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Change in carbon state by low-temperature aging in heavily drawn pearlitic steel wires
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Acta Materialia 60 (2012) 387–395 www.elsevier.com/locate/actamat
Change in carbon state by low-temperature aging in heavily drawn pearlitic steel wires Jun Takahashi a,⇑, Makoto Kosaka b, Kazuto Kawakami a, Toshimi Tarui b a
Advanced Technology Research Labs., Nippon Steel Corporation, 20-1 Shintomi, Futtsu-city, Chiba 293-8511, Japan b Steel Research Labs., Nippon Steel Corporation, 20-1 Shintomi, Futtsu-city, Chiba 293-8511, Japan Received 30 June 2011; received in revised form 8 September 2011; accepted 8 September 2011 Available online 4 November 2011
Abstract Atom probe tomography analysis of heavily drawn pearlitic steel wires was performed with and without low-temperature aging. In the as-drawn wire, the lamellar cementite hardly decomposed and remained in a sufficient amount. By contrast, almost homogeneous carbon atomic distribution of the concentration near the average carbon content was observed in the wire with maximum tensile strength aged at 150 °C for 30 min. In the 200 °C 30 min aging, carbon atoms were enriched at the boundary (prior lamellar cementite) and the carbon concentration in the lamellar ferrite was lower. The change in carbon state was explained by the presence of the high number density of vacancies that was introduced by heavy drawing. These results indicate that cementite decomposition occurred during the thermal aging after, and not during, drawing. The mechanism of the change in strength by low-temperature aging was discussed. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Atom probe; Steel wire; Cementite decomposition; Pearlite; Solute carbon
1. Introduction Pearlitic steel wires with high tensile strength of greater than 4000 MPa have been commercialized by the drawing of hypereutectoid steel rods. Developments of higher tensile wires are required to lighten automotive tires and thereby reduce carbon dioxide emissions. However, the lamellar cementite forming the pearlitic structure is decomposed partially or fully by heavy drawing, and the mechanical properties of the wires are changed by the cementite decomposition [1,2]. Cementite decomposition in cold-worked steels was detected by Mo¨ssbauer spectroscopy in the 1970s [3,4]. Furthermore, it was studied in heavily drawn pearlitic steel wires [5,6]. Cementite decomposition in heavily drawn wires has been directly observed using transmission electron microscopy (TEM) and atom probe field ion microscopy since the second half of the 1990s [7–17]. Elemental maps resulting ⇑ Corresponding author. Tel.: +81 439 80 2169; fax: +81 439 80 2746.
E-mail address: [email protected] (J. Takahashi).
from atom probe tomography (APT) analysis revealed the distribution of carbon atoms in drawn pearlitic lamellae after cementite decomposition [9–11,13,15–17]. It was reported, based on APT analysis, that most cementite was decomposed in drawn wires with true strains of 3.5–4.5 [8–11], and that the lamellar cementite was fully decomposed in heavily drawn wire with a true strain of 5.1 [13]. In contrast, Nam et al. [6] found, using Mo¨ssbauer measurements, that at most 50% of the cementite volume was decomposed in drawn wires with true strains of more than 3, and the ratio tended to saturate even if the true strain increased because of carbon saturation at dislocations located near the cementite/ferrite interface. Cementite decomposition has been reported not only in heavily drawn wires, but also in carbon steels subjected to high-pressure torsion [18–20]. Three mechanisms of cementite decomposition have been hypothesized. The first is the strong attractive interaction between carbon atoms and the high density of dislocations introduced by heavy deformation [1,6,14,20–22]. In this model, the interaction energy is larger than the binding energy between the carbon
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.09.014
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and iron atoms in cementite (0.8 and 0.5 eV atom1, respectively), thus, at temperatures at which the diffusion of carbon atoms is enabled, the carbon atoms move to the dislocations. The second possible mechanism is the Gibbs–Thomson effect [7,11,12]. Due to the fragmentation of cementite by heavy deformation, the free energy of the interface between cementite increases and matrix and the local solubility of carbon increases, which leads to the decomposition of the cementite. It has also been reported that off-stoichiometry cementite with carbon vacancies caused by heavy deformation is unstable, resulting in the decomposition of the cementite [19,20]. The third mechanism is the so-called “carbon drag effect”. In this model, moving dislocations during deformation carry the trapped carbon atoms because of the large attractive interaction between the carbon and the dislocation [18,23]. The dislocations generated from the interface of lamellar cementite move to the lamellar ferrite with trapped carbon atoms, and carbon atoms remain in the ferrite by the pair annihilation of the dislocations. Consequently, highly supersaturated solute carbon is formed. The mechanism of cementite decomposition has been discussed without a common consensus being reached. The carbon positions (trapping sites) after the cementite decomposition have also been discussed. Segregations at dislocations or interfaces between the lamellar ferrite and cementite, supersaturated solid solution in ferrite matrix and the formation of a partially amorphous phase have all been proposed [16,17]. There are two reasons for the difficulty in reaching a consensus. The first is that the heavily drawn wires have very fine microstructures, making it difficult to characterize the carbon state in the fine-drawn pearlite structure. The second is that the characterized wires did not undergo the same drawing and aging conditions. In particular, heavily drawn wires were significantly influenced by aging at room temperature (RT) after drawing. Further, in most of the previous reports, little attention was paid to any unintentional aging. In this study, we produced heavily drawn wires with the intention of separating aging during drawing and aging after drawing. We investigate the aging temperature dependence on the cementite decomposition and the carbon state of the wires using APT and TEM techniques. We discuss the mechanism of cementite decomposition and the sites of carbon atoms after the decomposition. We also discuss the relation between the tensile strength and the carbon state in terms of the strengthening mechanism. 2. Experimental 2.1. Preparation of the sample wires The material used in this study has the composition of a hypereutectoid steel: Fe–0.92C–0.48Mn–0.22Si wt.% or Fe–4.13C–0.47Mn–0.42Si at.%. The material was melted, continuously cast into blooms, broken down into billets, hot rolled and directly patented into wire rods 5.5 mm in
diameter. The wire rod was cold-drawn with powder lubricant to a diameter of 2.0 mm, and heated to 950 °C to be patented in a lead bath at 580 °C. Patented wires of both diameters were further cold-drawn to 0.2 mm in liquid lubricant to the true strain of e = 4.61, where true strain e is defined as ln(Ai/Af), where Ai and Af are the initial and final cross-sectional areas, respectively. The speed of the final drawing was set to 30 m min1, which is significantly lower than that of the commercial production process. As-drawn wires were immediately cooled down with dry ice or immersed in liquid nitrogen, then kept in a freezer below 30 °C to prevent them from aging at RT. It should be noted that the controlled production of sample wires is different from the commercial production process. The aim of the lower-speed drawing and sample freezing is to distinguish between the heavy deformation effect in the midst of the drawing and the thermal effect after drawing. Intentional post-drawing aging was conducted in air (6200 °C) or in a salt bath (>200 °C) for 30 min. The wire of e = 4.61 were aged at temperatures in 50 °C increments from 100 to 400 °C. The tensile test was conducted twice on both strain wires of all aged levels, including without aging, and the averaged value of the two tensile tests was used in this study. 2.2. Instruments for microstructure characterization A Hitachi HF2000 transmission electron microscope, operated at 200 kV, was applied for the characterization of microstructures of sample wires. Thin films for TEM were prepared using the following procedure. Several wires were embedded parallel in a matrix of epoxy resin. The embedded wires were then thinned to about 30 lm perpendicularly to the wire axis by mechanical polishing. Thin films of longitudinal sections of the wires were produced by electropolishing using an electrolyte of 5% perchloric acid in acetic acid. The contaminated surface and oxide layers were removed by argon ion milling. It should be noted that the specimen temperature was raised slightly to harden the resin (60 °C 90 min). Needle-shaped specimens for APT analysis were fabricated directly from the sample wires using the standard two-stage electropolishing technique [24]. The electrolytes used were 25% perchloric acid in acetic acid for the first stage and 2% perchloric acid in 1-butoxyethanol for the second stage. A needle specimen tip was fabricated by electropolishing for a few minutes. This short fabrication time effectively suppressed any unintentional aging even if the temperature of the electrolyte increased slightly. In the preparation method, the tip position of the needle specimen always corresponded to the center region of the wire. Consequently, APT analysis was conducted at the center region of the wire along the probing direction parallel to the drawing direction (the wire’s axis). It should be noted that the carbon concentration peaks of thin cementite lamellae appeared to be slightly lower and wider than the peaks by the local magnification effect associated with
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the local radius of curvature in the tip surface [15,24]. In contrast, the focused ion beam (FIB) milling method enables the fabrication of a needle tip from the specific region of the wire and probing in the direction perpendicular to the pearlitic lamellae with high spatial resolution. However, as the FIB milling may raise the specimen temperature when large beam currents are used [25], we used the electropolishing method in this experiment. An energy-compensated three-dimensional atom probe (3-DAP, Oxford NanoScience Ltd.) with a standard reflectron was applied for APT analysis. Measurements of 1 106 atoms were performed in an ultrahigh vacuum of less than 5 1011 torr at a specimen temperature of 65 K, with a total voltage for probing in the range of 10– 15 kV, a pulse fraction of 20% and a pulse frequency of 1.5 kHz. Several different positions in each sample wire were measured with needle specimens of more than three. In atom probe measurements, the lamellar cementite was located in the center of the field of view using FIM image. In order to obtain information on phase and crystallographic direction from a larger area, FIM images were observed using a neon gas of about 1 105 torr at 80 K, with a distance from the tip to the screen of 58 mm and a screen diameter of 73 mm. Atomic data sets were analyzed using Posap analysis software (version 1.76). Peaks of 27, 28, 28.5 and 29 amu were assigned to iron, a peak of 27.5 to manganese, and peaks of 14, 14.5 and 15 to silicon in the mass-to-charge spectrum. Carbon was also detected as molecular ions. In the carbon assignment, peaks of 6, 6.5, 12 and 13 were assigned to the monomers, peaks of 18, 18.5 and 36 to the trimers, a peak of 24.5 to the tetramer, and a peak of 24 to the dimer of carbon atoms [15,26]. We confirmed that the spherical cementite had a carbon concentration of 25–27 at.% by the peak assignment [26]. A small peak at 25 amu was excluded from the carbon assignment because of the overlap with the residual chromium peak. 3. Experimental results 3.1. Tensile test Fig. 1 shows the change in tensile strength of sample wires as a function of aging temperature. The tensile strength of the as-drawn wire was about 4480 MPa. The tensile strength increased with low-temperature aging, and showed a maximum strength with aging at 100–150 °C. The peak aging region is represented by the dark background in the figure. For aging above the temperature showing the maximum strength, the tensile strength decreased with increasing temperature. The maximum increment in tensile strength from the as-drawn wire was estimated to be approximately 150 MPa. In carbon steels, such strengthening through the lowtemperature aging is caused by dislocation locking by carbon atoms (so-called Cottrell atmosphere) [21], and the lowering of the strength above the peak temperature for aging is caused by recovery and/or recrystallization [27]. We
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Fig. 1. Change in the tensile strength of the sample wire (e = 4.61) as a function of the aging temperature. The arrows indicate the wires to study in detail.
discuss whether the interpretation is suitable in terms of the change in the microstructure and carbon state. Detailed characterization using TEM, FIM and APT was performed in the sample wires indicated by the arrows in Fig. 1. 3.2. TEM observation Fig. 2 shows microstructures of sample wires (e = 4.61) with and without aging in the longitudinal section by TEM bright-field observation. Cementite lamellae and ferrite lamellae in heavily drawn pearlitic wires are preferentially aligned along the drawing direction, i.e. the wire axis. The direction slightly inclined to the horizontal, parallel to the lamellae, corresponds to the drawing direction in the figures. The drawn pearlitic lamellar structure did not have uniform interlamellar spacing, and the thickness of the ferrite lamellae was in the range of 5–20 nm. Large internal strains and dislocations were observed in the lamellar ferrite. Distinct lamellar cementite was not observed with TEM in the wire with a large true strain. The lamellar ferrite exhibited a strong h1 1 0i texture along the wire axis but the each lamella showed a different contrast, indicating that each ferrite lamella had the different crystallographic direction around the axis. No difference in microstructure was observed between the wires aged at temperatures below 200 °C. In contrast, a recovered microstructure was observed in the wire aged at 350 °C, the contrast of the ferrite lamellae having changed to a slightly rounded shape. 3.3. FIM and APT analyses Fig. 3 shows the FIM and APT analysis results of an asdrawn wire without aging. The FIM image consists of distinct bright and dark areas (Fig. 3a), which correspond to the ferrite and cementite lamellae, respectively. The scale bar in the FIM image was calculated based on the tip radius estimated from applied voltage [24]. The two representative three-dimensional (3-D) carbon maps indicate that sufficient amounts of carbon atoms of the lamellar
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Fig. 2. TEM bright-field images of longitudinal sections of the sample wires with (a) no aging, (b) 150 °C aging, (c) 200 °C aging and (d) 350 °C aging.
Fig. 3. (a) FIM image at 10 kV and (b) 3-D carbon maps and concentration profile of the as-drawn sample wire without aging.
cementite remain and the lamellar ferrite maintains a low carbon concentration. The concentration profile across the lamellar cementite shows a high carbon concentration (>20 at.%) and a low silicon concentration in the center region of the cementite. The results indicate that hardly any decomposition of the cementite occurred despite the heavy drawing (e = 4.61). Fig. 4 shows the FIM and APT analysis results of a wire aged at 150 °C, which had the maximum tensile strength (peak aging). The carbon state differs significantly from that of the as-drawn wire. The FIM image consists of almost uniform contrast areas, where a few small dark portions remain. The 3-D carbon maps indicate that the
carbon atoms are distributed almost homogeneously. The concentration profile indicated a mostly uniform distribution of carbon atoms of about 4–5 at.% in the analyzed volume, which concentration is almost identical to the average carbon content in the wire. The maps did not show any distinct segregation of carbon atoms at the dislocations. A unique carbon distribution was shown in a heavily drawn pearlitic wire of e = 5.1 by Hono et al. [13] using APT, which they suggested was caused by the complete decomposition of cementite due to heavy deformation. APT analysis of the wire aged at 150 °C was performed four times using different needle specimens. Two of the samples showed the homogeneous carbon distribution
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Fig. 4. (a) FIM image at 13 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 150 °C.
Fig. 5. (a) FIM image at 9.5 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 200 °C.
portrayed in the upper map of Fig. 4b, while the other two samples showed the slightly uneven distribution in the lower map. The standard atom probe has a small analyzed volume of at most about 20 nm 20 nm 100 nm. Thus, the analyzed volume of Fig. 4b might be within the lamellar ferrite. However, such a uniform distribution of carbon atoms was observed in most of the analyzed volumes of the aged wire; furthermore, the FIM image, with its uniform contrast, also supports the APT result. The interlamellar spacing of the wire was estimated to be 5–20 nm, and thus it is considered that the analyzed volume mostly includes the region of the prior lamellar cementite and its interface. Therefore, the almost homogeneous carbon distribution of the APT result indicates that the local carbon concentration was almost the same throughout the lamellar ferrite, the lamellar cementite and their interface. Fig. 5 shows the FIM and APT results of a wire aged at 200 °C, which had a slightly lower tensile strength than that of the 150 °C aged wire. The FIM image shows a slightly inhomogeneous contrast, unlike the 150 °C aged wire (Fig. 5a). The bright and dark areas appear again. The 3D carbon maps indicate that carbon atoms are enriched in the slightly wide band region. The concentration profile across the enriched carbon band shows a much lower
carbon concentration than the 25 at.% stoichiometric carbon concentration of cementite. Thus, the enriched region is considered to be prior lamellar cementite or lamellar ferrite boundaries with dislocations. The carbon concentration in the ferrite matrix other than in the enriched region is much lower than that of the 150 °C aged wire. Fig. 6 shows the FIM and APT results of a wire aged at 350 °C, which had a tensile strength much lower than that of the as-drawn wire. The FIM image consists of clear bright and dark contrast areas, which is similar to the asdrawn wire. However, the morphology of the dark region differed significantly from that of the as-drawn wire. The as-drawn wire had a lamella-like structure, while the 350 °C aged wire contains some circular dark features in addition to the lamellar structure. The 3-D carbon maps indicate that the small features with a high carbon concentration appear along the ferrite boundary (prior lamellar cementite). The carbon concentration in the feature appears to be more than 20 at.%, while the silicon concentration is low within the feature and is slightly high near its edges. These results indicate that small cementite or carbide spheres was generated along the boundaries, where silicon atoms were discharged by the partitioning. In contrast, the carbon concentration in ferrite was much lower than that of the wires aged at lower temperatures.
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Fig. 6. (a) FIM image at 11 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 350 °C.
Fig. 7. Change in carbon state and tensile strength as a function of aging temperature. The yield strength is also represented in the figure.
4. Discussion 4.1. Mechanism of cementite decomposition Fig. 7 shows a summary of the change in carbon state and tensile strength of the wire with a true strain of 4.61 as a function of aging temperature. The yield strength is also shown in the figure, and is estimated to be 0.2% proof stress. The carbon distribution in drawn pearlitic lamellae in the wire changed significantly, depending on the aging temperature, i.e. the degree of aging. In our experiment, under the conditions of lower-speed drawing and sample freezing after drawing, cementite decomposition hardly occurred immediately after drawing (Fig. 3). However, through low-temperature aging at 150 °C, the lamellar cementite decomposed significantly and an almost homogeneous carbon distribution, with the high concentration corresponding to the average carbon content, was formed (Fig. 4). Since these are the results of sample wires produced by a controlled process, the influences of the temper-
ature rise during drawing and the unintentional aging after drawing are considered to be more in the wires produced by the commercial process. However, the results definitely indicate that the cementite decomposition proceeded during the aging after drawing, and not during the drawing itself. Therefore, the carbon drag effect, whereby the moving dislocations carry carbon atoms from cementite to ferrite lamellae during drawing, is not the main mechanism of cementite decomposition. Even if the wire temperature increases by high-speed drawing, the diffusion of carbon atoms is much lower than the velocity of the dislocation movement in the drawing. Thus, the carbon drag effect is considered not to be the main mechanism even in the commercial production process. The fact that carbon atoms moved during aging after drawing suggests the possibility that it is either the large dislocation–carbon interaction energy effect or the Gibbs–Thomson effect since both effects could be limited by the diffusion of carbon atoms [14].
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Furthermore, by increasing the aging temperature, the carbon concentration in the ferrite was decreased and carbon enrichment at the boundaries (prior cementite lamellae) appeared. This phenomenon is peculiar because the movement of carbon atoms changed to the opposite direction between 150 and 200 °C. In order to understand this phenomenon, temperature-dependent factors associated with the cementite decomposition must be taken into consideration. Based on the the Gibbs–Thomson effect, Hinchliffe and Smith [28] calculated the solubility of carbon in ferrite for hemicylindrical cementite particles of different radii and across a range of temperatures. The fragmented hemicylindrical cementite with a radius of 0.5 nm showed about 0.41 at.% at 150 °C. The value is one order lower than the observed carbon concentration of 4–5 at.% in lamellar ferrite (Fig. 4b). When a high concentration of defects was introduced in the cementite phase, the solubility of the carbon is predicted to increase more because the free energy of cementite is effectively raised to a higher level. Such a solid solution state is considered to be unstable because carbon clusters and/or carbides preferably precipitate [29]. However, no such clusters were observed with APT and TEM. Thus, the cementite decomposition and the change in carbon state cannot be explained by the Gibbs–Thomson effect alone. We therefore believe that the main mechanism of cementite decomposition is the interaction between carbon atoms and lattice defects introduced during heavy drawing. 4.2. Origin of the uniform distribution of carbon atoms We first discuss the origin of the homogeneous distribution of carbon atoms. We showed in this experiment that not only heavy drawing but also aging played an important role in the cementite decomposition. Two factors are required for the cementite decomposition; the first is the diffusion of carbon atoms for a distance substantially greater than the interlamellar spacing and the second is the large number density of more stable sites rather than the binding energy between carbon and iron in cementite [6,10]. The high density of dislocations introduced by heavy drawing and interfaces between lamellar cementite and ferrite have been proposed as the stable sites. The 3-D carbon maps in the wire aged at 150 °C showed an almost homogeneous carbon distribution of 4–5 at.% throughout the lamellar ferrite, lamellar cementite and their interfaces. High spatial resolution atom probe tomography can sufficiently discriminate carbon atomic positions when they are located at the dislocations and/or dislocation walls. Fig. 8a shows the magnified microstructure of the wire aged at 150 °C, which shows a homogeneous carbon distribution by APT analysis. Distinct cementite was not observed, but each ferrite lamella could be recognized as an image of different contrast due to misorientation and internal strain. Furthermore, a difference in contrast was also observed within each lamellar ferrite. Fig. 8b shows a schematic dia-
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Fig. 8. (a) Magnified bright-field micrograph and (b) schematic diagram showing the distribution of dislocations of the 150 °C aged sample wire. The region surrounded by the dashed line indicates the volume analyzed by APT.
gram explaining the distribution of dislocations in the wire. The diffraction analysis showed that the lamellar ferrite consisted of sub-grains intersected by small angle boundaries [14,16]. It is therefore considered that there were dislocation walls at interfaces of lamellar ferrite and at small angle boundaries across the lamellar ferrite. As mentioned previously, the homogeneous carbon distribution observed in the 150 °C aged wire indicated that the local carbon concentration was almost the same throughout the lamellar ferrite, lamellar cementite and their interface. Such a phenomenon is not explained by the carbon enrichment at the dislocation walls in ferrite and at the interfaces of the lamellar ferrite. Plastic deformation introduces excess vacancies in steel [30–32]. Vacancies of the order of 104 under a strain of about 20% were estimated by thermal desorption spectroscopy of tritium in 0.025 wt.% C steel, and the concentration of the vacancies were increased by lowering the temperature and by increasing the amount of strain [32]. In contrast, it is reported that the vacancies in steel had a large attractive interaction with carbon atoms and that the interaction energy (0.85 eV atom1) is sufficiently larger than the binding energy between carbon and iron in cementite [33]. The change in the carbon state in the experiment is explained by the presence of the high number density of excess vacancies. Migration for a sufficient distance was difficult at RT even though a high number density of vacancies was introduced into the drawing wire. In aging at 150 °C, the carbon atoms migrated a sufficient distance to become entrapped by the vacancies that were distributed homogeneously throughout the matrix ferrite. This is the reason why the cementite decomposition proceeded signif-
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icantly with the 150 °C aging. Carbon enrichment at the ferrite boundaries (prior lamellar cementite) was observed in the wire aged at 200 °C, which temperature was almost identical to the decomposition region of the carbon– vacancy complex [33,34]. Thus, the carbon atoms forming the complexes were decomposed and dissolved carbon atoms segregated at the boundaries, as shown in Fig. 5. X-ray diffraction measurements have already confirmed that the lattice parameter of the ferrite phase is not affected by cementite decomposition [12,13,17]. The experimental X-ray diffraction results can be explained by our proposed model whereby carbon atoms segregate at lattice defects in ferrite, such as vacancies, dislocations and boundaries. In the aging at higher temperatures, carbon atoms segregated at the boundaries and precipitated as small cementite spheres along the boundaries (Fig. 6). TEM observations indicated that the microstructure of the drawn pearlitic structure was recovered with aging at 350 °C, which causes the large decrement in tensile strength. 4.3. Strength mechanism of steel wires with low-temperature aging The very high tensile strength of heavily drawn pearlitic steel wires has been explained mainly by the refinement of the pearlitic lamellae by the heavy drawing [1,7,12–14,35]. However, various factors associated with dislocations, the state of the cementite, solute carbon and so on strongly affected the strength, and a different strengthening mechanism, dispersion of amorphous proto-cementite, has been proposed [17]. Here, we discuss only the change in yield and tensile strengths as a function of aging temperature. The yield strength was sufficiently lower than the tensile strength in the as-drawn wire, while it drastically increased through low-temperature aging at 100 °C. A large increment in the yield strength is normally considered to be due to strain aging [21,27], whereby dislocations are locked by carbon atoms (the so-called Cottrell atmosphere). Such behavior of the yield strength suggests that the dynamic strain aging during drawing and the static aging at RT after drawing were sufficiently suppressed in this experiment. In the experiment, the tensile strength reached a maximum when the carbon concentration in the lamellar ferrite also reached a maximum by aging. At such a low aging temperature the dislocations could not move, though carbon atoms could. Thus, the strength was influenced only by the change in the carbon state in the wire. By aging at 150 °C, the lamellar cementite was almost fully decomposed, and dissolved carbon atoms showed a homogeneous carbon distribution of 4–5 at.% throughout the lamellar cementite, lamellar ferrite and their interfaces. Such a change in the carbon state has two different influences on the strengthening mechanism. The first is the decomposition of lamellar cementite, which implies that the strength of the barriers to deformation decreases in pearlitic structure. As mentioned before,
the strength of heavily drawn pearlitc steel wires is explained mainly by the refinement of the pearlitic lamellae, i.e. the Hall–Petch relationship [35]. The Hall–Petch coefficient (slope) shows the measure of the strength of the barriers to deformation. If the lamellar cementite is mostly decomposed, the coefficient will decrease and the cementite’s contribution to strengthening will decrease. After heavy drawing, however, ferrite lamellae become misorientated with each other and with the dislocation walls, which suggests that the interfaces of ferrite lamellae act as grain boundaries. It is reported that the ferrite grain boundary has a large Hall–Petch coefficient (600 MPa lm1/2) when the carbon content is sufficiently high [36,37]. Therefore, it is considered that the decrease in strength due to the disappearance of lamellar cementite is not significant. The second influence is the large increment in carbon concentration in the lamellar ferrite, which causes the solid solution strengthening of carbon to increase in the wire. It was reported that the amount of solid solution strengthening pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ffi of carbon in martensite steel is about 1400 Cðwt:%Þ MPa, where C is the concentration of solute carbon [38]. The amount of solid solution strengthening in the concentration corresponding to an average carbon content of 0.92 wt.% (4.12 at.%) is estimated to be about 1300 MPa, although the solid solution strengthen of carbon in ferrite steel is not the same as that in martensite steel. According to our proposed model in wire aged at 150 °C, the carbon atoms were trapped at vacancies and formed carbon–vacancy complexes. In this case, the carbon atoms bonded with vacancies make a smaller contribution to the overall lattice strains than solid solution carbon in the ferrite matrix because the lattice strains caused by the carbon–vacancy complexes are smaller. Therefore, actual solid solution strengthening of carbon is considered to be smaller than the estimated amount. The observed increment in the tensile strength was actually about 150 MPa. This small value indicates that the decrement in the strength due to the decomposition of lamellar cementite and the increment in the solid solution strengthening of carbon was almost identical. The yield strength and tensile strength decreased rapidly with aging above 250 °C. This phenomenon is explained by the recovery of the drawn pearlitic structure. The APT analysis clearly showed that fine spherical cementite precipitated along the boundaries and further that the concentration of carbon in the ferrite decreased because of the reduction in lattice defects (vacancies and dislocations). 5. Conclusion APT analysis of heavily drawn pearlitic steel wires with and without low-temperature aging was performed to elucidate the mechanism of the cementite decomposition. A sample wire of true strain e = 4.6 was produced under the controlled conditions of low-speed drawing and sample freezing to distinguish between the effect of heavy deforma-
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tion during the drawing and the thermal effect after drawing. In the as-drawn wire, the lamellar cementite hardly decomposed and remained in a sufficient amount. In contrast, an almost homogeneous carbon distribution with average carbon content was observed in the 150 °C aging after drawing, which had the maximum tensile strength. In the 200 °C aging, carbon atoms segregated at the boundaries. In the 350 °C aging, small cementite or carbide spheres were generated along the boundaries. The experimental results led to the following conclusions. (1) Cementite decomposition proceeded during aging after drawing, but not during the drawing, which indicates that the carbon drag effect whereby the moving dislocations carry carbon atoms from cementite to ferrite during drawing is not the main mechanism of the cementite decomposition. (2) The high number density of vacancies introduced by the heavy drawing plays an important role in the decomposition of cementite in the wire. (3) The change in tensile strength by aging at 150 °C was explained by the decrement in strength due to the decomposition of lamellar cementite and the increment in solid solution strengthening due to dissolved carbon atoms in lamellar ferrite.
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Acta Materialia 60 (2012) 387–395 www.elsevier.com/locate/actamat
Change in carbon state by low-temperature aging in heavily drawn pearlitic steel wires Jun Takahashi a,⇑, Makoto Kosaka b, Kazuto Kawakami a, Toshimi Tarui b a
Advanced Technology Research Labs., Nippon Steel Corporation, 20-1 Shintomi, Futtsu-city, Chiba 293-8511, Japan b Steel Research Labs., Nippon Steel Corporation, 20-1 Shintomi, Futtsu-city, Chiba 293-8511, Japan Received 30 June 2011; received in revised form 8 September 2011; accepted 8 September 2011 Available online 4 November 2011
Abstract Atom probe tomography analysis of heavily drawn pearlitic steel wires was performed with and without low-temperature aging. In the as-drawn wire, the lamellar cementite hardly decomposed and remained in a sufficient amount. By contrast, almost homogeneous carbon atomic distribution of the concentration near the average carbon content was observed in the wire with maximum tensile strength aged at 150 °C for 30 min. In the 200 °C 30 min aging, carbon atoms were enriched at the boundary (prior lamellar cementite) and the carbon concentration in the lamellar ferrite was lower. The change in carbon state was explained by the presence of the high number density of vacancies that was introduced by heavy drawing. These results indicate that cementite decomposition occurred during the thermal aging after, and not during, drawing. The mechanism of the change in strength by low-temperature aging was discussed. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Atom probe; Steel wire; Cementite decomposition; Pearlite; Solute carbon
1. Introduction Pearlitic steel wires with high tensile strength of greater than 4000 MPa have been commercialized by the drawing of hypereutectoid steel rods. Developments of higher tensile wires are required to lighten automotive tires and thereby reduce carbon dioxide emissions. However, the lamellar cementite forming the pearlitic structure is decomposed partially or fully by heavy drawing, and the mechanical properties of the wires are changed by the cementite decomposition [1,2]. Cementite decomposition in cold-worked steels was detected by Mo¨ssbauer spectroscopy in the 1970s [3,4]. Furthermore, it was studied in heavily drawn pearlitic steel wires [5,6]. Cementite decomposition in heavily drawn wires has been directly observed using transmission electron microscopy (TEM) and atom probe field ion microscopy since the second half of the 1990s [7–17]. Elemental maps resulting ⇑ Corresponding author. Tel.: +81 439 80 2169; fax: +81 439 80 2746.
E-mail address: [email protected] (J. Takahashi).
from atom probe tomography (APT) analysis revealed the distribution of carbon atoms in drawn pearlitic lamellae after cementite decomposition [9–11,13,15–17]. It was reported, based on APT analysis, that most cementite was decomposed in drawn wires with true strains of 3.5–4.5 [8–11], and that the lamellar cementite was fully decomposed in heavily drawn wire with a true strain of 5.1 [13]. In contrast, Nam et al. [6] found, using Mo¨ssbauer measurements, that at most 50% of the cementite volume was decomposed in drawn wires with true strains of more than 3, and the ratio tended to saturate even if the true strain increased because of carbon saturation at dislocations located near the cementite/ferrite interface. Cementite decomposition has been reported not only in heavily drawn wires, but also in carbon steels subjected to high-pressure torsion [18–20]. Three mechanisms of cementite decomposition have been hypothesized. The first is the strong attractive interaction between carbon atoms and the high density of dislocations introduced by heavy deformation [1,6,14,20–22]. In this model, the interaction energy is larger than the binding energy between the carbon
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.09.014
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and iron atoms in cementite (0.8 and 0.5 eV atom1, respectively), thus, at temperatures at which the diffusion of carbon atoms is enabled, the carbon atoms move to the dislocations. The second possible mechanism is the Gibbs–Thomson effect [7,11,12]. Due to the fragmentation of cementite by heavy deformation, the free energy of the interface between cementite increases and matrix and the local solubility of carbon increases, which leads to the decomposition of the cementite. It has also been reported that off-stoichiometry cementite with carbon vacancies caused by heavy deformation is unstable, resulting in the decomposition of the cementite [19,20]. The third mechanism is the so-called “carbon drag effect”. In this model, moving dislocations during deformation carry the trapped carbon atoms because of the large attractive interaction between the carbon and the dislocation [18,23]. The dislocations generated from the interface of lamellar cementite move to the lamellar ferrite with trapped carbon atoms, and carbon atoms remain in the ferrite by the pair annihilation of the dislocations. Consequently, highly supersaturated solute carbon is formed. The mechanism of cementite decomposition has been discussed without a common consensus being reached. The carbon positions (trapping sites) after the cementite decomposition have also been discussed. Segregations at dislocations or interfaces between the lamellar ferrite and cementite, supersaturated solid solution in ferrite matrix and the formation of a partially amorphous phase have all been proposed [16,17]. There are two reasons for the difficulty in reaching a consensus. The first is that the heavily drawn wires have very fine microstructures, making it difficult to characterize the carbon state in the fine-drawn pearlite structure. The second is that the characterized wires did not undergo the same drawing and aging conditions. In particular, heavily drawn wires were significantly influenced by aging at room temperature (RT) after drawing. Further, in most of the previous reports, little attention was paid to any unintentional aging. In this study, we produced heavily drawn wires with the intention of separating aging during drawing and aging after drawing. We investigate the aging temperature dependence on the cementite decomposition and the carbon state of the wires using APT and TEM techniques. We discuss the mechanism of cementite decomposition and the sites of carbon atoms after the decomposition. We also discuss the relation between the tensile strength and the carbon state in terms of the strengthening mechanism. 2. Experimental 2.1. Preparation of the sample wires The material used in this study has the composition of a hypereutectoid steel: Fe–0.92C–0.48Mn–0.22Si wt.% or Fe–4.13C–0.47Mn–0.42Si at.%. The material was melted, continuously cast into blooms, broken down into billets, hot rolled and directly patented into wire rods 5.5 mm in
diameter. The wire rod was cold-drawn with powder lubricant to a diameter of 2.0 mm, and heated to 950 °C to be patented in a lead bath at 580 °C. Patented wires of both diameters were further cold-drawn to 0.2 mm in liquid lubricant to the true strain of e = 4.61, where true strain e is defined as ln(Ai/Af), where Ai and Af are the initial and final cross-sectional areas, respectively. The speed of the final drawing was set to 30 m min1, which is significantly lower than that of the commercial production process. As-drawn wires were immediately cooled down with dry ice or immersed in liquid nitrogen, then kept in a freezer below 30 °C to prevent them from aging at RT. It should be noted that the controlled production of sample wires is different from the commercial production process. The aim of the lower-speed drawing and sample freezing is to distinguish between the heavy deformation effect in the midst of the drawing and the thermal effect after drawing. Intentional post-drawing aging was conducted in air (6200 °C) or in a salt bath (>200 °C) for 30 min. The wire of e = 4.61 were aged at temperatures in 50 °C increments from 100 to 400 °C. The tensile test was conducted twice on both strain wires of all aged levels, including without aging, and the averaged value of the two tensile tests was used in this study. 2.2. Instruments for microstructure characterization A Hitachi HF2000 transmission electron microscope, operated at 200 kV, was applied for the characterization of microstructures of sample wires. Thin films for TEM were prepared using the following procedure. Several wires were embedded parallel in a matrix of epoxy resin. The embedded wires were then thinned to about 30 lm perpendicularly to the wire axis by mechanical polishing. Thin films of longitudinal sections of the wires were produced by electropolishing using an electrolyte of 5% perchloric acid in acetic acid. The contaminated surface and oxide layers were removed by argon ion milling. It should be noted that the specimen temperature was raised slightly to harden the resin (60 °C 90 min). Needle-shaped specimens for APT analysis were fabricated directly from the sample wires using the standard two-stage electropolishing technique [24]. The electrolytes used were 25% perchloric acid in acetic acid for the first stage and 2% perchloric acid in 1-butoxyethanol for the second stage. A needle specimen tip was fabricated by electropolishing for a few minutes. This short fabrication time effectively suppressed any unintentional aging even if the temperature of the electrolyte increased slightly. In the preparation method, the tip position of the needle specimen always corresponded to the center region of the wire. Consequently, APT analysis was conducted at the center region of the wire along the probing direction parallel to the drawing direction (the wire’s axis). It should be noted that the carbon concentration peaks of thin cementite lamellae appeared to be slightly lower and wider than the peaks by the local magnification effect associated with
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the local radius of curvature in the tip surface [15,24]. In contrast, the focused ion beam (FIB) milling method enables the fabrication of a needle tip from the specific region of the wire and probing in the direction perpendicular to the pearlitic lamellae with high spatial resolution. However, as the FIB milling may raise the specimen temperature when large beam currents are used [25], we used the electropolishing method in this experiment. An energy-compensated three-dimensional atom probe (3-DAP, Oxford NanoScience Ltd.) with a standard reflectron was applied for APT analysis. Measurements of 1 106 atoms were performed in an ultrahigh vacuum of less than 5 1011 torr at a specimen temperature of 65 K, with a total voltage for probing in the range of 10– 15 kV, a pulse fraction of 20% and a pulse frequency of 1.5 kHz. Several different positions in each sample wire were measured with needle specimens of more than three. In atom probe measurements, the lamellar cementite was located in the center of the field of view using FIM image. In order to obtain information on phase and crystallographic direction from a larger area, FIM images were observed using a neon gas of about 1 105 torr at 80 K, with a distance from the tip to the screen of 58 mm and a screen diameter of 73 mm. Atomic data sets were analyzed using Posap analysis software (version 1.76). Peaks of 27, 28, 28.5 and 29 amu were assigned to iron, a peak of 27.5 to manganese, and peaks of 14, 14.5 and 15 to silicon in the mass-to-charge spectrum. Carbon was also detected as molecular ions. In the carbon assignment, peaks of 6, 6.5, 12 and 13 were assigned to the monomers, peaks of 18, 18.5 and 36 to the trimers, a peak of 24.5 to the tetramer, and a peak of 24 to the dimer of carbon atoms [15,26]. We confirmed that the spherical cementite had a carbon concentration of 25–27 at.% by the peak assignment [26]. A small peak at 25 amu was excluded from the carbon assignment because of the overlap with the residual chromium peak. 3. Experimental results 3.1. Tensile test Fig. 1 shows the change in tensile strength of sample wires as a function of aging temperature. The tensile strength of the as-drawn wire was about 4480 MPa. The tensile strength increased with low-temperature aging, and showed a maximum strength with aging at 100–150 °C. The peak aging region is represented by the dark background in the figure. For aging above the temperature showing the maximum strength, the tensile strength decreased with increasing temperature. The maximum increment in tensile strength from the as-drawn wire was estimated to be approximately 150 MPa. In carbon steels, such strengthening through the lowtemperature aging is caused by dislocation locking by carbon atoms (so-called Cottrell atmosphere) [21], and the lowering of the strength above the peak temperature for aging is caused by recovery and/or recrystallization [27]. We
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Fig. 1. Change in the tensile strength of the sample wire (e = 4.61) as a function of the aging temperature. The arrows indicate the wires to study in detail.
discuss whether the interpretation is suitable in terms of the change in the microstructure and carbon state. Detailed characterization using TEM, FIM and APT was performed in the sample wires indicated by the arrows in Fig. 1. 3.2. TEM observation Fig. 2 shows microstructures of sample wires (e = 4.61) with and without aging in the longitudinal section by TEM bright-field observation. Cementite lamellae and ferrite lamellae in heavily drawn pearlitic wires are preferentially aligned along the drawing direction, i.e. the wire axis. The direction slightly inclined to the horizontal, parallel to the lamellae, corresponds to the drawing direction in the figures. The drawn pearlitic lamellar structure did not have uniform interlamellar spacing, and the thickness of the ferrite lamellae was in the range of 5–20 nm. Large internal strains and dislocations were observed in the lamellar ferrite. Distinct lamellar cementite was not observed with TEM in the wire with a large true strain. The lamellar ferrite exhibited a strong h1 1 0i texture along the wire axis but the each lamella showed a different contrast, indicating that each ferrite lamella had the different crystallographic direction around the axis. No difference in microstructure was observed between the wires aged at temperatures below 200 °C. In contrast, a recovered microstructure was observed in the wire aged at 350 °C, the contrast of the ferrite lamellae having changed to a slightly rounded shape. 3.3. FIM and APT analyses Fig. 3 shows the FIM and APT analysis results of an asdrawn wire without aging. The FIM image consists of distinct bright and dark areas (Fig. 3a), which correspond to the ferrite and cementite lamellae, respectively. The scale bar in the FIM image was calculated based on the tip radius estimated from applied voltage [24]. The two representative three-dimensional (3-D) carbon maps indicate that sufficient amounts of carbon atoms of the lamellar
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Fig. 2. TEM bright-field images of longitudinal sections of the sample wires with (a) no aging, (b) 150 °C aging, (c) 200 °C aging and (d) 350 °C aging.
Fig. 3. (a) FIM image at 10 kV and (b) 3-D carbon maps and concentration profile of the as-drawn sample wire without aging.
cementite remain and the lamellar ferrite maintains a low carbon concentration. The concentration profile across the lamellar cementite shows a high carbon concentration (>20 at.%) and a low silicon concentration in the center region of the cementite. The results indicate that hardly any decomposition of the cementite occurred despite the heavy drawing (e = 4.61). Fig. 4 shows the FIM and APT analysis results of a wire aged at 150 °C, which had the maximum tensile strength (peak aging). The carbon state differs significantly from that of the as-drawn wire. The FIM image consists of almost uniform contrast areas, where a few small dark portions remain. The 3-D carbon maps indicate that the
carbon atoms are distributed almost homogeneously. The concentration profile indicated a mostly uniform distribution of carbon atoms of about 4–5 at.% in the analyzed volume, which concentration is almost identical to the average carbon content in the wire. The maps did not show any distinct segregation of carbon atoms at the dislocations. A unique carbon distribution was shown in a heavily drawn pearlitic wire of e = 5.1 by Hono et al. [13] using APT, which they suggested was caused by the complete decomposition of cementite due to heavy deformation. APT analysis of the wire aged at 150 °C was performed four times using different needle specimens. Two of the samples showed the homogeneous carbon distribution
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Fig. 4. (a) FIM image at 13 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 150 °C.
Fig. 5. (a) FIM image at 9.5 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 200 °C.
portrayed in the upper map of Fig. 4b, while the other two samples showed the slightly uneven distribution in the lower map. The standard atom probe has a small analyzed volume of at most about 20 nm 20 nm 100 nm. Thus, the analyzed volume of Fig. 4b might be within the lamellar ferrite. However, such a uniform distribution of carbon atoms was observed in most of the analyzed volumes of the aged wire; furthermore, the FIM image, with its uniform contrast, also supports the APT result. The interlamellar spacing of the wire was estimated to be 5–20 nm, and thus it is considered that the analyzed volume mostly includes the region of the prior lamellar cementite and its interface. Therefore, the almost homogeneous carbon distribution of the APT result indicates that the local carbon concentration was almost the same throughout the lamellar ferrite, the lamellar cementite and their interface. Fig. 5 shows the FIM and APT results of a wire aged at 200 °C, which had a slightly lower tensile strength than that of the 150 °C aged wire. The FIM image shows a slightly inhomogeneous contrast, unlike the 150 °C aged wire (Fig. 5a). The bright and dark areas appear again. The 3D carbon maps indicate that carbon atoms are enriched in the slightly wide band region. The concentration profile across the enriched carbon band shows a much lower
carbon concentration than the 25 at.% stoichiometric carbon concentration of cementite. Thus, the enriched region is considered to be prior lamellar cementite or lamellar ferrite boundaries with dislocations. The carbon concentration in the ferrite matrix other than in the enriched region is much lower than that of the 150 °C aged wire. Fig. 6 shows the FIM and APT results of a wire aged at 350 °C, which had a tensile strength much lower than that of the as-drawn wire. The FIM image consists of clear bright and dark contrast areas, which is similar to the asdrawn wire. However, the morphology of the dark region differed significantly from that of the as-drawn wire. The as-drawn wire had a lamella-like structure, while the 350 °C aged wire contains some circular dark features in addition to the lamellar structure. The 3-D carbon maps indicate that the small features with a high carbon concentration appear along the ferrite boundary (prior lamellar cementite). The carbon concentration in the feature appears to be more than 20 at.%, while the silicon concentration is low within the feature and is slightly high near its edges. These results indicate that small cementite or carbide spheres was generated along the boundaries, where silicon atoms were discharged by the partitioning. In contrast, the carbon concentration in ferrite was much lower than that of the wires aged at lower temperatures.
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Fig. 6. (a) FIM image at 11 kV and (b) 3-D carbon maps and concentration profile of the sample wire with aging at 350 °C.
Fig. 7. Change in carbon state and tensile strength as a function of aging temperature. The yield strength is also represented in the figure.
4. Discussion 4.1. Mechanism of cementite decomposition Fig. 7 shows a summary of the change in carbon state and tensile strength of the wire with a true strain of 4.61 as a function of aging temperature. The yield strength is also shown in the figure, and is estimated to be 0.2% proof stress. The carbon distribution in drawn pearlitic lamellae in the wire changed significantly, depending on the aging temperature, i.e. the degree of aging. In our experiment, under the conditions of lower-speed drawing and sample freezing after drawing, cementite decomposition hardly occurred immediately after drawing (Fig. 3). However, through low-temperature aging at 150 °C, the lamellar cementite decomposed significantly and an almost homogeneous carbon distribution, with the high concentration corresponding to the average carbon content, was formed (Fig. 4). Since these are the results of sample wires produced by a controlled process, the influences of the temper-
ature rise during drawing and the unintentional aging after drawing are considered to be more in the wires produced by the commercial process. However, the results definitely indicate that the cementite decomposition proceeded during the aging after drawing, and not during the drawing itself. Therefore, the carbon drag effect, whereby the moving dislocations carry carbon atoms from cementite to ferrite lamellae during drawing, is not the main mechanism of cementite decomposition. Even if the wire temperature increases by high-speed drawing, the diffusion of carbon atoms is much lower than the velocity of the dislocation movement in the drawing. Thus, the carbon drag effect is considered not to be the main mechanism even in the commercial production process. The fact that carbon atoms moved during aging after drawing suggests the possibility that it is either the large dislocation–carbon interaction energy effect or the Gibbs–Thomson effect since both effects could be limited by the diffusion of carbon atoms [14].
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Furthermore, by increasing the aging temperature, the carbon concentration in the ferrite was decreased and carbon enrichment at the boundaries (prior cementite lamellae) appeared. This phenomenon is peculiar because the movement of carbon atoms changed to the opposite direction between 150 and 200 °C. In order to understand this phenomenon, temperature-dependent factors associated with the cementite decomposition must be taken into consideration. Based on the the Gibbs–Thomson effect, Hinchliffe and Smith [28] calculated the solubility of carbon in ferrite for hemicylindrical cementite particles of different radii and across a range of temperatures. The fragmented hemicylindrical cementite with a radius of 0.5 nm showed about 0.41 at.% at 150 °C. The value is one order lower than the observed carbon concentration of 4–5 at.% in lamellar ferrite (Fig. 4b). When a high concentration of defects was introduced in the cementite phase, the solubility of the carbon is predicted to increase more because the free energy of cementite is effectively raised to a higher level. Such a solid solution state is considered to be unstable because carbon clusters and/or carbides preferably precipitate [29]. However, no such clusters were observed with APT and TEM. Thus, the cementite decomposition and the change in carbon state cannot be explained by the Gibbs–Thomson effect alone. We therefore believe that the main mechanism of cementite decomposition is the interaction between carbon atoms and lattice defects introduced during heavy drawing. 4.2. Origin of the uniform distribution of carbon atoms We first discuss the origin of the homogeneous distribution of carbon atoms. We showed in this experiment that not only heavy drawing but also aging played an important role in the cementite decomposition. Two factors are required for the cementite decomposition; the first is the diffusion of carbon atoms for a distance substantially greater than the interlamellar spacing and the second is the large number density of more stable sites rather than the binding energy between carbon and iron in cementite [6,10]. The high density of dislocations introduced by heavy drawing and interfaces between lamellar cementite and ferrite have been proposed as the stable sites. The 3-D carbon maps in the wire aged at 150 °C showed an almost homogeneous carbon distribution of 4–5 at.% throughout the lamellar ferrite, lamellar cementite and their interfaces. High spatial resolution atom probe tomography can sufficiently discriminate carbon atomic positions when they are located at the dislocations and/or dislocation walls. Fig. 8a shows the magnified microstructure of the wire aged at 150 °C, which shows a homogeneous carbon distribution by APT analysis. Distinct cementite was not observed, but each ferrite lamella could be recognized as an image of different contrast due to misorientation and internal strain. Furthermore, a difference in contrast was also observed within each lamellar ferrite. Fig. 8b shows a schematic dia-
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Fig. 8. (a) Magnified bright-field micrograph and (b) schematic diagram showing the distribution of dislocations of the 150 °C aged sample wire. The region surrounded by the dashed line indicates the volume analyzed by APT.
gram explaining the distribution of dislocations in the wire. The diffraction analysis showed that the lamellar ferrite consisted of sub-grains intersected by small angle boundaries [14,16]. It is therefore considered that there were dislocation walls at interfaces of lamellar ferrite and at small angle boundaries across the lamellar ferrite. As mentioned previously, the homogeneous carbon distribution observed in the 150 °C aged wire indicated that the local carbon concentration was almost the same throughout the lamellar ferrite, lamellar cementite and their interface. Such a phenomenon is not explained by the carbon enrichment at the dislocation walls in ferrite and at the interfaces of the lamellar ferrite. Plastic deformation introduces excess vacancies in steel [30–32]. Vacancies of the order of 104 under a strain of about 20% were estimated by thermal desorption spectroscopy of tritium in 0.025 wt.% C steel, and the concentration of the vacancies were increased by lowering the temperature and by increasing the amount of strain [32]. In contrast, it is reported that the vacancies in steel had a large attractive interaction with carbon atoms and that the interaction energy (0.85 eV atom1) is sufficiently larger than the binding energy between carbon and iron in cementite [33]. The change in the carbon state in the experiment is explained by the presence of the high number density of excess vacancies. Migration for a sufficient distance was difficult at RT even though a high number density of vacancies was introduced into the drawing wire. In aging at 150 °C, the carbon atoms migrated a sufficient distance to become entrapped by the vacancies that were distributed homogeneously throughout the matrix ferrite. This is the reason why the cementite decomposition proceeded signif-
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icantly with the 150 °C aging. Carbon enrichment at the ferrite boundaries (prior lamellar cementite) was observed in the wire aged at 200 °C, which temperature was almost identical to the decomposition region of the carbon– vacancy complex [33,34]. Thus, the carbon atoms forming the complexes were decomposed and dissolved carbon atoms segregated at the boundaries, as shown in Fig. 5. X-ray diffraction measurements have already confirmed that the lattice parameter of the ferrite phase is not affected by cementite decomposition [12,13,17]. The experimental X-ray diffraction results can be explained by our proposed model whereby carbon atoms segregate at lattice defects in ferrite, such as vacancies, dislocations and boundaries. In the aging at higher temperatures, carbon atoms segregated at the boundaries and precipitated as small cementite spheres along the boundaries (Fig. 6). TEM observations indicated that the microstructure of the drawn pearlitic structure was recovered with aging at 350 °C, which causes the large decrement in tensile strength. 4.3. Strength mechanism of steel wires with low-temperature aging The very high tensile strength of heavily drawn pearlitic steel wires has been explained mainly by the refinement of the pearlitic lamellae by the heavy drawing [1,7,12–14,35]. However, various factors associated with dislocations, the state of the cementite, solute carbon and so on strongly affected the strength, and a different strengthening mechanism, dispersion of amorphous proto-cementite, has been proposed [17]. Here, we discuss only the change in yield and tensile strengths as a function of aging temperature. The yield strength was sufficiently lower than the tensile strength in the as-drawn wire, while it drastically increased through low-temperature aging at 100 °C. A large increment in the yield strength is normally considered to be due to strain aging [21,27], whereby dislocations are locked by carbon atoms (the so-called Cottrell atmosphere). Such behavior of the yield strength suggests that the dynamic strain aging during drawing and the static aging at RT after drawing were sufficiently suppressed in this experiment. In the experiment, the tensile strength reached a maximum when the carbon concentration in the lamellar ferrite also reached a maximum by aging. At such a low aging temperature the dislocations could not move, though carbon atoms could. Thus, the strength was influenced only by the change in the carbon state in the wire. By aging at 150 °C, the lamellar cementite was almost fully decomposed, and dissolved carbon atoms showed a homogeneous carbon distribution of 4–5 at.% throughout the lamellar cementite, lamellar ferrite and their interfaces. Such a change in the carbon state has two different influences on the strengthening mechanism. The first is the decomposition of lamellar cementite, which implies that the strength of the barriers to deformation decreases in pearlitic structure. As mentioned before,
the strength of heavily drawn pearlitc steel wires is explained mainly by the refinement of the pearlitic lamellae, i.e. the Hall–Petch relationship [35]. The Hall–Petch coefficient (slope) shows the measure of the strength of the barriers to deformation. If the lamellar cementite is mostly decomposed, the coefficient will decrease and the cementite’s contribution to strengthening will decrease. After heavy drawing, however, ferrite lamellae become misorientated with each other and with the dislocation walls, which suggests that the interfaces of ferrite lamellae act as grain boundaries. It is reported that the ferrite grain boundary has a large Hall–Petch coefficient (600 MPa lm1/2) when the carbon content is sufficiently high [36,37]. Therefore, it is considered that the decrease in strength due to the disappearance of lamellar cementite is not significant. The second influence is the large increment in carbon concentration in the lamellar ferrite, which causes the solid solution strengthening of carbon to increase in the wire. It was reported that the amount of solid solution strengthening pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ffi of carbon in martensite steel is about 1400 Cðwt:%Þ MPa, where C is the concentration of solute carbon [38]. The amount of solid solution strengthening in the concentration corresponding to an average carbon content of 0.92 wt.% (4.12 at.%) is estimated to be about 1300 MPa, although the solid solution strengthen of carbon in ferrite steel is not the same as that in martensite steel. According to our proposed model in wire aged at 150 °C, the carbon atoms were trapped at vacancies and formed carbon–vacancy complexes. In this case, the carbon atoms bonded with vacancies make a smaller contribution to the overall lattice strains than solid solution carbon in the ferrite matrix because the lattice strains caused by the carbon–vacancy complexes are smaller. Therefore, actual solid solution strengthening of carbon is considered to be smaller than the estimated amount. The observed increment in the tensile strength was actually about 150 MPa. This small value indicates that the decrement in the strength due to the decomposition of lamellar cementite and the increment in the solid solution strengthening of carbon was almost identical. The yield strength and tensile strength decreased rapidly with aging above 250 °C. This phenomenon is explained by the recovery of the drawn pearlitic structure. The APT analysis clearly showed that fine spherical cementite precipitated along the boundaries and further that the concentration of carbon in the ferrite decreased because of the reduction in lattice defects (vacancies and dislocations). 5. Conclusion APT analysis of heavily drawn pearlitic steel wires with and without low-temperature aging was performed to elucidate the mechanism of the cementite decomposition. A sample wire of true strain e = 4.6 was produced under the controlled conditions of low-speed drawing and sample freezing to distinguish between the effect of heavy deforma-
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tion during the drawing and the thermal effect after drawing. In the as-drawn wire, the lamellar cementite hardly decomposed and remained in a sufficient amount. In contrast, an almost homogeneous carbon distribution with average carbon content was observed in the 150 °C aging after drawing, which had the maximum tensile strength. In the 200 °C aging, carbon atoms segregated at the boundaries. In the 350 °C aging, small cementite or carbide spheres were generated along the boundaries. The experimental results led to the following conclusions. (1) Cementite decomposition proceeded during aging after drawing, but not during the drawing, which indicates that the carbon drag effect whereby the moving dislocations carry carbon atoms from cementite to ferrite during drawing is not the main mechanism of the cementite decomposition. (2) The high number density of vacancies introduced by the heavy drawing plays an important role in the decomposition of cementite in the wire. (3) The change in tensile strength by aging at 150 °C was explained by the decrement in strength due to the decomposition of lamellar cementite and the increment in solid solution strengthening due to dissolved carbon atoms in lamellar ferrite.
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