the effect of sheet processing on the microstructure, tensile, and creep behavior of inconel alloy...

7/27/2019 The Effect of Sheet Processing on the Microstructure, Tensile, And Creep Behavior of INCONEL Alloy 718

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The Effect of Sheet Processing on the Microstructure,Tensile, and Creep Behavior of INCONEL Alloy 718

C.J. BOEHLERT, D.S. DICKMANN, and N.C. EISINGER

The grain size, grain boundary character distribution (GBCD), creep, and tensile behavior of INCONEL*

*INCONEL is a registered trademark of Special Metals Corp., Huntington, WV.

alloy 718 (IN 718) were characterized to identify processing-microstructure-property relationships. Thealloy was sequentially cold rolled (CR) to 0, 10, 20, 30, 40, 60, and 80 pct followed by annealing attemperatures between 954 C and 1050 C and the traditional aging schedule used for this alloy. Inaddition, this alloy can be superplastically formed (IN 718SPF) to a significantly finer grain size andthe corresponding microstructure and mechanical behavior were evaluated. The creep behavior was eval-uated in the applied stress (a) range of 300 to 758 MPa and the temperature range of 638 C to 670 C.Constant-load tensile creep experiments were used to measure the values of the steady-state creep rateand the consecutive load reduction method was used to determine the values of backstress (0). Thevalues for the effective stress exponent and activation energy suggested that the transition between therate-controlling creep mechanisms was dependent on effective stresses (e a 0) and the transitionoccurred at e 135 MPa. The 10 to 40 pct CR samples exhibited the greatest 650 C strength, whileIN 718SPF exhibited the greatest room-temperature (RT) tensile strength (1550 MPa) and ductility(f 16 pct). After the 954 C annealing treatment, the 20 pct CR and 30 pct CR microstructures exhib-

ited the most attractive combination of elevated-temperature tensile and creep strength, while the mostseverely cold-rolled materials exhibited the poorest elevated-temperature properties. After the 1050 Cannealing treatment, the IN 718SPF material exhibited the greatest backstress and best creep resistance.Electron backscattered diffraction was performed to identify the GBCD as a function of CR andannealing. The data indicated that annealing above 1010 C increased the grain size and resulted in agreater fraction of twin boundaries, which in turn increased the fraction of coincident site lattice bound-aries. This result is discussed in light of the potential to grain boundary engineer this alloy.

I. INTRODUCTION

THE potential for improving the bulk properties of bothstructural and functional materials by manipulating the frequency

of special grain boundaries, defined here as low angle bound-aries (LABs; boundary misorientations less than 15 degrees)and coincident site lattice boundaries (CSLBs), was first intro-duced by Watanabe[1] and is considered to be the basis for mostcurrent grain boundary engineering (GBE) studies of metallicalloy systems. Over the past two decades, GBE efforts havedemonstrated significant improvements in stress corrosion crack-ing (SCC), fatigue, weldability, and creep of pure Ni and Ni-based superalloys.[215] In terms of Ni-based superalloys, effortshave been primarily focused on nonage-hardenable alloys suchas alloys 600 and 625, where the latter alloy has exhibitedsignificantly lower creep rates when it is processed to contain ahigher fraction of special boundaries. Although creep rates ofalloy 738 have also exhibited slower strain rates with increased

special boundary fractions,[2] few GBE studies have involvedthe age-hardenable series of Ni-based superalloys. Kruppet al.[16,17] have shown that the fraction of special boundariescan be significantly increased for IN 718 using cold rolling

deformation followed by recrystallization annealing whereincreased special boundary fractions reduced the sensitivity tooxygen-induced intergranular brittle fracture (dynamic embrit-tlement). In order to identify processing-microstructure-propertyrelations, the current work evaluated the GBCD of IN 718 as afunction of sheet processing using 10 pct cold-rolling incrementsfrom 0 to 80 pct followed by annealing temperatures rangingbetween 954 C to 1050 C, and the elevated-temperature tensileand tensile-creep properties were measured. The approach takenwas to characterize the creep behavior within the low-stresscreep regimes where diffusional creep and grain boundary slid-ing may be dominating the strain-rate response.

The creep mechanisms that operate during elevated-tem-perature deformation of pure metals and solid-solution-

strengthened alloys have been related to the value of the stressexponent, na, and the apparent activation energy, Qa, in theDorn steady-state creep rate equation:[18]

[1]

where Tis the creep temperature in degrees Kelvin, R is thegas constant, A is a constant, b is the Burgers vector, D0 isthe pre-exponential factor, is the shear modulus, and k isthe Boltzmann constant. However, the na and Qa values mea-sured for alloys containing dispersed second-phase particles

#

ss AD0 exp (Qa>RT) mb>kT(sa>m)na

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 37A, JANUARY 200627

C.J. BOEHLERT, Assistant Professor, is with the Department of Chem-ical Engineering, Michigan State University, East Lansing, MI 48824. Con-tact e-mail: [email protected] D.S. DICKMANN, ManufacturingEngineer, is with Electronic Ceramics, Ferro Corporation, Penn Yan, NY.N.C. EISINGER, Metallurgist, is with the Special Metals Corporation,Huntington, WV.

This article is based on a presentation made in the symposium entitledProcessing and Properties of Structural Materials, which occurred dur-ing the Fall TMS meeting in Chicago, Illinois, November 912, 2003, underthe auspices of the Structural Materials Committee.


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28VOLUME 37A, JANUARY 2006 METALLURGICAL AND MATERIALS TRANSACTIONS A

are generally considerably greater than those observed in puremetals and solid-solution-strengthened alloys.[1931] The n

a

values for precipitation-hardened alloys have ranged between5 and 15[1924] while, for dispersion-hardened alloys, includ-ing thoria-dispersed Ni-20Cr(wt pct), n

avalues have ranged

between 9 and 75.[27,30,31] The Qa

values have ranged from1 to 3 times those of the activation energy for self-diffu-sion.[29,31,32] These variations in the observed values ofn

aand

Qa

have been rationalized by introducing the concept of aback stress,

0, which is an internal stress opposing the dis-

location motion. In multiphase alloys such as IN 718, whichcontains an austenitic fcc phase matrix () and fine and strengthening precipitates, the applied stress during steady-state creep deformation is opposed by a backstress resultingfrom the presence of these strengthening particles and a defectstructure within the material.[19,27,3338] Therefore, the creepdeformation results from an effective stress (

e

a 0).

As a result, the steady-state creep rate can be represented by

[2]

where ne

is the effective stress exponent. Although tradi-tionally used to measure the back stress during high-stressdislocation power-law creep, it has been shown that the con-

secutive stress reduction method can also be applied to low-stress diffusional creep for IN 718, where low n

evalues (2)

are observed.[33,34]

II. EXPERIMENTAL

The IN 718 sheets used in this study were processed atSpecial Metals Corporation (Huntington, WV). The heats wereproduced by vacuum induction melting followed by elec-troslag remelting. The material was hot worked using con-ventional practices, and the as-processed condition includedmill annealing at 1066 C for all hot-rolling procedures thatpreceded the final CR and 982 C anneal. Subsequent ther-momechanical processing (TMP) treatments included CRbetween 10 and 80 pct. The CR steps were performed on sep-arate sheets each designated with 10, 20, 30, 40, 60, and 80 pctdeformation. Metallographic samples were prepared from thesections of the CR sheets prior to annealing, after annealingthen water quenching, and after annealing then aging. Theannealing treatment was performed at one of the followingtemperatures: 954 C, 1010 C, or 1050 C. In addition, oneset of CR samples was heat treated below the recrystalliza-tion temperature at 871 C. The aging treatment, used toprecipitate out the and strengthening phases, consistedof 718 C/8 h/furnace cool to 621 C then hold at 621 C fora total aging time of 18 hours. In addition, a separate heat ofIN 718SPF was produced in a similar fashion; however, thesheet cold working procedure, estimated to total between 55and 80 pct deformation, was altered to assure the productionof an ultrafine grain size product.[3942]

Electron backscattered diffraction (EBSD) orientation maps,obtained using an accelerating voltage of 25 keV and a stepsize of 0.5 m on a PHILIPS 515 SEM* with a LaB6 filament,

*PHILIPS is a trademark of Philips Electronic Instruments Corp.,Mahwah, NJ.

were obtained for the CR samples as well as CR samples whichwere subsequently annealed then water quenched. The final

#

ss A*(s

as0)

ne

preparatory polishing step included greater than 20 minutesusing 0.06-m colloidal silica. EDAX-TSL, Inc. (Draper, UT)manufactured the EBSD hardware and software. Brandonscriteria[43] were used to distinguish between the grain bound-ary types. The reported fractions of random general high-angleboundaries (GHABs), LABs, CSLBs, and twins (3) were theaveraged values taken from several orientation maps, performedon the cross sections, rolling faces, or longitudinal sections ofthe processed sheets, of areas typically greater than 1.4 mmby 0.7 mm.

Flat dogbone-shaped tensile and creep specimens with across section of approximately 1 12 mm and a gage lengthof 25 mm were machined, using either a mill or an elec-trodischarge machine, with the tensile axis parallel to therolling direction. The tensile experiments were performed inair at room temperature (RT) or 650 C using an Instron 8562(Norwood, MA) machine and a strain rate of approximately1.3 104s1. In most cases, multiple tests were performedand the reported strength and elongation values were aver-aged. Constant-load creep experiments were performed onApplied Test Systems Incorporated (Butler, PA) lever-armcreep apparati, using a 20:1 load ratio, in air at temperaturesranging between 638 C and 670 C and applied stresses rang-

ing between 300 and 758 MPa. The creep strain was moni-tored during the tests using a linear variable differentialtransformer attached to the gage section. The specimen tem-perature, monitored by three thermocouples located withinthe gage section during the creep experiments, was maintainedwithin 3 C using a single-zone furnace. The 0 values weredetermined at 638 C by the consecutive stress reductionmethod.[19,22,26] When the creep rate for a given

aremained

constant for at least 5 hours, it was assumed the steady-statecreep rate had been achieved. Thereafter, the sample was sub-jected to a small stress reduction (approximately 5 pct

a).

This resulted in an elastic contraction of the sample, followedby an incubation period with a zero creep rate. After a periodof time, creep began again at a lower rate. Once steady state

was reached, another stress reduction was performed. Thetime of the incubation period following each stress reductionwas recorded. The remaining stress vs the cumulative incu-bation time was plotted, and 0 was determined by taking theasymptotic value of the remaining stress when the cumula-tive incubation time appeared to be infinite. The 0 and

ss

values proved to be repeatable as duplicate samples weretested at the same temperature and

aand the measured

ss

values were within 5 pct of each other. In addition, the 0and

ssvalues were not dependent on strain history for total

creep strains less than 0.5 pct as several temperature/appliedstress conditions were performed, some in duplicate, beforethat of the backstress condition, and in each case, similar0 val-ues were recorded. Creep rupture experiments were performed

in air at 758 MPa and 649 C. All the creep and 650 C ten-sile experiments were initiated after soaking the samples atthe desired testing temperature for a minimum of 0.5 hours.

III. RESULTS

A. Microstructure

The chemical composition range of the IN 718 heats usedis shown in Table I. The annealed microstructures containedan equiaxed -phase (fcc) austenitic matrix and after aging


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Table I. Composition Range for the IN 718 Heats Used in This Study

Element Ni Ti Mo Co Cr Al P

Wt Pct 53.48 to 53.68 1.01 to 1.06 2.99 0.03 to 0.12 18.1 to 18.4 0.46 to 0.48 0.009 to 0.012Element C Fe Cu Si Mn Nb SWt Pct 0.03 17.99 to 18.3 0.02 to 0.17 0.01 to 0.17 0.04 to 0.12 5.07 to 5.11 0.001

(a)

(b)Fig. 1(a) Low-magnification and (b) high-magnification SEM photomi-crographs of the cross section of a 0 Pct cold rolled then 871 C annealedand aged microstructure illustrating the austenitic -phase matrix, fine and precipitates, and -phase precipitates (white).

fine (coherent spherical fcc (L12)) and (coherent ordereddisc-shaped body-centered tetragonal (DO22)) precipitatedthroughout (Figure 1). The average grain size was mea-sured through the line-intercept method (Table II).

1. Annealing temperature effectsAnnealing temperature had a significant effect on grain size.

Above 1010 C, grain growth occurs[44] and the 1050 C, 1-hannealed samples exhibited significantly larger grain sizes thanthose annealed at 954 C (Table II). The 1050 C annealedmicrostructures also exhibited an increased volume fraction of

twin boundaries compared with the 1010 C and 954 Cannealed microstructures; this will be discussed in Section 2.It is noteworthy that the RT hardness values significantlyincreased with CR deformation for the 871 C heat-treatedsamples, while hardness remained almost constant after 954 Cannealing (Tables III and IV). This indicated that 871 C isbelow the austenite (-phase) recrystallization temperature for0-40 pct CR IN 718, which is consistent with previous find-ings.[45] The relatively constant hardness values for the 954 Cannealed samples (Table II) indicated that the annealing tem-perature was above the recrystallization temperature and the

Table II. Grain Size of IN 718 as a Function of CR

and Annealing

954 C Annealed 1050 C Annealed

Standard StandardCR, Pct d, m Deviation d, m Deviation

0 36.2 6.6 92.0 11.010 37.5 8.2 77.5 14.020 30.6 6.5 89.1 15.330 28.9 3.4 76.4 10.140 28.7 2.5 83.9 8.960 23.0 4.0 101.8 13.380 20.0 3.2 99.8 13.3SPF 12.0 1.4 83.9 5.7

Table III. RT Hardness of 0 to 40 Pct CR then 954 CAnnealed-Then-Water-Quenched IN 718

CR, Pct Hardness, Rb Equivalent Hardness, Hv

0 95.1 22010 94.0 21520 93.7 21430 94.6 21840 94.2 216

Table IV. RT Properties of the 871 C Heat-Treated-Then-Aged Samples

CR, Pct YS, MPa UTS, MPa RT f, Pct Hardness, Hv0 1167 1435 20.7 434

10 1193 1433 22.5 43220 1346 1496 17.0 46030 1446 1557 15.1 47040 1497 1606 12.9 478


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Table V. GBCD Parameters of the Rolled and Rolled-Then-Annealed Samples

CR, Pct CSLBs GHABs LABs 3

As-rolled0 0.415 0.552 0.033 0.279

10 0.496 0.478 0.026 0.38220 0.215 0.665 0.120 0.19430 0.195 0.564 0.241 0.14140 0.158 0.483 0.359 0.109

954 C0 0.459 0.516 0.025 0.350

10 0.387 0.565 0.048 0.28720 0.342 0.563 0.095 0.26130 0.362 0.601 0.037 0.26140 0.321 0.636 0.043 0.22360 0.304 0.653 0.043 0.20980 0.213 0.755 0.032 0.1IN 718SPF 0.327 0.626 0.047 0.232

1010 C0 0.454 0.516 0.030 0.344

20 0.252 0.594 0.154 0.170

30 0.213 0.614 0.173 0.12040 0.152 0.606 0.241 0.073

1050 C0 0.350 0.616 0.034 0.233

10 0.476 0.496 0.028 0.33320 0.549 0.430 0.021 0.44030 0.412 0.556 0.032 0.31240 0.417 0.553 0.030 0.32860 0.543 0.426 0.031 0.43580 0.502 0.447 0.051 0.348IN 718SPF 0.563 0.399 0.038 0.342

quenching treatment successfully avoided formation of a sig-nificant volume fraction of the strengthening precipitates.[44]

Overall, the hardness values agreed well with the RT tensilestrength of the 871 C heat-treated-then-aged samples accord-ing to the established hardness-tensile strength relationship forIN 718[46] (data in Table IV).

2. Grain boundary character distributionTable V lists the GBCD parameters, including total frac-

tion of CSLBs, LABs, GHABs, and 3 boundaries, as a

function of CR and annealing temperature. The crystallo-graphic grain texture, measured using pole figure analysistaken from EBSD orientation maps (Figure 2), appeared toincrease slightly for increased rolling deformation where themaximum intensity for the preferential orientation of the40 pct CR condition was approximately twice that for thebaseline 0 pct CR condition. This corresponded to an increasein the LAB fraction of 0.03 in the 0 pct CR condition to0.36 in the 40 pct CR condition (Table V). However, afterannealing, the LAB fractions decreased to less than 0.3 forthe 40 pct CR condition, while the LAB fractions were evenlower for the 0 to 30 pct CR conditions. The maximum inten-sity for the preferred orientation of the IN 718SPF 954 Cannealed material was within 1.5 times that of the baseline

condition. In addition, no significant GBCD differences wereevident with respect to the sheet orientation, and therefore,the average CSLB, LAB, GHAB, and 3 boundary frac-tions used were taken from each of the three sheet orienta-tions. It is noted that GHABs dominated the GBCD of theas-cold-rolled samples as well as the samples annealed at

954 C and 1010 C (Table V). The LAB fractions were theleast prevalent and always less than 0.1. For the 0 to 80 pctCR materials annealed at 954 C, increased CR increasedthe fraction of GHABs and decreased the fraction of CSLBs,which were dominated by twins (Table V and Figure 3). TheGHABs ranged from 0.5 for the baseline 0 pct CR micro-structure to 0.75 for the 80 pct CR microstructure. Otherthan for 3 boundaries, no particular x(1x 29) bound-ary fraction was greater than 0.07.

The annealing temperature had a significant effect onGBCD. The fraction of twin boundaries increased signifi-cantly for the 1050 C annealed samples compared with the954 C and 1010 C annealed samples. For example, the twinboundary fractions increased from 0.21 to 0.26 (954 Cannealing) to 0.44 (1050 C annealing) for the 20 pct CRand 60 pct CR microstructures (Table V). Opposite to thatof the 954 C annealed samples, the GHABs decreased andthe CSLBs and twin boundary fractions increased withincreased CR for the 1050 C annealed samples (Figure 4).Representative EBSD data are illustrated in Figure 5, whichhighlights twin boundaries, LABs, and GHABS for a 1050 Cannealed 60 pct CR material. Figure 6 illustrates the distri-bution of CSLBs in this microstructure. It is noted that the

GBCD was not affected by aging or creep exposure, as suchsamples exhibited a similar distribution of GHABs, LABs,and CSLBs as those of the annealed-then-quenched samples.

B. Mechanical Behavior

1. Tensile behaviorFigure 7 illustrates the tensile stress/strain behavior of IN

718SPF in the 954 C annealed and 954 C annealed-then-aged conditions, where the significant strengthening effectof the fine precipitates is evident (Table VI). For both theannealed and annealed-then-aged conditions, half of the sam-ples represented in Figure 7 were machined in an orientation90 deg with respect to the rolling direction, indicating that

strength was not significantly effected by sheet orientation.This may be expected based on the lack of a strong texturein these severely deformed microstructures, as previouslydescribed. With aging, the tensile strength increased dramat-ically at the expense of elongation-to-failure (f), though thef values were always greater than 12 pct, and a ductile frac-ture was evident for all the samples tested (Figure 8). The RTstrengths of IN 718SPF approached those exhibited by the 20to 30 pct CR deformation-hardened material, which was notannealed but heat treated at 871 C prior to aging (compareTables IV and VI).

The 650 C yield strength (YS), ultimate tensile strength(UTS), and f values are presented in Table VII for theannealed-then-aged samples as a function of CR. The strength

of the IN 718SPF material decreased significantly from RTto 650 C. For the 954 C annealing treatment, the 40 pct CRmaterial exhibited the greatest YS value, and the 80 pct CRand IN 718SPF samples exhibited significantly lower YS val-ues. Thus, CR increases the elevated-temperature strength;however, the limit to this strengthening appears to be 40 pctCR. For the 1050 C annealed-then-aged samples, this was lessevident as the YS values ranged within a tighter band (865 to921 MPa), and the 10 pct CR materials YS was only slightlygreater than that for the IN 718SPF material. The 650 C fvalues exceeded 8 pct for nearly all the samples tested.


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Fig. 2Pole figures taken from EBSD data for a 1050 C annealed 60 pct CR sample. Strong texturing was not evident even after a large amount of CRdeformation.

Fig. 3The fraction of GHABs, LABs, CSLBs, and 3 boundaries as afunction of cold rolling for the 954 C annealed samples. The IN 718SPFdata were input as 55 pct CR.

Fig. 4The fraction of GHABs, LABs, CSLBs, and 3 boundaries as afunction of cold rolling for the 1050 C annealed samples. The IN 718SPFdata were input as 55 pct CR.

2. Creep behaviorDuring the creep experiments, the strain-time plots illus-

trated the three stages of creep: primary, secondary, and ter-tiary. The dependence of the steady-state creep rate on a is

illustrated in Figure 9. The na values were similar to thoseobserved previously for IN 718 by Han and Chaturvedi, [33]

whose data were interpolated to 638 C and included in Fig-ure 9(a) and Table VIII. Their data were for an as-processed


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Fig. 6CSLB distribution chart for the 1050 C annealed 60 pct CRsample represented in Figs. 2 and 5.

Fig. 7Stress vs strain plot for RT tensile tested IN 718SPF samples, whichwere either 954 C annealed then water quenched or 954 C annealedthen aged.

Table VI. RT Tensile Properties of IN 718SPF

Heat Treatment y, MPa UTS, MPa f, MPa f, Pct

954 C anneal thenwater quenched 682 1063 982 33.6

954 C annealedthen aged 1333 1554 1496 15.8

(a)

(b)

Fig. 5EBSD orientation map for a 1050 C annealed 60 pct CR sample. ( a) Normal direction EBSD inverse pole figure map where the colors representthe samples normal direction indexed to the fcc unit triangle. (b) Image quality map highlighting twin boundaries (yellow), LABs (red), and GHABs (blue).

condition, and they closely resembled the 0 pct CR data in thecurrent work. Note that the na values, which ranged between8 and 40, were similar to those measured for other particle-strengthened alloy systems[1931] and are considerably largerthan those generally observed for pure metals. For the 954 Cannealed-then-aged samples, it was apparent that there weretwo clusters of data, where the highest strain rates were exhib-ited by the samples CR to more than 40 pct and lower strainrates were exhibited by the 0 to 40 pct CR samples. This wasnot the case for the 1050 C annealed-then-aged samples.

a. Creep backstressFigure 10 illustrates plots used to determine 0, while

Figure 11 illustrates the corresponding vs e plots. Thevalues of ne, which are listed along with and 0 in

#

ss

#

ss


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Tables VIII and IX, were determined from Eq. [2]. For the954 C annealing treatment, the 20 pct CR and 30 pct CRsamples exhibited the greatest 0 values, 645 and 630 MPa,respectively, while the 60 pct CR, 80 pct CR, and IN 718SPFconditions exhibited the lowest 0 values, which were less

than half those of the highest 0 values. Correspondingly,the values for a given a were the lowest for the 20 pctCR and 30 pct CR materials (Table VII). For e less than135 MPa, the ne values were between one and two, whilefor e greater than 135 MPa, ne was greater than 3.5 (TableVIII and Figure 11). The Qa value (304 10 kJ/mol) wasdetermined for a 20 pct CR 954 C annealed-then-aged sam-ple using the ln ss vs 1/Tplot shown in Figure 12. For the1050 C annealed-then-aged condition, the 0 values fell ina narrower band, 450 to 630 MPa, than those for the 954 Cannealed-then-aged condition. The IN 718SPF sample exhib-

# ss

(a)

(b)

Fig. 8Fractographs for RT tensile tested IN 718SPF: (a) 954 C annealedand (b) 954 C annealed-then-aged samples.

Table VII. 650 C Tensile Properties of IN 718

CR Deformation, Pct 0.2 Pct YS, MPa UTS, MPa f, Pct

954 C anneal*0 988 1145 13.5

10 957 1115 10.820 1002 1184 11.430 959 1161 12.640 1054 1153 12.760 1050 1130 6.1**

80 948 1057 19.6SPF 991 1056 4.8**1050 C anneal*

0 899 1028 10.110 921 1041 8.120 878 1044 7.430 906 1054 14.040 888 1024 11.860 865 1036 12.780 876 1054 15SPF 913 1088 12

*All samples were aged according to the aging treatment 718 C for8 h followed by furnace cooling to 621 C and holding for total aging timeof 18 h.

**Sample broke out of the gage section.

ited the greatest 0 value, 630 MPa, while the 30 pct CRcondition exhibited the lowest 0 value, 505 MPa (Fig-ure 10(b)). Similar to that for the 954 C annealed-then-agedcondition, there was a transition in the ne values at e 135 MPa. Thus, this trend occurred independent of bothannealing temperature and CR deformation amount. Thisresult suggests that the active creep mechanisms may bemore dependent on e than on a. It is noteworthy that formost samples, the 0 value was over half of the YS at 650 C,and the 0 value was as high as 66 pct for the 30 pct CR954 C annealed-then-aged condition.

b. Creep rupture

The creep rupture data, listed in Table X, indicated thatincreased CR tended to increase Tr. For the 954 C annealed-then-aged samples, the greatest Tr and f values were exhib-ited by the 30 and 40 pct cold-rolled samples, where boththe Trand fvalues were greater than twice those for the base-line 0 pct and 10 pct CR conditions. Cold rolling below 30 pctdid not offer as significant of an increase in the creep rup-ture properties and both the 10 pct and 20 pct CR conditionsresulted in lower f values than that for the baseline 0 pctCR condition. Each of the ruptured samples exhibited ductiledimpling throughout the fractured surface (Figure 13). The871 C annealed samples also exhibited increased Tr valueswith increased CR, and the Trvalues were significantly greaterthan the 954 C annealed samples. This is considered to be

a result, in part, of the greater tensile strengths exhibited bythe 871 C heat-treated materials, which, as previously dis-cussed in Section III-A-1, were always maintained belowthe recrystallization temperature.

IV. DISCUSSION

A. Microstructure

The frequency of special boundaries can be significantlyincreased in Ni-based superalloys having a wide variety of


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compositions and strengthening mechanisms (i.e., precipita-tion, solution, etc.). For alloys 625 and 738, TMP hasincreased the special boundary fractions from nominally 0.1to 0.2 to 0.5 to 0.7, and the resulting microstructures haveexhibited sufficient thermal stability to sustain prolongedexposure to service operating temperatures of 850 C.[2,8] InAlloy 600, TMP has increased the special boundary fractionfrom 0.15 to 0.7 without significantly affecting grain size oradding texture.[9,47,48] In the current work, TMP increased the

fraction of special boundaries for IN 718 from less than 0.4to 0.6 without significantly adding texture, where both theannealing temperature and cold work played significant roles.

1. Grain boundary character distributionFor the as-rolled samples in the 10 to 40 pct CR defor-

mation range, increased CR tended to decrease the fractionof3 and CSLBs. Recrystallization annealing at temperaturesof 1010 C and 954 C did not change this trend, and GHABsmaintained the majority volume fractions for all the CR con-ditions. The overall result was that the GHABs increased and

the special boundary fractions subsequently decreased withincreased CR. On the other hand, the 1050 C annealing treat-ment, which is above the -phase grain growth temperature(1020 C[40]), had a different effect on the microstructureand in particular grain size and GBCD. The fraction of twinboundaries tended to increase with increased CR deformationfor the 1050 C annealed samples (Figure 4), and this resultedin increased CSLBs and decreased GHABs with increasedCR deformation. It is noted that some variability existed withinthis trend as the 20 pct CR data exhibited similar GHABs asthe 60 pct CR and IN 718SPF materials; these were the threemicrostructures that exhibited the greatest fractions of specialboundaries (0.6). Independent of annealing treatment theLAB fraction was always less than 0.1 and relatively unaf-fected by CR amount. Achieving a majority of special bound-aries therefore may be achieved for IN 718, where annealingtemperature is a vital factor. It is noted that the correspond-ing GBCD, and in particular twin boundary fractions, com-pared favorably to that measured by Krupp et al.[16,17] forcold-worked then 1050 C annealed IN 718. Overall, increas-ing the grain size through higher-temperature annealingresulted in increased twin boundary fractions, which is con-sistent with observations for other Ni-based superalloys, includ-

ing Waspaloy.

[49]

As most of the twins exhibited were completetwins (Figure 5), rather than island, edge, or incompletetwins,[50] this is an understandable result, as increased twinboundary fractions will result not only from the increasednumber of twins but also the increased boundary area of thetwins in the microstructure. For other GBE-processed Ni-basesuperalloys such as alloys 625 and 738, special boundary frac-tions similar to those found in the current work have led tosignificantly improved creep resistance including lower creeprates and increased Tr, without a deterioration in tensile strengthor ductility.[2] The creep data for the 1050 C annealed-then-aged samples indicate significantly decreased creep rates andincreased back stress values for the IN 718SPF condition.Thus, this processing condition may be optimal for creep,

indicating there may be fertile ground to GBE IN 718 as wellas other age-hardenable Ni-base superalloys. It is noted thatthe creep deformation did not alter the GBCD significantly,as deformed samples exhibited similar GHABs, LABs, CSLBs,and 3 values as the undeformed microstructures.

Although it has been shown here that IN 718 can be TMPto a majority of special boundary fractions similar to otherNi-based superalloys systems, the grain boundary statisticsindicate that these microstructures do not compare well withthe twin-limited ideal microstructure, which is composedentirely of special grain boundaries (i.e., 2/3 are twins) as dis-cussed by Davies and Randle.[51] The twin-limited microstructurehas been approached in other investigations of Ni-based alloysand austenitic steels but never achieved,[52,53] while the cur-

rent results, where the twin boundary fractions saturated at0.44, are similar to most findings including those for oxygen-free electronic (OFE) Cu and IN 600 alloy.[47] Thus, the desiredmaximum value of twins in the microstructure may not beachievable through commercially practical TMP treatments.

B. Mechanical Behavior

1. Tensile behaviorThe measured RT strength and fvalues of IN 718SPF were

intermediate to those published by Smith and Yates[39] for two

(a)

(b)

Fig. 9 vs e plot for the (a) 954 C annealed-then-aged creep sam-ples tested at 638 C. Also included are data from Han and Chaturvedi, [33]

which has been interpolated to 638 C. (b) vs e plot for the 1050 Cannealed-then-aged creep samples tested at 638 C.

#

ss

#

ss


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chemically different heats of IN 718SPF. However, similar totheir results, the RT tensile strengths of the IN 718SPF 954 Cannealed-then-aged samples were significantly greater thanthose for conventionally processed IN 718,[44] and this is mostlikely a result of the ultrafine grain size. In fact, the IN 718SPFmaterial exhibited the greatest RT properties of all the recrys-tallization-annealed materials studied. The only processed sam-

ples that exhibited greater strengths were the 30 pct and 40pct CR deformation-hardened samples, which were heat treatedat 871 C prior to aging (Table IV). However, the elevated-temperature properties of IN 718SPF degrade more rapidlywith temperature than the 0 to 40 pct CR samples as thestrength of IN 718SPF decreased30 pct from RT to 650 C,compare the data in Table VI and VII.

Table VIII. The 638 C Creep Data of IN 718 Samples Annealed at 954 C Then Aged

CR Deformation, Pct a, MPa 0, MPa a 0, MPass na ne

0 574 540 34 2.1 109 10.8 1.2574 540 34 1.9 109 10.8 1.2594 540 54 3.5 109 10.8 1.2595 540 55 3.8 109 10.8 1.2609 540 69 4.4 109 10.8 1.2611 540 71 4.6 109 10.8 1.2632 540 92 6.0 109 10.8 1.2

649 540 109 9.1 109

10.8 1.2674 540 134 1.1 108 10.8 1.210 591 570 21 1.0 109 18.3 1.7

596 570 26 1.1 109 18.3 1.7610 570 40 2.8 109 18.3 1.7634 570 64 5.1 109 18.3 1.7666 570 96 1.1 108 18.3 1.7694 570 124 1.9 108 18.3 1.7

20 666 645 21 3.8 109 39.8 1.7676 645 31 7.3 109 39.8 1.7683 645 38 1.0 108 39.8 1.7

30 648 630 28 1.5 109 36.0 1.7668 630 38 5.6 109 36.0 1.7680 630 50 8.4 109 36.0 1.7

40 578 550 28 2.4 109 13.7 1.5593 550 43 2.9 109 13.7 1.5596 550 46 2.8 109 13.7 1.5598 550 48 3.3 109 13.7 1.5610 550 60 4.7 109 13.7 1.5613 550 63 5.2 109 13.7 1.5613 550 63 4.8 109 13.7 1.5627 550 77 6.6 109 13.7 1.5

60 378 310 68 3.3 109 4.6 1.0399 310 89 4.2 109 4.6 1.0417 310 107 5.2 109 4.6 1.0436 310 126 7.2 109 10.4 3.6455 310 145 9.9 109 10.4 3.6471 310 161 1.2 108 10.4 3.6490 310 180 2.0 108 10.4 3.6501 310 191 3.3 108 10.4 3.6537 310 227 5.8 108 10.4 3.6

564 310 254 7.9 108 10.4 3.6598 310 288 1.2 107 10.4 3.6612 310 302 1.3 107 10.4 3.6

80 375 300 75 2.3 109 4.5 1.1397 300 97 3.0 109 4.5 1.1417 300 117 3.8 109 4.5 1.1457 300 157 6.0 109 10.3 3.9489 300 189 9.7 109 10.3 3.9523 300 223 2.4 108 10.3 3.9

Han and Chatervedi[34]

interpolated to 638 C 620 524 96 5.9 109 9.0 1.9673 524 149 1.2 108 9.0 1.9696 524 172 1.7 108 9.0 1.9720 524 196 2.3 108 9.0 1.9

SPF 334 305 29 2.1 109 5.3 1.1

374 305 69 3.3 109

5.3 1.1405 305 100 6.0 109 5.3 1.1


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(a)

(b)

Fig. 10Remaining stress vs cumulative incubation time for the cold-rolledand (a) 954 C annealed-then-aged and (b) 1050 C annealed-then-agedsamples.

(a)

(b)

Fig. 11 vs e plot for the cold-rolled and (a) 954 C annealed-then-aged and (b) 1050 C annealed-then-aged samples. The data were used tocalculate the listed ne values.

#

ss

Increasing the annealing temperature was also not bene-ficial to the elevated-temperature strengths, as the 1050 Cannealed samples tended to exhibit lower strengths thanthe 954 C annealed samples. This is expected to be a resultof the -phase grain-size discrepancy. Based on the similarYS values between the 10 pct and 40 pct CR conditions, itis expected that 10 to 40 pct CR does not have a signifi-cant effect on elevated-temperature strength, and therefore,

the creep behavior for the 10 to 40 pct CR conditions maybe compared without considering YS effects. It is noted,however, that CR 10 to 40 pct does significantly increasethe elevated-temperature strength from the baseline 0 pctCR condition.

2. Creep behaviora. Effective stress exponentHan and Chaturvedi[33,34] observed an incubation time after

stress reduction during creep in both the power-law and thediffusional-creep regimes for IN 718. According to Harris[54]

and Burton,[55] for diffusional creep to continue, the stressconcentration at the precipitate-matrix interface created bythe entrapment of diffusing vacancies can only be relievedby the formation of prismatic dislocation loops. However,Ansel and Weertman[56] have suggested that diffusional creepin two-phase alloys can only occur by the process of dislo-cation climb over the precipitate particles. Therefore, diffu-sional creep in two-phase alloys may not only involve vacancydiffusion in the matrix and dislocation motion in the grain

boundary region but also dislocation creation and motionwithin grains. Han and Chaturvedi[33] observed dislocationnetworks within grains, where it was suggested that theobserved incubation time could be due to the activation ofBardeenHerring sources.[57] Therefore they concluded thatthe consecutive stress reduction method can be used to deter-mine 0 during power-law creep as well as diffusional creep.

The experimental results of the current work indicate twoseparate regimes based on the effective stress level. The nevalues were measured to be between one and two for e lessthan 135 MPa, and ne was 3.0 to 4.3 for e greater than


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135 MPa. Based on this result, the creep mechanism may bedependent on e, and the low-stress regime may be dominatedby diffusional creep or grain boundary sliding. The ne valuesof the current work are in agreement with the findings of Anseland Weertman[56] and Han and Chaturvedi[34] and promote thesuggestion that differing creep mechanisms may be active anddependent on effective stress for particle-strengthened alloys.

b. Activation energyThe measured Qa value (304 kJ/mol) was greater than both

the activation energy for self-diffusion (265 to 280 kJ/mol)and the activation energy for the creep process (276 kJ/mol)of pure Ni[58] and the creep process of Ni-Cr in solid solu-tion (295 kJ/mol).[59] However, the creep rate expression usedto calculate these values considers neither the influence oftemperature on the value ofG nor the concept of back stress.Using the back stress, lower activation energies than thosecalculated using the applied stress have been calculated,[34]

and such values lie within the range expected for lattice self-

diffusion. Thus, a similar result would be expected basedon the Qa measured here, and self-diffusion is considered tobe more likely than grain boundary diffusion for creep ofIN 718 in the temperature range of 638 C to 670 C.

c. BackstressFor the 954 C annealed-then-aged samples, the significant

drop in 0 with increased CR deformation beyond 40 pct was

correlated with a decreasing average grain size. It also cor-responded to a significant decrease in the special boundaryfractions (Figure 3). It has been hypothesized that an increasein the special boundary fraction decreases the annihilation rateof dislocations in the grain boundary, leading to an increasein 0, a decrease in the e, and therefore a reduction in theoverall creep strain rate.[10,11] The CSLBs, and in particulartwin boundaries, are not considered to be effective sourcesand sinks, and a boundary close to a CSL may require a finiteexcess concentration of vacancies before it will reabsorbthem.[60,61] The closer the orientation to that of a CSL, the

Table IX. 638 C Creep Data of IN 718 Samples Annealed at 1050 C Then Aged

CR Deformation, Pct a, MPa 0, MPa a 0, MPass na ne

0 614 530 84 4.2 109 11.4 1.9638 530 108 5.2 109 11.4 1.9658 530 128 8.8 109 11.4 1.9668 530 138 1.1 108 11.4 1.9

10 591 560 31 2.6 109 22.7 2.0616 560 56 9.1 109 22.7 2.0639 560 79 1.8 108 22.7 2.0

642 560 82 1.6 108

22.7 2.020 627 590 37 3.1 109 19.5 2.0648 590 58 4.2 109 19.5 2.0668 590 78 9.0 109 19.5 2.0696 590 106 2.3 108 19.5 2.0715 590 125 3.5 108 19.5 2.0

30 579 505 74 4.8 109 8.1 1.3580 505 75 5.2 109 8.1 1.3601 505 96 7.7 109 8.1 1.3627 505 122 9.3 109 8.1 1.3

40 597 550 47 5.2 109 8.4 1.1617 550 67 7.2 109 8.4 1.1617 550 67 7.1 109 8.4 1.1638 550 88 9.7 109 8.4 1.1643 550 93 8.7 109 8.4 1.1681 550 131 1.6 108 8.4 1.1691 550 141 1.7 108 13.1 3.0733 550 183 3.7 108 13.1 3.0

60 598 520 78 4.3 109 6.3 1.0619 520 99 5.0 109 6.3 1.0620 520 100 4.8 109 6.3 1.0639 520 119 6.6 109 6.3 1.0682 520 162 1.8 108 16.0 4.3707 520 187 3.6 108 16.0 4.3740 520 220 6.8 108 16.0 4.3

80 592 570 22 2.35 109 11.4 1.0615 570 45 3.78 109 11.4 1.0634 570 64 6.22 109 11.4 1.0656 570 86 8.05 109 11.4 1.0678 570 108 1.09 108 11.4 1.0

SPF 655 630 25 3.5 109 31.6 1.3

661 630 31 5.1 109 31.6 1.3682 630 52 1.5 108 31.6 1.3688 630 58 1.6 108 31.6 1.3


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Table X. Creep Rupture Properties at 649 C/758 MPa for

the 954 C* and 871 C** Annealed Aged SamplesCR Deformation,Pct Tr,* h f,* Pct Tr,** h f,** Pct

0 24.9 9.3 73.6 26.110 27.5 6.7 65.7 25.220 38.5 7.9 99.8 22.030 64.8 28.9 103.8 18.240 57.9 38.0 128.9 23.8

Fig. 13Fractograph for a creep rupture sample which was 0 pct CR then954 C annealed-then-aged sample.

Fig. 12In vs 1/Tplot used to calculate Qa for a 20 pct CR 954 Cannealed-then-aged sample.

#

ss

higher this excess, and this may slow the rate of creep. [62] Forthe 1050 C annealed-then-aged samples, the IN 718SPFcondition exhibited both the highest 0 value and specialboundary fraction (0.6) of all the conditions studied. The 0

values observed in the current work suggest that 0 may bedependent on both special boundary fractions and grain size.

It has been suggested that the backstress arises from theinability of dislocations to bow between adjacent particles;i.e., at stresses below the Orowan bowing stress, disloca-tions will not be able to bow between adjacent particles andpositive strain will not occur. For example, according toEvans and Wilshire,[63] Ni20Cr2ThO2 single crystals willnot creep at stresses below its Orowan bowing stress, and

polycrystalline Ni20Cr2ThO2 will not creep until its Orowanbowing stress is exceeded locally. These results suggest thatan effective stress, which is equal to the applied stress minusthe Orowan bowing stress, governs the creep behavior.

The Orowan bowing stress, bow, can be approximated as

[3]

where G is the shear modulus, b is the burgers vector, and is the mean interparticle spacing. It is important to notethat varies with particle size. Using G 64 GPa,[44], b 0.3 nm,[64] and 44 to 56 nm[65], bow values range from340 to 440 MPa. The significantly higher back stress val-ues observed in this study (up to 645 MPa) suggest that theremust be other defect/dislocation interactions or diffusion bar-riers that contribute to the back stress in addition to theOrowan bowing stress.

d. Creep ruptureIncreased CR tended to increase Tr and f for 954 C

annealed-then-aged samples. The IN 718SPF material wasnot evaluated in creep in this study, but based on previouscreep rupture data,[39] its Tr value is expected to be similarto that exhibited by the 20 pct CR condition. Thus, thereappears to be a limit to the amount of CR deformation thatwill result in increased creep rupture life and f. This limitappears to be near 30 pct CR, as this condition exhibited

the maximumT

r value and a decrease inT

r was observedat 40 pct CR. It is noted that a significant decrease in thegrain size occurred with increased CR from the 40 pct CRcondition to the IN 718SPF condition, and this may be asignificant factor in the creep rupture discrepancy. Althoughprevious findings have correlated a significant increase increep rupture properties with increased special boundaryfractions,[2] further investigation of the effect of GBCD oncreep rupture properties is necessary to indicate the impor-tance of grain boundaries to the creep resistance under thesecreep stress and temperature conditions. Comparing the creepstrain rate and rupture properties, it appears that the 30 pctCR condition for the 954 C annealed-then-aged samplesresults in the most attractive overall creep behavior.

V. SUMMARY AND CONCLUSIONS

IN 718 was processed through sequential increments of CRbetween 0 to 80 pct followed by annealing between 954 Cto 1050 C then aging to evaluate processing-microstructure-property relationships. For annealing temperatures of 954 Cand 1010 C, increased CR led to decreased twin boundaryfractions. This resulted in decreased CSLBs and increasedGHABs with increased CR deformation. An opposite trend

tbow Gb>l


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was observed for 1050 C annealed microstructures, whichexhibited increased twin boundary fractions and CSLBs anddecreased GHABs with increasing CR deformation. This obser-vation was correlated to -phase grain size as the grain sizeincreased significantly with increased annealing temperaturefrom 954 C to 1050 C. The lowest GHAB fraction observedwas 0.4, and therefore, the greatest special boundary fractionrecorded was 0.6. This indicates the potential to GBE IN 718as a majority of special boundaries can be obtained throughcommercially practical TMP processing treatments.

The steady-state creep rate, backstress, and creep-ruptureproperties were measured for the 954 C annealed-then-agedsamples. For the 954 C annealed-then-aged microstructures,the greatest backstress and lowest creep rate values wereexhibited by the 20 pct and 30 pct CR microstructures, andthe 30 pct CR microstructure exhibited the greatest Tr value.For the 1050 C annealed-then-aged microstructures, thegreatest backstress and lowest creep rate values were exhib-ited by the IN 718SPF microstructure. The effective stressexponent values, which incorporated the backstress, suggestedthat the creep deformation mechanism is dependent on effec-tive stresses (e) where the transition point occurs at e 135 MPa. This was independent of both CR deformation and

annealing temperature. The significantly higher backstressvalues in relation to the estimated Orowan bowing stress sug-gests there must be other defect/dislocation interactions ordiffusion barriers that contribute to the backstress.

The RT tensile results revealed the exceptional strength andadequate elongation exhibited by the fine-grained IN 718SPFmaterial. For the 954 C annealed-then-aged samples, the 10to 40 pct CR microstructures exhibited the greatest 650 Ctensile strengths, while no clear trend in 650 C tensile strengthwas exhibited with CR deformation for the 1050 C annealed-then-aged condition. Regardless of processing condition, thetensile and creep specimens exhibited a ductile fracture. Over-all, the 20 pct and 30 pct CR samples exhibit the best ele-vated-temperature properties in the 954 C annealed condition,

while the IN 718SPF exhibited exceptional strength and creepresistance in the 1050 C annealed condition, on par withthe 20 to 30 pct CR 954 C annealed samples.

ACKNOWLEDGMENTS

This work was supported by the NSF Division of MaterialsResearch (Grant No. DMR-0533954). The IN 718SPF sheetswere processed under the supervision of Gaylord Smith (SMC),who along with James Crum (SMC) offered helpful guidancefor this work. The authors are grateful to Dr. Dingqiang Liand Messieurs Serkan Civelekoglu and Jay Spike for their tech-nical assistance.

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the effect of sheet processing on the microstructure, tensile, and creep behavior of inconel alloy...

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