The Effect of Cold Rolling on the Grain Boundary Character and Creep Rupture Properties of INCONEL alloy 718

Transcription

1 Journal Citation (to be inserted by the publisher) Copyright by Trans Tech Publications The Effect of Cold Rolling on the Grain Boundary Character and Creep Rupture Properties of INCONEL alloy 718 C.J. Boehlert 1, S. Civelekoglu 2, N. Eisinger 3, G. Smith 3, and J. Crum 3 1 Alfred University, School of Ceramic Engineering and Materials Science, 2 Pine Street, Alfred, NY 14802, USA, boehlecj@alfred.edu 2 Alfred University, School of Ceramic Engineering and Materials Science, 2 Pine Street, Alfred, NY 14802, USA, sc3@alfred.edu 3 Special Metals Corporation, 3200 Riverside Drive, Huntington, WV 25705, USA Keywords: superalloy, microstructure, electron backscatter diffraction, creep. Abstract. In order to evaluate the effects of sheet processing on the grain boundary character distribution (GBCD) of INCONEL alloy 718 (IN 718), electron backscattered diffraction (EBSD) mapping was performed on samples cold rolled between 0-40%. Increased cold rolling increased the fraction of low-angle boundaries at the expense of the coincident site lattice boundaries. The tensile-creep rupture life (T r ) and elongation-to-failure (ε f ) were evaluated at 649 C and 758MPa, and the data indicated that increased cold rolling significantly increased both the T r and ε f values. In addition the GBCD and room-temperature (RT) tensile properties were evaluated for superplastically formed INCONEL alloy 718 (IN 718SPF). The tensile results indicated the exceptional strength of the fine-grained IN 718SPF material, however the GBCD parameters were intermediate to those of the 10% and 20% cold rolled IN 718 materials. Introduction IN 718 is the most commonly used Ni-base superalloy for structural applications due to its good tensile, fatigue, creep and rupture strength at temperatures up to 700 C as well as its ease of fabrication [1]. A means to potentially improve the mechanical performance of superalloys such as IN 718 is through thermomechanical processing (TMP) to control the grain size and grain boundary structure. It has been demonstrated that IN 718 can be processed superplastically to a grain size of approximately 6µm and the resulting tensile and fatigue strength are superior to conventionally processed IN 718 [2-4]. Optimizing grain boundary structure in superalloys has provided for simultaneously improving creep, corrosion, fatigue, and weldability performance with virtually no degradation in either tensile strength or ductility [5]. Since altering the grain size and grain boundary structure does not involve variations in alloy chemistry, improvements in performance will not detrimentally affect thermal conductivity and/or phase stability. Advanced TMP schemes can be used to enhance the grain boundary character distribution (GBCD), which details the proportions of random and special grain boundaries as described by the coincident site lattice (CSL) model [6,7]. Fundamentally, TMP steps involving strain and recrystallization cause special boundaries to replace random boundaries in the boundary network when the appropriate conditions are obtained. It has been demonstrated that the frequency of special boundaries can be significantly increased in Ni-based alloys having a wide variety of compositions [5,8-11]. It has also been shown that the presence of a high fraction of special boundaries increases the creep resistance of Ni-base alloy such as Ni-16Cr-Fe, and INCONEL alloy 625 and alloy 738 [5,8,12-14]. In this work, the GBCD was characterized and quantified for IN 718 after cold rolling deformation between 0-40% and also after subsequent annealing at 921 C. EBSD was performed to INCONEL is a registered trademark of Special Metals Corporation.

2 Title of Publication (to be inserted by the publisher) identify the distribution of low-angle boundaries (LABs), high-angle boundaries (HABs), and coincident site lattice boundaries (CSLBs). Creep rupture experiments were conducted on coldrolled specimens, which were annealed and aged. In addition the GBCD and RT tensile properties of IN 718SPF were evaluated. Trends with respect to the relationships between GBCD and processing are presented. Experimental The IN 718 superalloy sheet was processed at Special Metals Corporation, Huntington, WV. The heats were produced by vacuum induction melting followed by electroslag remelting. The material was hot worked using conventional practices and the as-processed condition included mill annealing at 1066 C for all hot-rolling procedures which preceded the final cold roll and 982 C anneal. Subsequent TMP treatments included cold rolling between 0-40%. The rolling steps were performed on separate sheets each designated with 0%, 10%, 20%, 30%, and 40% cold rolling deformation. Metallographic samples were prepared from the sections of the rolled sheets prior to annealing, after annealing then water quenching, and after annealing then aging. The annealing treatment was performed at 921 C. The aging treatment consisted of 719 C/8h/furnace cool to 621 C then hold at 621 C for a total aging time of 18h. This aging treatment is intended to precipitate out the γ and γ strengthening phases in significant volumes, but it is not expected to change the GBCD. Sheet specimens with a dogbone geometry were machined with the tensile axis parallel to the rolling direction, and creep rupture experiments were performed in air at 758MPa and 649 C. In addition, a separate heat of IN 718SPF was produced in a similar fashion, however the sheet cold working procedure was altered to assure the production of a ultrafine grain size product. In addition to evaluating the GBCD, the IN 718SPF material was tensile tested at RT using a strain rate of approximately 1.3x10-3 s -1. EBSD orientation maps, obtained using an accelerating voltage of 25keV and a step size of 0.5µm on a Phillips 515 SEM with a LaB 6 filament, were obtained for the cold rolled samples as well as cold rolled samples which were subsequently annealed then water quenched. The final polishing step included several minutes using 0.06µm colloidal silica. EDAX-TSL, Inc., Draper, Utah manufactured the EBSD hardware and software. Brandon s criteria [15] were used to distinguish between the grain boundary types, and special boundary fractions were defined as LABs+CSLBs. The reported fractions of HABs, LABs, CSLBs, and twins (Σ3) were the average values taken from several orientation maps, performed on the cross-sections, rolling faces, or longitudinal sections of the processed sheets (see Fig. 1), of areas typically greater than 200µm by 400µm. It is noted that the minimum boundary tolerance angle selected was five degrees. Rolling Direction CS F L F Rolling Face CS Cross Section L - Longitudinal Figure 1. Schematic illustrating the sheet orientations from which the EBSD samples were sectioned. Results and Discussion Microstructure. The chemical composition range of the IN 718 heats used is shown in Table I. The annealed microstructures contained an equiaxed γ-phase (FCC) austenitic matrix and after

3 aging fine γ and γ precipitated throughout. The average γ grain diameter for the 0-40% cold rolled sheets ranged between µm as measured through the line intercept method as well as the EBSD analysis. Thus 0-40% cold rolling did not drastically change the equiaxed γ grain size. However, the superplastically formed microstructures exhibited a grain diameter between µm. For the annealed samples, the RT hardness values remained almost identical with respect to cold rolling deformation, see Table III. These values indicate that the annealing temperature was above the recrystallization temperature and the quenching treatment successfully avoided formation of a significant fraction of the strengthening precipitates [16]. Table II lists the GBCD parameters, including total fraction of CSLBs, LABs, HABs, and Σ3 boundaries, as a function of cold rolling, annealing, and sheet orientation. In general, there were small differences between the GBCD parameters of the different sheet orientations for a given cold rolling condition. However, no trends were evident with respect to the sheet orientation. The crystallographic texture, measured using pole figure analysis, appeared to increase slightly for increased rolling deformation where the maximum intensity for the preferential orientation of the 40% cold rolled condition was approximately twice that for the baseline 0% cold-rolled condition. The maximum intensity for the preferred orientation of the IN 718SPF material was within 1.5 times that of the baseline condition. Table1 Chemical composition range in weight percent for the IN 718 heats used in this study Nickel Titanium Molybdenum 2.99 Cobalt Chromium Aluminum Phosphorus Carbon 0.03 Iron Copper Silicon Manganese Niobium (plus Tantalum) Sulfur Table II. The grain boundary character parameters of the rolled and rolled-then-annealed samples Cold Rolling, % CSLBs HABs LABs Σ3 Σ3/CSLBs CSLBs+LABs GS, µm Rolled (not annealed) 0 F CS L F L F CS F CS F CS Rolled then annealled 20 F CS F CS F CS SPF then annealed IN 718SPF F IN 718SPF L GS: average grain diameter; F, CS, and L correspond to the sheet orientation (see Fig. 1). For the 0-40% cold rolled materials, increased rolling increased the fraction of LABs. The CSLBs, which were dominated by twins, increased slightly from the baseline material to the 10% cold rolled condition and then decreased with further rolling deformation. Overall the fraction of HABs was not significantly altered by the cold rolling and annealing operations where the HABs fractions typically hovered between This is unlike that for other Ni-base superalloys where

4 Title of Publication (to be inserted by the publisher) HAB fractions as low as 0.3 are achieved after TMP treatments [5,8-11]. It is noted however that those boundaries with misorientations below 5 degrees were eliminated from the analysis to avoid interpretation of subgrain boundaries as LABs. If such misorientations were included in the analysis the overall fraction of LABs would significantly increase at the expense of the HABs, and the resulting HAB fractions would be approximately 0.3. The authors have chosen to eliminate this ambiguity in an attempt to identify significant GBCD trends of the equiaxed γ grain boundaries as a function of cold rolling deformation. Overall, the main trends exhibited by cold rolling between 0-40% are increased cold rolling tends to decrease the fraction of CSLBs and increase the fraction of LABs. The respective averaged LAB and CSLB fractions for the 40% cold rolled then annealed samples were and 0.166, while the corresponding values for the baseline 0% cold rolled material were and There appears to be a limit to this trend as the IN 718SPF material, which is estimated to have cold rolling deformation between 55-80%, exhibited significantly larger CSLB fractions and smaller LAB fractions than the 40% cold rolled material. In fact, judging from the 0-40% cold rolled data (in particular the LAB and CSLB fractions), the IN 718SPF material GBCD parameters resemble those between the 10% and 20% cold rolled IN 718 materials. Creep Rupture. Fig. 2 illustrates the effect of cold rolling on the creep rupture life for the annealed then aged samples. The data, also listed in Table III, indicate that increased cold rolling deformation tends to increase the creep rupture life. The greatest rupture life and elongation values were exhibited by the 30% and 40% cold rolled samples, where both the T r and ε f values were greater than twice those for the as-processed and 10% cold rolled conditions. Cold rolling below 30% did not offer as significant of an increase in the creep rupture properties and both the 10% and 20% cold rolled conditions resulted in lower ε f values than that for the as-processed condition. Each of the ruptured samples exhibited ductile dimpling throughout the fractured surface, see Fig. 3a. The IN 718SPF material was not evaluated in creep in this study, but based on previous creep rupture data [2] its creep rupture life is expected to be similar to that exhibited by the 20% cold rolled condition. Thus there appears to be a limit to the amount of cold rolling deformation that will result in increased creep rupture life and ε f. This limit lies between 30% cold rolling and that estimated (55-80%) to be used for superplastic forming of IN 718. It is noted that a significant decrease in the grain size occurred from 40% cold rolling to superplastic forming, and this may be a significant factor in the creep rupture discrepancy. Increased cold rolling also corresponded to an increase in LAB fraction as mentioned previously. Further investigation of the effect of GBCD on the creep and creep rupture properties are necessary to indicate the importance of grain boundaries to the creep resistance under these creep stress and temperature conditions. In particular creep strain history will be evaluated to help delineate creep strain rate effects. Table III. The Creep Rupture Properties and RT Hardness of Cold Rolled IN 718 Cold Rolling Deformation, % T r *, hr Creep* ε f, % Hardness**, Rb Equivalent Hardness**, Hv *: annealed and aged samples crept at 649 C/758MPa; **: annealed then water quenched samples Tensile properties. Table IV lists the RT tensile properties of the IN 718SPF material as a function of heat treatment and orientation. With aging the tensile strength increased dramatically at the expense of ε f, though the ε f values were always greater than 12%. Only a small difference in the tensile properties was observed with respect to sample orientation. A ductile fracture was evident for all the samples tested, see Fig. 3b. The strength and elongation values lie intermediate to those

5 published by Gaylord and Yates [2] for two different heats of IN 718SPF. The RT tensile properties of IN 718SPF are significantly greater than those for conventionally processed IN 718 [17], and this is most likely a result of the ultrafine grain size. Creep Rupture Life, hrs ε f T r Cold Rolling Deformation, % Figure 2. Plot illustrating the relationship between creep rupture life and elongation-to-failure for the 0-40% cold rolled then annealed and aged IN 718 samples. (Creep conditions: 649 C and 758MPa). Table IV. RT Tensile Properties of the IN 718SPF material. Sample ID# Heat Treatment σ ys, MPa σ uts, MPa σ f, MPa ε f, % 3-L anneal L anneal T anneal T anneal L anneal+age L anneal+age T anneal+age T anneal+age L:tensile axis parallel to final rolling direction; T:tensile axis perpendicular to final rolling direction. Summary IN 718 was processed through sequential increments of cold rolling followed by annealing to evaluate the GBCD as a function of processing. The grain boundary character of the 0-40% cold rolled microstructures revealed that the fraction of LABs increases with rolling deformation while the fraction of CSLBs, which were dominated by twins, decreased with increased cold rolling. Neither the annealing treatment nor the sheet orientation significantly affected the GBCD parameters. Trends in the creep rupture data indicated that elongation-to-failure increases with amount of cold rolling from 10-40%. The same trend was also evident in the creep rupture life however the greatest rupture life was exhibited by the 30% cold rolled material. The as-processed and 10% cold rolled conditions exhibited the shortest creep rupture life and lowest ε f values. The RT tensile results revealed the exceptional strength exhibited by the fine-grained IN 718SPF material. Regardless of processing condition, the deformation behavior indicated a ductile fracture. Comparing the preliminary results of this investigation with that for other investigations of IN 718 and IN 718SPF, the processing route for obtaining an exceptional balance between creep rupture Elongation-to-Failure, %

6 Title of Publication (to be inserted by the publisher) and RT tensile properties is suggested to lie between 30% cold rolling deformation and superplastic formation. (a) (b) Figure 3. Photomicrographs of the fracture surfaces for a (a) creep rupture sample which was 0% cold rolled then annealed and aged and a (b) RT tensile tested IN 718SPF sample. Acknowledgments. This work was supported by the NSF Division of Materials Research (DMR ). The authors are grateful to Ronald Witt and Jeffrey Farrer of EDAX-TSL, Inc. Draper, Utah for performing some of the EBSD maps, and also Jay Spike for performing the tensile experiments. References [1] E.A. Lloria, The Status and Prospects of Alloy 718, Jour. of Metals (July 1988), [2] G.D. Smith and D.H. Yates: Proc. Advancements in Synthesis and Processes, Toronto, Canada. M207-M218 (Society for the Advancement of Material and Process Engineering, Covina, CA 1992). [3] G.D. Smith and H.L. Flower: Proc. Superalloys 718, 625, 706 and Various Derivatives, Pittsburgh, PA (The Minerals, Metals and Materials Society, Warrendale, PA 1994). [4] B.A. Baker: INCO Alloys International Technical Investigation Report No. BAB , Huntington, WV. September [5] E.M. Lehockey, G. Palumbo and P. Lin: Metall. and Mater. Trans. Vol. 29A (1998), [6] V. Randle: The Role of Coincident Site Lattice in Grain Boundary Engineering (The Institute of Materials, London, 1996). [7] V. Randle: Journal of Metals (February 1998), [8] G. Palumbo, E.M. Lehockey and P. Lin: Jour. of Metals (February 1998), [9] P. Lin, G. Palumbo, U. Erb and K.T. Aust: Scripta Materialia Vol. 33 No. 9 (1995), [10] M. Kumar, W.E. King and A.J. Schwartz: Acta Materialia Vol. 48 (2000) [11] A.J. Schwartz and W.E. King: Journal of Metals (February 1998), [12] G.S. Was, V. Thaveeprungsriporn and D.C. Crawford: Jour. of Metals (February 1998), [13] V. Thaveeprungsriporn and G. Was: Metall. and Mater. Trans. Vol. 28A (1997),

7 [14] E.M. Lehockey and G. Palumbo: Materials Sci. and Engineering Vol. A237 (1997), [15] D.G. Brandon: Acta Metallurgica Vol. 14 (November 1966), [16] W.C. Liu, Z.L. Chen and M. Yao: Metall. and Mater. Trans. Vol. 30A (1999), [17] Unpublished data from Special Metals Corporation, Huntington, West Virginia.