Transformation Characteristics of Ferrite/ Carbide Aggregate in Continuously Cooled, Low Carbon-Manganese Steels S. F. Di Martino & G. Thewlis Metallurgical and Materials Transactions A ISSN 173-5623 Metall and Mat Trans A DOI 1.17/s11661-13-235-x 1 23
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Transformation Characteristics of Ferrite/Carbide Aggregate in Continuously Cooled, Low Carbon-Manganese Steels S.F. DI MARTINO and G. THEWLIS Author's personal copy Transformation characteristics and morphological features of ferrite/carbide aggregate (FCA) in low carbon-manganese steels have been investigated. Work shows that FCA has neither the lamellae structure of pearlite nor the lath structure of bainite and martensite. It consists of a fine dispersion of cementite particles in a smooth ferrite matrix. Carbide morphologies range from arrays of globular particles or short fibers to extended, branched, and densely interconnected fibers. Work demonstrates that FCA forms over similar cooling rate ranges to Widmansta tten ferrite. Rapid transformation of both phases occurs at temperatures between 798 K and 973 K (525 C and 7 C). FCA reaction is not simultaneous with Widmansta tten ferrite but occurs at temperatures intermediate between Widmanstätten ferrite and bainite. Austenite carbon content calculations verify that cementite precipitation is thermodynamically possible at FCA reaction temperatures without bainite formation. The pattern of precipitation is confirmed to be discontinuous. CCT diagrams have been constructed that incorporate FCA. At low steel manganese content, Widmansta tten ferrite and bainite bay sizes are significantly reduced so that large amounts of FCA are formed over a wide range of cooling rates. DOI: 1.17/s11661-13-235-x Ó The Minerals, Metals & Materials Society and ASM International 213 I. INTRODUCTION PROCESS heat treatments for carbon and low alloy steels generally involve continuous cooling. Therefore, phases do not always form sequentially but rather can evolve simultaneously and under the influence of local temperature fluctuations. The nature of the phases formed can be difficult to identify. Microstructure classification schemes involving optical microscopy have been developed historically to facilitate this activity. Dubé et al. [1] and Aaronson [2] first characterized ferrite morphologies in isothermally transformed steels, but recognized that continuous cooling could render distinguishing morphological features indistinct. Allotriomorphic ferrite and various sideplate morphologies (often classed as bainite) were readily identified, but Widmansta tten ferrite was difficult to place and was regarded as a generically similar structure to bainite. A comprehensive classification scheme for weld metals (International Institute of Welding (IIW)) was devised in the 198s. [3] This scheme was applicable to weld metals where intragranular transformations dominate, but it could also be used for steels where austenite grain boundary transformations dominate. Ferrite sideplate and acicular ferrite were readily identified using the scheme, but associated principal structures e.g., Widmansta tten ferrite or bainite were difficult to define in S.F. DI MARTINO, Research Associate, is with the School of Materials, Loughborough University, Loughborough LE113TU, U.K. Contact e-mail: S.F.Di-Martino@lboro.ac.uk G. THEWLIS, formerly Principal Metallurgist, with the Steel Metallurgy Department, CORUS, Swinden Technology Centre, Moorgate, Rotherham S6 3AR, U.K., is now retired. Manuscript submitted May 7, 213. weld metals. The same was true of ferrite/carbide aggregate, where pearlite, bainite, or tempered martensite were not easy to recognize. Anelli and Di Nunzio [4] provided guidance on identifying transformation products associated with sideplate structures in steels with some success, but stereological effects i.e., the way constituents are orientated in space were not treated in depth. Thewlis [5] devised a comprehensive scheme (MiClass) for classifying and quantifying microstructure constituents and the associated principal structures in carbon and low alloy steels with the aim of aiding investigation of microstructure evolution during continuous cooling of industrial products and providing calibration data for theoretical models. Extensive evaluation exercises were carried out, which showed that a reasonable degree of consistency could be obtained between operators in identifying the principal structures primary ferrite, pearlite, and martensite, and also Widmansta tten ferrite and bainite constituting ferrite sideplate. The greatest discrepancy between operators occurred when attempting to determine the nature of ferrite/carbide aggregate. A major source of confusion was the presence in some steels of large grains of a dark etching phase containing widely dispersed and sometimes interconnected carbides that were difficult to resolve under the light microscope (see Figure 1). The phase (hereafter referred to as FCA) did not appear to fit any known stereotype and was prevalent in low carbon-manganese, continuously cooled steels forming at temperatures intermediate between the reconstructive (diffusion controlled with slow rates of reaction) and displacive (shear dominated with rapid rates of reaction) transformation regimes. Lemaire et al. [6] identified novel ferrite/carbide aggregate (FCA) in continuously cooled, low carbon-manganese steels that incorporated fibrous cementite precipitation. The
formation of carbides was thought to involve discontinuous inter-phase precipitation at temperatures intermediate between pearlite and bainite formation. Keehan et al. [7] and Pak et al. [8] reported a ferrite/carbide constituent that develops to large grain sizes in high strength, highly alloyed weld metals when the bainite start temperature is close to that of martensite. The constituent was recognized as coalesced bainite and was found to be detrimental to mechanical properties. Although much progress has been made with regard to classifying and quantifying steel microstructures, unusual FCA phases such as the FCA constituent [5] remain problematic. The accuracy of classification schemes and theoretical models may be significantly improved by a greater understanding of this phase. The object of the present work has therefore been to investigate the morphological features and transformation behavior of FCA utilizing the low carbon-manganese steels employed in evaluation exercises with the MiClass classification scheme. [5] The aim was to identify the temperature range of FCA formation relative to the known principal structures and locate the position of FCA in continuously cooled transformation (CCT) diagrams. II. EXPERIMENTAL PROCEDURE A. Materials The chemical compositions of steels investigated in the current work are given in Table I. Steels A and C Fig. 1 Ferrite/carbide aggregate (FCA) in.13 wt pct C, 1.2 wt pct Mn steel thermally cycled to a temperature of 12 C in a dilatometer and cooled at 1 C s 1 [173 K to 773 K (8 C to 5 C)] (after Ref. [5]). were 5 kg laboratory, vacuum furnace casts of base compositions.17 wt pct C, but with.52 wt pct Mn and 1.46 wt pct Mn, respectively. They were hot-rolled to 15-mm thick plates on a laboratory reversing plate mill. Steel B was a commercial Grade D ship plate of base composition.13 wt pct C with intermediate manganese content (.87 wt pct Mn). The steel was rolled to 2-mm thickness plate. Microstructures of steels A and C were representative of the commercial production material. B. Dilatometry Small cylindrical specimens 1 mm in length and 3 mm in diameter were machined from all plates in the longitudinal direction. The solid dilatometer samples were used to obtain both basic transformation data but also for interrupted quench experiments. Flats were polished on the surface of some dilatometer specimens to facilitate etching of austenite grain boundaries at the thermal cycle peak temperature. The prior austenite grain size (defined by thermal etched grooves) was measured at ambient temperature using a mean linear intercept method. Approximately, 1 intercepts were measured at a magnification of 95. The solid dilatometer specimens were subjected to controlled thermal cycles under vacuum in a Ba hr Type 85 high-speed quenching and deformation dilatometer. Dilatometer thermal cycles were chosen to investigate mutual interactions between the phase transformations taking place during cooling through the reconstructive and displacive transformation regimes. A wide range of microstructures was achieved by varying cooling rate but also by choosing specific peak temperature, which controlled the austenite grain size. Specimens from the low and high manganese steels A and C were heated to a peak temperature of 1373 K (11 C) at a rate of 5 C s 1, held for 2 minutes and then cooled at rates of typically 2, 5, 1, 25, 5, 75, 1, 15, and 2 C s 1 between 173 K and 773 K (8 C and 5 C). Thermal cycles were repeated for the intermediate manganese steel B using peak temperatures of 1373 K and 1573 K (11 C and 13 C) so that transformation from relatively coarse and fine austenite grain sizes could be investigated. The start transformation temperature T S, and finish transformation temperature T F for each steel thermal cycle were determined from the relative dilation curve. The 5 pct transformation temperature T 5pct was identified from computer generated graphs of percentage transformed against temperature. Dilation curves in each case were numerically differentiated to obtain information about Table I. Steel Compositions Steel Element (Wt Pct) Steel Code C Si Mn P S N Nb Ti V A.17.29.52 <.5 <.2.32 <.5 <.1 <.5 B.13.28.86 <.5 <.2.7 <.5 <.5 <.5 C.17.31 1.46 <.5 <.2.43 <.5 <.1 <.5
the rate of transformation at different temperatures. The temperature at which the peak rate of transformation T PRTT occurred for each steel thermal cycle was identified. The Ba hr dilatometer was capable of high cooling rates throughout the length and width of a 3-mm solid sample, thereby minimizing errors in thermal gradients and ensuring reliability of data. Errors in determining the T S and T 5pct temperatures were estimated at ~±5 K (±5 C). Errors in establishing T F may have been slightly greater. Errors owing to temperature gradients across the specimens may be expected to be typically ±5 K(±5 C) at the cooling rates used. [9] For interrupted quench experiments, work was focused on the thermally cycled, intermediate manganese steel B. Dilatometer samples were heated to peak temperature and cooled, but quenched with helium once a specific temperature had been reached. Thermal cycle peak temperatures of 1373 K and 1573 K (11 C and 13 C) were again employed but with cooling rates restricted to 1 and 5 C s 1 between 173 K and 773 K (8 C and 5 C). It was not practical to interrupt the cooling at faster rates. Quench temperatures were chosen such as to define the start and finish of the FCA phase transformation and the transformation sequence. The peak rate transformation temperature information obtained from the fully cycled specimens was used for guidance. C. Metallography Optical metallographic examination of microstructure features in transverse sections of dilatometer specimens was carried out after polishing and etching in 2 pct nital using a Reichert MEF-3 microscope. Phases were quantified according to the recently developed MiClass scheme. [5] Photographs were taken of microstructures at an appropriate magnification (59 or 19) andmon- tages were produced. Mesh grids inscribed on transparent acetate paper were overlaid on the microstructures so that constituents under or just touching grid cross-lines could be identified. FCA and principal structures primary ferrite (PF), pearlite (P), Widmanstätten ferrite (Wf), bainite (B), (with the distinction in upper bainite (UB) and lower bainite (LB) when relevant), and martensite (M) were quantified. Microphases associated with Widmansta tten ferrite were treated separately, while bainitic ferrite was quantified together with the carbide. A total of 5 points were counted of each microstructural condition. A LEO 153VP field emission gun scanning electron microscope (FEGSEM) with In-Lens detection was used to enable microstructure features to be viewed at high resolution (within the nanometer range) with good depth of field. The imaging conditions used were an accelerating voltage of 1 kv and a working distance of 3 mm. Images were linked to optical observations. Transmission electron microscopy (TEM) was performed using a JEOL 2 FX operating at 2 kv. Morphological characteristics of carbides in FCA and associated crystallographic information were obtained by means of carbon replicas. An EDAX Pegasus, combined energy dispersive X-ray analysis (EDX), and electron back scatter diffraction (EBSD) system attached to the LEO 153VP FEGSEM was used to investigate the orientation of FCA grains in the thermally cycled, intermediate manganese steel B, and identify low angle boundaries that could provide information as to the FCA transformation mechanism. The nominal step size was.1 lm. A low-angle boundary was defined as below a threshold limit of 2 deg. EBSD analysis involved the determination of crystal orientation maps of the sample. The crystallographic direction corresponding to the particular sample direction was calculated and grain orientations characterized by different colors. EBSD scanned regions were subsequently imaged by means of SEM (In-Lens detection) and optical microscopy in order to facilitate the identification of microstructural features of interest on the EBSD crystal orientation maps. D. Microhardness Microhardness measurements of FCA grains and principal structures in the thermally cycled, intermediate manganese steel B were made to help assess the FCA transformation mechanism. A Reichert microhardness tester was employed. A load of 2 g was applied for 6 seconds on each of the tested specimens. The microhardness numbers were then calculated from the following equation: HV ¼ 1854:4 P d 2 ½kg=mm2 Š; where HV is the microhardness number, P is the load in g, and d is the diagonal of the indentation in micrometers (i.e., the ocular reading multiplied by.157 lm). III. RESULTS A. FCA Characterization 1. Grains Micrographs of typical FCA grains in thermally cycled steels A, B, and C obtained using optical and high magnification FEGSEM are shown in Figure 2. Under the light microscope, the FCA appeared as a dark etching constituent usually adjacent to primary ferrite grains or Widmansta tten ferrite plates. It had a spongy appearance and consisted of a ferritic matrix interspersed with barely resolvable carbides. The FCA was comparable to that observed optically in.13 wt pct C, 1.2 wt pct Mn steel of Reference 5 heated to a peak temperature of 1473 K (12 C) and cooled at 1 C s 1 [173 K to 773 K (8 C to 5 C)] (see Figure 1). Viewed using a FEGSEM, FCA was revealed as a fine dispersion of carbides in a generally smooth ferrite matrix. There was no evidence of surface relief effects associated with displacive transformations. At low cooling rates, where the simultaneous formation of pearlite and FCA was possible, there was a correlation between the dimensions of the prior austenite grains and the nature of the transformation product formed. Measurements showed that pearlite regions were generally smaller than 5 lm, whereas FCA grains were typically greater than 5 lm and could reach dimensions of up to 6 lm.
(a) (b) PF WF FCA FCA 2 µm 2 µm PF Primary Ferrite; FCA Ferrite/Carbide Aggregate; WF Widmanstätten Ferrite Fig. 2 Typical ferrite/carbide aggregate (FCA) in steels A, B, and C thermally cycled to peak temperatures of 1373 K and 1573 K (11 C and 13 C) and continuously cooled (a) optical (b) FEGSEM. 2. Carbide morphology Different types of carbide dispersions in FCA grains of thermally cycled steels A, B, and C observed using FEGSEM are shown in Figure 3. In low manganese,.17 wt pct C steel A, austenitized at 1373 K (11 C) and cooled at slow rates, the FCA grain size was around 15 lm. Arrays of globular-shaped particles (average size 5 nm) were predominant in FCA regions. There were also small as well as more elongated fibers (5 to 2 nm), which were either formed discretely or in arrays. Overall fiber dimensions ranged between 2 nm and 1.5 lm while the average spacing between arrays was 5 nm. Irregular shaped particles were often seen, either formed discretely or as dense interconnected networks. Branched, elongated fibers up to 3 nm in size and multiple branched fibers were occasionally observed. The dimensions of irregular shaped particles and branched fibers varied between 4 nm and 1 lm. In some grains, geometrically shaped precipitate arrays or clusters were evident. A notable feature in some instances was pearlite at the boundaries of FCA grains. The pearlite lamellae growth direction was seemingly not reproduced by carbide precipitation within the FCA grains. In high manganese,.17 wt pct C steel C heat treated in an analogous manner to the low manganese material, the FCA grain dimensions were generally smaller (mainly around 1 lm). Carbides were much finer and the presence of small globular particles was limited. The morphology of the precipitates in the FCA was less regular and there were significantly greater numbers of both discrete and multiple branched fibres with a spacing of 1 to 2 nm. The lamellar pearlite growth encountered at the FCA boundary in the low manganese content steel was more marked at higher manganese contents. In intermediate manganese.13 wt pct C steel B, austenitized at 1373 K (11 C) and cooled at slow rates, FCA grain sizes were between 5 and 4 lm. Arrays of globular particles (4 to 5 nm), and also small (5 to 15 nm) and elongated fibers predominated, while the spacing between precipitate arrays was 1 to 4 nm. Increasing the austenitizing temperature to 1573 K (13 C) and cooling slowly resulted in generally coarser FCA grains (1 to 6 lm) with slightly finer internal carbides. As the cooling rate was increased in all the thermally cycled steels, FCA grains became finer (5 to 3 lm). Globular carbides and small fibers were observed inside the grains, but dense interconnected fibers and partially interconnected, irregularly shaped precipitates became more prevalent. The spacing between precipitate arrays was smaller at the faster cooling rates (5 to 3 nm). Carbide morphologies observed using FEGSEM were confirmed using TEM (see Figure 4). Arrays of globular precipitates and short fibers were evident as well as irregular and elongated branched fibers. Small particles appeared to develop from the boundary of grains and growth directions could be identified. 3. Carbide type TEM in conjunction with energy dispersive EDX of FCA regions in thermally cycled steels using thin foil indicated that the matrix was ferritic. Lattice parameters obtained from diffraction patterns were 2.86 A corresponding to the 2.8665 A of bcc alpha iron (.2 A accuracy) showing that the matrix was not significantly supersaturated with carbon. There was no evidence of laths or sideplates in FCA grains. Electron diffraction analysis of carbide particles on extraction replicas was limited by the interconnected nature of FCA precipitates. However, successful crystallographic analysis (see Figure 5) showed that particles were consistent with iron carbides, with many patterns indexing as cementite. 4. Grain orientation An EBSD misorientation map of FCA grains in thermally cycled, intermediate manganese steel B together with the corresponding FEGSEM image is shown in Figure 6. On the misorientation map, FCA regions (colored in light and dark blue) had uniform orientation throughout. The FCA grains pictured
Fig. 3 FEGSEM micrographs of typical carbides in FCA grains of thermally cycled steels A, B, and C. (a) globular particle arrays, (b) globular and short fiber arrays, (c) irregular shaped particles, (d) dense interconnected fibers, (e) elongated branched fibers, (f) multiple branched fibers, (g) geometric arrays, and (h) grain boundary pearlite. exhibited a slight change in orientation so that a low angle boundary (arrowed) was produced where the grains intersected. The FEGSEM micrograph suggests that FCA carbides nucleated at the adjacent prior austenite grain boundary within or below the image plane and precipitation followed a specific direction within the FCA grains. An EBSD misorientation map of FCA grains adjacent to pearlite in thermally cycled steel B together with the corresponding FEGSEM image is shown in
Fig. 4 TEM micrographs of typical carbides in FCA grains of thermally cycled steels A, B, and C: (a) globular and short fiber arrays, (b) elongated fiber arrays, (c) irregular shaped particles, and (d) elongated branched fibers. 1 nm (a) (b) Steel B Solution found for cementite. Z = [21] Fig. 5 Analysis of carbides in FCA grains: (a) carbide particles and (b) diffraction pattern.
Fig. 6 EBSD analysis of adjacent FCA grains in thermally cycled steel B showing homogenous orientation with precipitation direction and low angle grain boundary (arrowed): (a) misorientation map and (b) FEGSEM image. Fig. 7 EBSD analysis of adjacent FCA and pearlite grains in thermally cycled steel B showing distinct orientation difference: (a) misorientation map, and (b) FEGSEM image. Fig. 8 EBSD analysis of adjacent primary ferrite and pearlite grains in thermally cycled steel B, cooled at 5 C s 1, showing common orientation: (a) misorientation map, and (b) FEGSEM image. Figure 7. Pearlite exhibited a distinct orientation on the misorientation map (green colored) compared to the adjacent FCA (light blue colored), indicative of the fact that the two regions formed at different times and nucleated at different sites. The pearlite grew from an adjacent prior austenite grain boundary initially in a lamellar fashion, but subsequently degenerated into irregular length fibers seemingly associated with the same growth direction. FCA grains showed a clear orientation difference with adjacent ferrite grains. However, EBSD analysis of other regions of thermally cycled steel B containing only primary ferrite and pearlite showed that in most cases the pearlite islands had the same orientation as the adjacent ferrite grain. An example is shown in Figure 8. 5. Microhardness Microhardness data for microstructure constituents in intermediate manganese steel B thermally cycled to a peak temperature of 1573 K (13 C) and cooled at 1 and 5 C s 1 are shown in Figure 9. The hardness of FCA was found to be much less than bainite or martensite and consistently between that of primary ferrite and Widmansta tten ferrite.
6. Dilatometry Fig. 9 Microhardness data for microstructure constituents in steel B thermally cycled to a peak temperature of 1573 K (13 C) and cooled: (a) 1 C s 1 and (b) 5 C s 1 [173 K to 773 K (8 C to 5 C)]. a. Continuous cooling experiments. i..17 wt pct C (.52 wt pct, 1.46 wt pct Mn) steel. Heat treatment, transformation temperature, and microstructure data for.17 wt pct C steels A and C with low and high manganese content, austenitized at 1373 K (11 C) and continuously cooled, are shown in Table II. Relationships between microstructure constituents and cooling rate are shown in Figure 1. A low steel manganese content raised the transformation start temperature and substantially increased the amount of FCA formed while widening the range of cooling rates over which FCA was observed. In low manganese content steel, primary ferrite and pearlite were the main phases formed below 5 C s 1 cooling rate, while between 5 and 15 C s 1, Widmanstätten ferrite and FCA increased and decreased, peaking at cooling rates between 5 and 75 C s 1. The maximum amount of FCA (64 pct) was significantly greater than the highest Widmansta tten ferrite (29 pct). Bainite prevailed at cooling rates greater than 1 C s 1 followed by martensite. In the high manganese content steel, microstructure constituents formed in a similar sequence to the low manganese material, but primary ferrite and pearlite dominated at 2 C s 1 cooling rate, while FCA and Widmansta tten ferrite prevailed between 2 and 75 C s 1 with peaks between 1 and 25 C s 1. Table II. Heat Treatment and Transformation Data for Continuously Cooled.17 Wt Pct C Steels A and C with Low and High Mn Content Steel Code A (.52 Wt Pct Mn) C (1.46 Wt Pct Mn) Peak Temp. [K ( C)] Cooling Rate C s 1 [173 K to 773 K (8 C to 5 C)] Transformation Temp. ( C) Phase Amount (Pct) T S T 5pct T PRTT T F PF P WF MP FCA UB LB M 1373 (11) 2 79 7 75 575 77 23 5 788 68 7 578 62 13 12 3 12 1 783 685 685 577 58 18 4 2 25 764 66 655 564 3 22 6 42 5 731 63 62 539 12 19 5 64 75 713 63 62 54 12 29 8 51 1 676 57 575 468 4 28 7 32 3 15 633 43 39 36 1 3 15 25 47 2 616 425 42 276 33 67 1373 (11) 2 747 63 62 491 8 12 8 5 727 6 625 471 4 8 29 7 16 1 713 6 6 446 2 44 16 2 25 657 515 52 335 7 34 12 23 24 5 612 45 41 284 11 4 6 52 28 75 575 395 4 232 2 23 66 9 1 514 39 395 222 48 52 15 489 36 365 215 14 86 2 484 38 38 196 1 PF: primary ferrite; P: pearlite; WF: Widmanstätten ferrite; MP: microphase; FCA: ferrite/carbide aggregate; UB: upper bainite; LB: lower bainite; M: martensite.
Microstructure Constituent, % 1 9 8 PF 7 FCA 6 B M 5 4 3 P WF 2 1 2 4 6 8 1 12 14 16 18 2 Cooling Rate, K s -1 (8-5 C) (a) Microstructure Constituent, % 1 9 8 PF B M 7 6 WF 5 4 3 FCA 2 1 P 2 4 6 8 1 12 14 16 18 2 Cooling Rate, K s -1 (8-5 C) (b) Fig. 1 Relationship between volume percent microstructure constituents and cooling rate in thermally cycled,.17 wt pct C steels A and C, austenitized 1373 (11 C): (a).52 wt pct Mn and (b) 1.46 wt pct Mn. Table III. Steel Code B (.86 Wt Pct Mn) Heat Treatment and Transformation Data for Continuously Cooled.13 Wt Pct C Steel B with Intermediate Mn Content Peak Temp. [K ( C)] Cooling Rate C s 1 [173 K to 773 K (8 C to 5 C)] Transformation Temp. ( C) Phase Amount (Pct) T S T 5pct T PRTT T F PF P WF MP FCA UB LB M 1373 (11) 2 81 695 75 566 81.5 18.5 5 792 68 685 557 58 9 19 5 8 1 779 67 675 551 25 41 12 22 25 753 635 63 549 14 53 15 18 5 72 615 65 511 3 42 2 21 14 6 715 58 58 493 1 27 9 17 4 6 75 7 6 57 463.5 9 3.5 36 46 5 1 671 56 56 398.5 11.5 2 68 15 612 435 395 326 1 9 2 557 425 41 3 1 1573 (13) 2 761 53 79.6 2.4 5 757 512 52 13 12 4.5 11.5 1 756 63 63 516 16 45 13 26 25 735 52 12 48 18 22 5 694 58 6 48.5 16.5 1 18 51 4 6 679 383.5 15.5 4.5 16 34.5 28 75 655 355 7 3 22 49 21 1 64 321 4 28 68 15 536 395 39 286 9 91 2 487 269 1 In this case, the largest amount of FCA (23 pct) was significantly less than the maximum Widmansta tten ferrite (44 pct). ii..13 Wt pct C,.86 wt pct Mn steel. Heat treatment, transformation temperature, and microstructure data for.13 wt pct C steel B with intermediate manganese content, austenitized at low [1373 K (11 C)] and high [1573 K (13 C)] peak temperatures, and continuously cooled are presented in Table - III. Relationships between microstructure constituents and cooling rate are shown in Figure 11. Low austenitizing temperature raised the transformation start temperature but the volume fraction of FCA was little changed and the range of cooling rates over which FCA occurred was only slightly broadened. Irrespective of austenitizing temperature, the position of FCA in the transformation sequence was contemporary with Widmansta tten ferrite. Maximum volume percents of FCA were 22 and 26 pct at low and high austenitizing temperatures, respectively, which were significantly lower than for Widmansta tten ferrite (53 and 48 pct). iii. Interrupted quench experiments. Heat treatment and microstructure data for quenched,.13 wt pct C steel B with intermediate manganese content are shown in Table IV. Austenitizing steel B at 1373 K (11 C) and cooling at 1 C s 1 down to quench temperatures of 932 K and 91 K (659 C and 637 C) allowed significant proportions of primary ferrite and Widmansta tten to form. Pearlite was not seen as discrete
Microstructure Constituent, % 1 9 8 7 6 5 4 3 2 1 PF M B WF FCA P 2 4 6 8 1 12 14 16 18 2 Microstructure Constituent, % 1 9 M PF 8 B 7 WF 6 5 4 FCA 3 2 1 P 2 4 6 8 1 12 14 16 18 2 Cooling Rate, K s -1 (8-5 C) Cooling Rate, K s-1 (8-5 C) (a) (b) Fig. 11 Relationship between phase amounts and cooling rate in thermally cycled,.13 wt pct C steel B with.86 wt pct Mn, austenitized: (a) 1373 (11 C) and (b) 1573 K (13 C). Table IV. Heat Treatment and Transformation Data for Interrupted-Quenched.13 Wt Pct C Steel B with Intermediate Mn Content Steel Code Peak Temp. [K ( C)] Cooling Rate C s 1 [173 K to 773 K (8 C to 5 C)] Interrupted-Quench Temp. [K ( C)] Phase Amount (Pct) PF WF MP FCA UB LB M B (.86 Wt Pct Mn) 1373 (11) 1 932 (659) 18 1.5 1.5 7 91 (637) 23 22 4 63 888 (615) 28 39 9 24 5 938 (665) 3 26.3 7.7 928 (655) 3 29 68 879 (66) 4 35 1 6 853 (58) 3 46 7.5 1 33.5 1573 (13) 1 897 (624) 7 5 1 87 879 (66) 9 11 4 76 86 (587) 17 41 16 26 5 873 (6) 1 12 4 9 73 853 (58) 1 18 8 71 833 (56) 1 22 2 57 83 (53) 1 18 5 38 39 773 (5) 1 17 9.5 53.5 19 743 (47) 1 21 23 55 grains but was present in microphases associated with Widmansta tten ferrite. Remaining austenite transformed into martensite during the quench. Reducing the quench temperature to 888 K (615 C) resulted in little further growth of primary ferrite and Widmanstätten ferrite but there was significant FCA formation. An absence of martensite indicated that overall transformation was complete prior to the quench. Increasing the austenitizing temperature from 1373 K to 1573 K (11 C to 13 C) and cooling at 1 C s 1 down to quench temperatures of 897 K, 879 K, and 86 K (624 C, 66 C, and 587 C) resulted in a similar pattern of behavior. However, irrespective of austenitizing temperature, when the cooling rate was increased to 5 C s 1, overall transformation was not completed during cooling to the quench temperatures. In this case, the FCA transformation took place just before or after Widmanstätten ferrite growth was complete but it ended before or just after bainite transformation started. IV. DISCUSSION The object of the present work was to investigate the transformation characteristics and morphological features of ferrite/carbide aggregate, designated FCA, in continuously cooled low carbon-manganese steels. Unusual FCA phases have received little attention in the literature and their identification and classification is not straightforward. The results of the current work show that FCA grains consist of a ferrite matrix interspersed with cementite particles whose morphologies are complex and vary significantly with steel composition, heat treatment, and transformation tem-
perature. An important feature to emerge from the current work is the absence of laths or sideplates in FCA grains, indicative that FCA development is not a displacive transformation such as bainite or autotempered martensite. In the following discussion, FCA classification will be considered and attempts made to position FCA in the transformation sequence of principal structures. The thermodynamics and kinetics of the FCA reaction will be examined. Finally, CCT diagrams will be constructed and the influence of steel composition and austenitizing temperature on FCA development assessed. A. FCA Classification The current work confirms that under the light microscope it is difficult to classify FCA. Even using high-magnification FEGSEM or TEM techniques, results show that FCA has neither the classical lamellar structure of pearlite nor the lath structure of bainite and martensite. Carbide morphologies are seen to vary considerably from arrays of globular particles or short fibers to elongated, branched, and densely interconnected fibers. Work shows that FCA forms at much slower cooling rates than bainite or martensite and is a relatively soft phase. It is difficult therefore to interpret FCA as coalesced bainite, which develops in highly alloyed weld metals when the bainite start temperature is close to that of martensite. [7,8] Lemaire et al. [6] identified novel FCA in continuously cooled, low carbon-manganese steels that incorporated fibrous cementite precipitation, which was thought to involve discontinuous inter-phase precipitation. Arrays of carbides in FCA grains are suggestive of an inter-phase precipitation mechanism. However, it is not easy to explain the overall range of FCA carbide growth morphologies in this way or the relatively low hardness of FCA. It is notable from results that edges of FCA grains occasionally exhibit pearlite lamellae. It is possible therefore that FCA is a type of degenerate pearlite. Literature [1 12] shows that pearlite lamella growth can become discontinuous so that carbide precipitation becomes fragmented while still related to the original pearlite colony. Results show that the lamellae growth direction within the FCA grains is apparently not reproduced by carbide precipitation, although this would depend on the plane of observation. Furthermore, particles in FCA regions are considerably smaller and the morphologies more complex than the fibers usually observed in degenerate pearlite. [1] In the present study, FEGSEM, TEM, and EBSD analysis shows that in most cases there is a preferred precipitation direction in FCA grains, and generally an absence of laths. It is possible therefore that FCA may be thought of as a microstructural constituent in its own right forming at reconstructive transformation temperatures. A detailed understanding of the way in which different FCA carbide morphologies are produced is beyond the scope of the present work. Nevertheless, as will be shown later, discontinuous carbide precipitation may prove an important part of FCA development. B. Transformation Process Dynamics 1. Steel cooling rate Continuous cooling experiments carried out on steels in the current work show clear trends. At very slow cooling rates, large amounts of primary ferrite are formed accompanied by relatively small fractions of pearlite. As cooling rate increases, the amounts of primary ferrite and pearlite decrease rapidly, while FCA and Widmansta tten ferrite increase to a central peak or plateaux and then decrease. The decline in FCA and Widmanstätten ferrite is followed by a marked increase and decrease in bainite while martensite forms at very fast cooling rates. It is evident that FCA formation takes place over similar cooling rate ranges to Widmansta tten ferrite. However, low steel manganese content substantially increases the amount of FCA formed compared to Widmanstätten ferrite, while widening the range of cooling rates over which both are observed. It is well recognized that manganese is an austenite stabilizing element so that a decrease should increase the transformation start temperature as observed. In the intermediate manganese content steel, a decrease in austenitizing temperature raises the transformation start temperature commensurate with an increase in the surface area of austenite grain boundary for ferrite nucleation. There is some broadening of the cooling rate range over which FCA and Widmansta tten ferrite form, but the amount of FCA is significantly reduced compared to Widmansta tten ferrite. The effects of manganese and austenitizing temperature on FCA and Widdmanstatten ferrite formation will be discussed in more detail later in relation to CCT diagrams. 2. Steel transformation rate Although cooling rate is important in thermomechanical processing of steels, transformation temperature ultimately governs the type of microstructure constituent formed. In the slower cooled steels of the current work, transformation start temperature, T S, may be considered as sensitive to the formation of primary ferrite, while T 5pct provides information as to other principal structure development. Tables II and III show that maximum volume fractions of FCA correlate with T 5pct temperatures between 788 K and 93 K (515 C and 63 C), while the highest amounts of Widmanstätten ferrite are associated with T 5pct temperatures somewhat higher i.e., between 873 K and 98 K (6 C and 635 C). Important kinetic information can be obtained from dilatometer dilation curves to support classification of microstructures. [5] Numerical differentiation of the percent transformed against temperature curves can provide knowledge as to the overall rate of steel phase transformations. Transformation rate increases to a peak (T PRTT ) as the most prolific phase forms and then decreases, the process occurring over progressively narrower temperature ranges as transformations change from reconstructive to displacive regimes. At high temperatures, the driving force for transformation is low so that shallow peaks are evident. Conversely, at low temperatures where shear reactions prevail, the
Primary Ferrite, % 1 9 8 7 6 5 4 3 2 1 1 2 3 4 5 6 7 8 Pearlite, % 1 9 8 7 6 5 4 3 2 1 1 2 3 4 5 6 7 8 Peak Rate Transformation Temperature, C (a) Peak Rate Transformation Temperature, C (b) 1 1 Widmanstätten Ferrite, % 9 8 7 6 5 4 3 2 1 FCA, % 9 8 7 6 5 4 3 2 1 1 2 3 4 5 6 7 8 Peak Rate Transformation Temperature, C (c) 1 2 3 4 5 6 7 8 Peak Rate Transformation Temperatutre, C (d) 1 1 Bainite, % 9 8 7 6 5 4 3 Martensite, % 9 8 7 6 5 4 3 2 1 2 1 1 2 3 4 5 6 7 8 1 2 3 4 5 6 7 8 Peak Rate Transformation Temperature, C (e) Peak RateTransformation Temperature, C (f) Fig. 12 Relationship between microstructure constituents and peak rate transformation temperatures of thermally cycled steels A, B, and C. driving force for transformation is high and sharp peaks are observed. Relationships between volume percent microstructure constituents and steel peak rate transformation temperatures for the thermally cycled steels in the current work are shown in Figure 12. The highest rates of transformation to primary ferrite and pearlite are seen at temperatures approximately between 873 K and 973 K (6 C and 7 C). Maximum rates of transformation to both FCA and Widmansta tten ferrite are evident at temperatures generally between 798 K and 973 K (525 C and 7 C). Rapid bainite formation takes place at temperatures between 673 K and 873 K (4 C and 6 C), while peak rates of transformation to martensite are observed at temperatures around 673 K (4 C). Although peak rate transformation temperature and to some extent T 5pct temperature correlations with microstructure provide an indication as to when FCA and Widmanstätten ferrite reactions are maximized, it is not possible in continuously cooled steels to isolate transformation start and finish temperature of individual phases. It is necessary to interrupt cooling
at selected temperatures and determine the extent to which individual microstructural constituents form in order to identify a transformation sequence. 3. Transformation sequence The interrupted quench experiments carried out in the current work show quite clear patterns. Irrespective of austenitizing temperature [1373 K and 1573 K (11 C and 13 C)], cooling relatively slowly (1 C s 1 )to different quench temperatures results in FCA forming just before or after Widmansta tten ferrite growth finishes. When the cooling rate is increased sufficiently (5 C s 1 ), the FCA reaction takes place just before or after Widmansta tten ferrite growth is complete but ends before or just after bainite transformation starts. Thus, FCA formation is not simultaneous with Widmanstätten ferrite but occurs intermediate between Widmansta tten ferrite and bainite but slightly overlaps with each. The observed slight overlap may be explained by temperature gradients and local fluctuations in the carbon content of austenite. It is, therefore, reasonable to assume a transformation sequence of phases as follows: PF! P! WF! FCA! B! M: 4. Thermodynamics and kinetics Overall indications in the present study are that FCA is either a phase in its own right or is in some way related to pearlite. FEGSEM, TEM, and EBSD evidence suggests that FCA is a reconstructive transformation (diffusion controlled with slow rates of reaction). Interrupt quench experiments show that the FCA reaction takes place at temperatures after displacive transformation to Widmansta tten ferrite ceases and before displacive bainite transformation begins, although possibly overlapping slightly with each during continuous cooling. Widmansta tten ferrite is reportedly a shear transformation accompanied by the rapid diffusion of carbon atoms across the advancing austenite/ferrite interface, while carbon partitioning in bainite is claimed to occur after the diffusionless transformation of bainitic ferrite [5] Attempts therefore to isolate the thermodynamic and kinetic conditions under which FCA forms need to address parameters related to bainite and martensite transformations, as well as pearlite. Where FCA and pearlite are observed together in the thermally cycled steels of the current work, pearlite regions are generally smaller than 5 lm, while FCA grains can reach dimensions of 6 lm. The possibility exists therefore that when the overall carbon content of austenite grains is insufficient to sustain a lamellar type of pearlite growth, FCA forms by some kind of discontinuous precipitation that is dependent on the local carbon content of austenite. Some understanding of the latter in thermally cycled steels of the present study at different stages of cooling may be obtained from the interrupted quench experiments. Table V shows calculations of carbon in untransformed austenite after austenitizing steel B at 1373 K (11 C) and cooling at 1 C s 1 to quench temperatures of 932 K, 91 K, and 888 K (659 C, 637 C, and 615 C). The upper carbon limit considers that carbon in ferrite is zero, while the lower limit assumes that it is.2 wt pct. In samples quenched from 932 K and 91 K (659 C and 637 C), austenite carbon content is less than about.2 wt pct. Primary ferrite and Widmanstätten ferrite are observed but not FCA. On quenching at 888 K (615 C), austenite carbon content reaches a value of around.5 wt pct and FCA formation is marked. However, this carbon content is significantly below the.8 wt pct eutectoid carbon composition of pearlite. The above calculations assume a uniform carbon distribution in austenite. In practice, a carbon concentration gradient may occur at the transformation interface with the greatest segregation close to the interface. It may thus be thermodynamically possible for pearlite to form initially but then the carbon distribution in austenite is lowered to the point where reaction ceases. The remaining austenite transforms to FCA when the thermodynamic conditions allow. It is necessary to isolate these conditions from those for bainite formation. Figure 13 shows a predicted isopleth diagram for steel B constructed using an appropriate kinetic code, MUGC46 [13] in conjunction with a thermodynamic calculation software package, MTDATA [14] and a Fig. 13 Predicted isopleth diagram (MTData) for.13 wt pct C steel B with.86 wt pct Mn showing regions of possible austenite to ferrite phase transformations. Steel Code B (.86 Wt Pct Mn) Table V. Carbon in Austenite for Interrupted Quenched.13 Wt Pct C Steel B with Intermediate Mn Content Peak Temp. [K ( C)] Cooling Rate C s 1 [173 K to 773 K (8 C to 5 C)] Interrupted Quench Temp. [K ( C)] Carbon in Austenite (Wt Pct) Upper Limit Lower Limit 1373 (11) 1 932 (659).186.177 91 (637).26.192 888 (615).542.486
critically assessed thermodynamic database for steels. The phase diagram incorporates the T o curve for the bainite reaction and the extrapolated c/fe 3 C boundary. T o is considered an important boundary condition for bainite transformation in the literature [15] and is defined as the temperature at which c and a of the same composition have the same free energy. The diffusionless formation of bainitic ferrite at a given temperature is possible when austenite carbon concentrations are to the left of the T o curve while reaction ceases at T o. The intersection of the T o and c/fe 3 C lines in the isopleth diagram in Figure 13 can be seen to dissect it into three thermodynamic regions. In region I, the formation of bainitic ferrite is possible but not the precipitation of cementite. Region II allows supersaturated bainitic ferrite to form accompanied by precipitation of carbides. Region III allows nucleation of cementite but not the formation of bainitic ferrite. If the quench temperature [888 K (615 C)] for steel B and corresponding austenite carbon content (.5 wt pct) are superimposed on Figure 13, it is evident that conditions are thermodynamically commensurate with region III. The formation of FCA but not bainite on quenching at 888 K (615 C) is therefore consistent thermodynamically. Takahashi and Bhadeshia [16] have demonstrated how, depending on temperature and austenite carbon content, cementite precipitation is thermodynamically possible in steels without the formation of bainite. The pattern of precipitation should be discontinuous and not the cooperative mode found for the eutectoid reaction leading to pearlite. Compared to co-operative eutectoid pearlite transformation, discontinuous cementite precipitation is known to require a lower free energy and lower carbon concentration at the interface. [17] The free energy diagram in Figure 14 shows that in the case of discontinuous precipitation, every time second phase b with a composition richer than x forms, there is a corresponding lowering in free energy. This is true even when the amount of b phase is so small that the composition of the a phase is almost identical to the initial matrix composition x m. On the other hand, for the eutectoid reaction to be possible, the composition of the a phase must be lower than x, while the composition of the b phase has to be higher than x, so that the net free energy can be reduced. Thus compared to discontinuous precipitation, the eutectoid reaction needs a large fraction of the equilibrium segregation to be accomplished as cells grow. This would explain why in the current work FCA rather than pearlite forms in large austenite grains with carbon contents lower than the eutectoid composition and at temperatures intermediate between Widmansta tten ferrite and bainite. The pattern of cementite precipitation during FCA transformation is likely to be complex. The precipitation of cementite should deplete the surrounding austenite of carbon so that it transforms to ferrite. This in turn causes a local increase in carbon content of austenite so that the process is repeated. Alternate formation of ferrite and carbide would be strictly dependent on the local carbon concentration of the austenite and the transformation temperature. It is concluded, therefore, that FCA transformation is not a bainitic or martensitic type but is reconstructive in nature with the kinetics controlled by discontinuous precipitation of cementite. 5. CCT diagrams In the present study, CCT diagrams may be constructed for the thermally cycled steels A, B, and C using 1 T F % Ms T F (B) % B % Transformed Transformation sequence PF P WF FCA B M T s Transformation start T F Transformation finish Temperature % FCA % WF T s T F (FCA) T F (WF) T F (P) % P T F (PF) % PF Fig. 14 Schematic diagram showing free energy comparison of solute segregation in (a) discontinuous and (b) continuous precipitation (after Ref. [17]). Fig. 15 Schematic diagram showing microstructure constituent transformation sequence and finish transformation temperatures used for constructing CCT diagrams of thermally cycled steels A, B, and C.
Fig. 16 CCT diagrams for.17 wt pct C steels A and C, austenitized 1373 (11 C): (a).52 wt pct Mn and (b) 1.56 wt pct Mn. the phase transformation sequence identified earlier. Knowing the percent transformed against temperature curves for the thermally cycled steels, and knowing the volume percent of microstructure constituents (see Tables II and III) it is possible to estimate the start and finish temperatures during cooling for each phase formed sequentially as shown schematically in Figure 15. It is a relatively simple matter then to construct CCT diagrams. CCT diagrams for.17 wt pct C steels A and C with low (.52 wt pct) and high (1.56 wt pct) manganese contents austenitized at 1373 K (11 C) are shown in Figure 16. The effect of decreasing manganese content on increasing FCA formation is clearly evident. At the lower manganese content, transformation temperatures are raised so that the primary ferrite bay is substantially increased in size and extends over a wide range of cooling rates. The Widmansta tten ferrite and bainite bay sizes are significantly reduced enabling the FCA bay to extend to much faster cooling rates with the consequent marked increase in FCA formation. CCT diagrams for.13 wt pct C steel B with intermediate manganese content (.86 wt pct Mn) austenitized at 1373 K and 1573 K (11 C and Fig. 17 CCT diagram for.13 wt pct C steel B with.86 wt pct Mn austenitized: (a) 1373 (11 C) and (b) 1573 K (13 C). 13 C)areshowninFigure 17 The effect of austenitizing temperature on FCA formation is much less marked and more subtle than is the case with manganese. Decreasing austenitizing temperature increases transformation start temperatures analogous to decreasing manganese content (see above) and increases the size of the primary ferrite bay, particularly at slower cooling rates. However, the Widmanstätten ferrite bay is slightly increased in size and the FCA bay is decreased, although extending to faster cooling rates. A consequence is that at around the 5 C s 1 cooling rate both Widmansta tten ferrite and FCA formation is still possible at the expense of the bainite reaction. CCT diagrams are important since they summarize much of the information needed when processing steels. Heat treatments can be designed which achieve the desired microstructure and properties for a given application. Further work is required to determine the influence of FCA on mechanical properties. However, the CCT diagrams in the current work for the first time incorporate FCA. This should provide a greater accuracy in controlling product properties, and developing empirical or theoretical models for effective and reliable property predictions.