figure 1- Schematic illustration of FSW process. Figure 2- Geometry of CS4 tool used in the welding [26,27]. Figure 3- Configuration and dimensions of longitudinal tensile test specimens. Figure 4- Macroscopic images of the FSW samples: a) Sample-1; b) Sample-2; c) Sample-4; d) Sample-6. (HI: Heat Input; AS: Advancing Side; RS: Retreating Side). Figure 5- Microstructures in various zones of single-sided welded sample: a) BM; b) HZ; c) WN; d) HAZ.
Figure 6- TEM micrographs of (a,b) WN region and (c,d) HZ region of the single-sided welded sample (M/A: Martensite/Austenite constituents; LF: Lathy Ferrite; UB: Upper Bainite). Figure 7- Optical micrographs in the HZ regions of FSW samples with different heat input (HI: Heat Input). Figure 8- The PAG size measurement. The dashed lines in the figure show the grain boundaries of a PAG. Figure 9- PAG size (a) and bainite lath width (b) of the HZ regions in different heat input. Figure 10- TEM images of HZ region for a) Sample-2 (1151.5 J/mm) and b) Sample-6 (2848 J/mm) (M/A: Martensite/Austenite constituents; UB: Upper Bainite). Figure 11- Microhardness maps for a) Sample-1 (single-sided welded); b) Sample-2 (1151.5 J/mm); and c) Sample-5 (2600 J/mm). Figure 12- Peak hardness values in the HZ regions of the FSW samples versus heat input. Figure 13- (a) Peak hardness values versus PAG size; (b) Hall-Petch relation between hardness and PAG size; and (c) peak hardness values versus bainite lath width in the HZ region of the FSW samples.
Figure 14-
Strength values in the HZ regions of the FSW samples versus heat input.
Figure 15- Strength values in the HZ regions of the FSW samples versus microhardness. Figure 16- (a) Strength values versus PAG size; and (b) Hall-Petch relation between strength and PAG size in the HZ regions of the FSW samples.
Figure 17- Strength values in the HZ regions of the FSW samples versus bainite lath width.
Fig.1.
Fig.2.
Fig.3.
Fig.4. HI:1151 J/mm
Side-1
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RS HZ HAZ
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d) Carbide
Carbides
M/A M/A LF
Carbide
M/A Strip Carbide Carbide
a)a)
b)
M/A Carbides Strips
M/A Strips
UB UB Carbide
c)
d)
Fig.7. HI: 1400 J/mm
HI: 1151 J/mm
a)
b)
c)
HI: 2848 J/mm
HI: 2600 J/mm
Fig.8.
HI: 2000 J/mm
d)
e)
Prior Austenite Grain Size [micron]
Fig.9a
R2=0.9822
PAG Size [micron] = 1.3524 + 0.0108 (Heat Input) [J/mm]
Heat Input [J/mm]
Fig.9b
Bainite Lath Width [micron]
1
0.9
0.8 R2=0.9123 0.7
0.6 Bainite Lath Width [micron] = 0.5048 + 0.0001 (Heat Input) [J/mm] 0.5
Heat Input [J/mm]
Fig.10.
UB
UB
M/A
a)
b)
Fig.11.
a) HV0.5
b)
c)
Fig.12.
Microhardness [HV0.5]
350
R2=0.8971
325
300
Microhardness [HV0.5] = 352.62 - 0.0161 (Heat Input) [J/mm]
275 Heat Input [J/mm]
Fig.13a.
Microhardness [HV0.5]
375
350
R2=0.8207
325
300 Microhardness [HV0.5] = 349.04 – 1.4125 (PAG Size) [micron]
275 Prior Austenite Grain Size [micron]
Fig.13b.
Microhardness [HV0.5]
375
350
R2=0.8514
325
300 Microhardness [HV0.5] = 264.03 + 241.5 (PAG size)-1/2 [µm-1/2]
275 0.16
0.19
0.22
0.25
Prior Austenite Grain Size
0.28 -1/2
-1/2
[µm
]
0.31
Fig.13c. 375
Microhardness [HV0.5]
350
325 R2=0.8619
300
Microhardness [HV0.5] = 402.52 – 102.91 (Bainite Lath Width) [micron]
275 0.6
0.7
0.8
0.9
Bainite Lath Width [micron]
Fig.14.
Strength [MPa]
R2=0.777
R2=0.84
Tensile Strength [MPa] = 845.43 - 0.0328 (Heat Input) [J/mm] Yield Strength [MPa] = 802.19 - 0.0518 (Heat Input) [J/mm]
Heat Input [J/mm]
1
Fig.15.
Strength [MPa]
R2=0.9186 R2=0.8827
Tensile Strength [MPa] = 2.08 (Microhardness) + 114.29 Yield Strength [MPa] = 3.11 (Microhardness) - 297.96
Microhardness [HV0.5]
Fig.16a.
Strength [MPa]
R2=0.6583
R2=0.7344
Tensile Strength [MPa] = 835.93 - 2.7659 (PAG Size) [micron] Yield Strength [MPa] = 788.58 - 4.4417 (PAG Size) [micron]
Prior Austenite Grain Size [micron]
Fig.16b.
Strength [MPa]
R2=0.7328
R2=0.8224
Tensile Strength [MPa] = 665.41 + 15.5 (PAG size)-1/2 [mm-1/2] Yield Strength [MPa] = 514.19 + 24.97 (PAG size)-1/2 [mm-1/2]
Prior Austenite Grain Size-1/2 [mm-1/2]
Fig.17. 850
800 Strength [MPa]
R2=0.8362
750 R2=0.9284
700
650 Tensile Strength [MPa] = 955.35 - 219.73 (Bainite Lath Width) [micron] Yield Strength [MPa] = 981.86 - 355.04 (Bainite Lath Width) [micron]
600 0.6
0.7
0.8 Bainite Lath Width [micron]
0.9
1
Table 1. Chemical composition (wt.%) of API X80 pipe-line steel used in this investigation. Fe
C
Si
Mn
S
P
Ni
Cr
Cu
Ti
Mo
Nb
N
Al
V
Balance
0.04
0.135
1.7
0.001
0.013
0.147
0.41
0.263
0.014
0.005
0.102
0.006
0.031
0.002
Table 2. Welding parameters with corresponding heat input of FSW X80 welds. Side-1
Side-2
Sample No
Rotation Speed [rpm]
Welding Speed [mm/min]
1
350
127
2
550
284.5
550
228.6
1151.5
3
725
237.2
725
228.3
1400
4
550
124.5
550
124.5
2000
5
725
92.7
725
95.3
2600
6
550
80
550
70.9
2848
Rotation Speed [rpm]
Welding Speed [mm/min]
Heat Input [J/mm]
Single-Sided FSW Sample
Table 3. Longitudinal tensile properties in different weld zones of the single-sided welded sample (average values). (FSW zones as indicated in Figure 4a)
Yield Strength
Ultimate Tensile Strength
Elongation
Hardness
Rp0.2 [MPa]
Rm [MPa]
A [%]
[HV0.5]
WN
598
693
8.2
228
HZ
692
792
6.3
317
HAZ
552
667
13,7
207
BM
621
718
11.1
231
FSW Zones
Table 4. Longitudinal tensile properties of the HZ regions of the FSW samples in different heat input (average values). Heat Input
Yield Strength
Ultimate Tensile Strength
Elongation
Peak Hardness
Sample No
[J/mm]
Rp0.2 [MPa]
Rm [MPa]
A [%]
[HV0.5]
2
1151
758
814
4.5
333
3
1400
719
797
5.9
333
4
2000
679
766
7.0
316
5
2600
686
779
5.3
316
6
2848
650
744
6.7
304
Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input Hakan Aydin, Tracy W. Nelson
www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(13)00892-7 http://dx.doi.org/10.1016/j.msea.2013.07.090 MSA30190
To appear in:
Materials Science & Engineering A
Received date: 24 April 2013 Revised date: 20 July 2013 Accepted date: 22 July 2013 Cite this article as: Hakan Aydin, Tracy W. Nelson, Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2013.07.090 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure and mechanical properties of hard zone in friction stir welded X80 pipeline steel relative to different heat input Hakan AYDIN∗a, Tracy W. NELSONb
a)
Mechanical Engineering Department, Uludag University, 16059, Gorukle-Bursa, Turkey
b)
Mechanical Engineering Department, Brigham Young University, 435 CTB, Provo, UT 84602, USA.
Abstract The study was conducted to investigate the microstructure and mechanical properties of the hard zone in friction stir welded X80 pipeline steel at different heat inputs. Microstructural analysis of the welds was carried out using optical microscopy, transmission electron microscopy, and microhardness. Heat input during friction stir welding process had a significant influence on the microstructure and mechanical properties in the hard zone along the advancing side of the weld nugget. Based on the results, the linear relationships between heat input and post-weld microstructures and mechanical properties in the hard zone of friction stir welded X80 steels were established. It can be concluded that with decreasing heat input bainitic structure in the hard zone become finer and so hard zone strength increases. Keywords: X80 pipeline steel; Friction stir welding; Hard zone, Hardness; Tensile Properties.
∗
Corresponding author: Assistant Professor Dr. Hakan AYDIN Tel.: +90 224 294 06 52; fax.: +90 224 294 19 03; E-mail address: [email protected] Postal address: Uludag University, Engineering and Architecture Faculty, Mechanical Engineering Department, 16059, Gorukle-Bursa, Turkey.
1. Introduction The American Petroleum Institute (API) X80 steel has gained increasing attention from pipeline companies for use in oil and gas transmission pipelines [1]. The commercial X80 API-grade steels are produced by Thermo-mechanically Controlled Processing (TMCP), which produces uniform refined microstructures consisting of ferrite and bainite [2]. These refined microstructures provide superior combination of high-strength and excellent toughness. Welding is an essential process to make large-scale pipelines for long distance transportation of crude oil or natural gas under high pressure. During conventional fusion welding processes, linepipe steels often meet with some problems, such as significant loss of strength and toughness in the heat affected zone (HAZ). Moreover, in conventional fusion welding of linepipe steels, the HAZ is susceptible to hydrogen-assisted-cracking (HAC). The increasing relevance of X80 steel in oil and gas transportation operation requires research on more efficient and reliable joining processes. In recent years, friction stir welding (FSW) has shown considerable promise for joining aluminum alloys, as well as for magnesium, copper, titanium and steel [3-5]. The oil and gas industry has a number of potential applications for this technology. However, a more complete understanding of the technological process and the achieved mechanical behavior of FSW in steels is needed. FSW makes use of a nonconsumable welding tool to induce severe plastic deformation of the workpiece material, generating both heat and mixing across the joint [6,7]. FSW offers several advantages over conventional fusion welding processes, due to its low heat input and absence of melting and solidification process [5]. Advantages include better mechanical properties, low residual stress and distortion, and reduced occurrence of defects [3,4]. In addition, it is expected that hydrogen induced cracking in steels will be reduced or eliminated. These advantages are likely to make FSW an attractive process for joining linepipe steels.
While many studies have focused on FSW of the commercial Al alloys, such as 2XXX [8-11], 6XXX [12-15] and 7XXX [16-18], there is considerable interest in extending this technology to mild steel and other low to medium carbon ferrous alloys [2,19-22]. The mechanical properties of FSWed carbon steels are partially improved compared with the base metals, because FSW achieves grain refinement in the stir zone of the carbon steel, similar to Al alloys [19,20]. The microstructure and mechanical properties of friction stir welds are dependent on the chemical composition and the microstructure of the base metal (BM), the thermo-mechanical conditions during FSW process, and the cooling rate within the critical temperature range. The thermal cycles and high strain rates imposed during the FSW process of X80 steel lead to localized brittle hard zone (HZ) region in the weld. It has been shown that this localized HZ may play a significant role in the fracture toughness of FSW X80 steel [23]. Therefore, the microstructure and mechanical properties of this zone need to be characterized in more detail. Although there are some studies on microstructure and mechanical properties of FSW high strength low alloy (HSLA) steels in the literature [24-28], only a few focused on FSWed X80 steel [1,2,29,30]. Of these, there has been no work on the microstructural and mechanical properties of HZ region. Therefore, the objective of the present work was to evaluate the microstructure and mechanical properties of HZ region in FSWed X80 steel relative to different heat input (HI).
2. Experimental Details This study was performed on a commercial grade API X80 pipe-line steel with chemical composition (wt.%) listed in the Table 1. Before welding, oxide and surface scale were removed from both sides of the plate by grinding followed by degreasing with a methanol solvent. Test plates with dimensions of 1016 x 127 x 11 mm were friction stir welded parallel
to the rolling direction. Schematic illustration of the FSW process can be seen in Fig.1. A Polycrystalline cubic boron nitride (PCBN) CS4 tool was used for all welds (Fig.2). An argon gas atmosphere, at a flow rate of 1.1m3/h, was used to prevent surface oxidation of the weld and tool during the weld cycle. All welds were made under a depth-controlled process with a head tilt of 0.5º. Both single and double-sided welds were evaluated. When double sided, plates were flipped longitudinally to maintain advancing and retreating side symmetry. The FSW welding parameters with corresponding HI used in this investigation are listed in Table 2. Samples were removed transverse to the weld using a water-jet cutter from each weld for post-weld metallographic and microhardness characterization. After that, samples were mounted in bakelite then ground and polished successively through 1µm diamond paste, etched with 2% Nital reagent for 20 seconds, and analyzed optically using an Olympus GX51 microscope. Vickers micro-hardness measurements were performed on a LECO LM 100AT microhardness tester using a diamond pyramid indenter with a 500 gram load, 15 second dwell time and 400μm inter-indentation spacing. For micro-tensile test specimens, the FSW samples were cut parallel to the welding direction using a water-jet cutter. Then, longitudinal through-thickness tensile test specimens were removed from the weld zone of the samples by an electrical discharge cutting machine (EDM). The configuration and dimensions of the longitudinal tensile test specimens can be seen in Fig.3. Tensile tests were carried out on an Instron universal testing machine (Model No. 123), at room temperature and with a crosshead speed of 0.762 mm/min. Transverse samples for transmission electron microscopy (TEM) studies were removed as 3 mm diameter cylinders by cutting with the EDM from selected areas of the samples. 300 µm thick discs were sliced by a low speed diamond saw from these cylinders. Discs were
mechanically thinned from 300µm to 70µm by grinding with 1200 grit SiC water-proof abrasive paper. Finally, the electron transparent thin sections for TEM analysis were prepared by means of double jet electro-polishing, using a solution of 10% (volume fraction) perchloric acid and 90% glacial acetic acid with 20V below 0 ºC in the icy water. The thin foil specimens were examined in a Philips Tecnai F30 TEM FEG model transmission electron microscope at 300 kV to observe the microstructural details.
3. Results and Discussion
3.1. Microstructure Fully consolidated welds X80 steel were produced at all weld conditions. No macroscopic defects, such as voids, cracks or distortion, have been observed in the inspection of the welds. The macroscopic views of the welds at low magnification reveal distinct changes in microstructures at the different parameters investigated (Fig. 4). The various friction stir weld zones are indicated in Fig. 4a. The HZ is located at the top surface and in the weld nugget (WN) along the advancing side (AS). As can be seen in Fig. 4b-d, HI has a strong effect on as-welded microstructure, especially in the HZ region. The HZ and HAZ regions expand with increasing HI. In addition, higher HI produced a more swirl-like microstructure pattern that was asymmetric about the weld centerline. The optical microscopy observations performed on the cross-section of the single-sided welded sample revealed the distinct changes in microstructures of the weld zones (Fig.5). As can be seen in Fig.5a, the BM microstructure is mainly composed of elongated fine-grained polygonal ferrite, with the average grain size about 6 µm, and a small amount of refined upper bainite islands which are almost parallel to the rolling direction of the plate. The post-weld microstructures in the weld zone are a function of the thermo-mechanical cycles experienced
during the welding process. There is no evidence of BM microstructures in the WN and HZ regions (Fig. 5 b, c). The elongated grains in BM have completely transformed to the bainitic structure with well-defined laths in the HZ region. This bainitic microstructure consists of thin, relative straight and long, parallel ferrite laths with discontinuous carbide particles (second phases) at the lath boundaries. This microstructure indicates that the HZ region has reached a peak temperature in excess of A3, which is the temperature at which ferrite transforms to austenite on heating, and the cooling rate is sufficient to form lath bainite. In addition, Wei et. al. [26] reported that the HZ region experienced high strain and/or strain rate which aided in the transformation to austenite. Fairchild et al. [31] have also reported that the temperature and strain rates were very high at the top surface and on the AS compared to other regions in the weld. The WN microstructures are predominantly coarse bainitic structure with carbides along the lath boundaries (Fig.5c). Bainite lath boundaries in the WN region are more irregular that those observed in the HZ region. This is likely the result of deformation during the transformation in the high heat input welds. Additionally, the prior austenite grain (PAG) boundaries are discontinues in both the WN and HZ microstructure [26,27]. The elongated ferrite grains in the BM underwent recrystallization forming more refined equiaxed polygonal ferrite in the HAZ. It is likely that temperatures in the HAZ did not exceed the A3 temperature, but the temperature was sufficient to cause significant coarsening and spheroidization of the carbides to occur. The TEM images of the WN and HZ microstructures of the single-sided welded sample are shown in Fig.6. The microstructure in WN consists of nonequiaxed and interwoven nonparallel lathy ferrite with dense density dislocations, and coarse carbide phases distributed intra-lath and along the lath boundaries (Fig.6a,b). The dark phases in this matrix are
Martensite/Austenite (M/A) constituents or retained austenite [32]. The formation of M/A constituents may be attributed to the partitioning of carbon during the transformation to bainite and the post-transformation of carbon-enriched austenite. Blocky M/A constituents are generally beneficial to the improvement of toughness compared with M/A strip [33]. Thin M/A strips are observed in the well-defined lath boundary locations in the HZ microstructure (Fig.6c,d). Give this; the HZ may exhibit lower toughness than the WN region as it has more M/A strips as reported by Tribe [23]. The optical micrographs of the HZ region of the welded samples with different HI are shown in Fig.7. These figures show that the microstructures in the HZ are similar, consisting of predominantly classical lath-bainitic morphologies with more well defined PAG boundaries. It is clear from these micrographs that HI has a significant influence on microstructure in the HZ region. The HZ regions of all samples display relatively coarse PAGs which will have an effect on the resulting transformation product. The PAG size measurement was performed using a method described by Wei [35]. In this method, as seen in Fig.8, a circle is placed on the top of each PAG and the intercepted length on each diameter is measured. The grain size of a PAG is the average value of these measurements. The traditional ASME intercept or planimetric methods are not suitable to use due to inadequate numbers of PAGs in the micrographs. Wei [35] stated that this method was good to measure the large grain size in the image with limited number grains. Roughly 60% decrease was observed in the PAG size at the HZ in the lowest HI (1151 J/mm) compared to the highest HI (2848 J/mm). In addition, as can be seen in Fig.9a, PAG size increases linearly with increasing HI. Higher HI produces higher peak temperatures, which coarsens PAGs size. In addition, higher HI results in slower cooling rate [27,34]. It is well known that cooling rate plays a crucial role in determining weld microstructure. The polygonal and allotriomorphic grain boundary ferrite may occur at relatively slow cooling rates, while the intermediate transformation products, i.e., bainite, may
be observed at relatively high cooling rates. The bainitic lath structure becomes thinner with increasing cooling rate. Cooling rates in the HZ regions of the welds investigated were fast enough for the formation of lath bainite. Cota et al. [36] has reported that when the cooling rate increases, the reduction in the bainite start temperature led to an increase in the driving force for the nucleation rate of sub-units of ferrite and, consequently, a decreased in the width of bainitic laths. Higher HI and slower cooling rate provides favorable diffusion conditions promoting bainite coarsening in the HZ region. The lath width of the bainitic structure in the HZ region is finest as the lowest HI (fastest cooling rate) compared to the samples with the highest HI. This is shown in the TEM images in Fig.10. The bainite lath widths were measured using a mean intercept method on the TEM micrographs: the lath width at lowest HI is 0,64 µm compared to the 0,90 µm lath width at the highest HI. As expected, similar to the PAGs size, bainite lath width in the HZ regions also increases almost linearly with increasing HI (Fig.9b).
3.2. Microhardness In order to examine the property variations across the welds, microhardness was measured on transverse cross sections of each weld. Fig.11 shows the hardness maps from three different heat input FSW in X80 steel. These microhardness maps show several features in the weld zones: HAZ softening on both sides of the weld; slightly higher hardness of the WN compared to the BM; and the HZ that varies in size and distribution as a function of HI. The non-uniform distribution of hardness correlates to a non-uniform heat gradient generated in the stir zone. Higher strains on the AS of the weld result in an asymmetric heat generation across the stir zone. This also means that there is a larger difference in material properties between the AS and retreating side (RS) of the weld. In Fig.11, the hardness peaks commonly follow a line on the AS of the weld close to the location of the shoulder and edge of the weld
tool. This is likely the region of greatest shear strain which would result in greater heating [37]. If this were the case, higher peak temperatures would produce larger PAG size and subsequent transformation to completely lath-bainitic microstructure, as observed in the microstructure examination. The HAZ has the lowest hardness values due to the predominantly overtempered polygonal ferrite microstructure at elevated temperatures during FSW process. The minimum hardness value in the HAZ decreases due to the higher amount of heat transferred to the weld (Fig.11b,c). The influence of the HI on the microhardness of weld zone can be seen in Fig.11b,c. The hardness of the WN decreases with increasing HI. Although the area of HZ region expands almost to the RS of the weld with increasing HI (Fig.11b,c), the peak hardness in the HZ regions decreases with increasing HI similar to the rest of the WN (Fig.12). The increase in HZ area is likely the result of a large fraction of re-austenitization due to the increased heat input. The decrease in peak hardness in the HZ ad WN are the result of coarsening laths and PAG structure with increasing HI and/or slower cooling rate (Fig.13). Decrease in the HI leads to an increase of the peak hardness values in the HZ region from about 304 HV0.5 to about 333 HV0.5. Peak hardness values in the HZ regions of the samples decrease almost linearly with increasing HI, bainite lath width and PAG size (Figs.12 and 13). Additionally, it is well known that the strength/hardness of macroscopic material is proportional to the reciprocal of the square root of grain size, according to the Hall-Petch relation [38]. The linear behavior between hardness and PAG size follows a Hall- Petch type relationship too (Fig.13b). Both the expansion and decreasing peak hardness of the HZ region with increasing HI provides evidence of strong dependence of the HZ microstructure and properties on weld HI.
3.3. Tensile Properties The mechanical properties of each microstructural distinct region of the single-sided welded sample were evaluated. Table 3 shows the tensile properties and hardness values from the different weld zones of the single-sided welded sample. The HZ region exhibits the highest yield and ultimate tensile strength, but this is accompanied by the lowest elongation: 43% lower than the BM and 23% lower than the WN. Additionally, there is no evidence of grain coarsening in the HAZ. Given the lack of grain coarsening and slight tempering observed the FSW HAZ; it is likely that the HAZ in the FSW will exhibit much greater toughness than that observed in arc weldments of the same steel. The longitudinal tensile properties of the HZ regions for the various heat inputs investigated are summarized in Table 4. Fig.14 shows the effect of the HI on the strength values in the HZ regions of the samples, i.e. the yield strength and ultimate tensile strength. This figure shows that there is a fairly strong linear correlation between tensile properties and HI. Over the range of parameters investigated, the yield strength in the HZ decreased roughly 15% when the HI increased from 1151 J/mm to 2848 J/mm. Overall the tensile properties appear consistent with the microhardness values in the HZ regions (Fig.15). Variation in strength in the HZ regions of FSW X80 steels can be explained in terms of microstructure. FSW samples with lower HI had higher strength as a result of finer lath microstructures resulting from higher cooling rates. This tendency of increase in strength and hardness with decrease in HI is related with peak temperature and cooling rate. Cota et. al. [36] have reported that when the cooling rate increased, the reduction in the transformation start temperature led to finer bainitic structure and an increase dislocation density, producing an increase in the ultimate tensile and yield
strengths. In HSLA steels, the precipitation of fine microalloying carbides and carbonitrides are not the predominant factor for increasing strength. The predominant factor is fine bainitic microstructure. These results are in agreement with those from the literature [39-42] for HSLA low-carbon bainitic steels. The relationships between bainite lath width, PAG size and strength values in the HZ regions of the samples can be seen in Figs.16 and 17. The strength values in the HZ regions is highly dependent upon the bainite lath width and PAG size; yield and ultimate tensile strength in the HZ regions of the samples increase almost linearly with decreasing bainite lath width and PAG size. Similar to the linear Hall- Petch type relationship between hardness and PAG size, here again, the linear behavior between strength and PAG size follows strongly a Hall- Petch type relationship (Fig.16b). Elongation values in the HZ regions of the samples essentially tend to decrease with decreasing bainite lath width and PAG size, but a significant linear correlation has not been established between elongation and bainite lath width and PAG size. In short, lower the heat input produces smaller the PAG and finer grain/lath structure resulting in higher strength and lower ductility.
4. Conclusions The present study has focused on the microstructure and mechanical properties of HZ region in FSWed X80 steel relative to HI. Correlations between HI and post-weld microstructure/mechanical properties of HZ in FSW HSLA X80 line-pipe steels have been investigated. From this investigation, the following conclusions can be derived: • The HZ microstructure consists of upper bainitic structure with well-defined laths while BM presents typical rolled fine-grained polygonal ferrite microstructure with refined upper bainite islands.
• HI had a significant influence on the microstructure of the HZ. With increasing the HI, (1) the HZ width become larger and (2) the PAGs in the HZ become coarser. The coarser PAGs resulted in the coarser the bainitic structure, and the coarser the lath structure. Roughly 60% and 30% decrease were observed in the PAG size and lath width at the HZ in the lowest HI (1151 J/mm), respectively, compared to the highest HI (2848 J/mm). • The mechanical properties of HZ were higher than those of the other zones, but this zone had the lowest elongation. • Using welding parameters, which cause a higher HI, induced lower microhardness values in the weld zone. • The HI had an obvious effect on the mechanical properties of the HZ. With increasing the HI, (1) the strength values decreased and (2) elongation values had the tendency to increase. When the HI is 1151 J/mm, the yield and ultimate tensile strength were above 16.6% and 9.4% of the welded joint with 2848 J/mm HI, respectively, but the lowest elongation with 4.5%. • The linear relationships between HI and post-weld microstructures and mechanical properties in the HZ of FSW samples were established. With increasing the heat input, lath width and PAG size in the HZ increased linearly, and the peak hardness and strength values in the HZ decreased linearly. Elongation values in the HZ essentially tend to decrease with decreasing HI, lath width and PAG size, but a significant linear correlation has not been established between elongation and HI and post-weld microstructure.
Acknowledgments The authors want to acknowledge the support provided by Science Fellowships and Grant Programmes Department (BIDEB) of The Scientific and Technological Research Council of Turkey (TUBITAK). The authors are also grateful to Dr. J. Farrer for technical assistance, Brigham Young University.
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Figure 1- Schematic illustration of FSW process. Figure 2- Geometry of CS4 tool used in the welding [26,27]. Figure 3- Configuration and dimensions of longitudinal tensile test specimens. Figure 4- Macroscopic images of the FSW samples: a) Sample-1; b) Sample-2; c) Sample-4; d) Sample-6. (HI: Heat Input; AS: Advancing Side; RS: Retreating Side). Figure 5- Microstructures in various zones of single-sided welded sample: a) BM; b) HZ; c) WN; d) HAZ.
Figure 6- TEM micrographs of (a,b) WN region and (c,d) HZ region of the single-sided welded sample (M/A: Martensite/Austenite constituents; LF: Lathy Ferrite; UB: Upper Bainite). Figure 7- Optical micrographs in the HZ regions of FSW samples with different heat input (HI: Heat Input). Figure 8- The PAG size measurement. The dashed lines in the figure show the grain boundaries of a PAG. Figure 9- PAG size (a) and bainite lath width (b) of the HZ regions in different heat input. Figure 10- TEM images of HZ region for a) Sample-2 (1151.5 J/mm) and b) Sample-6 (2848 J/mm) (M/A: Martensite/Austenite constituents; UB: Upper Bainite). Figure 11- Microhardness maps for a) Sample-1 (single-sided welded); b) Sample-2 (1151.5 J/mm); and c) Sample-5 (2600 J/mm). Figure 12- Peak hardness values in the HZ regions of the FSW samples versus heat input. Figure 13- (a) Peak hardness values versus PAG size; (b) Hall-Petch relation between hardness and PAG size; and (c) peak hardness values versus bainite lath width in the HZ region of the FSW samples.
Figure 14-
Strength values in the HZ regions of the FSW samples versus heat input.
Figure 15- Strength values in the HZ regions of the FSW samples versus microhardness. Figure 16- (a) Strength values versus PAG size; and (b) Hall-Petch relation between strength and PAG size in the HZ regions of the FSW samples.
Figure 17- Strength values in the HZ regions of the FSW samples versus bainite lath width.
Fig.1.
Fig.2.
Fig.3.
Fig.4. HI:1151 J/mm
Side-1
RS
AS
AS
RS HZ HAZ
HZ
WN
HAZ
BM Side-2
a) HI: 2000 J/mm
b)
Side-1
RS
AS
Side-2
c)
HI: 2848 J/mm RS
Side-2 AS
Side-1
d)
Fig.5.
b)
Fig.6.
c)
d) Carbide
Carbides
M/A M/A LF
Carbide
M/A Strip Carbide Carbide
a)a)
b)
M/A Carbides Strips
M/A Strips
UB UB Carbide
c)
d)
Fig.7. HI: 1400 J/mm
HI: 1151 J/mm
a)
b)
c)
HI: 2848 J/mm
HI: 2600 J/mm
Fig.8.
HI: 2000 J/mm
d)
e)
Prior Austenite Grain Size [micron]
Fig.9a
R2=0.9822
PAG Size [micron] = 1.3524 + 0.0108 (Heat Input) [J/mm]
Heat Input [J/mm]
Fig.9b
Bainite Lath Width [micron]
1
0.9
0.8 R2=0.9123 0.7
0.6 Bainite Lath Width [micron] = 0.5048 + 0.0001 (Heat Input) [J/mm] 0.5
Heat Input [J/mm]
Fig.10.
UB
UB
M/A
a)
b)
Fig.11.
a) HV0.5
b)
c)
Fig.12.
Microhardness [HV0.5]
350
R2=0.8971
325
300
Microhardness [HV0.5] = 352.62 - 0.0161 (Heat Input) [J/mm]
275 Heat Input [J/mm]
Fig.13a.
Microhardness [HV0.5]
375
350
R2=0.8207
325
300 Microhardness [HV0.5] = 349.04 – 1.4125 (PAG Size) [micron]
275 Prior Austenite Grain Size [micron]
Fig.13b.
Microhardness [HV0.5]
375
350
R2=0.8514
325
300 Microhardness [HV0.5] = 264.03 + 241.5 (PAG size)-1/2 [µm-1/2]
275 0.16
0.19
0.22
0.25
Prior Austenite Grain Size
0.28 -1/2
-1/2
[µm
]
0.31
Fig.13c. 375
Microhardness [HV0.5]
350
325 R2=0.8619
300
Microhardness [HV0.5] = 402.52 – 102.91 (Bainite Lath Width) [micron]
275 0.6
0.7
0.8
0.9
Bainite Lath Width [micron]
Fig.14.
Strength [MPa]
R2=0.777
R2=0.84
Tensile Strength [MPa] = 845.43 - 0.0328 (Heat Input) [J/mm] Yield Strength [MPa] = 802.19 - 0.0518 (Heat Input) [J/mm]
Heat Input [J/mm]
1
Fig.15.
Strength [MPa]
R2=0.9186 R2=0.8827
Tensile Strength [MPa] = 2.08 (Microhardness) + 114.29 Yield Strength [MPa] = 3.11 (Microhardness) - 297.96
Microhardness [HV0.5]
Fig.16a.
Strength [MPa]
R2=0.6583
R2=0.7344
Tensile Strength [MPa] = 835.93 - 2.7659 (PAG Size) [micron] Yield Strength [MPa] = 788.58 - 4.4417 (PAG Size) [micron]
Prior Austenite Grain Size [micron]
Fig.16b.
Strength [MPa]
R2=0.7328
R2=0.8224
Tensile Strength [MPa] = 665.41 + 15.5 (PAG size)-1/2 [mm-1/2] Yield Strength [MPa] = 514.19 + 24.97 (PAG size)-1/2 [mm-1/2]
Prior Austenite Grain Size-1/2 [mm-1/2]
Fig.17. 850
800 Strength [MPa]
R2=0.8362
750 R2=0.9284
700
650 Tensile Strength [MPa] = 955.35 - 219.73 (Bainite Lath Width) [micron] Yield Strength [MPa] = 981.86 - 355.04 (Bainite Lath Width) [micron]
600 0.6
0.7
0.8 Bainite Lath Width [micron]
0.9
1
Table 1. Chemical composition (wt.%) of API X80 pipe-line steel used in this investigation. Fe
C
Si
Mn
S
P
Ni
Cr
Cu
Ti
Mo
Nb
N
Al
V
Balance
0.04
0.135
1.7
0.001
0.013
0.147
0.41
0.263
0.014
0.005
0.102
0.006
0.031
0.002
Table 2. Welding parameters with corresponding heat input of FSW X80 welds. Side-1
Side-2
Sample No
Rotation Speed [rpm]
Welding Speed [mm/min]
1
350
127
2
550
284.5
550
228.6
1151.5
3
725
237.2
725
228.3
1400
4
550
124.5
550
124.5
2000
5
725
92.7
725
95.3
2600
6
550
80
550
70.9
2848
Rotation Speed [rpm]
Welding Speed [mm/min]
Heat Input [J/mm]
Single-Sided FSW Sample
Table 3. Longitudinal tensile properties in different weld zones of the single-sided welded sample (average values). (FSW zones as indicated in Figure 4a)
Yield Strength
Ultimate Tensile Strength
Elongation
Hardness
Rp0.2 [MPa]
Rm [MPa]
A [%]
[HV0.5]
WN
598
693
8.2
228
HZ
692
792
6.3
317
HAZ
552
667
13,7
207
BM
621
718
11.1
231
FSW Zones
Table 4. Longitudinal tensile properties of the HZ regions of the FSW samples in different heat input (average values). Heat Input
Yield Strength
Ultimate Tensile Strength
Elongation
Peak Hardness
Sample No
[J/mm]
Rp0.2 [MPa]
Rm [MPa]
A [%]
[HV0.5]
2
1151
758
814
4.5
333
3
1400
719
797
5.9
333
4
2000
679
766
7.0
316
5
2600
686
779
5.3
316
6
2848
650
744
6.7
304