A

CrcemnTr©

K

1

atoiA(rt(dtNbXpTi

0d

Materials Science and Engineering A 473 (2008) 233–237

The effect of thermomechanical processing on the creepbehavior of Alloy 690

C.J. BoehlertDepartment of Chemical Engineering and Materials Science, Michigan State University, East Lansing, MI 48824, USA

Received 22 February 2007; received in revised form 17 March 2007; accepted 20 March 2007

bstract

The effect of thermomechanical processing on the microstructure and elevated-temperature creep behavior of Alloy 690 was investigated.ommercially available sheet was subjected to four cycles of cold rolling to 25% deformation followed by annealing at 1000 ◦C for 1 h. Both the

esultant microstructure and the original microstructure were characterized using electron backscattered diffraction. The thermomechanically pro-essed microstructure exhibited a slightly lower fraction of twins and a smaller average grain size than the original microstructure. Tensile–creepxperiments were performed in an open-air environment at temperatures between 650 and 690 ◦C and stresses between 75 and 172 MPa. Theeasured creep stress exponents (4–5) activation energies (320–368 kJ/mol) suggested that dislocation creep with lattice self-diffusion was domi-

ant. The thermomechanically processed microstructure exhibited significantly worse creep resistance than the original as-processed microstructure.hus, cyclic strain and annealing processing, which has been shown to improve the ductility-dip cracking susceptibility of Alloy 690, is not

ecommended for enhancing the creep resistance. 2007 Elsevier B.V. All rights reserved.

ostruc

rftooaedtaatTs(rr

eywords: Nickel-based alloy; Electron backscattered diffraction; Creep; Micr

. Introduction

Alloy 690, a commercially available nickel-based alloy withnominal composition close to Ni–30Cr–10Fe (wt.%), is attrac-

ive for pressurized-water nuclear reactor components becausef its superior corrosion resistance [1–3]. Due primarily to itsntergranular stress corrosion cracking resistance (IGSCC),lloy 690 is intended to replace Alloy 600 (Ni–16Cr–9Fe

wt.%)) as a steam generator tube material in pressurized-watereactors [4]. One potential means to improve the IGSCC resis-ance of Alloy 690 is through thermomechanical processingTMP) treatments which alter the grain boundary characteristribution (GBCD). Strain-recrystallization-based TMPreatments have resulted in improved IGSCC resistance of purei and Ni-based alloys [5–7]. However, strain-recrystallization-ased TMP investigations of Alloy 690 have been limited [8,9].ia et al. [8] have evaluated the effects of thermomechanical

rocessing on the distributions of twin boundaries in Alloy 690.hey found that the strain and annealing processes significantly

nfluenced the distributions of twins [8]. They performed cold

E-mail address: boehlert@egr.msu.edu.

rtpcid

921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2007.03.082

ture

olling between 5 and 50% followed by annealing at 1100 ◦Cor 5 min. With small strains (5% cold rolled material), thewin boundaries were parts of clusters, and the overall fractionf special boundaries was 0.73. However, with larger amountsf cold rolling deformation almost no twin clusters existed,nd the overall special boundary fraction was 0.47 [8]. Davet al. [9] evaluated the effect of TMP on the microstructure anductility-dip cracking susceptibility of Alloy 690. In their work,he as-received wrought mill microstructure was subjected to

repeated cycle of 25% cold rolling followed by an annealt 1000 ◦C for 1 h. This cycle was repeated four times and theotal reduction of the initial sheet was approximately 67%.he TMP material exhibited a slightly greater percentage ofpecial boundaries (50–55%) than the as-received material40% special boundaries), and there were regions in which theandom boundary network was effectively disrupted. The strain-ecrystallization processed material exhibited a higher ductilityecovery temperature and a higher minimum ductility thanhe material that did not undergo cyclic strain-recrystallization

rocessing. Thus the additional TMP treatment had a benefi-ial impact on the alloy’s resistance to cracking and therebymproved its fracture behavior. A correlation between intact ran-om boundary networks and intergranular brittle fracture modes

2 nd En

wwTtmos

ralwmoepIgaat

2

arcmwiowcauawftstwTdicetbd2bdoat

sm

ubOcfztbtiwhdp5asfticws

smas

3

3

bFTeigawomicrostructures. Fine M23C6 carbides were observed to dec-orate the grain boundaries in both microstructures, see Figs. 1band 2a, while larger carbides were distributed throughout themicrostructures.

Table 1

34 C.J. Boehlert / Materials Science a

as observed, and in regions where this random boundary net-ork had been disrupted, transgranular ductile fracture occurred.heir work suggested that the GBCD and, more specifically, the

opological connectivity of random boundaries, had an effect onaterial resistance to ductility-dip cracking, although there were

ther intervening microstructural factors mentioned in theirtudy.

A complete understanding of the physical mechanismsesponsible for the elevated-temperature creep behavior andssociated microstructure–property relationships of Alloy 690 isacking. In particular, it has yet to be established if the GBCD,hich has been shown to have a significant influence on theechanical deformation behavior, including creep resistance,

f pure Ni and Ni-based superalloy systems [10–14], has anffect on the creep behavior. This work was intended to evaluaterocessing–microstructure–property relationships of Alloy 690.n particular the effect of TMP on the microstructure (GBCD,rain size, etc.) and high-temperature creep behavior was evalu-ted. The material was chosen because it is a commercially usedlloy and, as such, the aim of this study was to show relevanceo actual engineering materials in use.

. Experimental procedure

The as-processed (AP) Alloy 690 sheet was mill annealedt 1066 ◦C. To produce the TMP sheet, the AP sheet mate-ial was subjected to the following strain annealing sequence:old rolling to 25% deformation followed by a solution treat-ent at 1000 ◦C for 1 h followed by air cooling. This sequenceas repeated four times. This sequence was chosen based on

dentical TMP treatments performed by Dave et al. [9]. Theriginal thickness of the AP sheet was approximately 8 mmhile the final thickness of the TMP sheet was 2.7 mm. This

onstituted a total reduction of the AP sheet’s thickness ofpproximately 66%. Bulk chemical analysis was performedsing inductively coupled plasma optical emission spectroscopynd inert gas fluorescence. Each AP and TMP sheet materialas sectioned and metallographically polished to prepare it

or imaging. Spatially resolved electron backscattered diffrac-ion (EBSD) orientation maps were obtained from polishedections using a FEI XL-30 Field Emission Gun scanning elec-ron microscope (SEM). The EBSD hardware and softwareere manufactured by EDAX-TSL, Inc. (Draper, UT, USA).he specimens were ground mechanically by 15, 6, and 1 �miamond suspension for 10 min, respectively, and then pol-shed by 0.06 �m colloidal silica for 60 min. The GBCD washaracterized using the orientation maps to determine the ori-ntation relationships between grains. For each map, morehan 500 grain boundaries were analyzed using a step size ofetween 0.5 and 2 �m. Low-angle boundaries (LABs) wereefined as those boundaries containing misorientations betweenand 15◦. Brandon’s criteria [15] were used to distinguish

etween general high-angle boundaries (GHABs) and coinci-

ent site lattice boundaries (CSLBs). The reported fractionsf GHABs, LABs, CSLBs, and twins (�3) were the aver-ged values taken from several orientation maps, performed onhe cross-sections, rolling faces, or longitudinal sections of the

M

N

6

gineering A 473 (2008) 233–237

heets. Grain size was determined using the mean line interceptethod [16,17].Blanks from the AP and TMP sheet materials were machined,

sing either electrodischarge machining or a mill, into a flat dog-one geometry used for tensile and tensile–creep specimens.pen-air tensile–creep experiments were performed on verti-

al Applied Test System, Incorporated (Butler, PA, USA) loadrames with a 20:1 lever-arm ratio. Applied Test System singleone furnaces were used to heat the specimens to within ±3 ◦C ofhe target temperature. Specimen temperatures were monitoredy three chromel–alumel type K thermocouples located withinhe specimen’s reduced section. Creep strain was monitored dur-ng the tests using a linear variable differential transformer thatas connected to a 25.4 mm gage length Applied Test Systemigh-temperature extensometer. The extensometer was attachedirectly to the gage section of each sample. The testing tem-eratures and stresses ranged between 650 and 690 ◦C and0–172 MPa, respectively. All creep specimens were loaded par-llel to the rolling direction, and the experiments were conducteduch that the specimens were soaked at the creep temperatureor at least 60 min prior to applying load in order to minimizehe thermal stresses. After the creep strain had proceeded wellnto the secondary regime, either the load or temperature washanged or the creep test was discontinued. The tested specimensere cooled under load to minimize recovery of the deformed

tructures.Room-temperature (RT) tensile tests were performed at a

train rate of 10−3 s−1 using an Instron 4206 tensile testingachine. Strain was measured during the tensile tests with

n extensometer attached directly to the gage section of theample.

. Results and discussion

.1. Microstructure

Table 1 lists the measured alloy composition. Fig. 1 illustratesackscattered electron SEM images of the AP microstructures.ig. 2 illustrates backscattered electron SEM images of theMP microstructures. The TMP material exhibited an averagequiaxed grain size of 8.9 �m and the AP material exhib-ted an average grain size of 16.1 �m. It is noted that therain boundaries were intact and uncracked in both the APnd TMP microstructures. The AP microstructure was some-hat banded which has also been observed in previous studiesf this alloy [9]. Banding was also exhibited in the TMP

easured Composition of the IN 690 alloy in wt.%

i Cr Fe Mn Si C Si Al Mo Co O N

0.3 29.2 10.1 0.18 0.05 0.04 0.05 0.2 0.09 0.02 0.01 0.01

C.J. Boehlert / Materials Science and Engineering A 473 (2008) 233–237 235

Fm

mwfitTiwt0tamCcwtb[

Fn

3

A(m7ci[mwat

tobserved even for samples deformed to over 15% strain. Fig. 7a

ig. 1. Backscattered electron SEM images (a) and (b) of the as-processed (AP)icrostructure.

Fig. 3 illustrates an EBSD orientation map for the APicrostructure, which exhibited an equiaxed microstructurehere the grain orientations were distributed fairly evenly. Polegure analysis indicated that the microstructure was not strongly

extured. Fig. 4 illustrates an EBSD orientation map for theMP microstructure. The GHAB fractions were the majority

n both microstructures. The volume fraction of twins observedas 0.10 and 0.05 for the AP and TMP microstructures, respec-

ively. The volume fraction of LABs observed was 0.25 and.30 for the AP and TMP microstructures, respectively. Thus,he TMP material exhibited a slightly lower fraction of twinsnd a slightly higher fraction of LABs compared to the APaterial. If we consider that the combination of the LAB andSLB fractions (<�29) constitute the overall fraction of spe-ial boundaries, then the maximum special boundary fractionas approximately 0.35 for both the AP and TMP microstruc-

ures. This value is significantly less than the maximum special

oundary fraction value, 0.5–0.55, measured by Dave et al.9].

ibc

ig. 2. Backscattered electron SEM images (a) and (b) of the thermomecha-ically-processed (TMP) microstructure.

.2. Tensile and creep behavior

Representative RT tensile curves are illustrated in Fig. 5. TheP microstructure exhibited a slightly greater yield strength

364 MPa) than the TMP microstructure (344 MPa). The ulti-ate tensile strengths for both microstructures were close to

25 MPa and significant work hardening was exhibited. Thereep strain-life history resembled that for most metals exhibit-ng three stages of creep: primary, secondary, and tertiary18,19]. Fig. 6 illustrates strain versus time plots for each of theaterials at T = 650 ◦C and σ = 172 MPa. Once the creep stressas reached there appeared to be a short incubation period beforepositive creep strain was achieved. This was common for both

he AP and TMP conditions.The AP material exhibited significantly greater creep resis-

ance than the TMP material. No grain boundary cracking was

llustrates the minimum creep strain rate versus applied stressehavior at T = 650 ◦C, while Fig. 7b illustrates the minimumreep strain versus inverse temperature response at σ = 125 MPa.

236 C.J. Boehlert / Materials Science and Engineering A 473 (2008) 233–237

Fig. 3. EBSD orientation map for the as-processed (AP) microstructure.

Fig. 4. EBSD orientation map for the thermomechanically processed (TMP)microstructure.

Fi

TAa0cevamsocc(

ctm

Fp

ig. 5. RT stress vs. strain behavior for as-processed (AP) and thermomechan-cally processed (TMP) specimens.

he TMP condition exhibited lower creep resistance than theP material at all the applied stress and temperatures evalu-

ted. In particular, the minimum creep rate was approximately.5 orders of magnitude faster for the TMP microstructureompared with the AP condition. The measured creep stressxponents were between 4 and 5 at T = 650 ◦C, and the Qappalues were between 320 kJ/mol (TMP) and 368 kJ/mol (AP)t σ = 125 MPa. Such creep parameters are similar to thoseeasured for other similar nickel-based superalloys [14] and

uggest a dislocation climb mechanism [20,21] is active. Basedn the measured exponents and activation energies, dislocationreep with lattice self-diffusion was the suggested dominantreep mechanism over the entire applied stress range examined75 MPa < σ < 172 MPa).

It is felt that the smaller grain size and the lower twin fractions

ould have contributed to the worse creep resistance exhibited byhe TMP material. The presence of the grain boundary carbides

ay have also contributed to the creep behavior discrepancy as

ig. 6. Creep strain vs. time plots for as-processed (AP) and thermomechanicallyrocessed (TMP) specimens at T = 650 ◦C and σ = 172 MPa.

C.J. Boehlert / Materials Science and

Fmr

ti

4

aimwi

mtptmrfct

A

tNU

R

[[

[

[

[

[[[

[

ig. 7. Creep plots used to determine the creep parameters n and Qapp: (a)inimum creep rate vs. applied stress at T = 650 ◦C and (b) ln minimum creep

ate vs. 1/T at σ = 125 MPa.

he TMP material exhibited more grain boundary area, due tots fine grain size, than the AP material.

. Summary

This work evaluated the effect of TMP on the microstructurend creep behavior of Alloy 690. The AP microstructure exhib-

ted a slightly larger volume fraction of twins than the TMP

icrostructure, yet the overall fraction of special boundariesas similar for both microstructures. The AP material exhib-

ted a grain size almost twice that of the TMP material. The AP

[

[[

Engineering A 473 (2008) 233–237 237

icrostructure exhibited significantly greater creep resistancehan the TMP microstructure. For the applied stresses and tem-erature evaluated, the measured creep parameters suggestedhat dislocation creep was the dominant creep deformation

echanism. Overall, the TMP sequence degraded the creepesistance, refined the equiaxed grain size and promoted a lowerraction of twins in the microstructure. Thus TMP involvingyclic cold rolling and annealing is not suggested for improvinghe creep resistance of Alloy 690.

cknowledgements

This work was supported by the National Science Founda-ion through grant DMR-0533954. The author is grateful to Mr.athan Eisinger (Special Metals Corporation, Huntington, WV,SA) for directing the alloy processing effort.

eferences

[1] R.S. Dutta, R. Tewari, Br. Corros. J. 3 (34) (1999) 201–205.[2] A.J. Sedricks, J.W. Schultz, M.A. Cordovi, Boshoku Gijutsu 28 (1979)

82–95.[3] C. Cheung, U. Erb, G. Palumbo, Mater. Sci. Eng. 185A (1994) 39–43.[4] M. Thuvander, K. Stiller, Mater. Sci. Eng. 281A (2000) 96–103.[5] G. Palumbo, K.T. Aust, Acta. Metall. Mater. 11 (38) (1990) 2343–2352.[6] G. Palumbo, U.S. Patent 5,817,193 (1998).[7] B. Alexandreanu, B.M. Capell, G. Was, Mater. Sci. Eng. 300A (2001)

94–104.[8] S. Xia, B.X. Zhou, W.J. Chen, W.G. Wang, Scripta Mater. 54 (2006)

2019–2022.[9] V.R. Dave, M.J. Cola, M. Kumar, A.J. Schwartz, G.N.A. Hussen, Welding J.

(January 2004) 1-S–5-S (American Welding Society and Welding ResearchCouncil).

10] E.M. Lehockey, G. Palumbo, Mater. Sci. Eng. 237A (1997) 168–172.11] G.S. Was, V. Thaveeprungsriporn, D.C. Crawford, J. Met. (1998) 44–

49.12] V. Thaveeprungsriporn, G. Was, Metall. Mater. Trans. A 28A (1997)

2101–2112.13] G. Palumbo, K.T. Aust, in: D. Wolf, S. Yip (Eds.), Special Properties of

� Grain Boundaries, Materials Interfaces: Atomic Level Structure andProperties, Chapman and Hall, NY, 1989, pp. 190–211.

14] C.J. Boehlert, D.S. Dickmann, N.C. Eisinger, Metall. Mater. Trans. A 37A(1) (2006) 27–40.

15] D.G. Brandon, Acta Metall. 14 (1966) 1479–1484.16] J.E. Hilliard, Met. Prog. 78 (1964) 99–100.17] Standard Test Methods for Determining Average Grain Size, ASTM Des-

ignation E112-96e3, American Society for Testing and Materials, WestConshohocken, PA.

18] R.W. Evans, B. Wilshire, Creep of Metals and Alloys, The Institute of

Metals, New York, NY, 1985.

19] R.W. Hertzberg, Deformation and Fracture Mechanics of EngineeringMaterials, fourth ed., John Wiley and Sons, New York, NY, 1996.

20] T.G. Langdon, P. Yavari, Acta Metall. 30 (1982) 881–887.21] J. Weertman, Trans. Am. Soc. Met. 61 (1968) 681–694.