Materials and Design

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1 Materials and Design 53 (2014) Contents lists available at SciVerse ScienceDirect Materials and Design journal homepage: Effects of Mo, Cr and Nb on microstructure and mechanical properties of heat affected zone for Nb-bearing X80 pipeline steels Xiao-wei Chen a,b, Gui-ying Qiao c, Xiu-lin Han b, Xu Wang a,b, Fu-ren Xiao a,, Bo Liao a a Key Lab of Metastable Materials Science & Technology, College of Materials Science & Engineering, Yanshan University, Qinhuangdao , China b CNPC Bohai Equipment Manufacturing Co. Ltd., Qingxian , China c School of Environmental and Chemical Engineering, Yanshan University, Qinhuangdao , China article info abstract Article history: Received 21 March 2013 Accepted 6 July 2013 Available online 20 July 2013 Keywords: High-Nb steel Microstructure Mechanical property The microstructure and mechanical properties of welding heat affected zone (HAZ) of three typical X80 pipeline steels, i.e. Mn Cr Nb, Mn Mo Nb and/or Mn Cr Mo Nb steels, have been studied using the welding thermal simulation method on a Gleeble-3500 thermal simulator. The results show that the chemical compositions and welding process parameters have significant effects on the microstructure and properties of HAZ. With increase of the cooling rate, the amount of microstructure transformed at lower temperature increases, and the microstructure becomes finer, furthermore, the strength and toughness of HAZ show a rising trend. Within the range of the peak temperature from 600 to 1350 C, there are two brittle zones and one major strength weakening zone in HAZ, especially for high-nb added Cr steel, weaken of HAZ is very serious. Effects of the key chemical elements such as Mo, Cr and Nb on microstructure transformation of HAZ have been discussed. Although these three alloy system steels in this work display good weldability, the Mn Mo Nb and/or Mn Cr Mo Nb steels show better combination of strength and toughness of HAZ compared with the Mn Cr Nb steels. Therefore, for high-nb steels, it is necessary to add suitably Mo to improve the toughness and strength of HAZ, especially, increase the strength. Crown Copyright Ó 2013 Published by Elsevier Ltd. All rights reserved. 1. Introduction Ever since the new concept of high-nb pipeline steels was proposed in the early 1990s [1], the high-nb steels have shown up many advantages, such as the balance of high strength and excellent toughness, good weldability, and less cost. Therefore, the high-nb pipeline steels have been widely applied in many long-distance gas transportation projects, e.g., the Cantarell Project in Mexico, the Cheyenne Plains Project in North America and the second West-to-East Gas Transportation Pipeline Project in China [2 5]. The strongest attractions of these widely applied steels are that the austenite recrystallization temperature increases greatly with the increasing of dissolved Nb in austenite, and the steel can be produced by high temperature process (HTP) [6 9], meanwhile, the dissolved Nb in austenite can restrain the ferrite transformation and promote acicular ferritic (or low-carbon bainite) transformation. Thereby, the high strength pipeline steels without Mo (or little Mo) can be produced by the high temperature processing (HTP), consequently, the cost of the steel plates or coils can be decreased greatly [4,5]. Comparing with the conventional Mn Mo Nb high-strength pipeline steel, the high-nb steels add Corresponding author. Tel./fax: addresses: frxiao@ysu.edu.cn, cyddjyjs@263.net (F.-r. Xiao). Cr to substitute partly or wholly for high-cost Mo, and many laboratory researches revealed that the Cr-added high-nb steels have good strength, toughness and weldability [10,11]. Therefore, the high-strength X80 pipeline steels of two types of alloy systems, i.e., Mn Cr Nb and Mo Cr Nb of low-mo content, have been widely applied in the Second West-to-East Gas Transportation Pipeline Project [4,5]. However, during the actual production process of the X80 steel pipe, the reject ratio of steel pipes coiled by Mn Cr Nb steels is higher, which is always considered that it is because of the degeneration of toughness in HAZ [12]. Whereas, our statistic analysis results for actual steel pipes show that the pipe body and HAZ of the pipes coiled by Mn Cr Nb steels, compared with pipes coiled by Mn Cr Nb with low-mo steels, have higher impact toughness and lower strength. Therefore, it is necessary to systematically study the difference of variation of microstructure and mechanical properties in HAZ between steels of different alloy systems, and deeply understand the effectiveness of alloy elements such as Cr, Mo and/or Nb in HAZ, furthermore, improve the mechanical properties of the X80 steel pipe. In this work, three typical X80 steels, i.e. Mn Cr Nb, Mn Mo Nb and/or Mn Cr Mo Nb, were selected, and the microstructure and mechanical properties of HAZ were studied using the welding thermal simulation method, furthermore, the effects of alloys and /$ - see front matter Crown Copyright Ó 2013 Published by Elsevier Ltd. All rights reserved.

2 X.-w. Chen et al. / Materials and Design 53 (2014) weld heat inputs were discussed. The results will be helpful for developing of high-nb pipeline steels. 2. Experimental procedures Test pieces of these three experimental steels were cut from the central section of economical API X80 steel plates, the chemical compositions are shown in Table 1. The test pieces were machined into specimens of mm for HAZ simulation. The welding thermal simulation experiments were carried out on a Gleeble 3500 thermal mechanical simulator. For the welding CCT diagrams, the specimens were uniformly and rapidly heated up at a rate of 100 C/s to a peak temperature of 1350 C and held for 1 s, then cooled down with different cooling rates to the room temperature. The cooling rates were in a range from 60 C/s to 0.2 C/s. In order to study the effects of peak temperature on microstructure and properties of HAZ, the heat inputs of 2.5 kj/mm and 4.0 kj/ mm, and the peak temperatures from 1350 C to 650 C were selected to simulate the reheated HAZ from CGHAZ region to the sub-intercritically reheated region of O.D mm pipe under the multi-wire tandem submerged arc welding condition. The thermal cycle parameters were determined by ANSYS analysis. e.g. For the CGHAZ at the peak temperature of 1350 C, the specimens were uniformly and rapidly heated up at a rate of 100 C/s to a peak temperature of 1350 C and held for 1 s, then cooled down from 800 C to 500 C in 12 s and 44 s respectively. The other cycles were conducted at different peak temperatures, and the cooling times below 800 C were the same as those in the above thermal cycles. After simulation, the specimens were further machined into standard Charpy V-notch samples for impact toughness measurements at 10 C according to ASTM:E23. Some type specimens after thermal cycle were selected for microstructure observation. The microstructure observations were conducted using optical microscopy and transmission electron microscopy (TEM). The content of dissolved Nb of the specimens after treatment at different experimental conditions was determined. The specimens were dissolved in hydrochloric acid, the solution was filtrated. Then the Nb atomic emission spectrometry in solution was measured by inductively coupled plasma-atomic emission spectrometry (ICP-AES). According to these data, the content of Nb in solution and/or precipitates at different experimental conditions was obtained. 3. Results 3.1. Effect of cooling rate on microstructure of CGHAZ Figs. 1 3 show typical optical microstructures of three steels after reheat at 1350 C and cooling down at different cooling rates. For the low-carbon microalloyed steels, in order to identify the microstructure, the classified method proposed in Ref. [13,14] was used in this work, i.e., the microstructure is classified as polygonal ferrite (PF), quasi-polygonal ferrite (QF or massive ferrite MF), granular bainite (GB), bainite (BF). For Steel A of typical Mn Cr Nb system steel, as the cooling rate is lower than 0.5 C/s, coarse non-polygonal ferrite with some black island constituents distributing freely in ferrite grains or at the ferrite grain boundaries appears in the microstructure, and the prior austenite grain boundary is not observed. The microstructure shows up a typical QF (or MF) characteristic (Fig. 1a), which is usually called as acicular ferrite (AF) in pipeline steels [15]. With the increase of the cooling rate, the microstructure becomes the mixture of QF with GB. For the GB, it is considered that some island constituents are distributed in matrix, and the prior austenite grain boundary network exists in the microstructure. With the increase of the cooling rate, the fraction of GB increases, and the size of ferrite and island constituent decreases, meanwhile, the prior austenite grain boundaries appear in the microstructure. As the cooling rate increases to 7 C/s, the microstructure mainly consists of GB, which appears that the fine islands tend to distribute linearly, and the prior austenite grain boundaries become a clear network (Fig. 1b). As the cooling rate increases to 20 C/s, the fine islands change to bundle shape distributing between ferrite laths, which appears BF characteristic (Fig. 1c). However, with the further increase of the cooling rate, although the bainite microstructure becomes finer, some quasi-polygonal acicular ferrite still can be found in the microstructure (Fig. 1d). For the Steel B of typical Mn Mo Nb system steel, the microstructure transformation law is same as the Steel A, however, some difference can be found (Fig. 2). As the cooling rate is 0.5 C/s, the microstructure is mainly QF (Fig. 2a), while, it is finer than that of Steel A. As the cooling rate increases to 7 C/s, the microstructure changes and mainly consists of BF and GB. The microstructure is finer, and the film-like islands distribute between the ferrite laths, furthermore, the prior austenite grain boundaries become a clear network (Fig. 2b). With further increase of the cooling rate, the microstructure becomes dominant BF, while some acicular ferrite segmented between the BF colonies can be observed (Fig. 2c and d). In contrast to the microstructures of Steel A and B, the PF and QF transformation of Steel C with Mn Cr Mo Nb system is restrained due to complex addition of Cr Mo. For example, as the cooling rate is 0.5 C/s, the microstructure mainly consists of GB, which appears some dot-line and needle-line islands distributing in the ferrite matrix, meanwhile, the prior austenite grain boundary can be observed clearly (Fig. 3a). As the cooling rate is increased to 7 C/s, the microstructure changes and mainly consists of bainite, and the ferrite matrix converts to lathy form, meanwhile, the islands change to a film-like form distributing between the ferrite laths (Fig. 3b). With further increase of the cooling rate, the microstructure becomes finer (Fig. 3c and d). As stated above, the transformed products of the steels in the tested cooling rate range are the intermediate transformation microstructures, possibly including QF (or MF), GB, BF. The bainitic ferrite is very complex, which appears to consist of different forms of ferrite and island constituents scattered throughout the matrix, while the ferritic forms can not be identified clearly by means of optical microscopy. Therefore, by examining in detailed TEM of the microstructures obtained at typical cooling rates for two steels, the difference of ferritic forms and island constituents can be revealed clearly, which are shown in Figs. 4 and 5. The morphology Table 1 Chemical composition of experimental steels (wt.%). Steels C Si Mn P S Mo Ni + Cu Cr Nb Ti Ceq Pcm A B C Note: Ceq ¼ C þ Mn 6 þ NiþCu 15 þ CrþMoþV 5 ; Pcm ¼ C þ Si 30 þ MnþCuþCr 20 þ Ni 60 þ Mo 15 þ V 10 þ 5B.

3 890 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 1. Optical microstructure of Steel A at different cooling rates (a) 0.5 C/s; (b) 7 C/s; (c) 20 C/s; (d) 40 C/s. Fig. 2. Optical microstructure of Steel B at different cooling rates (a) 0.5 C/s; (b) 7 C/s; (c) 20 C/s; (d) 40 C/s. of QF (or MF) obtained at cooling rate of 0.5 C/s for Steel A is composed of an assemblage of interwoven nonparallel ferrite laths and nonequiaxed fine grained structures (Fig. 4a and b), meanwhile, some M/A island constituents distribute at the nonequiaxed ferrite grain boundary (Fig. 4a), even some bainite contained carbide can be observed (Fig. 4c). Moreover, a fairly high dislocation density presents in the nonparallel ferrite laths and nonequiaxed fine ferrite grain under TEM, which is an essential characteristic of intermediate transformation products. However, for GB obtained at the cooling rate of 0.5 C/s for Steel C, the morphology of ferrite is composed of an assemblage of interwoven nonparallel ferrite laths (Fig. 5a) and parallel ferrite laths structures (Fig. 5b). Some refined M/A island constituents are entrapped among ferrite laths (Fig. 5a), and some large size M/A island constituents distribute at the bainite colony boundaries and prior austenite grain boundaries (Fig. 5c). As the cooling rate is at 7 C/s, the detailed morphology of the ferrite and M/A island constituents in TEM for Steel A (Fig. 4d f)

4 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 3. Optical microstructure of Steel C at different cooling rates (a) 0.5 C/s; (b) 7 C/s; (c) 20 C/s; (d) 40 C/s. is similar to the microstructure of Steel C obtained at the cooling rate of 0.5 C/s, the results indicate that the microstructure is mainly composed of GB. While, some differences can be found, i.e., the amount and size of the M/A constituent decrease, moreover, the M/A constituents change to dot-like and rod-like forms, and are entrapped among the ferrite laths or at the prior austenite grain boundary (Fig. 4f). In contrast, the microstructure of Steel C at the cooling rate of 7 C/s mainly consists of mixture of parallel lathy ferrite (Fig. 5d) with interwoven nonparallel ferrite laths (Fig. 5e). A few thin film-like M/A constituents exist between the parallel ferrite laths (Fig. 5d), and the rod-like M/A constituents exist in the nonequiaxed ferrite matrix (Fig. 5e). Meanwhile, the parallel lathy ferrite and/or nonequiaxed ferrite are transformed from the prior austenite grain boundary, and some small M/A constituents remain at the prior austenite grain boundary (Fig. 5f). However, as the cooling rate is 20 C/s, the microstructures of these two steels are mainly dominated by the lathy ferrite, and the rod-like M/A constituents distributing between the ferrite laths (Figs. 4g and h and 5g i). Meanwhile, some massive ferrite distributed dot-like M/A constituents freely can be found in Steel A (Fig. 4i), which is deemed as the acicular ferrite in the optical micrograph (Fig. 1a). From the results as stated above, addition of Mo in high-nb steel can restrain PF and QF (or MF) transformation and promote GB and BF transformation, and affect the mechanical properties of HAZ Welding CCT diagrams of CGHAZ The welding CCT diagrams are believed to be useful in gaining an insight into microstructures and designing an optimal welding process during the industrial processing of the welded steel pipes. The welding CCT diagrams of the three steels are constructed using the linear cooling rate dilatometric method and the aforementioned microstructural observations. As shown in Fig. 6, the three steels all exhibit multilayer transformation curves in the welding CCT diagrams. The major transformation curves of these steels almost lie in the temperature range from 410 to 695 C for Steel A, from 396 to 675 C for Steel B and from 350 to 640 C for Steel C, respectively. Because the three steels have different amounts of added alloy elements, such Cr, Mo, Nb etc., the phase transformation temperatures and critical cooling rates are different. For Steel A (Mn Cr Nb), in the welding CCT diagram (Fig. 6a), the nonisothermal undercooled austenite transformation could occur at C, and the full intermediate transformed microstructures could be achieved when the cooling rates are higher than 0.5 C/s. Cooling rates higher than 10 C/s would promote the predominant formation of a mixture of bainite with parallel lath ferrite (BF) and granular bainite with the dot-like and/or rod-like M/A constituents distributing in the ferrite matrix (GB). Cooling rates lower than 10 C/s would promote the predominant formation of a mixture of GB and QF (or MF). While, as the cooling rates are lower than 0.5 C/s, the polygonal ferrite transformation field appears in up right corner in CCT diagram. For Steel B (Mn Mo Nb), the welding CCT diagram shows similarities in the shapes as Steel A, while, due to Mo is added in steel to replace Cr, the temperature of undercooled austenite transformation decreases by about 20 C (Fig. 6b). The result indicates that Mo is a stronger element than Cr in decreasing the austenite transformation. Moreover, the compound addition of Cr and Mo can strengthen the effectiveness of decrease austenite transformation and restraining the PF of QF transformation. The effect illustrates clearly by comparing Fig. 6c with Fig. 6a and b. For Steel C (Mn Cr Mo Nb), the nonisothermal undercooled austenite transformation could occur at C, which is below about 50 C or30 C than the Steel A or Steel B, respectively. Moreover, the PF transformation region do not appear in the tested cooling rate range, the QF transformation can be avoided at a cooling rate of slightly lower than 0.5 C/s, and full BF with lathy ferrite structure is obtained at the cooling rate of slightly lower than 20 C/s Effects of cooling rate on mechanical properties Fig. 7 shows the effects of cooling rates on mechanical properties of HAZ of the steels. As the cooling rate is below 2 C/s, the strength and impact energy appears a slight increase trend with the increase of the cooling rate. Comparing the three steels, the impact energy of Steel A is slightly higher than Steel B and Steel C. In

5 892 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 4. TEM micrographs of the samples of Steel A at cooling rates of (a c) 0.5 C/s; (d f) 7 C/s; (g i) 20 C/s. contrast, the strength of Steel A is lower than Steel B and Steel C obviously, and the yield strength of Steel A is even lower than the technical standard value of 555 MPa. However, as the cooling rate is over 2 C/s, the strength and impact energy increase obviously with the increase of the cooling rate. While, as the cooling rate is over 20 C/s, the increasing trends of strength and impact energy decrease, even begin to decrease. Moreover, the strength and impact energy of Steel C are higher than those of Steel B and Steel A. The changes of mechanical properties attribute to difference of the transformed microstructures at different cooling rates for the three steels. The Steel A with Mn Cr Nb system has low strength at lower cooling rates, which would results in the weakening strength in CGHAZ during the welding process with high welding heat input Effects of peak temperature on mechanical properties of HAZ The welding CCT diagrams are believed to be useful in gaining an insight into microstructures and forecasting the mechanical properties of HAZs during the industrial processing of welding steel pipe [16], however, the actual welding cooling processes are non-linear, which will result in some difference of the microstructures and the mechanical properties of HAZ forecasted by welding CCT diagrams between the actual welded pipes. On the other hand, the microstructure of HAZ of the actual welded pipe is made up of grad structures under different peak temperatures. Therefore, it is necessary to further investigate the changes of microstructures and mechanical properties of HAZs at the different peak temperatures during the actual pipe welding processes, so the microstructure and mechanical properties of HAZs of the tested steels were investigated by thermal simulation with two kinds of welding heat inputs in order to simulate the typical two-wire and four-wire submerged arc welding. The deterioration of HAZ is one of many concerned issues for pipeline steels [11,12]. The effects of the welding heat input and the peak temperature on the impact toughness of HAZs of the steels are shown in Fig. 8. There are two brittle zones presented at the peak temperatures of 1350 C and 750 C, respectively. With the increase of the heat inputs, the impact toughness of those zones decrease. While, favorite impact toughness of HAZ can be obtained when the peak temperature is within the range form 1000 C to 900 C or below 700 C, which comes near to or even higher than those of the base metals. However, for the brittle zone of coarse grain heat affected zone (CGHAZ), the impact energy values are higher than 100 J. The results show that the high-nb steels have good weldability. However, from the results of effects of the cooling rates on mechanical properties (Fig. 7), as the cooling rates are below 5 C/s, the CGHAZs have the lowest impact toughness and strength, especially, for steel A, the strength is lower than the standard value of 555 MPa. The results indicate that the strength becomes another key issue for high-nb X80 pipeline steels, so the effects of the peak

6 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 5. TEM micrographs of the samples of Steel C at cooling rates of (a c) 0.5 C/s; (d f) 7 C/s; (g i) 20 C/s. temperatures on strength of HAZs for the two typical steels were investigated, as shown as in Fig. 9. It can be seen that the strength decreases with the decrease of peak temperature, and the lowest strength appears when the peak temperature decreases to C. With further decrease of the peak temperature, the strength begins to increases slightly. While, as the peak temperature is below C, the strength decreases slightly. Of course, the heat input and the chemical composition have great effects on mechanical properties of the HAZs. For the steels with different chemical compositions, as the heat input is 2.5 kj/mm, the strength of Steel A is lower than Steel C, and the lowest value of yield strength of Steel A is even lower than the standard value of 555 MPa when the peak temperature is within the range from 800 to 1050 C (Fig. 9a). In addition, with the increase of the heat input, the strength further decreases remarkably (Fig. 9b), while, the strength of Steel C meet the requirement of the standard. Those results as stated above indicate that the weakness of strength in HAZs would become a key issue for high-nb X80 pipeline steels Effects of peak temperature on microstructure of HAZ Figs. 10 and 11 show the microstructures of the simulated samples of typical Steels A and C under different peak temperatures and heat input conditions. The microstructures of the samples after the weld thermal cycle changed significantly, although the changes of different steels are similar, some difference still can be found. For example, as the peak temperature is 1350 C and the heat input is 2.5 kj/mm, the microstructure of Steel A is mainly composed of bainite and little granular bainite, plenty of film-like, rod-like M/ A and little dot-like constituents distributing in the matrix, and prior-austenite grain boundary is clearly visible (Fig. 10a). Meanwhile, with the increase of the heat input, the microstructure changes to mainly composing of granular bainite and little bainite, and microstructure coarsens (Fig. 10b). While, the microstructures of Steel C are mainly composed of bainite (Fig. 11a and b), the microstructure mainly consists of parallel lathy ferrite, a few thin film-like M/A constituents exist between the parallel ferrite laths (Fig. 12a and b), meanwhile, some nonparallel ferrite laths exist, and some dot-like M/A constituents distribute at the nonequiaxed ferrite and prior austenite grain boundaries Fig. 12c. With the increase in the heat input, the microstructure coarsens (Fig. 11b). Those results are similar to the CCT diagrams, and indicate that the weld CCT diagrams could describe the microstructural changes clearly. With the decrease of peak temperature to 1150 C, the microstructures are refined (Figs. 10c and d and 11c and d), the amount and size of the M/A constituents are decreased, and the prior-austenite grain boundaries becomes weak. However, as the peak temperature further decreases to C, the microstructures change greatly, which appear to consist predominantly of QF ferrite, and the size of the QF grain and the amount of M/A constituents is decreased clearly (Figs. 10e and f and 11e and f). TEM micrographs show that the QF ferrite

7 894 X.-w. Chen et al. / Materials and Design 53 (2014) Discussion Fig. 6. Welding CCT diagrams of CGHAZ of the steels (a) Steel A; (b) Steel B; (c) Steel C. grain is very fine, and some dot-like M/A constituents distribute in the ferrite grain and at ferrite grain boundaries (Fig. 12d f). As the peak temperature decreases to 850 C, great variation in the microstructure was observed. The ferrite grain displayed a banded structure, which showed the as-rolled microstructure feature, and the size of the ferrite grain and the amount of M/A constituents are increased, meanwhile, the M/A constituents appear to distribute along the prior-austenite grain boundaries (Figs. 10g and h and 11g and h). TEM micrographs show that the microstructure of the as-rolled ferrite matrix has little change after the welding heat cycles (Fig. 12g i). While, as the peak temperature further decreases below 700 C, the microstructures have no obvious change anymore. For the high-nb pipeline steel, the strongest attractions of the steels widely applied are the lowest cost of the steel plates or coils, because the austenite recrystallization temperature increases greatly with the increase of dissolved Nb in austenite, the steel can be produced by high temperature process (HTP), moreover, the dissolved Nb in austenite can restrain the ferrite transformation, and promote acicular ferritic (or low-carbon bainite) microstructure transformation, thereby, the high-strength pipeline steels without Mo can be produced [2,3]. Many research results show that the economical high-nb steels without Mo can be produced using grade X80 pipeline steel plates of thickness over 22 mm by the optimized thermal mechanical controlled processes, the strength and low-temperature impact toughness are over 555 MPa and 300 J [4]. However, for the welded pipe, the microstructure and mechanical properties of HAZ would be deteriorated during welding processes, because the refined microstructures produced by optimized thermal mechanical controlled processes are damaged, furthermore, the deterioration of microstructure and mechanical properties of HAZ may be more serious with increase of the heat input [17]. While, for the same size pipe, the welding process parameters are similar at the same heat input condition, therefore, the chemical compositions of the pipeline steels become one of the most important roles to affect the mechanical properties of HAZ. From the results of effects of cooling rates on the impact toughness of CGHAZ as stated above, the three steels appear good weldability, the impact energy of CGHAZ is over 100 J when the cooling rate is over 7 C/s (Fig. 7b), which indicates that the critical heat input is about 4.0 kj/mm (Fig. 6), which is agreed with the results of the welding heat simulation (Fig. 8). While, some difference appears in different alloy system steels, the mechanical properties of HAZ of the Steel A with the Mn Cr Nb system (1.81%Mn 0.314%Cr 0.91%Nb) are lower than the Steel B with the Mn Mo Nb system (1.80%Mn 0.26%Mo 0.70%Nb) and Steel C with the Mn Mo Cr Nb system (1.83%Mn 0.073%Mo 0.269%Cr 0.57%Nb), and the mechanical properties of HAZ of the Steel B are lower than the Steel C (Figs. 7 and 9). Especially, the strength of HAZ of Steel A is notably lower than Steel B and Steel C, and the strength even lower than the standard value when the cooling rate is below 5 C/s (Fig. 7). Those results indicate that the Mn Mo Nb system and the Mn Cr Mo Nb system steels may be more suited to produce X80 grade pipeline steel than Mn Cr Nb system steels. While, comparing the chemical compositions of the tested steels (Table 1), some alloy elements of Ni, Cu, Nb and so on were added in the steels, the alloy elements also have certain effect on the microstructure and mechanical properties of HAZ, which results that the effects are complex. Therefore, the parameters of the carbon equivalent, such as Ceq and Pcm, are usually used to evaluate the weldability of the pipeline steel. Many research results show that the weldability of the low carbon pipeline steels decrease with the increase of the Ceq and Pcm, which results toughness decrease and strength increase of the HAZ [18]. Hence, the critical values of the Ceq and Pcm are restricted below 0.43 and 0.23 in standard, respectively [19]. However, from the results as stated above, the mechanical properties of HAZ seem not meeting the laws. Comparing Steel A and Steel C, the toughness and strength of HAZ seem to increase with the increase of Ceq and Pcm (Table 1, Figs. 7 and 8). The toughness and strength of HAZ depend on the prior austenite grain size and the transformed microstructure, especially, the microstructure. Many previous research results indicate that the optimal microstructure with the high toughness is relation to the phase transformation temperature [20]. Therefore, suitable addition of Mo, Cr in the low-c high-nb steels can fit to phase transfor-

8 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 7. Effects of cooling rates on the mechanical properties of the tested steels (a) strength; (b) impact toughness. Fig. 8. Effects of peak temperatures on impact toughness of HAZs of the steels at heat inputs of (a) 2.5 kj/mm and (b) 4.0 kj/mm. Fig. 9. Effects of peak temperatures on strength of HAZs of the steels at heat inputs of (a) 2.5 kj/mm and (b) 4.0 kj/mm. mation temperature (Fig. 6), and obtain the optimal microstructure, furthermore, improve the toughness of HAZ. On the other hand, the Ceq and Pcm do not consider the effects of Nb, while, for high-nb steel, the amount of Nb in solution or precipitate increases, therefore, the effects of Nb on microstructure and toughness of HAZ must be considered, which would be discussed in detail below. From the results of the cooling rate on mechanical properties of HAZ, as the cooling rate is lower than 5 C/s, the strength and impact energy have a little change with the increase of the cooling rate. However, when the cooling rate is over 7 C/s, the strength and impact energy increase remarkably with the increase of critical cooling rate, and the critical cooling rate of about 20 C/s appears, with the further increase of the cooling rate, the impact energy appears a decreasing trend. The prior austenite grain size and its transformed microstructure are considered as two of the most important factors affecting the impact toughness of HAZ [21,22]. For the steels in this work, the average size of the prior austenite grain is lower than 63 lm for Steel A, 64 lm Steel B and 72 lm for Steel C, respectively, when the cooling rate is higher than

9 896 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 10. Effect of heat input and peak temperature on microstructure of Steel A, (a) 2.5 kj/mm, 1350 C; (b) 4.0 kj/mm, 1350 C; (c) 2.5 kj/mm, 1150 C; (d) 4.0 kj/mm, 1150 C; (e) 2.5 kj/mm, 950 C; (f) 4.0 kj/mm, 950 C; (g) 2.5 kj/mm, 750 C; (h) 4.0 kj/mm, 750 C. 5 C/s. The prior austenite grain in HAZ has lower coarse tendency, because the (TiNb)(CN) precipitates are not dissolved completely even at the thermal cycle [11], and the undissolved (TiNb)(CN) particles in Nb steels can effectively restrain the austenite grain growth [11,23]. While, from Fig. 7b, the Steel C has higher impact toughness, therefore, the transformed microstructure becomes a key factor affecting the impact toughness of HAZ, which strongly depends on the chemical compositions of the steels, because all alloying elements added in steels would affect the phase transformation kinetics, sequentially, have a great effect on the microstructure and the mechanical properties of the HAZ. From the Figs. 1 5, the microstructures of all steels in the range of the tested cooling rates are mainly medium-temperature transformed products, which consist of different forms of ferrite and M/ A constituents. Many research results show that the size of ferrite, bainite colony and M/A constituent have great effect on the mechanical properties of the HAZ [24]. From the results as stated above, as the cooling rate is lower, the microstructures mainly consist of large size QF, PF, GF and block M/A constituent (Fig. 1a 3a), which result the lower strength and impact toughness (Fig. 7); With the increase of the cooling rate, the microstructures change to consist of mixture of GF and BF, meanwhile, the amount and

10 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 11. Effect of heat input and peak temperature on microstructure of Steel C, (a) 2.5 kj/mm, 1350 C; (b) 4.0 kj/mm, 1350 C; (c) 2.5 kj/mm, 1150 C; (d) 4.0 kj/mm, 1150 C; (e) 2.5 kj/mm, 950 C; (f) 4.0 kj/mm, 950 C; (g) 2.5 kj/mm, 750 C; (h) 4.0 kj/mm, 750 C. the size of M/A constituents decrease and its form change to dotlike and film-like forms, the microstructures are refined (Fig. 1c 3c), as a result, the strength and impact toughness of the HAZ are improved remarkably (Fig. 7); With the further increase of the cooling rate, the microstructures mainly consist of BF and little GB(Figs. 1d 3d), the refined lathy ferrite and film-like M/A constituents are obtained, and the amount and the size of M/A constituents decrease, so the strength further increases (Fig. 7a), While, the size of bainite colony with the same orientation lathy ferrite increases. The bainite colony with low angle grain boundary lathy ferrite has a little inhibition crack propagation [25], therefore, the impact toughness decreases (Fig. 7b). However, comparing the three steels, at the same cooling rate, the microstructures appear some differences, which are attributed to the different effects of the chemical compositions of the steels on phase transformations. From Fig. 6, the Steel A appears higher phase transformation temperature, while the Steel C has lower phase transformation temperature. For alloy additions, it is well known that Mn, Mo, Cr, Ni, Cu, Nb, and Ti can promote the formation of inter medium transformed microstructure and restrain the formation of polygonal ferrite and/or pearlite [26]. It is suggested that there exists a Mn concentration spike in austenite in front of the a/c interface when the polygonal ferrite grows into austenite [27]. The presence of this spike demonstrates that the polygonal

11 898 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 12. TEM micrographs of Steel C after heat input 2.5 kj/mm and peak temperature at (a c) 1350 C; (d f) 950 C; (g i) 750 C. ferrite is controlled by the diffusion of Mn in austenite. This is a likely reason why the increase of Mn content prolongs the incubation time for polygonal ferrite formation. It can be well understood that the steels have lower critical cooling rate occurred PF transformation (Fig. 6). While, comparing the weld CCT diagrams (Fig. 6), although the amount of Mn is similar, the critical cooling rate and phase transformation temperature are different for the steels, it is attributed to additions of Cr, Mo, Cu, Ni, Nb. The Cr, Mo additions apparently suppress or delay the formation of QF, PF and pearlite, and promotes the formation of bainitic ferrite [13,28], and increase the transformed temperature of QF and pearlite, furthermore, decrease the transformed temperature of BF. While, from the results, the effectiveness of Mo may be stronger than Cr because that the transformed temperatures of Steel A are higher than Steel B (Fig. 6). Moreover, the Cr, Mo complex additions can enhance the effectiveness to suppress the PF transformation, and decrease the transformed temperature. Therefore, the transformed products are refined, and the mechanical properties are improved. In addition, the Nb addition is considered to suppress the phase transformation, however, the effect of Nb on suppressing the transformed temperature depends on the amount of Nb in solution, and any precipitation of Nb in NbC will accelerate the transformation and increase the transformed temperature [3,5]. Fig. 13 shows the effects of the cooling rates on the amount of Nb in solution of Steel A and Steel C. As the cooling rate is higher than 2 C/s, the amount of dissolved Nb changes little with the variation of cooling rate. While, as the cooling rate is below 2 C/s, with the decrease of the cooling rate, the amount of dissolved Nb decreases notably. In addition, the amount of dissolved Nb of Steel A is higher than Steel C, however, the amount of dissolved Nb is only about 0.032% and 0.028% for Steel A and Steel C, respectively, which is lower than the amount of Nb added in steels. The results indicate that the Nb do not dissolve largely during the welding circles. Moreover, the dissolved Nb do not precipitate at high cooling rate, while, with the decrease of the cooling rate, the dissolved Nb precipitates during the cooling process. Our previous work shows that each addition of 0.01% Nb in the solution will lead to a decrease of phase transformation temperature by about 11 C; while each addition of 0.01 wt.% undissolved Nb would lead to an increase of phase transformation temperature by about 23 C. Thus, it can be well understood that although there is a higher Nb addition in Steel A (0.91%Nb) than Steel C (0.57%Nb), and the phase transformation temperatures of Steel A are higher than Steel C (Fig. 6). On the other hand, the peak temperatures and weld heat inputs have great effects on the microstructure of HAZ, furthermore, effect on the mechanical properties [29,30]. From the results as stated above, the variation of mechanical properties with the peak temperature shows that the lowest strength occurs at the peak temperature of C, and the lowest impact toughness occurs at the peak temperature of 750 C and 1350 C. The results indicate

12 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 13. Effect of the cooling rate on the amount of Nb in solution. that the peak temperature affects the austenite state during the heat process, and then affects the phase transformation during the cooling process, furthermore, affects the room temperature microstructures. According to the chemical compositions of tested steels, the austenite transformation temperatures (A c1 and A c3 ) are calculated [31], they are about 710 and 870 C, respectively. For this reason, as the peak temperature is below 700 C, the acicular ferrite (AF) mixed QF with GB(or BF) microstructure of base metal in the HAZ is only tempered during the weld thermal circle. Therefore, the microstructures have no great change, and the strength and the toughness are at the same level, which is similar to the result in reference [32], that shows the acicular ferrite microstructure has higher temper stability, the strength and toughness can be improved when the tempering temperature is even higher C [32]. As the peak temperature increases to 750 C, which is just over A c1, the prior austenite grain boundaries in the base metal become preferential nucleation sites, as a result, some austenite forms along the prior austenite grain boundaries, and the austenite grains transform to M/A island constituents during the quickly cooling process (Figs. 10g and h and 11g and h), which can be illustrated clearly in Fig. 14. Fig 14 shows that the effects of the peak temperatures on M/A island constituents of Steel C. It can be seen that the M/A island constituents distribute aggregately along the grain boundaries as-rolled microstructure when the peak temperature is at 750 C(Fig. 14e and f). The M/A island constituents are considered as the hard and brittle phase [33], which can increase the strength, so the strength increases and the impact energy decreases (Fig. 9). Moreover, the block brittle M/A constituent distributed on the grain boundaries aggregately, are usually regarded as the site of the crack initiation, and thus the intersections of the austenite grain boundaries with M/A particles provide a good site for fracture initiation. The cleavage cracks propagate from one cleavage facet to another [33]. This is why the impact energy decreases sharply (Fig. 8). While, as the peak temperature increases, the amount of austenite formed during the heat process increases, the amounts of C and alloy elements concentrated in austenite decrease, the austenite will transform to fine BF and/or GB even QF. As a result, the toughness improves, and the strength decreases. In addition, the untransformed AF structure is tempered strongly, which results in further decrease of the strength (Figs. 8 and 9). As the peak temperature increases to 900 C, the austenitization nearly completes. Yet the carbon nitride cannot dissolve, austenite grain size is very small and the alloy elements distribution is heterogeneous, so the phase transformation temperature is higher. The microstructure is mainly QF consisting of refined ferrite and M/A constituents (Figs. 10e and f and 11e and f), which can be confirmed clearly by TEM observation (Fig. 12). The M/A constituents distribute uniformly in matrix (Fig. 14c and d), so the toughness is improved and the lowest strength is obtained (Figs. 8 and 9). While, with the increase of the peak temperature, the austenite homogeneity increases, some bainite transforms during the cooling process, and the microstructure changes to a mixture of QF with BF, therefore, the strength and toughness are improved (Figs. 8 and 9). As the peak temperature increases to over 1150 C, the carbonitride begins to dissolve quickly [34], and the austenite grain grows quickly, the sub-haz is called CGHAZ. The microstructure also changes to bainite (Figs. 10a d and 11a d). While, with the increase of the heat peak temperature, the size of the prior-austenite grain and the bainite coarsen, so the strength increases and the toughness decrease (Figs. 8 and 9). In addition, comparing the results of the mechanical properties of the weld CCT diagrams (Fig. 7) with the welding thermal circles (Figs. 8 and 9), although the cooling processes are different, and which results that certain difference exists in the values of mechanical properties, the change rule of mechanical properties with the chemical compositions of steels display strong comparability. The results indicate that the welding CCT diagrams can be used to estimate the microstructure and mechanical properties of HAZ. From the results as stated above, the Steel B with the Mn Mo Nb system and Steel C with the Mn Cr Mo Nb system have higher toughness and strength in CGHAZ; While, the Steel A with Mn Cr Nb system appears lower strength in CGHAZ, because it has higher phase transformation temperature. Not only the additions of Cr, Mo and Cr Mo play important role, but also the higher Nb content in the steel plays an important role. The undissolved NbC precipitates suppress effectively the growth of the austenite grain, refine the prior austenite grain, and decrease the deterioration of the impact toughness of HAZ. While, the amount of Nb in solution decreases, and the effectiveness of Nb in decreasing the phase transformation temperature weakens, conversely, the undissolved NbC precipitates would lead to an increase of the phase transformation temperature, which results the coarsening of the final microstructure of HAZ, and decrease of the mechanical properties. From this point of view, with decrease of the peak temperature, the amount of dissolved Nb decreases during the heat process, meanwhile, the NbC precipitation growth tendency increases [9,34], which would accelerate the PF and QF transformation, furthermore, further decrease the strength. Therefore, the lowest strength appears in the sub-haz of peak temperature at the range from 900 to 1000 C, moreover, with the increase of the heat input, the strength further decreases, and the peak temperature range for the strength below the standard value increases (Fig. 9). From all results as stated above, although high-nb steels have good weldability, the HAZs have higher impact toughness, on the contrary, the strength of HAZ of the high Nb steel only with Cr degenerates observably. Nevertheless, the addition of Mo and/or Cr Mo, especially complex addition of Cr Mo can not only improve the toughness of HAZ, but also improve the strength. Therefore, for the high-nb X80 pipeline steels, suitable addition of a little Mo can preferably improve the mechanical properties of welded steel pipes. 5. Conclusions Three typical X80 steels, i.e. Mn Cr Nb, Mn Mo Nb and/or Mn Cr Mo Nb, were selected, and the effects of main alloy elements of Mo, Cr and Nb on the microstructure and mechanical properties of HAZ were investigated. The following cases can be concluded as the main findings of this study:

13 900 X.-w. Chen et al. / Materials and Design 53 (2014) Fig. 14. Distribution M/A constituents in sample of Steel C after different heat input and peak temperature, (a) 2.5 kj/mm, 1350 C; (b) 4.0 kj/mm, 1350 C; (c) 2.5 kj/mm, 950 C; (d) 4.0 kj/mm, 950 C; (e) 2.5 kj/mm, 750 C; (f) 4.0 kj/mm, 750 C; (g) 2.5 kj/mm, 650 C; (h) 4.0 kj/mm, 650 C. Although the cooling processes of the welding CCT diagrams and the welding heat cycles exist certain differences, and the difference certainly exists in values of the mechanical properties. While, the change rules of the mechanical properties with different chemical compositions of steels and cooling rate display strong comparability, therefore, the welding CCT diagrams can be used to estimate the microstructure and mechanical properties of HAZ. The three alloy system steels in this work display good weldability, the brittlement zones appear in the sub-hazs of the peak temperature at about 750 C and 1350 C respectively, while, as the heat input is 4.0 kj/mm and/or the cooling rate is over 5 C/ s, the Charpy impact energy of CGHAZ is over 100 J. The reason of brittlement of CGHAZ is attributed to the coarse prior austenite grain and its transformed bainite, and the reason of brittlement of the sub-hazs of the peak temperature at about 750 C is attributed to the large size and amount of M/A constituents along the prior austenite grain boundary. On the contrary, the strength weakening zones appear in the sub-hazs of the peak temperature from 900 to 1000 C, and with the increase of the heat input or decrease of the cooling rate. The strength decreases to a value even below the standard value of 555 MPa.

14 X.-w. Chen et al. / Materials and Design 53 (2014) The addition of Mo in high-nb X80 grade pipeline steels is more effective to improve the toughness and strength of HAZ than the addition of Cr, while, the complex additions of Cr and Mo strongly enhance the effectiveness to increase the toughness and strength of HAZ. For the high-nb steels, the Nb in solution can suppress the transformation of PF and QF and decrease the phase transformation temperature, meanwhile, the Nb in precipitates can restrain the growth of the prior austenite grain, improve the toughness of CGHAZ. Nevertheless, the undissolved Nb precipitates would lead to an increase of phase transformation temperature, and coarsen the transformed microstructure, furthermore, decrease the toughness and strength. Therefore, for the high-nb steels, it is necessary to further consider the effectiveness of Nb, and add suitably Cr and Mo in steel to improve the toughness and strength of HAZ, especially, increase the strength. 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