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Microstructure and mechanical properties of a flash butt welded pearlitic rail
Journal of Materials Processing Tech. 270 (2019) 20–27
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Journal of Materials Processing Tech. journal homepage: www.elsevier.com/locate/jmatprotec
Microstructure and mechanical properties of a flash butt welded pearlitic rail
T
Rodrigo Rangel Porcaroa, , Geraldo Lúcio Fariaa, Leonardo Barbosa Godefroida, Gabriela Ribeiro Apolonioa, Luiz Cláudio Cândidoa, Elisângela Silva Pintob ⁎
a b
Universidade Federal de Ouro Preto, Ouro Preto, MG, Brazil Instituto Federal de Educação Ciência e Tecnologia de Minas Gerais, Ouro Preto, MG, Brazil
ARTICLE INFO
ABSTRACT
Associate Editor: C.H. Caceres
The structural changes resulting from the Flash Butt Welding (FBW) of pearlitic rails have been associated with wear/premature failures, despite this, there are no studies applying dilatometry to correlate the welding thermal cycles with the microstructural development of such material. The microstructural evolution of the heat affected zone is clarified with the aid of dilatometry. The increase in the steel hardenability associated with a larger austenitic grain size promotes the austenite-pearlite transformation at lower temperatures in the grain growth region. This explains why this region has larger pearlite colony size but smaller interlamellar pearlite spacing and higher hardness than the grain refined region. Partial cementite spheroidization in the heat affected zone is responsible for significant decrease in hardness and tensile strength and is correlated to localized dipping, rolling contact fatigue and failures. A dilatometry based methodology is proposed to define a process window and control the post-weld cooling rate at the rail head in order to improve the weld performance due to a better hardness profile, without increasing costs or welding time. For the steel evaluated, a 20% increase in the hardness of the softened area at the HAZ was obtained by dilatometric simulation of a safe accelerated cooling (5 °C/s).
Keywords: Flash butt welding Dilatometry Rail steels Mechanical properties of welded joints Welding metallurgy
1. Introduction The increasing demand for rail transportation around the world has been accompanied by increased axle load, higher train speed and frequency of railroad use, which requires more carefulness with structural integrity to reduce rail failures and to increase the transport availability (Girsch et al., 2009). The use of continuous welded rails (CWR) by the flash butt welding (FBW) process has also grown in order to assure higher continuity and to improve the dynamic behavior of the railway system (Mansouri and Monshi, 2004). The microstructure of steels used in railroad tracks worldwide is mainly pearlitic (Boer and Masumoto, 2001). The increase on surface hardness of the rails decreases the wear on the head, and the mechanisms used for this purpose have mainly been the decrease of the pearlite interlamellar spacing by heat treatments and/or use of alloying elements (Olivares et al., 2011). Garnham and Davis (2009) showed how aspects related to the previous austenitic grain size and the pearlite interlamellar spacing influence the mechanical properties of railroad tracks. However, Micenko et al. (2013) reported that there are no studies in the literature describing the structural evolution of pearlitic rails ⁎
during the FBW process. The authors also suggested the use of dilatometry techniques to correlate aspects of welding thermal cycles, such as peak temperature and cooling rates, with the structure development and the resulting properties. Mansouri and Monshi (2004) and Porcaro et al. (2017) describe the structural changes resulting from the FBW of pearlitic rails. These authors presented the heat affected zone (HAZ) of joints composed of three regions: grain growth, grain refined due to recrystallization and partial transformation. The structural changes, also reported by Micenko et al. (2013), include partial decarburization and large variation of hardness in the partial transformation region due to the partial cementite spheroidization. These variations in the rail head hardness after the FBW may cause localized wear or plastic strain during the railway lifetime (Steenbergen and Dollevoet, 2013), which in turn, increases the vibrations associated with fatigue. In a recent work, Mutton et al. (2016) evaluated the microstructure influence on the development of RCF in flash butt welded joints of pearlitic rails. Fatigue cracks developed readily in the spheroidization bands on both sides of the welding centerline and propagated further at these locations. Severe plastic deformation in the softened zones of FBW
Corresponding author. E-mail address: [email protected] (R.R. Porcaro).
https://doi.org/10.1016/j.jmatprotec.2019.02.013 Received 31 August 2018; Received in revised form 11 January 2019; Accepted 10 February 2019 Available online 16 February 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.
Journal of Materials Processing Tech. 270 (2019) 20–27
R.R. Porcaro, et al.
was pointed out as RCF crack initiation sites at the gauge corner. The authors suggest investigating in order to establish a correlation between the welding conditions and the cementite spheroidization process in different rail steel grades, considering the need to better understand the involved mechanisms to promote process improvements. There is a knowledge lack on the mechanisms that control the microstructure development and it has hindered the setting of FBW process parameters in order to improve weld performance. In this context, this paper presents detailed structural and mechanical characterizations of FBW joints of a pearlitic rail steel. Results of measurements of pearlite interlamellar spacing and colonies sizes are presented for all HAZ regions, associated with dilatometric studies to describe steel phase transformations due to the welding process, including the effects of austenitic grain size and partial cementite spheroidization. The results allow understanding, in detail, how the pearlite morphology is affected in terms of phase transformations by the weld thermal cycles of the FBW process and its relation with the properties and some failure modes. In addition, the dilatometric results presented have the potential to provide systematic guidance for turning FBW parameters and/or applying controlled accelerated cooling to achieve improved metallurgical features.
Table 2 Controlled parameters for the flash butt welding of the TR57 intermediate steel rail.
The evaluated material was a pearlitic steel for railroad application classified as intermediate according to AREMA (2013) Standard, profile TR57. The mechanical properties specified for this class of steel rail are shown in Table 1. The rails correspond to a single batch of material and were welded in a Brazilian stationary flash butt welding facility, according to the parameters presented in Table 2. No accelerated cooling was applied after welding. The base metal was subjected to chemical analysis by optical emission spectrometry. Metallographic analysis of the welded joints and the base metal were made at the rail head, 10 mm below the rolling surface (Fig. 1 (a)). The metallographic preparation of samples followed the usual procedures (ASTM E3, 2015) and a 2% Nital etchant was used to reveal the structure. Microstructural analysis and measurements of the pearlite interlamellar spacing using atomic force microscope (AFM) and scanning electron microscope (SEM), were performed. Measurements of pearlite interlamellar spacing were made in different positions: grain growth region (2 mm from the central line of the weld); grain refined region (8 mm from centerline); partial transformation region (12 mm from centerline); base metal. To perform the measurements, those colonies with smaller interlamellar spacing were looked over the micrographs, which are the ones with lamellae perpendicular to the polishing cut (Krauss, 2005). Fifteen different images were used in each region to measure the pearlite interlamellar spacing. The spacing measurements from the topographic images obtained by AFM were performed with the aid of Gwyddion software (Fig. 1 (b)), and in the case of the SEM, the ImageJ software was used. Thermal etching, by oxidation at 700 °C in a controlled atmosphere furnace (low oxygen partial pressure controlled with continuous argon flux), was performed for 10 min in the section shown in Fig. 1 (a). The objective was to measure the pearlite colony sizes in all joint regions, Table 1 Mechanical properties specifications for intermediate steel rail. AREMA (2013). Minimum required
Yield strength (MPa) Tensile strength (MPa) Elongation (%) Superficial Hardness (HB)
551 1013 8 325
Values
Intensity and duration of the initial flash Number of preheat current pulses and pulse duration Intensity of the preheat pulses Force intensity during the preheating pulses Intensity and duration of the final flash Intensity of final force Total displacement
77.4kA 10un. 45-70kA 106kN 38.3kA 477kN 37-45 mm
20s 3.8s – – 14.4s – –
using an optical microscope and the intercept method (ASTM E1382, 2015). For the measurements, five images of each analyzed area were considered. Also in the section 10 mm below the rail head surface, Vickers microhardness profile was drawn with 200gf and 0.5 mm indentation spacing along the entire joint. For the tensile tests, three specimens from the base metal and three from the welded joint were used according to ASTM E8M (2013). The tensile test specimens were machined with the dimensions presented in Fig. 2 (a). The used gage length was larger than the total width of the HAZ (about 30 mm). All specimens were longitudinally sampled from the rail heads at 10 mm below the rolling surface. The tensile specimens from welding joints were etched with Nital 4% to ensure that the full extent of the HAZ was contained in the gage length (Fig. 2 (b)). Considering the Weingrill et al. (2016) results related to welding thermal cycles, besides the measured pearlite colony sizes and interlamellar spacing, two groups of dilatometric tests were performed and evaluated in the present study. Group (i): austenitizing temperature of 900 °C (60 s), used to obtain the continuous cooling transformation (CCT) diagram of the base metal, since the temperature and time of austenitization did not allow the growth of austenitic grains. Group (ii): austenitizing temperature of 1300 °C for 1 s, used to obtain the CCT diagram equivalent to the grain growth region in the HAZ. In both cases, the initial cooling temperature to obtain the austenite decomposition was 900 °C and the applied cooling rates can be seen in Fig. 3 (a). Another group of dilatometric tests was performed in order to simulate the partially transformed region, Fig. 3 (b). The three conditions, showed in Fig. 3 (b), were determined based on the results of Weingrill et al. (2016) and on optical pyrometric measurements performed by this paper authors at the welding plant regarding to the minimum residence times between Ac1 and Ac3. The condition Spheroidization 1 in Fig. 3 (b) is equivalent to the actual FBW process applied to the intermediate steel rail. Spheroidization 2 and Spheroidization 3 aimed to evaluate the influence of higher post-weld cooling rates on the residence time between Ac1 and Ac3 and on the resulting pearlite morphology/microhardness. The specimens for dilatometry (solid cylinders with 10 mm x 3 mm) were machined in the base metal, longitudinal to the rail and at 10 mm below the head rolling surface. The tests were performed in a dilatometer R.I.T.A. L78, with induction heating and helium flux. The specimens underwent metallographic analysis to confirm events obtained in the thermal cycles applied and for comparison to the welded joint. Furthermore, Vickers microhardness was measured in the dilatometry specimens (fifteen measurements on each specimen).
2. Materials and methods
Mechanical Properties
Parameter controlled during welding
3. Results and discussion 3.1. Chemical analysis Table 3 presents the chemical composition of the studied steel. It is possible to conclude that the base metal agrees with the AREMA (2013) Standard requirements. However, as the amount of carbon is relatively 21
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Fig. 1. (a) Schematic representation of the cut section for structural characterization of rails welded joints, indicated by an arrow. (b) Example of a test-line drawn in a direction perpendicular to the lamellae at a pearlite colony imaged by AFM and the corresponding topographic profile used to measure the spacing.
Fig. 2. Specimens for tensile tests. (a) Specimen dimensions (mm); (b) specimen sampled from welded joint etched with Nital 4% highlighting the HAZ.
Fig. 3. Dilatometry tests. (a) Thermal cycles applied to the material to obtain CCT diagrams of the base metal and coarse grained HAZ. (b) Thermal cycles applied to the material to simulate the partially transformed region of the HAZ with different post-weld cooling rates and resultant residence times between Ac1 and Ac3.
22
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Table 3 Results of chemical analysis of the base metal, intermediate steel rail. Chemical composition (wt. %) Average
C 0.715
Mn 0.843
Si 0.240
P 0.016
S 0.008
Cr 0.078
Mo 0.009
Ni 0.012
V 0.002
Nb 0.004
Fig. 4. Micrographs of flash butt welded joint of the intermediate steel rail etched with Nital 2% near the center of the rail head, 10 mm below the rolling surface (Fig. 1 (a)). Original magnification 100 × .
of the pearlite morphology, with partial cementite spheroidization. The lower hardness values of the partial transformation and fine grain regions compared to the base metal and grain growth region, shown in Fig. 5, can be related to the morphological aspects of the pearlite, mainly the interlamellar spacing and the spheroidization. The decrease in hardness due to the partial spheroidization of cementite, especially in the region of partial austenitization, is one of the major technological challenges for rails manufacturers and for railways (Micenko et al., 2013). As presented by Mutton et al. (2016) when investigating different flash butt welded premium steel rails, the spheroidized microstructure shows a significantly different mechanical behavior at the railwheel interface and this is associated with the RCF cracks nucleation on the rail surface. Their experimental work also indicated a higher crack growth rate at the spheroidized region. Other aspects worth mentioning are the lack of studies on the kinetics of cementite spheroidization in the FBW process of pearlitic rails and the relationships between the thermomechanical cycles of the process, the pearlite morphology and the mechanical properties (Porcaro et al., 2017). The characterization results of the welded joints revealed that the grain growth region of the HAZ has larger average size of pearlite colonies than the grain refined region, however, the pearlite interlamellar spacing is lower in the first one compared to the second one. Similar results were presented by Jilabi (2015), but the author could not explain the reasons and the resultant hardness profile. Considering that the post-weld cooling rates are approximately the same over the entire joint for temperatures lower than 900 °C (Messler, 1999), the differences in the interlamellar spacing of pearlite can be due to changes in the hardenability of the material. In other words, the CCT diagram of the grain growth region is shifted to the right and down due to the larger austenitic grain size. Therefore, for the same cooling rate exhibited at temperatures below 900 °C, the grain growth region presents austenite-pearlite transformation at lower temperatures than the grain refined region. This hypothesis will be confirmed in a later section with results from dilatometric tests.
low, a higher amount of manganese was expected, since this type of steel is generally eutectoid and the manganese lowers the eutectoid carbon content (Krauss, 2005). Due to the low amount of manganese and carbon, the steel structure showed some amount of proeutectoid ferrite (less than 5%). 3.2. Microstructural characterization The metallographic analysis is shown in Fig. 4, where it is possible to observe that the entire extent of the welded joint is mostly pearlitic. In the bonding region, partial decarburization occurred with formation of a “line” composed by free ferrite, as indicated in Fig. 4. A partial decarburization area is common in flash butt welded rails joints and it generally does not represent a problem (Micenko et al., 2013). Also in Fig. 4, the grain growth region, grain refined region, and partial transformation region of the HAZ, as well as the base metal, can be identified. A schematic figure relating the peak temperatures during welding with pearlite colonies sizes (revealed by thermal etching), pearlite morphology of each region obtained by SEM and Vickers microhardness is in Fig. 5. It is observed in Fig. 5 that the pearlite colony size changes as a function of the position from the centerline of the weld up to the base metal. The pearlite colony size can be directly related to the previous austenitic grain size at the end of the welding process (Garnham and Davis, 2009). Due to the deformation at the end of the flash butt welding process, dynamic recrystallization of austenite occurs in HAZ (Mansouri and Monshi, 2004). Therefore, the size of the previous austenitic grain and, consequently, the pearlite colonies sizes in each region of the joint depend on the peak temperature: in the grain growth region, the temperature was high enough to allow grain growth and originated the coarser pearlite colonies in the Fig. 5 (45 μm). On the other hand, in the grain refined region, the average pearlite colony size (9 μm) is refined because it derived from a recrystallized/fine austenite. Topographic images obtained by AFM in the various regions of the FBW joint are presented in Fig. 6 and the respective measurements of pearlite interlamellar spacing are in the Fig. 5. It can be seen that the grain growth region of the HAZ showed significantly smaller interlamellar spacing than the grain refined region, with average values of 0.17 μm and 0.20 μm, respectively. It is also observed from Figs. 5 and 6 that the partial transformation region showed a significant modification
3.3. Dilatometry The dilatometry results, obtained for different cooling rates are presented in Fig. 7, with two overlapping CCT diagrams applying different austenitizing temperatures (900 °C and 1300 °C). In addition, the 23
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Fig. 5. Schematic image relating the peak temperature (a) and the various microstructures, pearlite morphology and colony size, pearlite interlamellar spacing and Vickers microhardness of the FBW joint. (b) Thermal etching, (c) SEM. Section represented in Fig. 1 (a).
should be done with care and cannot be based on the CCT diagram of the base metal, since the grain growth can significantly increase the hardenability of the HAZ. However, the Fig. 7 also shows that a postweld accelerated cooling up to 5 °C/s is safer for the steel evaluated here, since there is no martensite formation even in the coarse grained region. Considering a cooling rate around 1 °C/s as a representative value of the welding process applied to the rail steel in this work, characterization results of the dilatometric samples from this cooling rate are shown in Fig. 8 for the two peak temperatures. The results corroborate the hypothesis raised in the previous section regarding the formation of the microstructure in the HAZ, that is, the smaller interlamellar spacing of pearlite in the grain growth region compared to the fine grain region is due to the displacement (to the right and down) of the CCT diagram of the steel. As can be seen in Fig. 8, for the same cooling rate (1 °C/s), the pearlite that originated from the larger austenite grains resulted in a bigger average colony size, but had a smaller interlamellar spacing and higher hardness in comparison to those that originated from the smaller austenite grains. Moreover, analysis of the variance of the data obtained
results of average previous austenitic grain sizes and microhardness are also presented. It is observed in the CCT diagrams that the peak temperature and, consequently, the previous austenitic grain size had a significant influence on the phase transformations during the continuous cooling of the rail steel and in the Vickers microhardness values. Another key point in Fig. 7 is the beginning of the martensitic transformation from a cooling rate equal to 10 °C/s for the largest austenitic grain size. On the other hand, for a smaller austenitic grain size (equivalent to the base metal), martensitic transformation only occurs for cooling rates above 15 °C/s. The natural cooling rate of flash butt welded rails is around 1 °C/s (Weingrill et al., 2016). Saita et al. (2013) and Tawfik et al. (2008) reported that some companies apply accelerated cooling using forced convection and/or water spray after FBW of rails in order to reduce the hardness differences over the HAZ or to increase the productivity. Cai et al. (2011) reported that typical cooling rates of rails after FBW are between 1 °C/s and 10 °C/s. From the results presented in Fig. 7, it is clear that the choice of accelerated cooling rates after FBW of rails
Fig. 6. Microstructures of the FBW joint obtained by AFM near the center of the rail head, 10 mm below the running surface (Fig. 1 (a)). (a) Grain growth, 2 mm from central line; (b) Grain refined, 8 mm from central line; (c) Partial transformation, 12 mm from central line; (d) Base metal, 20 mm from central line. 24
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Fig. 7. Results of the CCT diagrams of the steel from different austenitizing temperatures and average previous austenitic grain sizes. Vickers microhardness for each cooling rate are also shown.
(interlamellar spacing and hardness) indicated different average values for the two austenitizing temperatures. Student’s t-tests performed on the paired data also indicated, with 95% confidence, that there were different means, with p < 0.05. The results from the dilatometric group that simulated the partially transformed region are presented in Fig. 9 and compared to those from the weld joints. As can be seen, the results from the sample spheroidization 1, that simulated the natural welding cooling rate, is very close to the partially transformed region from the weld joint, including pearlite morphology (partial cementite spheroidization) and Vickers microhardness. The time decrease at the intercritical temperature due to the simulated post-weld accelerated cooling rates, along with the higher cooling rates, was effective in inhibiting the partial cementite spheroidization and increasing the hardness of this region. As can be seen in Fig. 9, the tested accelerated cooling rate of 5 °C/s (spheroidization 3) resulted in hardness close to that obtained in the base
metal and in the coarse grained region, representing an increase of 20% (323 HV) compared to the softened region in natural cooled welded joints (269 HV). The combination of the results shown from Figs. 7–9 is very important to understand the structural evolution of the pearlite at the various regions of the HAZ during the FBW of pearlitic rails. Another highlight is that a post-weld controlled accelerated cooling in the rail head is recommended to reduce the hardness drop in the partially transformed region. The BS EN 14587-1 (2005) recommends a postweld treatment just after the burr removal when welding the grade R350HT. The Indian Railways (2012) recommends an air quenching treatment for head hardened rails. To date, no studies have been found showing the results presented here. Besides, the methodology presented using the results from dilatometric tests could be used to determine improved welding conditions for different rail steel grades. For the rail steel presented here, an accelerated cooling up to 5 °C/s can be applied
Fig. 8. Pearlite morphology, austenite grain size and hardness of dilatometric samples cooled at 1 °C/s from different austenitizing temperatures: (a) 900 °C; (b) 1300 °C. 25
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Fig. 9. Pearlite morphology and microhardness in the partial transformation region of the FBW joint in comparison to the dilatometric samples that simulated different post-weld cooling rates on the partially transformed region of the HAZ. (a) Optical microscope; (b) SEM; (c) Hardness.
safely without martensite formation even in the coarse grained region, in other words, there is a process window and it does not mean higher costs or longer welding times and it is effective to minimize the hardness loss in the partial cementite spheroidization region. The lack of knowledge on the mechanisms that control the microstructure development has hindered the setting of FBW process parameters in order to achieve welds with better performance. Among the possibilities found in the literature, one can cite the work of Micenko et al. (2013) that suggest the reduction of the preheating cycles and the formation of joints with narrow HAZ. However, Mutton et al. (2016) warn that this procedure can significantly increase the residual stress levels. The dilatometric results presented in this paper have the potential to provide systematic guidance for turning FBW parameters and/ or applying controlled accelerated cooling to achieve improved metallurgical features.
Fig. 10. Vickers microhardness profile in the head of FBW joint, 10 mm below the surface. Regions of the joint: GG – Grain growth; GR – Grain refined; PA – Partial austenitization; BM – Base metal.
3.4. Mechanical tests
Table 4 Tensile tests results of specimens from base metal (BM) and flash butt welded joints (WJ) of intermediate steel rail.
Results of Vickers microhardness profile in the entire HAZ of the welded rail head is shown in Fig. 10. Due to the smaller interlamellar pearlite spacing, the grain growth region is the only one to present microhardness values closest to the base metal; on the other hand, the microhardness values decrease continuously from the grain refined region up to a minimum of 250 HV in the region of partial transformation. These hardness results are related to the pearlite morphology resulting from the thermomechanical cycles of the flash butt welding process, especially interlamellar spacing and partial cementite spheroidization. As reported by Li et al. (2011), the hardness loss of FBW rails implies the formation of surface irregularities that could increase dynamic impact loads between wheels and rails. The results of tensile tests of the base metal and the welded joints
Specimen
Yield Strength (MPa)
Tensile Strength (MPa)
Elongation (%)
BM WJ
768 ± 32 639 ± 13
1185 ± 31 1035 ± 19
9±3 10.5 ± 0.7
are presented in Table 4. The base metal fulfill the specifications proposed by the AREMA (2013). However, regarding the welds, an average decrease of 17% in the yield strength and 13% in the tensile strength was observed, accompanied by a small increase in the percent elongation. In addition, as illustrated in Fig. 11, all specimens from welded joints fractured in the partial transformation region, between 10 mm 26
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Fig. 11. Tensile specimen of the flash butt welded joint of intermediate steel rail. The local strain and fracture at the partial austenitization region are highlighted.
and 12 mm of the weld centerline. The results were similar to those obtained by Godefroid et al. (2015). The decrease of the yield and tensile strengths of the material due to the FBW process can be attributed to the structural changes previously indicated in the metallography, such as the partial globulization of the cementite in pearlite and the increase in the interlamellar spacing. As can be seen in Fig. 11, both sides of the welded specimen at the end of the HAZ (partial austenitization region) presented local deformation. Again, the controlled post-weld cooling rate based on dilatometry could improve the mechanical results on the rolling surface. Fegredo et al. (1993) showed that partial cementite spheroidization of pearlite increases the wear rate and modify the plastic flow on the surface of rails. The effect of ratchetting (exhausted ductility by the incremental accumulation of surface plastic deformation) on surface cracking would increase with the decrease of yield stress and hardness due to cementite spheroidization. Mutton et al. (2016) suggested this as responsible for the readily formation and rapid propagation of RCF cracks in FBW rails.
Acknowledgements The authors thank the Brazilian VLI Company for kindly providing materials for this study. References ASTM E1382-97, 2015. Standard Test Methods for Determining Average Grain Size Using Semiautomatic and Automatic Image Analysis. ASTM International, West Conshohocken, PA. ASTM E3, 2015. Standard Guide for Preparation of Metallographic Specimens. ASTM International, West Conshohocken, PA. ASTM E8/E8M-13, 2013. Standard Test Methods for Tension Testing of Metallic Materials. ASTM International, West Conshohocken, PA. Boer, H., Masumoto, H., 2001. Niobium in rail steel. Proceedings of the International Symposium Niobium, Orlando, USA. pp. 1–24. BS EN 14587-1, 2005. Railway Applications. Track. Flash Butt Welding of Rails. New R220, R260, R260Mn and R350HT Grade Rails in a Fixed Plant. Cai, Z., Nawafune, M., Ma, N., Qu, Y., Cao, B., Murakawa, H., 2011. Residual stresses in flash butt welded rail. Transactions of Joining and Welding Research Institute 40, 79–87. Fegredo, D.M., Kalousek, J., Shehata, M.T., 1993. The effect of progressive minor spheroidization on the dry-wear rates of a standard carbon and a Cr-Mo alloy rail steel. Wear 161, 29–40. Garnham, J.E., Davis, C.L., 2009. Rail materials. In: In: Lewis, R., Olofsson, U. (Eds.), Wheel-Rail Interface Handbook, vol. 1. CRC Press, Boston, pp. 125–171. Girsch, G., Keichel, J., Gehrmann, R., Zlatnik, A., Frank, N., 2009. Advanced rail steels for heavy haul application – track performance and weldability. Proceedings of the 9th International Heavy Haul Conference, Shangai, China. pp. 171–179. Godefroid, L.B., Faria, G.L., Cândido, L.C., Viana, T.G., 2015. Failure analysis of recurrent cases of fatigue fracture in flash butt welded rails. Eng. Fail. Anal. 58 (2), 407–416. Indian Railways, 2012. Manual for Flash Butt Welding of Rails. Government of India, Ministry of Railways. Research, Design and Standards Organization, Lucknow-11. Jilabi, A.S.J.A.Z., 2015. Welding of Rail Steels. PhD Thesis. University of Manchester, UK, pp. 1–243. Krauss, G., 2005. Steels: Processing, Structure and Performance, first ed. ASM International – Materials Park, Ohio, pp. 281–296. Li, W., Xiao, G., Wen, Z., Xiao, X., Jin, X., 2011. Plastic deformation of curved rail at rail weld caused by train-track dynamic interaction. Wear 271, 311–318. Mansouri, H., Monshi, A., 2004. Microstructure and residual stress variations in Weld Zone fo flash-butt welded railroads. Sci. Technol. Weld. Join. 9 (3), 237–246. AREMA, 2013. American Railway Engineering and Maintenance-of-Way Association Manual of railway engineering. Messler, R.W., 1999. Principles of Welding: Processes, Physics, Chemistry and Metallurgy, first ed. John Wiley & Sons Inc., New York, pp. 161–172. Micenko, P., Muruganant, A., Huijun, Li, Xiaofeng, Xu, 2013. Double dip hardness profiles in rail weld heat-affected zone – literature and research review report. Final Report, Project name: Improvements to Railway Welding. CRC for Rail Innovation, Brisbane, Australia, pp. 1–49. Mutton, P., Cookson, J., Qiu, C., Welsby, D., 2016. Microstructural characterization of rolling contact fatigue damage in Flashbutt Welds. Wear 366-367, 368–377. Olivares, R.O., Garcia, C.I., Deardo, A., Kalay, S., Hernández, F.F.R., 2011. Advanced metallurgical alloy design and thermomechanical processing for rails steels for north american heavy haul use. Wear 271 (1-2), 364–373. Porcaro, R.R., Lima, D.A.P., Faria, G.L., Godefroid, L.B., Cândido, L.C., 2017. Microestrutura e propriedades mecânicas de um aço para trilhos ferroviários soldado por centelhamento. Soldag. Inspeã§ã£o 22 (1), 59–71. Saita, K., Ueda, M., Yamamoto, T., Karimine, K., Iwano, K., Hiroguchi, K., 2013. Trends in rail welding technologies and our future approach. Nippon Steel & Sumitomo Metal Technical Report 105, 84–92. Steenbergen, M., Dollevoet, R., 2013. On the mechanism of squat formation on train rails – part two: growth. Int. J. Fatigue 47, 373–381. Tawfik, D., Mutton, P.J., Chiu, W.K., 2008. Experimental and numerical investigations: alleviating tensile residual stress in flash-butt welds by localized rapid post-weld heat treatment. J. Mater. Process Tech 196 (1-3), 279–291. Weingrill, L., Krutzler, J., Enzinger, N., 2016. Temperature field evolution during flash butt welding of railway rails. Mater. Sci. Forum 879, 2088–2093.
4. Conclusions Detailed mechanical and structural characterizations of flash butt welded joints of a pearlitic steel rail were made, including dilatometric analysis to investigate the effects of welding thermal cycles on the HAZ. The following conclusions can be drawn:
• The HAZ consists of three regions: (i) grain growth region, with
• • • •
larger pearlitic colony size (44.9 μm), but smaller interlamellar pearlite spacing (0.17 μm), which resulted in microhardness close to the base metal (323 HV); (ii) grain refined region, with smaller pearlite colonies (9.2 μm) originated from recrystallized austenite, greater interlamellar spacing of pearlite (0.20 μm), and lower microhardness (286 HV); (iii) partial transformation region, with partial cementite spheroidization in pearlite and the lowest microhardness (262 HV). The CCT diagram of the grain growth region is shifted to the right and down due to the larger austenitic grain size, presenting austenite-pearlite transformation at lower temperatures than the grain refined region, explaining the smaller interlamellar pearlite spacing in the grain growth region. The grain growth region limits the maximum post-weld cooling rate due to the increased hardenability. Dilatometry tests indicated a process window that allows the increase of the post-weld cooling rate up to 5 °C/s without martensite formation for the steel evaluated. The partial spheroidization promoted an average decrease of 17% in the yield strength and 13% in the tensile strength, when compared to the base metal, besides the hardness loss. Dilatometry tests successfully simulated the partial cementite spheroidization at the HAZ and showed that post-weld accelerated cooling is an effective way to reduce the spheroidization, improve hardness distribution and welding conditions. A 20% increase in the hardness of the softened HAZ region was obtained by dilatometric simulation of a safe accelerated cooling (5 °C/s). 27
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Journal of Materials Processing Tech. journal homepage: www.elsevier.com/locate/jmatprotec
Microstructure and mechanical properties of a flash butt welded pearlitic rail
T
Rodrigo Rangel Porcaroa, , Geraldo Lúcio Fariaa, Leonardo Barbosa Godefroida, Gabriela Ribeiro Apolonioa, Luiz Cláudio Cândidoa, Elisângela Silva Pintob ⁎
a b
Universidade Federal de Ouro Preto, Ouro Preto, MG, Brazil Instituto Federal de Educação Ciência e Tecnologia de Minas Gerais, Ouro Preto, MG, Brazil
ARTICLE INFO
ABSTRACT
Associate Editor: C.H. Caceres
The structural changes resulting from the Flash Butt Welding (FBW) of pearlitic rails have been associated with wear/premature failures, despite this, there are no studies applying dilatometry to correlate the welding thermal cycles with the microstructural development of such material. The microstructural evolution of the heat affected zone is clarified with the aid of dilatometry. The increase in the steel hardenability associated with a larger austenitic grain size promotes the austenite-pearlite transformation at lower temperatures in the grain growth region. This explains why this region has larger pearlite colony size but smaller interlamellar pearlite spacing and higher hardness than the grain refined region. Partial cementite spheroidization in the heat affected zone is responsible for significant decrease in hardness and tensile strength and is correlated to localized dipping, rolling contact fatigue and failures. A dilatometry based methodology is proposed to define a process window and control the post-weld cooling rate at the rail head in order to improve the weld performance due to a better hardness profile, without increasing costs or welding time. For the steel evaluated, a 20% increase in the hardness of the softened area at the HAZ was obtained by dilatometric simulation of a safe accelerated cooling (5 °C/s).
Keywords: Flash butt welding Dilatometry Rail steels Mechanical properties of welded joints Welding metallurgy
1. Introduction The increasing demand for rail transportation around the world has been accompanied by increased axle load, higher train speed and frequency of railroad use, which requires more carefulness with structural integrity to reduce rail failures and to increase the transport availability (Girsch et al., 2009). The use of continuous welded rails (CWR) by the flash butt welding (FBW) process has also grown in order to assure higher continuity and to improve the dynamic behavior of the railway system (Mansouri and Monshi, 2004). The microstructure of steels used in railroad tracks worldwide is mainly pearlitic (Boer and Masumoto, 2001). The increase on surface hardness of the rails decreases the wear on the head, and the mechanisms used for this purpose have mainly been the decrease of the pearlite interlamellar spacing by heat treatments and/or use of alloying elements (Olivares et al., 2011). Garnham and Davis (2009) showed how aspects related to the previous austenitic grain size and the pearlite interlamellar spacing influence the mechanical properties of railroad tracks. However, Micenko et al. (2013) reported that there are no studies in the literature describing the structural evolution of pearlitic rails ⁎
during the FBW process. The authors also suggested the use of dilatometry techniques to correlate aspects of welding thermal cycles, such as peak temperature and cooling rates, with the structure development and the resulting properties. Mansouri and Monshi (2004) and Porcaro et al. (2017) describe the structural changes resulting from the FBW of pearlitic rails. These authors presented the heat affected zone (HAZ) of joints composed of three regions: grain growth, grain refined due to recrystallization and partial transformation. The structural changes, also reported by Micenko et al. (2013), include partial decarburization and large variation of hardness in the partial transformation region due to the partial cementite spheroidization. These variations in the rail head hardness after the FBW may cause localized wear or plastic strain during the railway lifetime (Steenbergen and Dollevoet, 2013), which in turn, increases the vibrations associated with fatigue. In a recent work, Mutton et al. (2016) evaluated the microstructure influence on the development of RCF in flash butt welded joints of pearlitic rails. Fatigue cracks developed readily in the spheroidization bands on both sides of the welding centerline and propagated further at these locations. Severe plastic deformation in the softened zones of FBW
Corresponding author. E-mail address: [email protected] (R.R. Porcaro).
https://doi.org/10.1016/j.jmatprotec.2019.02.013 Received 31 August 2018; Received in revised form 11 January 2019; Accepted 10 February 2019 Available online 16 February 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.
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was pointed out as RCF crack initiation sites at the gauge corner. The authors suggest investigating in order to establish a correlation between the welding conditions and the cementite spheroidization process in different rail steel grades, considering the need to better understand the involved mechanisms to promote process improvements. There is a knowledge lack on the mechanisms that control the microstructure development and it has hindered the setting of FBW process parameters in order to improve weld performance. In this context, this paper presents detailed structural and mechanical characterizations of FBW joints of a pearlitic rail steel. Results of measurements of pearlite interlamellar spacing and colonies sizes are presented for all HAZ regions, associated with dilatometric studies to describe steel phase transformations due to the welding process, including the effects of austenitic grain size and partial cementite spheroidization. The results allow understanding, in detail, how the pearlite morphology is affected in terms of phase transformations by the weld thermal cycles of the FBW process and its relation with the properties and some failure modes. In addition, the dilatometric results presented have the potential to provide systematic guidance for turning FBW parameters and/or applying controlled accelerated cooling to achieve improved metallurgical features.
Table 2 Controlled parameters for the flash butt welding of the TR57 intermediate steel rail.
The evaluated material was a pearlitic steel for railroad application classified as intermediate according to AREMA (2013) Standard, profile TR57. The mechanical properties specified for this class of steel rail are shown in Table 1. The rails correspond to a single batch of material and were welded in a Brazilian stationary flash butt welding facility, according to the parameters presented in Table 2. No accelerated cooling was applied after welding. The base metal was subjected to chemical analysis by optical emission spectrometry. Metallographic analysis of the welded joints and the base metal were made at the rail head, 10 mm below the rolling surface (Fig. 1 (a)). The metallographic preparation of samples followed the usual procedures (ASTM E3, 2015) and a 2% Nital etchant was used to reveal the structure. Microstructural analysis and measurements of the pearlite interlamellar spacing using atomic force microscope (AFM) and scanning electron microscope (SEM), were performed. Measurements of pearlite interlamellar spacing were made in different positions: grain growth region (2 mm from the central line of the weld); grain refined region (8 mm from centerline); partial transformation region (12 mm from centerline); base metal. To perform the measurements, those colonies with smaller interlamellar spacing were looked over the micrographs, which are the ones with lamellae perpendicular to the polishing cut (Krauss, 2005). Fifteen different images were used in each region to measure the pearlite interlamellar spacing. The spacing measurements from the topographic images obtained by AFM were performed with the aid of Gwyddion software (Fig. 1 (b)), and in the case of the SEM, the ImageJ software was used. Thermal etching, by oxidation at 700 °C in a controlled atmosphere furnace (low oxygen partial pressure controlled with continuous argon flux), was performed for 10 min in the section shown in Fig. 1 (a). The objective was to measure the pearlite colony sizes in all joint regions, Table 1 Mechanical properties specifications for intermediate steel rail. AREMA (2013). Minimum required
Yield strength (MPa) Tensile strength (MPa) Elongation (%) Superficial Hardness (HB)
551 1013 8 325
Values
Intensity and duration of the initial flash Number of preheat current pulses and pulse duration Intensity of the preheat pulses Force intensity during the preheating pulses Intensity and duration of the final flash Intensity of final force Total displacement
77.4kA 10un. 45-70kA 106kN 38.3kA 477kN 37-45 mm
20s 3.8s – – 14.4s – –
using an optical microscope and the intercept method (ASTM E1382, 2015). For the measurements, five images of each analyzed area were considered. Also in the section 10 mm below the rail head surface, Vickers microhardness profile was drawn with 200gf and 0.5 mm indentation spacing along the entire joint. For the tensile tests, three specimens from the base metal and three from the welded joint were used according to ASTM E8M (2013). The tensile test specimens were machined with the dimensions presented in Fig. 2 (a). The used gage length was larger than the total width of the HAZ (about 30 mm). All specimens were longitudinally sampled from the rail heads at 10 mm below the rolling surface. The tensile specimens from welding joints were etched with Nital 4% to ensure that the full extent of the HAZ was contained in the gage length (Fig. 2 (b)). Considering the Weingrill et al. (2016) results related to welding thermal cycles, besides the measured pearlite colony sizes and interlamellar spacing, two groups of dilatometric tests were performed and evaluated in the present study. Group (i): austenitizing temperature of 900 °C (60 s), used to obtain the continuous cooling transformation (CCT) diagram of the base metal, since the temperature and time of austenitization did not allow the growth of austenitic grains. Group (ii): austenitizing temperature of 1300 °C for 1 s, used to obtain the CCT diagram equivalent to the grain growth region in the HAZ. In both cases, the initial cooling temperature to obtain the austenite decomposition was 900 °C and the applied cooling rates can be seen in Fig. 3 (a). Another group of dilatometric tests was performed in order to simulate the partially transformed region, Fig. 3 (b). The three conditions, showed in Fig. 3 (b), were determined based on the results of Weingrill et al. (2016) and on optical pyrometric measurements performed by this paper authors at the welding plant regarding to the minimum residence times between Ac1 and Ac3. The condition Spheroidization 1 in Fig. 3 (b) is equivalent to the actual FBW process applied to the intermediate steel rail. Spheroidization 2 and Spheroidization 3 aimed to evaluate the influence of higher post-weld cooling rates on the residence time between Ac1 and Ac3 and on the resulting pearlite morphology/microhardness. The specimens for dilatometry (solid cylinders with 10 mm x 3 mm) were machined in the base metal, longitudinal to the rail and at 10 mm below the head rolling surface. The tests were performed in a dilatometer R.I.T.A. L78, with induction heating and helium flux. The specimens underwent metallographic analysis to confirm events obtained in the thermal cycles applied and for comparison to the welded joint. Furthermore, Vickers microhardness was measured in the dilatometry specimens (fifteen measurements on each specimen).
2. Materials and methods
Mechanical Properties
Parameter controlled during welding
3. Results and discussion 3.1. Chemical analysis Table 3 presents the chemical composition of the studied steel. It is possible to conclude that the base metal agrees with the AREMA (2013) Standard requirements. However, as the amount of carbon is relatively 21
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Fig. 1. (a) Schematic representation of the cut section for structural characterization of rails welded joints, indicated by an arrow. (b) Example of a test-line drawn in a direction perpendicular to the lamellae at a pearlite colony imaged by AFM and the corresponding topographic profile used to measure the spacing.
Fig. 2. Specimens for tensile tests. (a) Specimen dimensions (mm); (b) specimen sampled from welded joint etched with Nital 4% highlighting the HAZ.
Fig. 3. Dilatometry tests. (a) Thermal cycles applied to the material to obtain CCT diagrams of the base metal and coarse grained HAZ. (b) Thermal cycles applied to the material to simulate the partially transformed region of the HAZ with different post-weld cooling rates and resultant residence times between Ac1 and Ac3.
22
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Table 3 Results of chemical analysis of the base metal, intermediate steel rail. Chemical composition (wt. %) Average
C 0.715
Mn 0.843
Si 0.240
P 0.016
S 0.008
Cr 0.078
Mo 0.009
Ni 0.012
V 0.002
Nb 0.004
Fig. 4. Micrographs of flash butt welded joint of the intermediate steel rail etched with Nital 2% near the center of the rail head, 10 mm below the rolling surface (Fig. 1 (a)). Original magnification 100 × .
of the pearlite morphology, with partial cementite spheroidization. The lower hardness values of the partial transformation and fine grain regions compared to the base metal and grain growth region, shown in Fig. 5, can be related to the morphological aspects of the pearlite, mainly the interlamellar spacing and the spheroidization. The decrease in hardness due to the partial spheroidization of cementite, especially in the region of partial austenitization, is one of the major technological challenges for rails manufacturers and for railways (Micenko et al., 2013). As presented by Mutton et al. (2016) when investigating different flash butt welded premium steel rails, the spheroidized microstructure shows a significantly different mechanical behavior at the railwheel interface and this is associated with the RCF cracks nucleation on the rail surface. Their experimental work also indicated a higher crack growth rate at the spheroidized region. Other aspects worth mentioning are the lack of studies on the kinetics of cementite spheroidization in the FBW process of pearlitic rails and the relationships between the thermomechanical cycles of the process, the pearlite morphology and the mechanical properties (Porcaro et al., 2017). The characterization results of the welded joints revealed that the grain growth region of the HAZ has larger average size of pearlite colonies than the grain refined region, however, the pearlite interlamellar spacing is lower in the first one compared to the second one. Similar results were presented by Jilabi (2015), but the author could not explain the reasons and the resultant hardness profile. Considering that the post-weld cooling rates are approximately the same over the entire joint for temperatures lower than 900 °C (Messler, 1999), the differences in the interlamellar spacing of pearlite can be due to changes in the hardenability of the material. In other words, the CCT diagram of the grain growth region is shifted to the right and down due to the larger austenitic grain size. Therefore, for the same cooling rate exhibited at temperatures below 900 °C, the grain growth region presents austenite-pearlite transformation at lower temperatures than the grain refined region. This hypothesis will be confirmed in a later section with results from dilatometric tests.
low, a higher amount of manganese was expected, since this type of steel is generally eutectoid and the manganese lowers the eutectoid carbon content (Krauss, 2005). Due to the low amount of manganese and carbon, the steel structure showed some amount of proeutectoid ferrite (less than 5%). 3.2. Microstructural characterization The metallographic analysis is shown in Fig. 4, where it is possible to observe that the entire extent of the welded joint is mostly pearlitic. In the bonding region, partial decarburization occurred with formation of a “line” composed by free ferrite, as indicated in Fig. 4. A partial decarburization area is common in flash butt welded rails joints and it generally does not represent a problem (Micenko et al., 2013). Also in Fig. 4, the grain growth region, grain refined region, and partial transformation region of the HAZ, as well as the base metal, can be identified. A schematic figure relating the peak temperatures during welding with pearlite colonies sizes (revealed by thermal etching), pearlite morphology of each region obtained by SEM and Vickers microhardness is in Fig. 5. It is observed in Fig. 5 that the pearlite colony size changes as a function of the position from the centerline of the weld up to the base metal. The pearlite colony size can be directly related to the previous austenitic grain size at the end of the welding process (Garnham and Davis, 2009). Due to the deformation at the end of the flash butt welding process, dynamic recrystallization of austenite occurs in HAZ (Mansouri and Monshi, 2004). Therefore, the size of the previous austenitic grain and, consequently, the pearlite colonies sizes in each region of the joint depend on the peak temperature: in the grain growth region, the temperature was high enough to allow grain growth and originated the coarser pearlite colonies in the Fig. 5 (45 μm). On the other hand, in the grain refined region, the average pearlite colony size (9 μm) is refined because it derived from a recrystallized/fine austenite. Topographic images obtained by AFM in the various regions of the FBW joint are presented in Fig. 6 and the respective measurements of pearlite interlamellar spacing are in the Fig. 5. It can be seen that the grain growth region of the HAZ showed significantly smaller interlamellar spacing than the grain refined region, with average values of 0.17 μm and 0.20 μm, respectively. It is also observed from Figs. 5 and 6 that the partial transformation region showed a significant modification
3.3. Dilatometry The dilatometry results, obtained for different cooling rates are presented in Fig. 7, with two overlapping CCT diagrams applying different austenitizing temperatures (900 °C and 1300 °C). In addition, the 23
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Fig. 5. Schematic image relating the peak temperature (a) and the various microstructures, pearlite morphology and colony size, pearlite interlamellar spacing and Vickers microhardness of the FBW joint. (b) Thermal etching, (c) SEM. Section represented in Fig. 1 (a).
should be done with care and cannot be based on the CCT diagram of the base metal, since the grain growth can significantly increase the hardenability of the HAZ. However, the Fig. 7 also shows that a postweld accelerated cooling up to 5 °C/s is safer for the steel evaluated here, since there is no martensite formation even in the coarse grained region. Considering a cooling rate around 1 °C/s as a representative value of the welding process applied to the rail steel in this work, characterization results of the dilatometric samples from this cooling rate are shown in Fig. 8 for the two peak temperatures. The results corroborate the hypothesis raised in the previous section regarding the formation of the microstructure in the HAZ, that is, the smaller interlamellar spacing of pearlite in the grain growth region compared to the fine grain region is due to the displacement (to the right and down) of the CCT diagram of the steel. As can be seen in Fig. 8, for the same cooling rate (1 °C/s), the pearlite that originated from the larger austenite grains resulted in a bigger average colony size, but had a smaller interlamellar spacing and higher hardness in comparison to those that originated from the smaller austenite grains. Moreover, analysis of the variance of the data obtained
results of average previous austenitic grain sizes and microhardness are also presented. It is observed in the CCT diagrams that the peak temperature and, consequently, the previous austenitic grain size had a significant influence on the phase transformations during the continuous cooling of the rail steel and in the Vickers microhardness values. Another key point in Fig. 7 is the beginning of the martensitic transformation from a cooling rate equal to 10 °C/s for the largest austenitic grain size. On the other hand, for a smaller austenitic grain size (equivalent to the base metal), martensitic transformation only occurs for cooling rates above 15 °C/s. The natural cooling rate of flash butt welded rails is around 1 °C/s (Weingrill et al., 2016). Saita et al. (2013) and Tawfik et al. (2008) reported that some companies apply accelerated cooling using forced convection and/or water spray after FBW of rails in order to reduce the hardness differences over the HAZ or to increase the productivity. Cai et al. (2011) reported that typical cooling rates of rails after FBW are between 1 °C/s and 10 °C/s. From the results presented in Fig. 7, it is clear that the choice of accelerated cooling rates after FBW of rails
Fig. 6. Microstructures of the FBW joint obtained by AFM near the center of the rail head, 10 mm below the running surface (Fig. 1 (a)). (a) Grain growth, 2 mm from central line; (b) Grain refined, 8 mm from central line; (c) Partial transformation, 12 mm from central line; (d) Base metal, 20 mm from central line. 24
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Fig. 7. Results of the CCT diagrams of the steel from different austenitizing temperatures and average previous austenitic grain sizes. Vickers microhardness for each cooling rate are also shown.
(interlamellar spacing and hardness) indicated different average values for the two austenitizing temperatures. Student’s t-tests performed on the paired data also indicated, with 95% confidence, that there were different means, with p < 0.05. The results from the dilatometric group that simulated the partially transformed region are presented in Fig. 9 and compared to those from the weld joints. As can be seen, the results from the sample spheroidization 1, that simulated the natural welding cooling rate, is very close to the partially transformed region from the weld joint, including pearlite morphology (partial cementite spheroidization) and Vickers microhardness. The time decrease at the intercritical temperature due to the simulated post-weld accelerated cooling rates, along with the higher cooling rates, was effective in inhibiting the partial cementite spheroidization and increasing the hardness of this region. As can be seen in Fig. 9, the tested accelerated cooling rate of 5 °C/s (spheroidization 3) resulted in hardness close to that obtained in the base
metal and in the coarse grained region, representing an increase of 20% (323 HV) compared to the softened region in natural cooled welded joints (269 HV). The combination of the results shown from Figs. 7–9 is very important to understand the structural evolution of the pearlite at the various regions of the HAZ during the FBW of pearlitic rails. Another highlight is that a post-weld controlled accelerated cooling in the rail head is recommended to reduce the hardness drop in the partially transformed region. The BS EN 14587-1 (2005) recommends a postweld treatment just after the burr removal when welding the grade R350HT. The Indian Railways (2012) recommends an air quenching treatment for head hardened rails. To date, no studies have been found showing the results presented here. Besides, the methodology presented using the results from dilatometric tests could be used to determine improved welding conditions for different rail steel grades. For the rail steel presented here, an accelerated cooling up to 5 °C/s can be applied
Fig. 8. Pearlite morphology, austenite grain size and hardness of dilatometric samples cooled at 1 °C/s from different austenitizing temperatures: (a) 900 °C; (b) 1300 °C. 25
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Fig. 9. Pearlite morphology and microhardness in the partial transformation region of the FBW joint in comparison to the dilatometric samples that simulated different post-weld cooling rates on the partially transformed region of the HAZ. (a) Optical microscope; (b) SEM; (c) Hardness.
safely without martensite formation even in the coarse grained region, in other words, there is a process window and it does not mean higher costs or longer welding times and it is effective to minimize the hardness loss in the partial cementite spheroidization region. The lack of knowledge on the mechanisms that control the microstructure development has hindered the setting of FBW process parameters in order to achieve welds with better performance. Among the possibilities found in the literature, one can cite the work of Micenko et al. (2013) that suggest the reduction of the preheating cycles and the formation of joints with narrow HAZ. However, Mutton et al. (2016) warn that this procedure can significantly increase the residual stress levels. The dilatometric results presented in this paper have the potential to provide systematic guidance for turning FBW parameters and/ or applying controlled accelerated cooling to achieve improved metallurgical features.
Fig. 10. Vickers microhardness profile in the head of FBW joint, 10 mm below the surface. Regions of the joint: GG – Grain growth; GR – Grain refined; PA – Partial austenitization; BM – Base metal.
3.4. Mechanical tests
Table 4 Tensile tests results of specimens from base metal (BM) and flash butt welded joints (WJ) of intermediate steel rail.
Results of Vickers microhardness profile in the entire HAZ of the welded rail head is shown in Fig. 10. Due to the smaller interlamellar pearlite spacing, the grain growth region is the only one to present microhardness values closest to the base metal; on the other hand, the microhardness values decrease continuously from the grain refined region up to a minimum of 250 HV in the region of partial transformation. These hardness results are related to the pearlite morphology resulting from the thermomechanical cycles of the flash butt welding process, especially interlamellar spacing and partial cementite spheroidization. As reported by Li et al. (2011), the hardness loss of FBW rails implies the formation of surface irregularities that could increase dynamic impact loads between wheels and rails. The results of tensile tests of the base metal and the welded joints
Specimen
Yield Strength (MPa)
Tensile Strength (MPa)
Elongation (%)
BM WJ
768 ± 32 639 ± 13
1185 ± 31 1035 ± 19
9±3 10.5 ± 0.7
are presented in Table 4. The base metal fulfill the specifications proposed by the AREMA (2013). However, regarding the welds, an average decrease of 17% in the yield strength and 13% in the tensile strength was observed, accompanied by a small increase in the percent elongation. In addition, as illustrated in Fig. 11, all specimens from welded joints fractured in the partial transformation region, between 10 mm 26
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Fig. 11. Tensile specimen of the flash butt welded joint of intermediate steel rail. The local strain and fracture at the partial austenitization region are highlighted.
and 12 mm of the weld centerline. The results were similar to those obtained by Godefroid et al. (2015). The decrease of the yield and tensile strengths of the material due to the FBW process can be attributed to the structural changes previously indicated in the metallography, such as the partial globulization of the cementite in pearlite and the increase in the interlamellar spacing. As can be seen in Fig. 11, both sides of the welded specimen at the end of the HAZ (partial austenitization region) presented local deformation. Again, the controlled post-weld cooling rate based on dilatometry could improve the mechanical results on the rolling surface. Fegredo et al. (1993) showed that partial cementite spheroidization of pearlite increases the wear rate and modify the plastic flow on the surface of rails. The effect of ratchetting (exhausted ductility by the incremental accumulation of surface plastic deformation) on surface cracking would increase with the decrease of yield stress and hardness due to cementite spheroidization. Mutton et al. (2016) suggested this as responsible for the readily formation and rapid propagation of RCF cracks in FBW rails.
Acknowledgements The authors thank the Brazilian VLI Company for kindly providing materials for this study. References ASTM E1382-97, 2015. Standard Test Methods for Determining Average Grain Size Using Semiautomatic and Automatic Image Analysis. ASTM International, West Conshohocken, PA. ASTM E3, 2015. Standard Guide for Preparation of Metallographic Specimens. ASTM International, West Conshohocken, PA. ASTM E8/E8M-13, 2013. Standard Test Methods for Tension Testing of Metallic Materials. ASTM International, West Conshohocken, PA. Boer, H., Masumoto, H., 2001. Niobium in rail steel. Proceedings of the International Symposium Niobium, Orlando, USA. pp. 1–24. BS EN 14587-1, 2005. Railway Applications. Track. Flash Butt Welding of Rails. New R220, R260, R260Mn and R350HT Grade Rails in a Fixed Plant. Cai, Z., Nawafune, M., Ma, N., Qu, Y., Cao, B., Murakawa, H., 2011. Residual stresses in flash butt welded rail. Transactions of Joining and Welding Research Institute 40, 79–87. Fegredo, D.M., Kalousek, J., Shehata, M.T., 1993. The effect of progressive minor spheroidization on the dry-wear rates of a standard carbon and a Cr-Mo alloy rail steel. Wear 161, 29–40. Garnham, J.E., Davis, C.L., 2009. Rail materials. In: In: Lewis, R., Olofsson, U. (Eds.), Wheel-Rail Interface Handbook, vol. 1. CRC Press, Boston, pp. 125–171. Girsch, G., Keichel, J., Gehrmann, R., Zlatnik, A., Frank, N., 2009. Advanced rail steels for heavy haul application – track performance and weldability. Proceedings of the 9th International Heavy Haul Conference, Shangai, China. pp. 171–179. Godefroid, L.B., Faria, G.L., Cândido, L.C., Viana, T.G., 2015. Failure analysis of recurrent cases of fatigue fracture in flash butt welded rails. Eng. Fail. Anal. 58 (2), 407–416. Indian Railways, 2012. Manual for Flash Butt Welding of Rails. Government of India, Ministry of Railways. Research, Design and Standards Organization, Lucknow-11. Jilabi, A.S.J.A.Z., 2015. Welding of Rail Steels. PhD Thesis. University of Manchester, UK, pp. 1–243. Krauss, G., 2005. Steels: Processing, Structure and Performance, first ed. ASM International – Materials Park, Ohio, pp. 281–296. Li, W., Xiao, G., Wen, Z., Xiao, X., Jin, X., 2011. Plastic deformation of curved rail at rail weld caused by train-track dynamic interaction. Wear 271, 311–318. Mansouri, H., Monshi, A., 2004. Microstructure and residual stress variations in Weld Zone fo flash-butt welded railroads. Sci. Technol. Weld. Join. 9 (3), 237–246. AREMA, 2013. American Railway Engineering and Maintenance-of-Way Association Manual of railway engineering. Messler, R.W., 1999. Principles of Welding: Processes, Physics, Chemistry and Metallurgy, first ed. John Wiley & Sons Inc., New York, pp. 161–172. Micenko, P., Muruganant, A., Huijun, Li, Xiaofeng, Xu, 2013. Double dip hardness profiles in rail weld heat-affected zone – literature and research review report. Final Report, Project name: Improvements to Railway Welding. CRC for Rail Innovation, Brisbane, Australia, pp. 1–49. Mutton, P., Cookson, J., Qiu, C., Welsby, D., 2016. Microstructural characterization of rolling contact fatigue damage in Flashbutt Welds. Wear 366-367, 368–377. Olivares, R.O., Garcia, C.I., Deardo, A., Kalay, S., Hernández, F.F.R., 2011. Advanced metallurgical alloy design and thermomechanical processing for rails steels for north american heavy haul use. Wear 271 (1-2), 364–373. Porcaro, R.R., Lima, D.A.P., Faria, G.L., Godefroid, L.B., Cândido, L.C., 2017. Microestrutura e propriedades mecânicas de um aço para trilhos ferroviários soldado por centelhamento. Soldag. Inspeã§ã£o 22 (1), 59–71. Saita, K., Ueda, M., Yamamoto, T., Karimine, K., Iwano, K., Hiroguchi, K., 2013. Trends in rail welding technologies and our future approach. Nippon Steel & Sumitomo Metal Technical Report 105, 84–92. Steenbergen, M., Dollevoet, R., 2013. On the mechanism of squat formation on train rails – part two: growth. Int. J. Fatigue 47, 373–381. Tawfik, D., Mutton, P.J., Chiu, W.K., 2008. Experimental and numerical investigations: alleviating tensile residual stress in flash-butt welds by localized rapid post-weld heat treatment. J. Mater. Process Tech 196 (1-3), 279–291. Weingrill, L., Krutzler, J., Enzinger, N., 2016. Temperature field evolution during flash butt welding of railway rails. Mater. Sci. Forum 879, 2088–2093.
4. Conclusions Detailed mechanical and structural characterizations of flash butt welded joints of a pearlitic steel rail were made, including dilatometric analysis to investigate the effects of welding thermal cycles on the HAZ. The following conclusions can be drawn:
• The HAZ consists of three regions: (i) grain growth region, with
• • • •
larger pearlitic colony size (44.9 μm), but smaller interlamellar pearlite spacing (0.17 μm), which resulted in microhardness close to the base metal (323 HV); (ii) grain refined region, with smaller pearlite colonies (9.2 μm) originated from recrystallized austenite, greater interlamellar spacing of pearlite (0.20 μm), and lower microhardness (286 HV); (iii) partial transformation region, with partial cementite spheroidization in pearlite and the lowest microhardness (262 HV). The CCT diagram of the grain growth region is shifted to the right and down due to the larger austenitic grain size, presenting austenite-pearlite transformation at lower temperatures than the grain refined region, explaining the smaller interlamellar pearlite spacing in the grain growth region. The grain growth region limits the maximum post-weld cooling rate due to the increased hardenability. Dilatometry tests indicated a process window that allows the increase of the post-weld cooling rate up to 5 °C/s without martensite formation for the steel evaluated. The partial spheroidization promoted an average decrease of 17% in the yield strength and 13% in the tensile strength, when compared to the base metal, besides the hardness loss. Dilatometry tests successfully simulated the partial cementite spheroidization at the HAZ and showed that post-weld accelerated cooling is an effective way to reduce the spheroidization, improve hardness distribution and welding conditions. A 20% increase in the hardness of the softened HAZ region was obtained by dilatometric simulation of a safe accelerated cooling (5 °C/s). 27