Volume 04 - Heat Treating Part 5 docx

Fig. 25 Strain- induced precipitation of NbCN in austenite in a steel containing 0.9 wt% C and 0.07 wt% Nb. Specimen was reheated at 1250 °C (2280 °F), rolled 25%, and held a t 950 °C (1740 °F), and then air cooled to room temperature. Centered dark field electron micrograph was obtained using a (111) NbC zone axis. Source: Ref 74 Effect of Niobium in Suppressing Recrystallization. The second question comparing the suppression of recrystallization by solute and precipitate is addressed in Table 5. In this case, the time to 20% static softening was measured for a variety of conditions, including those where niobium was largely in solution and those where niobium was present as a precipitate. The data for the strain rate of 10 s -1 are perhaps the more interesting because they have been generated under conditions approximating industrial plate rolling. When the niobium is in solution, the measured delay was less than 1.5 s at 1000 °C (1830 °F), whereas the delay caused by precipitates was on the order of several minutes. Furthermore, a comparison of the curves presented in Fig. 23 and 24 indicates that the precipitation-start time at 900 °C (1650 °F) corresponds to about 34% recrystallization in the niobium-free reference steel. A close examination of Fig. 24 reveals that this means the niobium in solution has retarded recrystallization for about 8 s at 900 °C (1650 °F) in the niobium steel. Table 5 Times required for 20% softening in three selected steels at strain rates of 0.1 and 10 s -1 Temperature Time to 20% static softening (t 20% ), s Strain rate (ε), s -1 °C °F Alloy K-0 Alloy K-1 Alloy K-2 900 1650 5 2000 (a) 3000 (a) 1000 1830 2 15 200 (a) 0.1 1100 2010 0.6 3 3.5 900 1650 0.2 300 (a) 1000 (a) 10 1000 1830 0.1 1 1.5 Source: Ref 69 (a) Condition where extensive precipitation occurs. Recrystallization Controlled Rolling An alternate form of thermomechanical processing exists in addition to controlled rolling. This new process is called recrystallization controlled rolling. The metallurgical basis of the RCR process and how it is distinguished from the CCR or CHR processes and intensified controlled rolling (ICR) can be understood with the help of Fig. 22. Whereas there is a clear distinction in processing (that is, finish rolling temperatures) between recrystallization controlled rolling and conventional controlled rolling, the difference between recrystallization controlled rolling and conventional hot rolling is not initially obvious because they both rely on high-temperature deformation. There are in fact two major differences between recrystallization controlled rolling and conventional hot rolling. First, all of the deformation in recrystallization controlled rolling is completed in the region of full recrystallization. This is in contrast to the CHR method, where a portion of the deformation takes place in the region of partial recrystallization. Second, only certain steels are amenable to the RCR method since a mechanism to inhibit grain coarsening must be present in the steel. Alloys not configured for recrystallization controlled rolling undergo static recrystallization following each high-temperature deformation. This is followed by rapid grain coarsening in the remaining interpass time. Hence, this leads to coarse, as- rolled austenite grain sizes. This austenite structure is characterized by a small S v and is therefore indicative of poor thermomechanical processing. The central feature of steels exhibiting RCR behavior is the presence of a grain coarsening inhibition system. Recent research has indicated that this inhibition system can take one of two forms (Ref 44, 94, 96). The first relies on the presence of stable particles whose distribution provides an adequate pinning force to suppress coarsening following post-deformation static recrystallization. These particles must have sufficient stability to resist Ostwald ripening (solution-crystallizer phenomenon in which small crystals, more soluble than large ones, dissolve and reprecipitate onto larger particles) themselves and must be present in a sufficiently fine dispersion. As was shown above, the pinning force exerted by an array of particles on a boundary is given by Equation 18 assuming a rigid boundary model. As was shown by Ashby (Ref 65), the driving force for normal grain coarsening is nearly three orders of magnitude smaller than that for recrystallization. Hence, the particle pinning force required to suppress coarsening is much smaller than that required to suppress recrystallization. Because the particle pinning force varies with f v /r, even particle distributions with rather small f v and large r can still be effective barriers to grain coarsening. This, of course, would not be true in the suppression of recrystallization, where only large, local volume fractions of fine particles are effective. Particles such as titanium nitride have been shown to be particularly well suited for the suppression of grain coarsening (Ref 44, 96). In this regard, the titanium nitride particles are formed via solid state precipitation immediately following solidification and, hence, exist prior to deformation. An alternate grain coarsening inhibition mechanism has recently been discovered (Ref 110). This mechanism involves intense solute drag as the critical element. It has been found that high levels of soluble molybdenum and niobium can be very effective in retarding boundary motion at high reheating temperatures (see Fig. 26). Hence, these solute effects act to retard grain coarsening both during reheating and after post-deformation static recrystallization. In summary, recrystallization controlled rolling offers an attractive alternative to conventional controlled rolling in several applications where low-temperature rolling is impractical (for example, older underpowered plate mills and hot forgings) (Ref 112, 113). That recrystallization controlled rolling can lead to excellent austenite conditioning is evidenced by values of S v well in excess of 100 mm -1 which can be easily realized (Ref 11). Fig. 26 Effect of molybdenum addition on austenite grain size a fter reheating at 1150 °C (2100 °F) and holding at this temperature for various times. Source: Ref 111 Intensified Controlled Rolling A variation in conventional controlled rolling, known as intensified controlled rolling, is shown in Fig. 22. Note that intensified controlled rolling differs from conventional controlled rolling in that finish rolling not only occurs below the T RXN , but also extends to temperatures below the Ar 3 . The goal of intensified controlled rolling is to increase the strength and toughness of the as-rolled MA steel over what could be achieved through conventional controlled rolling. The combination of the lower reheat and lower rolling temperatures leads to finer as-rolled austenite. Furthermore, because a portion of the rolling occurs in the intercritical or two-phase (α + γ) region, a certain amount of proeutectoid ferrite would be deformed during this rolling. This deformed ferrite is partially responsible for the higher strengths observed with intensified controlled rolling (Ref 93, 114). High Toughness Rolling Processes Heat treatment processes incorporating high toughness rolling have been developed by Sumitomo and Broken Hill Proprietary. Sumitomo High Toughness Rolling. An additional rolling practice shown in Fig. 22 is known as the Sumitomo high toughness (SHT) process (Ref 93, 115, 116). The goal of the SHT process is to produce steels with very high resistance to brittle fracture at the expense of strength when compared to a steel processed by the CCR process. In the SHT process, extra grain refinement of the austenite is achieved by the introduction of a cooling and low-temperature reheating cycle positioned between rough and finish rolling. The finish rolling is performed in the intercritical region as with intensified controlled rolling. Broken Hill Proprietary High Toughness Rolling. There are some conditions where neither the CCR process nor the ICR process will achieve a resistance to brittle fracture in plate steels. Such conditions exist when the plate thickness exceeds ≈ 18 mm ( ≈ 23 32 in.) or when the plate width causes the mill loads to exceed acceptable levels. These shortcomings have been overcome by Broken Hill Proprietary (BHP) through a variation of both the alloy design and controlled rolling practice of their plate steels intended for line pipe and offshore applications (Ref 117, 118, 119). During conventional controlled rolling, the total reduction is fairly evenly distributed between the roughing and finishing passes. In the BHP process, a larger portion of the deformation takes place in the roughing passes and, consequently, a smaller portion in the finishing passes. Hence, while the low-temperature deformation in conventional controlled rolling and intensified controlled rolling is used to generate high values of S v through extensive pancaking, the high-temperature passes and the associated recrystallization and grain refinement are responsible for the high S v found in the BHP process. This emphasis on high-temperature rolling is beneficial for the rolling of both heavy sections and large widths. Because the BHP process involves high-temperature rolling, the phenomenon of grain coarsening must be avoided. This is accomplished by a microtitanium addition to a niobium MA steel. The titanium nitride that forms is effective in suppressing grain coarseing. References cited in this section 11. C.I. Garcia, A.K. Lis, and A.J. DeArdo, 31st Mechanical Working and Steel Processing Conference, Iron and Steel Society/AIME, 1989, p 505 21. A. Arrowsmith, J. Iron Steel Inst., Vol 110, p 317 44. Y. Zheng, G. Fitzsimons, and A.J. DeArdo, HSLA Steels: Technology and Applications, American Society for Metals, 1984, p 85 53. L.J. Cuddy, Thermomechanical Processing of Microalloyed Austenite, A.J. DeArdo et al., Ed., The Metallurgical Society of AIME, 1982, p 129 65. M.F. Ashby, Recrystallization and Grain Growth of Multi- Phase and Particle Containing Materials: 1st RISO International Symposium on Metallurgy and Material Science, N. Hansen, A.R. Jones, and T. Leffers, Ed., RISO National Laboratory, Roskilde, Denmark, 1980, p 325 68. S.S. Hansen, J.B. Vander Sande, and M. Cohen, Metall. Trans. A, Vol 11A, 1980, p 387 69. O. Kwon, Ph.D. thesis, University of Pittsburgh, 1985, p 208 74. A.J. DeArdo, J.M. Gray, and L. Meyer, Niobium, H. Stuart, Ed., The Metallurgical Society of AIME, 1984, p 685 81. E.J. Palmiere, Ph.D. thesis, University of Pittsburgh, 1991 93. H. Gondoh and T. Osuka, Niobium, H. Stuart, Ed., The Metallurgical Society of AIME, 1984, p 833 94. T. Siwecki et al., Thermomechanical Processing of Microalloyed Austenite, A.J. DeArdo, G.A. Ratz, and P.J. Wray, Ed., The Metallurgical Society of AIME, 1982, p 163 95. Sekine et al., Thermomechanical Processing of Microalloyed Austenite, A.J. DeArdo, G.A. Ratz, and P.J. Wray, Ed., The Metallurgical Society of AIME, 1982, p 141 96. R.M. Fix, Y.Z. Zheng, and A.J. DeArdo, HSLA Steels: Technology and Applications, American Society for Metals, 1984, p 219 97. J.J. Jonas, C.M. Sellars, and W.J. McG. Tegart, Met. Rev., Vol 14, 1969, p 1 98. H.J. McQueen and J.J. Jonas, Plastic Deformation of Materials, R.J. Arsenault, Ed., Academic Press, 1975, p 393 99. J.D. L'Ecuyer and G. L'Esperance, Acta Metall., Vol 37, 1989, p 1023 100. J.J. Jonas, International Confer ence on Physical Metallurgy of Thermomechanical Processing of Steel and Other Metals, I. Tamura, Ed., The Iron and Steel Institute of Japan, 1988, p 59 101. O. Kwon and A.J. DeArdo, Acta Metall., 1990, in press 102. A.J. DeArdo, Mathematical Modelling of Hot Rolling of Steel, S. Yue, Ed., CIM, 1990, p 220 103. M.J. Luton, R. Dorvel, and R.A. Petkovic, Metall. Trans. A, Vol 11A, 1980, p 411 104. A.B. Rothwell, Mem. Sci. Rev. Met., Vol 69, 1972, p 413 105. T. Greday and M. Lamberigts, The Hot Deformation of Austenite, J. Ballance, Ed., The Metallurgical Society of AIME, 1976, p 75 106. B.L. Phillippo and F.A.A. Crane, J. Iron Steel Inst., Vol 211, 1973, p 653 107. G.A. Wilber, J.R. Bell, J.H. Bucher, and W.S. Childs, Trans. AIME, Vol 242, 1968, p 2305 108. P.L. Mangonon, Jr. and W.E. Heitmann, International Conference on Steel Rolling, The Iron and Steel Institute of Japan, Tokyo, 1980, p 59 109. K. Relander and E. Tyni, Low Alloy High Strength Steels, Frese-Druck, Düsseldorf, 1970, p 81 110. A.J. DeArdo, HSLA Steels, The Chinese Society of Metals, Beijing, 1990, in press 111. C.I. Garcia and A.J. DeArdo, University of Pittsburgh, unpublished research, 1991 112. T. Siwecki, Microalloyed Vanadium Steels, M. Korchynsky et al., Ed., Association of Polis h Metallurgical Engineers, Krakow, Poland, 1990, p 63 113. J.R. Paules, 31st Mechanical Working and Steel Processing Conference, Iron and Steel Society/AIME, 1989, p 131 114. T. Hashimoto et al., Thermomechanical Processing of Microalloyed Austenite, A.J . DeArdo, G.A. Ratz, and P.J. Wray, Ed., The Metallurgical Society of AIME, 1982, p 501 115. H. Takeuchi et al., Steel Rolling, Vol 2, The Iron and Steel Institute of Japan, 1980, p 957 116. T. Tanaka et al., The Sumitomo Search, No. 19, May, 1978 117. C.R. Killmore, G.R. Harris, and J.G. Williams, High Strength Low Alloy Steels, D.P. Dunne and T. Chandra, Ed., South Coast Printers, Port Kembla, Australia, 1985, p 57 118. J.G. Williams et al., HSLA Steels: Metallurgy and Applications, J.M. Gray et al., Ed., American Society for Metals, 1986, p 567 119. R.H. Phillips, 3.6. Williams, and J.E. Croll, Microalloyed HSLA Steels, ASM International, 1988, p 235 Applications of Thermomechanical Processing to Heat Treat Low-Alloy Steels Up until this point, the benefits of thermomechanical processing have been examined for as-hot rolled strip and plate microalloyed steels. By definition, these steels contain less than about 2 wt% in total alloying and usually exhibit F-P microstructures. Recently, the benefits of thermomechanical processing have been extended to the low-alloy steels, including the HY steels, the HSLA steels, the ultralow carbon bainitic steels, and the multiphase steels. These steels differ from the as-rolled microalloyed steels in two ways: • Steels contain >3 wt% in total alloying • Steels may be heat treated after rolling These steels are often bainitic, martensitic, or multiphase in microstructure. HSLA Steels. Thermomechanical processing has been shown to greatly improve the resistance to brittle fracture in the new HSLA-80 and HSLA-100 steels (Ref 120). Figure 27 compares the notch toughness of quenched and aged steels that had been either hot rolled or controlled rolled prior to heat treatment. The benefit of the controlled rolling is obvious. Note that the benefits derived from the controlled rolling existed even after the post-rolling reheating treatment. Fig. 27 Comparison of dynamic tear toughness at -40 °C (-40 °F) of 25 mm (1 in.) thick HSLA- 80 plates produced either by controlled rolling or conventional rolling prior to reheating, water quenching, and aging. Source: Ref 120 ULCB Steels. Austenite conditioning has also been shown to be critical in achieving low-temperature toughness in as-hot rolled ULCB steels (Ref 11, 121). A study of the factors that control the resistance to brittle fracture in these steels indicated the importance of the prior austenite grain size (see Fig. 4). In a similar study (Ref 121), an important relationship was found between low-temperature toughness and the S v of the controlled rolled austenite that existed prior to transformation. While S v improves the toughness of F-P steels through the refinement of the ferrite grain size, the role of S v in ULCB steels is somewhat different. As was shown in a later study (Ref 11), the only significant barrier to a growing cleavage crack in a high-strength bainitic matrix is the prior austenite grain boundary. Hence, the higher the S v , the smaller the growth phase of the cleavage cracking event. It was shown that when these cleavage cracks could be kept relatively small, their existence did not preclude an otherwise high-toughness value at low temperatures. Multiphase steels intended for long product applications have also benefitted from thermomechanical processing. Low- carbon steels based on the manganese-molybdenum-niobium system have been shown to exhibit ferrite-lower bainite microstructures after controlled rolling. These steels display excellent ductility and work hardening characteristics and, hence, are ideal steels for cold forging applications. In this case, the controlled rolling not only leads to grain refinement but also to establishing the proper balance between the phases in the final microstructure (Ref 11, 122, 123). This same steel exhibits austenite conditioning via recrystallization controlled rolling during high-temperature forging processing. In the as-direct quenched condition, these forgings display unusually high levels of strength, toughness, and fatigue resistance (Ref 124). References cited in this section 11. C.I. Garcia, A.K. Lis, and A.J. DeArdo, 31st Mechanical Working and Steel Processing Conference, Iron and Steel Society/AIME, 1989, p 505 120. A.D. Wilson et al., Microalloyed HSLA Steels, ASM International, 1988, p 259 121. C.I. Garcia and A.J. DeArdo, Microalloyed HSLA Steels, ASM International, 1988, p 291 122. C.I. Garcia, A.K. Lis, and A.J. DeArdo, Wire Association international Conference Proceedings, Guilford, CT, 1990, in press 123. C.I. Garcia, A.K. Lis, and A.J. DeArdo, Trans. Iron Steel Soc. AIME, in press 124. C.I. Garcia, A.K. Lis, and A.J. DeArdo, Wire J. Int., in press Future Outlook The thermomechanical processing of austenite is often associated with significant improvements in final properties and offers a degree of flexibility in achieving various packages of properties that are quite impressive. While the elements of thermomechanical processing (reheating, hot rolling, and cooling) are easily distinguishable and individually amenable to thorough investigation, the response of a given MA steel to a given processing step will, to a large measure, depend on what had occurred in previous stages. It is this interdependence that makes the rationalization of commercial hot rolling so very complex and challenging. This is specially true in the MA steels, where the dissipation or retention of microalloy solute supersaturation in a given process step might have drastically different effects on microstructural evolution in subsequent steps. There is little doubt that thermomechanical processing offers a very cost-effective route to good properties in steel. It is interesting to consider where the concepts of thermomechanical processing and microalloying technology may lead us in the future. Introduction to Surface Hardening of Steels S. Lampman, ASM International Introduction SURFACE HARDENING, a process which includes a wide variety of techniques (Table 1), is used to improve the wear resistance of parts without affecting the more soft, tough interior of the part. This combination of hard surface and resistance to breakage upon impact is useful in parts such as a cam or ring gear that must have a very hard surface to resist wear, along with a tough interior to resist the impact that occurs during operation. Further, the surface hardening of steel has an advantage over through hardening because less expensive low-carbon and medium-carbon steels can be surface hardened without the problems of distortion and cracking associated with the through hardening of thick sections. Table 1 Engineering methods for surface hardening of steels Layer Additions Hardfacing Fusion hardfacing (welded overlay) Thermal spray (nonfusion bonded overlay) Coatings: Electrochemical plating Chemical vapor deposition (electroless plating) Thin films (physical vapor deposition, sputtering, ion plating) Ion mixing Substrate treatment Diffusion methods: Carburizing Nitriding Carbonitriding Nitrocarburizing Boriding Titanium-carbon diffusion Toyota diffusion process Selective hardening methods: Flame hardening Induction hardening Laser hardening Electron beam hardening Ion implantation Selective carburizing and nitriding Use of arc lamps There are two distinctly different approaches to the various methods for surface hardening (Table 1): • Methods that involve an intentional buildup or addition of a new layer • Methods that involve surface and subsurface modification without any intentional buildup or increase in part dimensions The first group of surface hardening methods includes the use of thin films, coatings, or weld overlays (hardfacings). Films, coatings, and overlays generally become less cost effective as production quantities increase, especially when the entire surface of workpieces must be hardened. The fatigue performance of films, coatings, and overlays may also be a limiting factor, depending on the bond strength between the substrate and the added layer. Fusion-welded overlays have strong bonds, but the primary surface-hardened steels used in wear applications with fatigue loads include heavy case- hardened steels and flame- or induction-hardened steels. Nonetheless, coatings and overlays can be effective in some applications. With tool steels, for example, TiN and Al 2 O 3 coatings are effective not only because of their hardness but also because their chemical inertness reduces crater wear and the welding of chips to the tool. Overlays can be effective when the selective hardening of large areas is required. This introductory article on surface hardening focuses exclusively on the second group of methods, which is further divided into diffusion methods and selective hardening methods (Table 1). Diffusion methods modify the chemical composition of the surface with hardening species such as carbon, nitrogen, or boron. Diffusion methods allow effective hardening of the entire surface of a part and are generally used when a large number of parts are to be surface hardened. In contrast, selective surface hardening methods allow localized hardening. Selective hardening generally involves transformation hardening (from heating and quenching), but some selective hardening methods (selective nitriding, ion implantation and ion beam mixing) are based solely on compositional modification. Factors affecting the choice of these surface hardening methods are discussed in the section "Process Selection" in this article. Diffusion Methods of Surface Hardening As previously mentioned, surface hardening by diffusion involves the chemical modification of a surface. The basic process used is thermochemical because some heat is needed to enhance the diffusion of hardening species into the surface and subsurface regions of a part. The depth of diffusion exhibits a time-temperature dependence such that: Case depth ∝ K time (Eq 1) where the diffusivity constant, K, depends on temperature, the chemical composition of the steel, and the concentration gradient of a given hardening species. In terms of temperature, the diffusivity constant increases exponentially as a function of absolute temperature. Concentration gradients depend on the surface kinetics and reactions of a particular process. Methods of hardening by diffusion include several variations of hardening species (such as carbon, nitrogen, or boron) and of the process method used to handle and transport the hardening species to the surface of the part. Process methods for exposure involve the handling of hardening species in forms such as gas, liquid, or ions. These process variations naturally produce differences in typical case depth and hardness (Table 2). Factors influencing the suitability of a particular diffusion method include the type of steel (Fig. 1), the desired case hardness (Fig. 2), and the case depth (Fig. 3). Table 2 Typical characteristics of diffusion treatments Process Nature of case Process temperature, Typical case Case hardness, Typical base Process characteristics °C (°F) depth HRC metals Carburizing Pack Diffused carbon 815-1090 (1500- 2000) 125 μm- 1.5 mm (5-60 mils) 50-63 (a) Low-carbon steels, low- carbon alloy steel Low equipment costs, difficult to control case depth accurately Gas Diffused carbon 815-980 (1500- 1800) 75 μm- 1.5 mm (3-60 mils) 50-63 (a) Low-carbon steels, low- carbon alloy steels Good control of case depth, suitable for continuous operation, good gas controls required, can be dangerous Liquid Diffused carbon and possibly nitrogen 815-980 (1500- 1800) 50 μm- 1.5 mm (2-60 mils) 50-65 (a) Low-carbon steels, low- carbon alloy steels Faster than pack and gas processes, can pose salt disposal problem, salt baths require frequent maintenance Vacuum Diffused carbon 815-1090 (1500- 2000) 75 μm- 1.5 mm (3-60 mils) 50-63 (a) Low-carbon steels, low- carbon alloy steels Excellent process control, bright parts, faster than gas carburizing, high equipment costs Nitriding Gas Diffused nitrogen, nitrogen compounds 480-590 (900- 1100) 125 μm- 0.75 mm (5-30 mils) 50-70 Alloy steels, nitriding steels, stainless steels Hardest cases from nitriding steels, quenching not required, low distortion, process is slow, is usually a batch process Salt Diffused nitrogen, nitrogen compounds 510-565 (950- 1050) 2.5 μm- 0.75 mm (0.1- 30 mils) 50-70 Most ferrous metals including cast irons Usually used for thin hard cases <25 μm (1 mil), no white layer, most are proprietary processes Ion Diffused nitrogen, nitrogen compounds 340-565 (650- 1050) 75 μm- 0.75 mm (3-30 mils) 50-70 Alloy steels, nitriding, stainless steels Faster than gas nitriding no white layer, high equipment costs, close case control Carbonitriding Gas Diffused carbon and nitrogen 760-870 (1400- 1600) 75 μm- 0.75 mm (3-30 mils) 50-65 (a) Low-carbon steels, low- carbon alloy steels, stainless steel Lower temperature than carburizing (less distortion), slightly harder case than carburizing gas control critical Liquid (cyaniding) Diffused carbon and nitrogen 760-870 (1400- 1600) 2.5-125 μm (0.1- 5 mils) 50-65 (a) Low-carbon steels Good for thin cases on noncritical parts, batch process, salt disposal problems Ferritic nitrocarburizing Diffused carbon and nitrogen 565-675 (1050- 1250) 2.5-25 μm (0.1- 1 mil) 40-60 (a) Low-carbon steels Low-distortion process for thin case on low-carbon steel, most processes are proprietary Other Aluminizing (pack) Diffused aluminum 870-980 (1600- 1800) 25 μm-1 mm (1-40 mils) <20 Low-carbon steels Diffused coating used for oxidation resistance at elevated temperatures Siliconizing by chemical vapor deposition Diffused silicon 925-1040 (1700- 1900) 25 μm-1 mm (1-40 mils) 30-50 Low-carbon steels For corrosion and wear resistance, atmosphere control is critical Chromizing by chemical vapor deposition Diffused chromium 980-1090 (1800- 2000) 25-50 μm (1-2 mils) Low-carbon steel, <30; high-carbon steel, 50-60 High- and low- carbon steels Chromized low-carbon steels yield a low-cost stainless steel, high-carbon steels develop a hard corrosion- resistant case Titanium carbide Diffused carbon and titanium, TiC compound 900-1010 (1650- 1850) 2.5-12.5 μm (0.1- 0.5 mil) >70 (a) Alloy steels, tool steels Produces a thin carbide (TiC) case for resistance to wear, high temperature may cause distortion Boriding Diffused boron, boron, compound 400-1150 (750- 2100) 12.5-50 μm (0.5- 2 mils) 40->70 Alloy steels, tool steels, cobalt and nickel alloys Produces a hard compound layer, mostly applied over hardened tool steels, high process temperature can cause distortion Source: Ref 1 (a) Requires quench from austenitizing temperature. [...]... fuel gas Acetylene 53 .4 1433 31 05 5620 23 25 42 15 1.0 26.7 716 53 5 21 14,284 15, 036 12 City gas 11.233 .5 300900 254 0 4600 19 85 36 05 (b) (b) (b) (b) (b) (b) (b) (b) Natural gas (methane) 37.3 1000 27 05 4900 18 75 34 05 1. 75 13.6 364 280 11 3,808 4, 004 9.0 Propane 93.9 252 0 26 35 47 75 19 25 34 95 4.0 18.8 50 4 3 05 12 5, 734 6 ,048 25. 0 MAPP 90 2406 2927 53 01 1760 3200 3 .5 20.0 53 5 381 15 7,620 8,0 25 22 (a) Product... G Parrish, The Influence of Microstructure on the Properties of Case-Carburized Components, American Society for Metals, 1980, p 15 9-1 60, 16 4-1 65 Flame Hardening of Steels Revised by Thomas Ruglic, Hinderliter Heat Treating, Inc Introduction FLAME HARDENING is a heat- treating process in which a thin surface shell of a steel part is heated rapidly to a temperature above the critical point of the steel... International, 1990, p 5- 1 2 5 B Edenhofer, M.H Jacobs, and J.N George, Industrial Processes, Applications and Benefits of Plasma Heat Treatment, in Plasma Heat Treatment, Science and Technology, PYC Édition, 1987, p 39 9-4 15 6 W.L Grube and J.G Gay, High-Rate Carburizing in a Glow-Discharge Methane Plasma, Metall Trans A, Vol 91, 1987, p 142 1-1 429 7 K.-E Thelning, Steel and Its Heat Treatment, 2nd ed.,... cylinders Depth of Heating Shallow hardness patterns (less than 3.2 mm, or 0.1 25 in., deep) can be attained only with oxy-gas fuels The high-temperature flames obtained with oxy-gas fuels provide the fast heat transfer necessary for effective localization of the heat pattern Deeper hardness patterns permit the use of either oxy-gas fuels or air-gas fuels Oxy-gas fuels will localize the heat, but care is... controlled by the quench rather than by the heating The deeper-seated heat produced by air-gas flames may preclude the use of air-gas mixtures because excessive distortion may occur In consideration of these factors, the use of air-gas heating will depend primarily on the shape of the part insofar as the configuration favors heat localization and a lower rate of heat transfer Gas Consumption, Time, and... hardened Fig 9 Setups for flame-hardening gears, idler wheels, and sprockets (a) Radiant burners (b) High-velocity convection burners Wide-face parts can be heated with double or staggered rings of burners The high-velocity convection burner is basically a miniature refractory-lined furnace in which heat is released at rates as high as 4 15 MJ/m3 · s (4 × 107 Btu/ft3 · h) The air-gas premixture, supplied... Engineers, 19 85 5 B Edenhofer, M.H Jacobs, and J.N George, Industrial Processes, Applications and Benefits of Plasma Heat Treatment, in Plasma Heat Treatment, Science and Technology, PYC Édition, 1987, p 39 9-4 15 11 P.A Molian, Engineering Applications and Analysis of Hardening Data for Laser Heat Treated Ferrous Alloys, Surf Eng., Vol 2, 1986, p 1 9-2 8 12 R Zenker and M Mueller, Electron Beam Hardening, Part. .. p 67 5- 6 90 18 A.H Deutchman et al., Ind Heat. , Vol 42 (No.1), Jan 1990, p 3 2-3 5 Process Selection The benefits of the most common methods of surface hardening are compared in Table 3 Flame and induction hardening are generally limited to certain families of steels such as medium-carbon steels, medium-carbon alloy steels, some cast irons, and the lower-alloy tool steels There is no size limit to parts... 4 Comparison of heating times for MAPP, acetylene, and propane Flame velocity, 170 m/s (55 0 ft/s); port size, No 69 drill (0.74 mm, or 0.0292 in.); coupling distance, 9 .5 mm ( to-fuel ratios: MAPP, 5. 0; acetylene, 1.33; propane, 4 .5 3 in.); material, 1036 steel Oxygen8 The ratio of oxygen to fuel is very important in obtaining maximum heating efficiency from the fuel However, oxygento-fuel ratios should... apparatus that automatically indexes, heats, and quenches parts Large parts such as gears and machine tool ways, with sizes or shapes that would make furnace heat treatment impractical, are easily flame hardened With improvements in gas-mixing equipment, infrared temperature measurement and control, and burner design, flame hardening has been accepted as a reliable heat- treating process that is adaptable . nitrocarburizing Diffused carbon and nitrogen 56 5- 6 75 (1 05 0- 1 250 ) 2. 5- 2 5 μm (0. 1- 1 mil) 4 0-6 0 (a) Low-carbon steels Low-distortion process for thin case on low-carbon steel, most processes are. nitrogen compounds 51 0 -5 65 ( 95 0- 1 050 ) 2 .5 μm- 0. 75 mm (0. 1- 30 mils) 5 0-7 0 Most ferrous metals including cast irons Usually used for thin hard cases < 25 μm (1 mil), no white. chromium 98 0-1 090 (180 0- 2000) 2 5- 5 0 μm ( 1-2 mils) Low-carbon steel, <30; high-carbon steel, 5 0-6 0 High- and low- carbon steels Chromized low-carbon steels yield a low-cost stainless