EFFECT OF HEATING RATE AND MICROSTRUCTURAL SCALE ON AUSTENITIZATION

Transcription

1 Proceedings of the International Conference on Solid-Solid Phase Transformations in Inorganic Materials 2015 Edited by: Matthias Militzer, Gianluigi Botton, Long-Qing Chen, James Howe, Chadwick Sinclair, and Hatem Zurob EFFECT OF HEATING RATE AND MICROSTRUCTURAL SCALE ON AUSTENITIZATION 1 Kester D. Clarke, 1 Amy J. Clarke, 1 Robert E. Hackenberg, 1 Chastity J. Vigil, 2 Chester J. Van Tyne 1 Los Alamos National Laboratory, P.O. Box 1663, Los Alamos, NM USA 2 Colorado School of Mines, 1500 Illinois St., Golden, CO USA Keywords: Induction heating, Austenitization, Steel Abstract The effect of heating rate and prior microstructure on austenitization kinetics has been evaluated for induction hardenable steels with ferrite-pearlite, quench and tempered martensite, or ferritespheroidized carbide initial microstructures. Microstructural scale was varied by initial heattreatment and dilatometry was used to assess austenitization as a function of heating rate and maximum temperature. As-quenched hardness was determined as a function of heating rate and maximum temperature for each microstructural condition. The initial microstructural scale and alloy composition are shown to cause significant variation in the as-quenched microstructure and material properties, particularly for the highest heating rates. Diffusion simulations support the observed microstructural and material property differences between the various initial microstructures and low-alloy and Cr-alloyed steel compositions. Introduction After solidification and homogenization, austenitization is a critical thermal cycle influencing microstructure, and therefore performance. Austenitization is particularly important for induction hardening applications. Differences in initial microstructure have been shown to result in significant variations in final microstructure [1-4] after hardening. We have recently examined the effects of microstructure and heating rate in three induction-hardenable steels [5-10], with the goal of understanding the effect of alloy content and microstructure on properties. Here, we summarize the overall effects of heating rate on low-alloy and Cr-alloyed medium carbon steels, and show results from a hypereutectoid alloy with significant Cr content. Materials and Experimental Procedures Three commercial steel alloys were selected as the experimental materials: 1045, 5150, and Compositions and equilibrium A 3 (or A c for the hypereutectoid composition) temperatures, calculated with ThermoCalc, are presented in Table I. The as-received condition for 1045 and 5150 was hot-rolled, and the was fine spheroidized. For each alloy, half of the as-received material was further heat-treated to produce additional initial conditions: normalized 1045, quench and tempered 5150, and coarse spheroidized The goal of the subsequent heat-treatments was to produce microstructures with different length-scales than the as-received material and to evaluate the effect of microstructural scale using material with the same composition. Thus, a total of six initial microstructural conditions were examined to characterize the austenitization and austenite homogenization response as a function of heating rate. Starting microstructures for each condition are presented in Fig

2 Table I Composition of experimental alloys (wt. pct.) Alloy C Mn Cr P S Si Ni Mo Cu V Al A 3/c, C (a) (c) (e) (b) (d) (f) Figure 1. Starting microstructural conditions: (a) 1045 hot-rolled, (b) 1045 normalized, (c) 5150 hot-rolled, (d) 5150 quench and tempered, (e) fine spheroidized, (f) coarse spheroidized. Optical micrographs, nital etch. Induction dilatometry was used to simulate continuous induction heat-treatments and to evaluate the material response both on-heating and on-cooling. Dilatometry was performed using an MMC quenching and deformation dilatometer with solid cylindrical samples nominally 10 mm in length and 3 mm in diameter with flat and parallel ends. Heat-treatments consisted of heating at an approximately constant heating rate to a maximum temperature. Experimental heating rates used were nominally 0.3, 3.0, 30, and 300 C/s. For the 0.3 and 3.0 C/s heating rates, in order to reduce the overall testing time, samples were heated at a rate of approximately 30 C/s to 700 C, followed by heating at the experimental heating rates. Upon reaching the maximum temperature, the samples were immediately quenched using helium gas. The helium gas quench achieved an 800 to 500 C cooling time of ~2.4 s, which translates to an average cooling rate of approximately 125 C/s. The microstructures of the as-received and experimentally heat-treated samples were evaluated using standard metallographic sampling and preparation techniques. The experimentally heattreated dilatometer cylinders were mounted in epoxy with the axis of the cylinder parallel to the polish plane. The samples were then ground to the axial center plane of the cylinders. Metallographic preparation was performed by rotary grinding with successively finer (60 grit to 800 grit) papers. Ground samples were then polished through 6, 1, and 0.25 μm diamond polishing paste on metallographic preparation cloths. The initial microstructural scale was evaluated by optical image analysis, with details presented in Table II. 106

3 Table II Starting Condition Microstructural Measurements Microstructural Feature 1045 HR 1045 N 5150 HR 5150 QT F C Pearlite lamellar spacing, μm Pearlite colony size, μm Ferrite grain size, μm Ferrite volume fraction Carbide size, μm Carbide volume fraction Knoop (500 g) microhardness indentations were used to determine the bulk hardness of the experimentally heat-treated samples after metallographic preparation. A minimum of five indentations were placed along the axis of each metallographically prepared dilatometer sample, oriented with the long axis of the indentation perpendicular to the axis of the cylinder. Knoop microhardness indentations were performed on at least two identically heat-treated cylinders for each material and heat-treatment condition. Reported microhardness results are an average of all indentations performed on all samples tested for a given heat-treat condition, with error bars indicating the span of the data. Rockwell C values are reported, which have been converted from Knoop 500 g microhardness values using ASTM standard E140. Thermo-calc and DICTRA TM software was used to simulate the microstructure and composition development as a function of maximum temperature and heating rate. Simulation development was based on previous work [11,12] for similar initial microstructures for continuous heating and isothermal heat-treatments. The goal of the diffusion simulations was to evaluate the austenitization and austenite homogenization kinetics, and to provide further understanding of the austenite composition development as a function of heating rate. Two simulation approaches were used, including one for the spheroidized carbide + ferrite microstructure and the quench and tempered microstructure (52100 fine and coarse spheroidized, and 5150 quench and tempered), and another for the ferrite + pearlite microstructures (5150 hot-rolled, and 1045 hotrolled and 1045 normalized). These approaches are discussed in detail in [5], and results from only the ferrite + pearlite simulations are presented here. The compositions evaluated for all cases are for representative values of iron, carbon, chromium, and manganese. Results and Discussion The material response to heat-treatment was measured by as-quenched hardness, as shown in Figure 2. For all continuous heat-treatments, higher heating rates require higher maximum temperatures to reach maximum alloy hardness, which was found to be dependent on alloy carbon level but independent of prior microstructure. With respect to microstructural scale, finer starting microstructures (i.e., 1045 normalized, 5150 quench and tempered, and fine spheroidized) generally required lower maximum temperatures to reach a given hardness. For example, 1045 hot-rolled material required heating to 875 C, even at the slowest heating rate, 0.3 C/s, whereas the 1045 normalized material reached maximum hardness with a peak temperature of 825 C, and was within 1 HRC of maximum with a peak temperature of 775 C. For these types of continuous heat-treatments, this represents a significant difference in thermal cycle required to reach maximum hardness, which increases with increasing heating rates. The 5150 material has higher hardenability, finer starting microstructures, and significantly less ferrite in the starting hot-rolled material than the 1045 material, and thus reached hardnesses above 57 HRC even at the lowest maximum temperatures tested (750 C, which is just above the 107

4 calculated A 3, Table I). Hardness variation between the hot-rolled and quench and tempered starting microstructures were also much less than for the coarser 1045 starting microstructures. The largest hardness differences are found for the highest heating rates. The hypereutectoid material also shows some difference in hardness as a function of starting microstructure at heating rates of 30 /s or more up to approximately 900 C. With increasing maximum temperature, the hardness reaches a maximum of roughly 65 HRC, which corresponds to a martensite carbon content of approximately 0.8 wt. pct. [13]. At heat-treatment temperatures above the peak hardness, the hardness decreases due to increasing amounts of carbon-stabilized retained austenite. The temperature at which the hardness reaches a maximum increases with heating rate, but is independent of the two microstructural scales evaluated here. (a) (c) (e) (b) (d) (f) Figure 2. Rockwell C hardness, converted from Knoop microhardness, versus maximum heattreatment temperature for (a) 1045 hot-rolled, (b) 1045 normalized, (c) 5150 hot-rolled, (d) 5150 quench and tempered, (e) fine spheroidized, and (f) coarse spheroidized. The as-quenched microstructures after heating at 300 C/s to specific maximum temperatures for each alloy, along with selected other heating rate and dilatometry results are presented in Fig. 3. Figures 3a-3c show the significant effect of microstructure on austenitization for the 1045 composition after heating at 300 C/s to 950 C. The hot-rolled starting microstructure (Fig. 3a) results in a much lower hardness (41 vs. 55 HRC) than the normalized starting microstructure (Fig. 3b). Although on-heating dilatometry results [12] from both thermal cycles suggest that the samples were fully austenitized, the on-cooling dilation response (Fig. 3c) and resulting microstructures indicate that the local hardenability in the austenite varied sufficiently to produce substantial regions of non-martensitic transformation products, particularly in the material with the hot-rolled prior microstructure. Figures 3d and 3e highlight microstructural features from the 5150 composition after heating at 300 C/s to 850 C. Relative to the low-alloyed 1045 material, both the hot-rolled and quench and 108

5 tempered microstructures of the 5150 composition appear to be fully austenitized and sufficiently hardenable to produce a fully martensitic microstructure, even though the maximum heattreatment temperature was 100 C lower. This is partially due to the slightly higher carbon content, which lowers the A 3 temperature, but is also a result of the significantly finer starting microstructures. Comparing the hot-rolled to the quench and tempered microstructure for the 5150 material, there appears to be little effect of starting microstructure, since the as-quenched material for both starting microstructures are fully martensitic and equivalent in hardness. However, the hot-rolled material shows remnants of the prior pearlite structure (Fig. 3d), either cementite lamellae that have not been fully dissolved, or substitutional alloying elements that were previously partitioned to the cementite [14] and have not had sufficient time to diffuse away. After performing a quench and temper treatment (austenitize at 830 C, oil quench, 500 C temper) on the hot-rolled 5150, these remnants no longer appear (Fig 3e). Figure 3f shows a region of a nearly fully martensitic microstructure from a 5150 HR sample heated to 775 C at 0.3 C/s. The substructure in the darker region on the left side of the image suggests this has transformed to martensite. Since the dilatometric results [12] show no evidence of nonmartensitic transformation products but do suggest full austenitization temperatures were not reached, the remaining portion of the microstructure is likely retained pearlite. Interestingly, the prior pearlite lamellae are clearly visible in the newly formed martensite, giving insight into the growth of austenite in the prior pearlitic structure, and suggests martensite forms before the cementite is fully dissolved or that the alloy content of the carbides leaves a region with altered etching response. Figures 3g and 3h show the microstructures formed from the fine and coarse spheroidized materials after heating to 1100 C at 300 C/s. As shown in Figure 2, at maximum temperatures of approximately 1000 C and above, the hardness of the material appears to be roughly independent of prior microstructure, and it is difficult to find any significant differences in the microstructures. The microstructures presented in Figs. 3g and 3h show some retained carbides, even though these samples were heat-treated well beyond the maximum hardness temperature for 300 C/s, which would suggest most of the carbon would be in solution. Upon closer inspection, many of the bumps where carbides appear may actually be indications of where the carbides were. We have suggested that these regions may be high in substitutional solid solution alloy elements that diffuse slowly relative to carbon, and continue to show an indication of the prior carbide due to locally higher alloy content, which locally changes the etching response [5,10]. This concept may also apply to the 5150 pearlitic steels, where the etch response indicating retained pearlite lamellae may only be due to increased local alloy content. Further work would be required to confirm local alloy segregation due to prior pearlite regions (5150) or prior spheroidized carbide regions (52100), but thermodynamic and kinetic simulations suggest that is may be plausible [5,10], much like solidification banding that can cause local variations in hardening response during induction processing. This type of local composition gradient may be important when considering induction cycles in alloy material or during the development of precise thermal cycles for the production of advanced high strength steels, which may require higher alloy content. Figure 3f presents a microstructure after heating at 0.3 C/s to 875 C, and shows that carbides near grain boundaries are the most likely to dissolve first, and regions near grain boundaries are also most likely to contain regions of very high carbon content, and therefore, retained austenite. Local variations in composition are likely quite significant in determining the microstructural constituent present at any given location, and it may be problematic to consider solid solution alloying additions such as chromium and manganese as homogenously distributed. 109

6 (a) ((d)) (g) (b) (e) (h) (c) (f) (i) Figure 3. As-quenched microstructures of dilatometer samples with the following material, heating rate, maximum temperature, and hardness combinations: (a) hot-rolled 1045, 300 C/s, 950 C, 41 HRC, (b) 1045 normalized, 300 C/s, 950 C, 55 HRC, (c) offset on-cooling dilation versus temperature for 1045 HR and 1045N samples from (a) and (b), (d) hot-rolled 5150, 300 C/s, 850 C, 60 HRC, (e) quench and tempered 5150, 300 C/s, 850 C, 61 HRC (f) hot-rolled 5150, 30 C/s, 775 C, 57 HRC, (g) fine spheroidized 52100, 300 C/s, 1100 C, 60 HRC, (h) coarse spheroidized 52100, 300 C/s, 1100 C, 60 HRC, (i) coarse spheroidized 52100, 0.3 C/s, 875 C, 62HRC. Secondary electron images, nital etch, except (g) and (h), picral etch. Figure 4 presents results from DICTRATM simulations for the hot rolled 1045 starting microstructure. Figures 4(a-c) show continuous heating simulations of a prior pearlite (now austenite) region next to a ferrite region to determine the kinetics of ferrite dissolution. The approach of assuming prior pearlite regions transform almost immediately to austenite seems reasonable when considering the results presented in Figs. 4(d-f), which show that the dissolution of a pearlitic cementite lamella requires very little time. Although the alloy content is low in the 1045 material, these calculations support the concept that, while carbon may be quick to diffuse over substantial microstructural distances during rapid heating cycles, solid solution alloying elements require significantly longer time and may be inhomogenously distributed due to prior coarse pearlite or coarse spheroidized alloy carbide microstructures. In these simulations, the cementite, prior pearlite, and ferrite compositions were calculated based on equilibrium formation during cooling. 110

7 (a) (b) (c) (d) (e) (f) Figure 4. Results from DICTRA TM simulations based on the 1045 hot rolled starting microstructure. (a-c) Carbon, chromium, and manganese diffusion from austenite to ferrite (transforming to austenite) regions as a function of time above 790 C and a heating rate of 300 C/s. Dashed lines indicate the initial austenite ferrite interface. The austenite region represents the prior pearlite, which transforms to austenite more quickly than prior ferrite regions. (d-f) Carbon, chromium, and manganese diffusion from cementite lamellae to ferrite (transforming to austenite) simulating the pearlite transformation process. Summary and Conclusions The austenitization of three induction hardenable steels with variations in initial microstructure scale was evaluated to determine the effect of heating rate on microstructure and properties for continuous heat-treatments. Controlled dilatometric heat-treatments were performed to simulate thermal cycles with heating rates from 0.3 C/s to 300 C/s. Hardness measurements provided an indication of the heat-treatment response of each material-microstructure combination. Examination of as-quenched microstructures provided some insight into the condition of the austenite immediately prior to quenching. As might be expected, there are significant effects of starting microstructure, particularly for microstructures with the longest length scales. Thermodynamic and kinetic simulations suggested that substitutional solid solution alloying elements may not be homogenously distributed within a given starting microstructure, and highrate heat-treatments do not provide sufficient time for substantial homogenization during an austenitization and quench cycle. Based on the results presented, the following conclusions can be reached: 1. Starting microstructure scale can have a significant effect on subsequent high heating rate thermal treatments. Increasing microstructural scale in continuous heat-treatments may require substantially higher maximum temperatures to achieve equivalent hardness and similar final microstructures. 111

8 2. Local composition gradients may exist in as-quenched materials, particularly in alloys that contain significant substitutional solid solution alloying elements. These composition gradients may occur due to prior microstructural features, such as enriched alloy carbides or ferrite with depleted alloy composition. 3. The heating rates examined here, up to 300 C/s, were intended to be relevant for induction processing of medium-carbon steels. However, these rates are likely also relevant to the complex thermal cycles that may be required to process the next generation of advanced high strength steels. For these complex thermal cycles, consideration of prior microstructural scale and local composition variation may be important. Acknowledgements KDC and CJVT gratefully acknowledge the support from the sponsors of the Advanced Steel Processing and Products Research Center (ASPPRC), an industry/university cooperative research center. KDC, AJC, REH, and CJV gratefully acknowledge support from Los Alamos National Laboratory, operated by Los Alamos National Security, LLC under Contract No. DE-AC52-06NA25396 for the U.S. Department of Energy. KDC and CJVT gratefully acknowledge the Timken Co. for supplying material for this study. References [1] Wever, F., Rose, A., Peter, W., Strassburg, W., Rademacher, L., Atlas zur Wärmebehandlung der Stähle, Band 1, 1961, Düsseldorf, Germany. [2] Rose, A., Hougardy, H., Atlas zur Wärmebehandlung der Stähle, Band 2, 1972, Düsseldorf, Germany. [3] Orlich, J., Rose, A., Wiest, P., Atlas zur Wärmebehandlung der Stähle, Band 3, 1973, Düsseldorf, Germany. [4] Orlich, J., Pietrzeniuk, H-J., Atlas zur Wärmebehandlung der Stähle, Band 4, 1976, Düsseldorf, Germany. [5] K.D. Clarke, Ph.D. Thesis, Colorado School of Mines, Golden, CO, [6] S.-J. Lee, K.D. Clarke, Met. and Mat. Trans. A, vol. 41, no. 12, Dec. 2010, pp [7] S.-J. Lee, K.D. Clarke, C.J. Van Tyne, Met. and Mat. Trans. A, vol. 41, no. 9, Sep. 2010, pp [8] K.D. Clarke, C.J. Van Tyne, R.E. Hackenberg, C.J. Vigil, Journal of Materials Engineering and Performance, vol. 20, no. 2, March, 2011, pp [9] C.J. Van Tyne, K.D. Clarke, ASM Induction Handbook, article 6B, Materials Park, OH: ASM International, June [10] K.D. Clarke, Comprehensive Materials Processing, Volume 12: Thermal Engineering of Steel and Other Alloy Systems, George Krauss, Editor, Elsevier, April [11] Katsamas, A.I., A Computational Study of Austentite Formation Kinetics in Rapidly Heated Steels. Surface and Coatings Technology, : p [12] Liu, Z.K., Ågren, J., Morphology of Cementite Decomposition in an Fe-Cr-C Alloy. Metallurgical Transactions A, : p [13] Kelly, P.M., Nutting, J., Proc. of the Royal Society of London. Series A, Mathematical and Physical Sciences, (1296): p [14] A.J. Clarke, D.R. Coughlin, M.K. Miller, R.D. Field, K.D. Clarke, K.A. Powers, P.J.Gibbs, P.A. Papin, V.K. Judge, D.T. Pierce, D.J. Alexander, E. De Moor, J.G. Speer, G. Krauss Microstructural evolution in 4340 steel after quenching and tempering, International Conference on Solid-Solid Phase Transformations in Inorganic Materials, June 28-July 3, 2015, Whistler, BC, Canada. 112