Mechanical Behaviour of Intercritically Reheated

Proceedings of the 2018 12th International Pipeline Conference IPC2018 September 24-28, 2018, Calgary, Alberta, Canada

IPC2018-78317 MECHANICAL BEHAVIOUR OF INTERCRITICALLY REHEATED COARSE-GRAINED HEAT AFFECTED ZONE IN HIGH STRENGTH LINE PIPE STEELS Madhumanti Mandal Department of Materials Engineering, The University of British Columbia Vancouver, BC, Canada

Thomas Garcin Department of Materials Engineering, The University of British Columbia Vancouver, BC, Canada

Warren J. Poole Department of Materials Engineering, The University of British Columbia Vancouver, BC, Canada

Matthias Militzer Department of Materials Engineering, The University of British Columbia Vancouver, BC, Canada

Laurie Collins Evraz Inc NA Regina, SK, Canada

representative of CGHAZ have excellent impact properties. The impact toughness of the microstructures typical of ICCGHAZ is strongly dependent on the composition as well as morphology and spatial distribution of the resulting martensite-austenite (M/A) islands transformed from inter-critically formed austenite. This zone can play a significant role in fracture initiation and thus needs to be considered in alloy and welding process designs.

ABSTRACT Multipass welding of high strength steels used for fabrication and joining of transmission pipelines presents a number of metallurgical challenges. A key concern is both the strength and toughness of the heat affected zone (HAZ) adjacent to both seam and girth welds. In this work, a systematic study has been conducted on regions of the heat affected zone in the base metal where the first welding pass produces a thermal excursion which results in a coarse-grained heat affected zone (CGHAZ). The subsequent weld pass involves intercritical annealing of this region, i.e. a microstructure associated with intercritically reheated coarse grain heat affected zone (ICCGHAZ). The small ICCGHAZ region is often identified as being particularly susceptible to crack initiation. This work was undertaken to understand microstructure development in this zone and how the ICCGHAZ may affect the overall performance of the HAZ. Gleeble thermomechanical simulations have been conducted to produce bulk samples representative of different welding scenarios. Charpy impact tests and tensile tests have been performed over a range of temperatures. It was found that when a continuous necklace of martensite-austenite islands form on the prior austenite grain boundaries (i.e. for a M/A fraction of ≈10%), the Charpy impact toughness energy is dramatically decreased and the ductile brittle transition temperature is significantly raised. Detailed studies on the secondary cracks have been conducted to examine the fracture mechanisms in the different microstructures. The results show that the lower bainite microstructures obtained after the 1st thermal treatment,

INTRODUCTION The potential formation of low toughness regions in the heat affected welded zone is a concern in the welding of line pipe steels. There are two regions of particular concern, i.e. the coarse grained heat affected zone (CGHAZ) next to the fusion zone [1]– [4] and the intercritical coarse grained heat affected zone (ICCGHAZ) which is formed during multi-pass welding [5]–[9]. Earlier investigations have shown that the toughness of the ICCGHAZ is mainly controlled by the size and volume fraction of high carbon martensite-austenite (M/A) constituents [8], [10]. However, it has also been reported that the loss in toughness is not only related to the presence of M/A, but also the morphology and distribution of the M/A constituents [11]. For example, Matsuda et al. showed that while the morphology of the M/A constituents only had a small influence on the crack initiation energy, the morphology strongly influenced the crack propagation energy [12]. In addition, the matrix microstructure may also have an important role in the local toughness. For example, the presence of low temperature transformation products such as lower bainite may significantly improve the

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710, 725 and 760 oC holding for 60 s followed by cooling at a rate of ≈50 oC/s was employed. The temperatures and times for the intercritical annealing treatment were selected based on the dilation measurements to form approximately 5%, 10% and 20% austenite which upon cooling would correspond to a similar fraction of M/A. After the thermal treatment, i) sub-size tensile tests (see Figure 1a) and ii) standard Charpy V-notch samples (see Figure 1b) were machined from the Gleeble samples.

toughness by increasing the density of high angle grain boundaries which act as obstacles to crack propagation [13]. It is challenging to study the effect of the different microstructural features observed in the heat affected zone of an actual weld due to the complexity of the graded microstructure found in the heat affected zone. Thus, this study has evaluated the mechanical properties of bulk samples produced in laboratory simulations which have microstructures representative of CGHAZ and ICGGHAZ regions. EXPERIMENTAL METHODOLOGY A hot-rolled X80 line pipe steel supplied by Evraz Inc. NA (Regina, SK) has been used in this study. The chemical composition of the steel is shown in Table 1. The microstructure of the as-received material mainly consisted of fine irregular ferrite with a random distribution of M/A islands. Three basic types of precipitates, i.e. Mo2C, Nb(C,N) and TiN, have been observed by transmission electron microscopy (TEM) studies on the as-received microstructure [14]. Table 1: Chemical composition of X80 steel (key elements)

wt% at%

C 0.06 0.28

Mn 1.65 1.7

Nb 0.034 0.020

Ti 0.012 0.014

Mo 0.24 0.14

N 0.005 0.020

Figure 1 shows the two types of test samples were machined from the as-received material for heat treatment using the Gleeble 3500 thermo-mechanical simulator. The heat treatment paths were designed to produce microstructures representative of CGHAZ and ICCGHAZ for a welding scenario relevant to gas metal arc welds (GMAW). During the Gleeble tests, the temperature was measured using a K-type thermocouple, spot welded to the centre of the specimens. A dilatometer was attached to the centre of the specimen to measure the change in width during the heating and cooling cycles. The dilation data was analyzed using the Kop method to obtain the transformation kinetics [15]. A two-step heattreatment path, as shown schematically in Figure 2, was used to produce specimens for the tensile and Charpy tests. The first step involved heating to 1300 oC at a heating rate of 100 oC/s, holding for 5 s and then cooling at ≈50 oC/s. The cooling rate was chosen based on experimental measurements (using embedded thermocouples) of the average cooling rate for laboratory based GMAW weld trials of Gaudet et al [16]. In these weld trials, a heat input of 0.57 kJ/mm, a travel speed of 8 mm/s and no preheat was employed. For the second step in the thermal history which involved, intercritical annealing, the highly nonisothermal temperature spike observed in the weld trials was simplified to an isothermal hold using the metallurgical principle of addtivitiy, i.e. an isothermal treatment can be determined with is equivalent to the thermal spike [17], [18]. In this case, heating rate of 50 oC/s to intercritical temperatures of

Figure 1: Schematic showing sample geometries for heattreatment and mechanical tests (a) sub-size tensile tests and (b) standard Charpy V-notch tests. All dimensions are in mm.

Figure 2: Schematic diagram of the thermal cycles used in the experiments.

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been colored based on their misorientation, i.e. green for 2-15o (low angle), blue for 15-50o (medium angle) and red for 50-63o (high angle). In this map, it can be seen that there is a high density of low and high angle boundaries but relatively fewer medium angle boundaries. In fact, the high angle boundaries have a special orientation relationship with the prior austenite grain.

Samples for microstructural characterization were prepared using standard metallographic preparation followed by chemical etching with a 2% Nital solution. Scanning electron microscope images were acquired with the in-lens detector (secondary electrons) using a Zeiss Sigma field emission gun scanning electron microscope (FEGSEM). Samples for characterization by electron backscatter diffraction (EBSD) were also prepared by standard metallographic procedures but in this case, the final step was done by electropolishing the sample in a solution of 95% acetic acid 5% perchloric acid. EBSD maps were acquired using a step size of 50 nm. Post processing and analysis of the EBSD data were conducted with TSL Orientation Imaging Microscopy (OIM) data analysis following the procedure suggested by Reichert et al. [19]. Uniaxial tensile tests were conducted on sub-size tensile specimens (see Figure 1a) with a gauge length of 12.5 mm using an Instron 8872 testing machine at a strain rate of 2 x 10-3 s-1. The axial strain was calculated from measurements on the change of length of the sample with a clip-on extensometer. The yield stress was determined by the 0.2% offset method. Standard Charpy V-notch specimens (10x10x55 mm) were machined from the Gleeble sample (see Figure 1b). The Charpy impact tests were conducted at temperatures of (-60, -45, -20, 0 and 21 °C) using an Instron 400 J instrumented impact machine. For the low temperature tests, the temperature was maintained within ±2 °C in a medium containing a mixture of methanol and dry ice. The fracture surfaces were examined by secondary electron imaging in the FEGSEM. In some cases, the broken Charpy samples were sectioned perpendicular to the fracture surface so that the relationship between the microstructure and the sub-surface damage could be examined.

Figure 3: Microstructure as observed in FEGSEM after the 1st thermal treatment using an in-lens detector

Figure 4c shows a pole figure where the orientation data has been taken from a single prior austenite grain (marked by the dashed line in Figure 4a). In the (001) pole figure, the 24 variants based on the Kurdjumov-Sachs Orientation Relationship (KSOR) with the austenite grain have been plotted (24 small black circles). The orientation of the prior austenite grain was determined from the small amount of retained austenite in the grain. The experimental results are superimposed on the ideal orientation relationships and have been colored brown, blue and yellow to distinguish the 3 main Bain groups each containing 8 variants. It can be observed that the experimentally determined orientations correspond closely to the theoretical values. In Figure 4d, the prior austenite grain has been colored based on the 3 Bain groups. This shows that the prior austenite grain has subdivided into an almost equal mixture of the 3 Bain groups. The orientation relationships between austenite and bainite ferrite have been extensively discussed by Takayama et al. [20]. In the following, we will refer to this type of microstructure with a high density of high angle special boundaries as lower bainite.

RESULTS AND DISCUSSION Microstructure produced in first thermal cycle Figure 3 is a FEGSEM image of the microstructure produced after the first thermal cycle shown in Figure 2. The thermal cycle approximates the heating and cooling cycle which would be found near the fusion line, i.e. a region with substantial austenite grain growth due to dissolution of NbC precipitates followed by relatively fast cooling; the CGHAZ. The etching with Nital reveals two important microstructural features in this image. First, one can observe fine bright particles distributed throughout the microstructure and these correspond to M/A islands and second, the prior austenite grain boundaries are also revealed as the solid white lines. Based on a series of images, the prior austenite grain size was measured to be ≈50 µm and EBSD analysis characterized the fraction of M/A particles to be ≈2-3 %. This is consistent with the results of Reichert et al. on the same steel with a similar heat treatment [19]. Figure 4 presents results from high resolution EBSD characterization of the microstructure. Figure 4a is an inverse pole figure (IPF) map showing the wide variation of crystal orientations in the microstructure. Figure 4b plots the same area where the boundaries between different regions of the microstructure have

As mentioned earlier, the thermal path was chosen to produce a microstructure similar to what one would obtain in the CGHAZ of a GMAW weld. However, in the current Gleeble tests, the heating rate was well below that observed in welding (i.e. > 1000 o C/s) and the steel was held at 1300 oC for 5 s whereas in a weld the time at peak temperature is < 1 s. This was done for two

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microstructure evolution to a GMAW weld where the peak temperature was ≈ 1350 oC, i.e. austenite grain size of ≈ 50 µm and all the Nb and Mo in solid solution. Turning to the cooling after the hold at 1300oC for 5s, the analysis of dilation measurements using the Kop method [15], indicated that the transformation start temperature (as defined by 5% transformed) was ≈ 530 oC. The detailed experimental measurements of Gaudet [16] on GMAW welds found that the cooling rate at 530 o C was ≈50 oC/s, i.e. the same as the current experiments. A transformation start temperature of 530 oC is consistent with previous measurements on the same steel under similar conditions [13], [21], [22].

reasons, first, this thermal cycle could be well controlled in the Gleeble and second, experiments conducted above 1300 oC lead to several technical problems (sagging of samples, loss of the control thermocouple and excessive oxidation). On the other hand, thermodynamic calculations on the stability of the Mo, Nb and Ti precipitates using Thermo-Calc (TCFE7 database) predict that only TiN is stable at 1300 oC, i.e the Mo and Nb precipitates will dissolve. The kinetics of dissolution was examined using the model of Garcin et al. which indicated that all the Nb and Mo precipitates would dissolve after holding at 1300 oC for 5 s [21]. Using the model of Garcin et al., it is estimated that the current experimental thermal cycle would be similar in terms of

Figure 4: EBSD maps after 1st thermal cycle showing (a) crystal orientations; (b) grain boundary misorientations; (c) (001) pole figure in comparison to 24 calculated KS variants (black circles); (d) Bain map showing three Bain groups according to (a).

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M/A particles increasing with the intercritical temperature. Analysis of the dilation data determined that approximately 5, 10 and 20 % austenite were formed at 710, 725 and 760 oC, respectively. The dilation measurements also showed that there was no evidence of the austenite to ferrite (or bainite) transformation upon cooling suggesting that the M/A observed in these images reflects the austenite that was formed during the intercritical hold. Finally, it appears qualitatively that the fraction of the prior austenite grain boundary with M/A particles increases significantly from Figure 5a (i.e. 5 % austenite) to Figure 5b where the prior austenite grain boundary is almost completely covered with M/A particles. Mechanical Tests Tensile tests - Figure 6 shows the results from the tensile tests conducted at ambient temperature using the sub-size tensile specimens machined from the Gleeble sample after the heat treatment. The results for the yield stress, ultimate tensile stress and total elongation are summarized in Table 2. It can be observed that the yield stress is similar for the sample after the first thermal cycle (i.e. lower bainite) and the samples with 10 and 20 % M/A at 680-700 MPa. The sample with 5% M/A has a higher yield stress of 738 MPa. On the other hand, the ultimate tensile stress increases from 808 MPa to 893 MPa as the fraction of M/A increases to 20%. Finally, the biggest effect of the fraction of M/A is on the total elongation (note: the samples have a sub-size gauge length). The total elongation increases from 4% to 15% as the intercritical M/A is increased to 5% and then decreases to 12% and 10% for 10 and 20% M/A, respectively. The role of M/A fraction has been extensively studied in dualphase and complex phase steels. The M/A fraction in a ferrite matrix plays a complex role in modifying the elastic-plastic transition, the work hardening behaviour and the initiation/growth of damage in these type of steels. In the current case, the situation is even more complex as the matrix is bainitic and the matrix may undergo tempering (softening) during the intercritical anneal. A detailed analysis of the competing processes is the subject of on-going research for these steels. Charpy Impact Tests – Figure 7 illustrates the results for the Charpy impact energy as function of test temperature for the 4 conditions considered in this work. The mean impact energy for each condition is indicated with the solid symbols (the vertical lines represent the range of the impact energies measured). The impact energy, Y, has been fit as a function of temperature, T, to the following equation [23]:

Figure 5: Microstructures of ICCHAZ having a lower bainite matrix with (a) 5% M/A, (b) 10% M/A and (c) 20% M/A

𝑇 − 𝑇0 𝑌 = 𝐴 + 𝐵 𝑡𝑎𝑛ℎ ( ) 𝐶

Microstructure produced after the second thermal cycle

where A, B, C are defined as: Lower shelf energy = A – B, Upper shelf energy = A + B, Transition temperature = T0 and Slope at T0 = B/C

Figure 5 shows the FEGSEM images of the microstructure after holding at 710, 725 and 760 oC, respectively. The bright white M/A islands are arranged in a necklace fashion along the prior austenite grain boundaries with the fraction and size of the

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The transition temperature in this paper is referred to as the temperature corresponding to 50% of the difference between upper –and lower shelf energy. It can be observed that for the case of the lower bainite microstructure (after the first thermal cycle), the upper shelf energy is ≈ 270 J and the transition temperature is below -52 oC. As the volume fraction of M/A increases to 5% (i.e. the decomposition product of the intercritical austenite from the second thermal cycle), the upper shelf energy increases to ≈300 J with a transition temperature of -45 oC. The increase in the upper shelf energy is consistent with the higher yield and tensile stress and the increase in total elongation for the sample with 5% M/A, i.e. the energy absorbing capability is increased. However, in the case of 10% M/A, a dramatic degradation occurs in the impact energy behavior. There is no evidence of a transition or an upper shelf (at least over the range of temperatures examined here) and while the impact energy increases slightly from 10 J at -20 oC to 31 J at 21 oC, the impact energy is below 50 J for all cases, i.e. an industry rule of thumb for low impact toughness. Finally, for the case of 20% M/A, the ductile brittle transition returns with an upper shelf energy of ≈ 200 J and a transition temperature of -15 o C.

Figure 6: Room temperature engineering stress-strain results for different conditions (note: sub-size tensile samples with a gauge length of 12.5 mm were used).

The results from the Charpy impact tests can be rationalized as follows. The good behavior (low transition temperature and high upper shelf energy) for the lower bainite samples can be attributed to the presence of a fine distribution of bainite packets with different Bain group variants (see Figure 4). In particular, the high density of high angle boundaries is considered to make the stress for macroscopic cleavage fracture more difficult as the cleavage cracks must change their propagation direction as they cross each high angle boundary. The fracture process can be viewed as a competition between ductile fracture and brittle failure with the higher macroscopic cleavage stress favoring ductile fracture. In contrast, the samples with 10% and 20% M/A form a nearly continuous layer of M/A along the prior austenite grain boundaries. An analysis based on the Fe-C phase diagram (in this case using ThermoCalc with the TCFE7 database and assuming the carbon content of the M/A is the same as the intercritical austenite) predicts that the carbon content in 10% and the 20% M/A is 0.6 and 0.3 wt.%, respectively. A simple estimate of the yield stress for the martensite can be made using the data of Leslie on bulk martensite samples [24]. This would suggest that yield stress for the martensite in the 10% and 20% microstructures is 2000 and 1300 MPa, respectively (note; these might be slight overestimates since the role of retained austenite is not considered). Thus, in the case of 10% M/A, it is unlikely that plastic deformation of the martensite can occur at any of the temperatures of the impact tests and thus, there will be a significant strain incompatibility between the matrix and the M/A as the material is loaded. Figure 8a shows a FEGSEM image of the fracture surface where it appears that fracture occurs on the prior austenite grain boundaries. Figure 8b shows a crosssection FEGSEM image perpendicular to the fracture surface and just below the fracture surface. Here, it appears based on secondary cracks below the fracture surface, that the fracture

Figure 7: Charpy impact behavior for different microstructures tested at different temperature

Table 2: Tensile results

Lower Bainite IA-5% M/A IA-10% M/A IA-20% M/A

0.2% YS (MPa) 682 738 694 698

UTS (MPa) 808 834 868 893

Total Elongation 4% 15 % 12.2 % 9.7 %

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path follows the boundary between M/A and the bainitic matrix. The combination of a nearly continuous layer of M/A on the prior austenite grains and the high yield stress of the M/A produces a low energy path for crack propagation to occur. In the case of 20% M/A, there is also a nearly continuous layer of M/A on the prior austenite grain boundaries but in this case the yield stress of the M/A would be substantially lower than the 10% case, i.e. 1300 MPa vs. 2000 MPa. Given the tensile strength for this condition is ≈900 MPa and the load transfer process from the matrix to second phase particles, it may be possible that in this case, the M/A constituents may undergo plastic deformation which shifts the balance towards ductile fracture. This would be consistent with the observations of an upper shelf energy of ≈200 J and a transition temperature of -25 oC which is much improved over the case of 10% M/A. Finally, the case of 5% M/A can be rationalized by the observation that the M/A particles only cover approximately 50% of the prior austenite grain boundaries, i.e. in this case, there is not a continuous crack path for fracture to occur along. This microstructure shows the best Charpy impact energy behavior. Future work will involve examining the effect of reducing the carbon content of the steel on the Charpy impact behavior. The hypothesis being that a reduction in the bulk carbon content will also reduce the carbon content of the intercritical austenite and resulting M/A particles thereby lowering the propensity of the crack path to follow the M/A network.

SUMMARY The current study shows that the impact energy behavior is strongly dependent on the microstructures developed when intercritical temperature excursions are found in the final pass of multi-pass weld scenarios. 1.

2.

3.

4.

Figure 4: (a) Fracture surface and (b) Secondary surface perpendicular to and below fracture surface of IA-10% M/A, broken at -20 oC

For cooling rates of 50 oC/s after a high temperature thermal cycle with a peak temperature of 1300 oC, the austenite grains transform to a variety of Bain groups leading to a final bainite structure which has a high density of high angle grain boundaries. This microstructure has excellent impact fracture properties. The yield stress is observed to be only weakly dependent on the fraction of M/A but the ultimate tensile strength increases almost linearly with the fraction of M/A. For volume fraction of ≥10%, an almost continuous film of M/A is found on the prior austenite grain boundaries. This leads to a continuous low energy, fracture path. Fracture appears to occur along the boundary between M/A and the bainitic matrix. In the case of ≈20% M/A, some recovery in impact fracture behavior is observed. It is proposed that this may be related to the possibility of plastic deformation of the M/A constituents.

ACKNOWLEDGMENTS The financial support received from the Natural Sciences and Engineering Research Council of Canada, Evraz Inc. NA and TransCanada Pipelines Ltd. is duly acknowledged with gratitude. We also acknowledge the Charpy impact testing facility provided to us by Evraz Inc. NA, Regina. REFERENCES

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