Characterization of low cycle fatigue damage in 9Cr–1Mo ferritic steelusing X-ray diffraction technique

Sanjay Rai, B.K. Choudhary, T. Jayakumar, K. Bhanu Sankara Rao, Baldev Raj*

Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, India

Received 23 November 1998; accepted 3 December 1998

Abstract

X-ray diffraction (XRD) technique has been used to characterize the low cycle fatigue (LCF) damage in 9Cr–1Mo ferritic steel. In thisstudy, full-width at half-maximum (FWHM) of the XRD peak has been measured for assessing the fatigue damage. Fully reversed total-axial-strain controlled fatigue tests have been performed at ambient temperature (300 K) at strain amplitudes of^ 0.25%,^ 0.50% and^0.75%. FWHM measurement of {3 1 0} plane has been carried out on specimens interrupted at different fatigue life fractions, whichrepresent the various stages of deformation and fracture such as cyclic hardening, cyclic softening, saturation and surface crack initiationand propagation. The cyclic hardening, which occurred in the early stages of fatigue deformation, exhibited broadening of diffraction profileand a rapid increase in the FWHM at all strain amplitudes. Also, large oscillations were observed in thed(2u) vs. sin2C curves. With furthercycling, FWHM remained almost constant in the softening and saturation stages. Finally, at the onset of rapid stress drop in the cyclic stressresponse and cusp formation in the compression portions of stress–strain hysteresis loops, which indicate surface crack initiation andpropagation, a rapid decrease in FWHM was observed. This drop in FWHM is attributed to relieving of microstresses owing to surfacecrack initiation and propagation. Relieving of micro stresses has also been confirmed by a significant reduction in oscillations in thed(2u ) vs.sin2C curves.q 1999 Elsevier Science Ltd. All rights reserved.

Keywords:Diffraction; Low cycle fatigue; Stress–strain

1. Introduction

Non-destructive evaluation (NDE) techniques are beingused for the assessment of initial microstructure andmechanical properties and their subsequent degradationduring service on exposure to elevated temperature understatic and/or cyclic loading. Advanced NDE techniquessuch as acoustic emission [1,2], ultrasonic attenuation andvelocity measurements [3,4], magnetic Barkhausen emis-sion [5,6], positron annihilation technique [7,8] and smallangle X-ray and neutron scattering [9–14] have shown theirpotential for the characterization and evaluation of micro-structure and creep/fatigue damage. Analysis of X-raydiffraction (XRD) peak profiles indicated that full-width athalf-maximum (FWHM) is sensitive to the variation inmicrostructure and stress–strain accumulation in the mate-rial [9,12–14]. Residual stresses of different types arepresent in the materials [15]. The first order stresses,commonly known as macro stresses, are homogeneous

over large scale involving many grains i.e., a few hundredsof microns. The main effect of these stresses on the XRDpeak is to shift the peak position. The second and third orderstresses (microstresses) are homogeneous over smalldomains such as a part of a grain, with dimensions ofabout tens of microns. Stresses of the second and thirdorder manifest themselves in terms of broadening of theXRD peak i.e., increase in the FWHM.

The high cycle fatigue (HCF) behavior of 17-4PH steel inshort-peened condition and low cycle fatigue (LCF) beha-vior in Cr–Mo–V steel were investigated using XRD tech-nique [12]. In 17-4PH steel, the HCF strength was correlatedwith FWHM and residual stress. Both FWHM and highcompressive residual stress decreased gradually withincreasing number of cycles followed by a rapid decreaseat the end of fatigue life. Under LCF loading in Cr–Mo–Vsteel, FWHM remained nearly constant in the early stage ofcyclic loading followed by a rapid decrease at highernumber of cycles [12]. Quesnel et al. [9] observed adecrease in FWHM with cyclic softening in a cold rolledhigh strength low alloy (HSLA) steel. In hot rolled con-dition, the alloy exhibited a stable stress response accom-panied with an increase in FWHM after fatigue cycling [9].

International Journal of Pressure Vessels and Piping 76 (1999) 275–281

IPVP 1928

0308-0161/99/$ - see front matterq 1999 Elsevier Science Ltd. All rights reserved.PII: S0308-0161(98)00140-9

* Corresponding author. Tel.:191-4114-40234/40301; fax:191-4114-40301/40360.

E-mail address:dmg@igcar.ernet.in (B. Raj)

9Cr–1Mo ferritic steel finds wide application inpetroleum, chemical processing and thermal power plants.This is also a candidate material for steam generator appli-cations in liquid metal cooled fast breeder reactors(LMFBRs). The components of steam generators that oper-ate at elevated temperatures are often subjected to repeatedthermal stresses as a result of temperature gradients whichoccur on heating and cooling during start-ups and shut-downs or during variations in operating conditions of areactor. These transient thermal stresses give rise to LCFdamage. Detailed investigation on the LCF behavior of9Cr–1Mo steel forging indicated that the alloy exhibitsinitial cyclic hardening followed by a gradual softeningand a saturation stage [16,17]. The stage-I crack initiationand stage-II crack propagation remained transgranular. Thepresent investigation is aimed at examining and understand-ing the influence of LCF deformation and fracture in 9Cr–1Mo steel at ambient temperature on FWHM measured byXRD technique using a portable back reflection type X-raydiffractometer. Use of portable XRD system on laboratoryLCF specimens would help in extending this technique forin-service fatigue damage assessment of the components.More emphasis has been given to identify and correlatethe fatigue crack initiation and propagation with FWHM,which limits the useful fatigue life of the material.

2. Experimental details

A nuclear grade 9Cr–1Mo ferritic steel plate of 20 mmthickness was used in this investigation. The chemicalcomposition of the steel is given in Table 1. The plate wassubjected to normalizing at 1223 K for 15 min followed byair cooling and then tempering at 1023 K for 2 h followedby air cooling. The microstructure at the end of normalizingand tempering (N1 T) treatment was composed oftempered lath martensite. The average grain size measuredby linear intercept method was found to be,22mm.Following the heat treatment, axial-cylindrical-ridge typeLCF specimens of 10 mm gauge diameter and 25 mmgauge length were machined in the rolling direction of theplate. LCF tests were conducted in air under fully reversedaxial-total-strain control mode in an Instron 1343 servohydraulic testing system using triangular waveform. Testswere performed at room temperature at strain amplitudes of^ 0.25%,^ 0.50% and^ 0.75% employing a constantstrain rate of 1× 1023 s21. Stress–strain hysteresis loopswere recorded continuously during the LCF tests. The testswere interrupted at various life fractions to characterize

the various stages of LCF deformation and fracture. LCFtests were performed first in tension followed by compres-sion cycle, and in each interruption, specimens wereremoved after completing compression cycle. The FWHMmeasurement of XRD peaks was carried out in the heat-treated virgin as well as LCF tested specimens after eachinterruption. In order to avoid the scatter in the data, onlyone specimen was used at each strain amplitude.

XRD studies were carried out using a portable XRDsystem (Model RIGAKU MSF-2M) having a back reflectiontype goniometer with a 2u scan range from 1408 to 1708.The radiation used for the study was Co–Ka and the {h k l}plane considered for analysis was {3 1 0}. For FWHMmeasurements, the counting time was optimized to about100 s to reduce statistical errors. The error in peak locationwas less than 0.018 and the uncertainty in the peak widthafter correcting for instrumental broadening was less than4%. Lorentzian curve fitting was used to calculate peakwidth [15]. Macro residual stress measurements werecarried out using multipleC method. The fourC anglesused for the measurements were 08, 108, 208 and 308 [15].Location of the peak top was carried out using parabolicpeak top method after incorporating background andabsorption Lorentz Polarization (LP) corrections. Measure-ments were carried out at four circumferential locations inthe center of gauge length and their average value has beenreported in each interrupted condition. After crack initiationFWHM measurement was carried out around the crack. TheX-ray spot size used in this investigation was 2× 2 mm.

3. Results and discussion

Typical XRD intensity profiles of {3 1 0} planes forvirgin heat treated material and LCF specimen tested at atotal strain amplitude of^ 0.75% after various interrup-tions are shown in Fig. 1(a)–(e). Fig. 2(a)–(c) shows changein shape of hysteresis loop with number of cycles. In theN 1 T condition, diffraction from {3 1 0} plane exhibitedsharp Ka1–Ka2 doublet (Fig. 1(a)). At all strain amplitudes,broadening of peaks was observed after initial cycling(e.g., at strain amplitude of̂ 0.75% after 7 cycles inFig. 1(b)). With further fatigue cycling to larger numberof cycles, no significant change in the diffraction peakshape was noticed (e.g., at strain amplitude of^ 0.75%after 7 and 107 cycles in Fig. 1(b) and (c), respectively).After 400 cycles, an appreciable increase in peak width wasobserved (Fig. 1(d)). The broadening of peaks is manifestedas overlapping of the Ka1–Ka2 doublet as a result of fatigue

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Table 1Chemical composition of 9Cr–1Mo steel

Element C Si Mn S P Cr Mo Fe

Amount (wt.%) 0.110 0.490 0.455 0.002 0.008 8.363 0.930 Balance

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Fig. 1. Typical X-ray intensity profiles obtained in the (a) heat treated virgin, and LCF specimens tested at total strain amplitude of^0.75% after variousinterruptions of (b) 7, (c) 107, (d) 400 and (e) 603 cycles.

cycling (Fig. 1(b)–(d). Relatively sharper diffraction peaksare observed at all strain amplitudes after cusp appeared inthe compression portions of the stress–strain hysteresisloops (e.g., at strain amplitude of̂ 0.75% after 603 cyclesin Fig. 1(e)).

The variation in the FWHM values with number of cyclesin the various stages of cyclic loading together with thevariation in peak tensile stress amplitude with number ofcycles is shown in Fig. 3. FWHM measurements werecarried out at four circumferential locations in the centerof gauge length. Variation in FWHM values at differentlocations was less then 5%. Hence, an average value hasbeen reported in each interrupted condition. 9Cr–1Mo steelin the normalized and tempered condition exhibited initialcyclic hardening followed by a progressive cyclic softeningand a saturation stage. The cyclic softening occurred by agradual decrease in stresses in both tension and compressionportions of the cycle. Further, the rate of reduction in stressvalues during cyclic softening stage decreased with increas-ing number of cycles giving rise to a saturation stage athigher number of cycles. In the saturation stage, the alloyexhibited a nearly stable stress–strain hysteresis loop with aconstant stress value over a large number of cycles. Thesaturation stage persisted until the surface crack initiationand propagation impaired the load bearing capacity of thespecimens as indicated by a rapid fall in the stresses (Fig. 3)and cusp formation in the stress–strain hysteresis loops (Fig.2). It can be seen from Fig. 3 that there is a rapid increase inFWHM corresponding to initial cyclic hardening followedby a nearly constant value in the softening and saturationstages. Finally, with the onset of surface crack initiation andpropagation, a rapid decrease in FWHM was observed at allthe strain amplitudes.

All the virgin LCF specimens exhibited high compressivemacro residual stress (approximately 300 MPa) on the

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Fig. 2. Stress–strain hysteresis loops showing cusp formation at total strainamplitudes of (a)̂ 0.75% after 603 cycles, (b)̂ 0.5% after 1660 cyclesand (c)^0.25% after 6253 cycles.

Fig. 3. Variation of stress amplitude and the corresponding FWHM as a function of number of cycles at total strain amplitudes of^0.75%,^0.5% and^ 0.25%.

surface. The high compressive residual stresses present inthe surface of the specimens were caused by machining ofthe specimens. These stresses were relieved in the initialcycling, and low (approximately zero) residual stresseswere present throughout the fatigue cycling. Further,d(2u) vs. sin2C curves exhibited large oscillations owingto fatigue cycling compared to that observed in the virginspecimen (Fig. 4). The oscillations ind(2u ) vs. sin2C alsopersisted in the softening and saturation stages. Similaroscillations ind(2u ) vs. sin2C plot were observed by Ques-nel et al. [9] in cold rolled HSLA steel as a result of fatiguecycling. It was suggested that the strong oscillations ind(2u) vs. sin2C indicate an orientation dependent micros-tress distribution. The oscillations can arise from the plasticdeformation of favourably oriented grains, while thesurrounding grains deform only elastically at small strains.After release of applied stress, a microstress distributiondevelops as elastically deformed grains attempt to relaxby exerting stress on those grains, which have already gotplastically deformed. At high strains, favourably orientedgrains deform at the largest strain amplitudes, while othersdeform at lower strain amplitudes resulting in the develop-ment of inhomogeneous structure and microstress distribu-tion after stress release. Following fatigue crack initiationand propagation, the oscillations ind(2u) vs. sin2C curvedecreased considerably indicating relaxation of micro-stresses (Fig. 4).

The initial cyclic hardening observed in the presentinvestigation is in agreement with that reported for9Cr–1Mo steel forging [16,17], modified 9Cr–1Mosteel [18] and 12Cr–Mo–V steel [19]. The rapidincrease in FWHM values corresponding to initial cyclichardening can be ascribed to dislocation–dislocationand precipitate-dislocation interactions [16,17] resultingin the increase in number of dislocation tangles anddislocations entangled with matrix precipitates. Suchinteractions would enhance the defect density in thestructure thereby resulting in an increase in FWHM in

the cyclic hardening stage. The initial cyclic hardeningincreased with increase in strain amplitude. Theinfluence of strain amplitude on FWHM can be seenin Fig. 3.

The progressive cyclic softening observed in the presentinvestigation (Fig. 3) is in agreement with the reportedobservation in 9Cr–1Mo [16,17,20,21] and modified 9Cr–1Mo [18,22–24] steels and 12Cr–Mo–V steel at ambient[21] and elevated [25] temperatures. The gradual decreasein the stress values with increasing number of cycles at roomtemperature can arise from the annihilation of dislocationsintroduced during martensitic transformation, developmentof cell structure within the martensite laths and transforma-tion of original lath structure to cell structure [16,17,19–24].Such a dislocation arrangement owing to fatigue cyclingwould lead to regions of high dislocation density i.e., cellwalls and regions of low dislocation density i.e., cell inter-ior. In this softening stage, FWHM exhibited a constantvalue. In the saturation stage, where constant stress responseover a large number of cycles results from a balancebetween cyclic hardening and cyclic softening in 9Cr–1Mo steel, a marginal increase in FWHM was observed athigher strain amplitudes.

Quesnel et al. [9] observed cyclic softening and adecrease in FWHM in cold rolled HSLA steel owing todecrease in the dislocation density through the formationof dislocation cells. In the hot rolled condition, the alloyexhibited a constant stress response and an increase inFWHM during cyclic loading. It was suggested that thebroadening of intensity profile could result from microstruc-tural change. This result brings out an important point thatconstancy in mechanical properties alone cannot representthe microstructural state of the material [9]. Under LCFloading, accumulation of a small amount of plastic straintakes place in each cycle, which increases with increasingnumber of cycles. The observed marginal increase inFWHM in the saturation stage can be attributed to the accu-mulation of plastic strain.

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Fig. 4. d(2u) vs. sin2C curves for virgin and LCF specimen tested at strain amplitude of^ 0.75% before (400 cycles) and after surface crack initiation andpropagation (603 cycles).

At all strain amplitudes, a rapid decrease in the FWHMvalues has been observed following surface crack initiationand propagation, indicated by cusp formation in thecompression portion of stress–strain hysteresis loop (Fig.2). Goto [12] reported a rapid decrease in FWHM towardsthe end of the fatigue life in 17-4PH steel tested at stresseshigher than the fatigue limit. Decrease in the FWHMresponse near the end of LCF tests in Cr–Mo–V steel hasalso been observed [14]. The rapid decrease in FWHMfollowing surface crack initiation and propagation in 9Cr–1Mo steel observed in the present investigation is attributedto the relaxation of microstresses. Decrease in the FWHMhas been reflected in the sharpening of the diffraction profile(Fig. 1(e)) and a considerable decrease in the oscillations inthe d(2u ) vs. sin2C curves (Fig. 4).

Among the many criteria used to define the fatigue life ofa laboratory specimen in strain controlled low cycle fatiguecondition, Rao et al. [26] observed that the number of cyclesto 20% stress reduction in saturation stress and number ofcycles for cusp formation provide the two best criteria todefine fatigue life. Also, the fatigue life defined using thesetwo criteria were found to be similar, which remains within10% of the number of cycles for separation of specimen intotwo pieces [26]. In the present investigation, the specimenswere interrupted at the number of cycles corresponding todetectable cusp formation in the stress–strain hysteresisloops, which in turn corresponds to stage-I crack initiationand propagation. Once a crack of sizeable length formed,the stage-II crack propagation would be rather faster andwill take smaller number of cycles for separation of speci-men into two pieces. Therefore, the definition of fatigue lifeas the number of cycles for cusp formation would provide asafe and useful fatigue life of a component. The rapiddecrease in FWHM caused by crack initiation and propaga-tion observed in the present investigation suggests that thisparameter can be used to indicate the impending failurecaused by fatigue damage and would also help in preventingthe catastrophic failures of components.

4. Conclusions

XRD technique was used to assess the LCF damage in9Cr–1Mo ferritic steel. The various stages of LCF deforma-tion and damage such as cyclic hardening, cyclic softening,saturation and surface crack initiation and propagation wereidentified using FWHM. The cyclic hardening, whichoccurred in the early stage of LCF cycling, exhibited broad-ening in the XRD profile and a rapid increase in the FWHM.Increase in FWHM has been associated with the oscillationsin the d(2u) vs. sin2C curves at all strain amplitudes. TheFWHM increased with increase in the amount of hardeningand strain amplitude. In the cyclic softening stage, no sig-nificant change in FWHM was observed. The saturation stage,where the stress value remained constant for a large numberof cycles owing to the formation of stable dislocation

substructure, displayed a marginal increase in FWHM. Atall strain amplitudes, a resharpening of the diffractionprofile (i.e., a rapid decrease in FWHM) was observedwith the onset of fatigue crack initiation and propagation.Also, a considerable decrease in oscillations in thed(2u) vssin2C curves was observed after crack initiation and propa-gation. This suggests that the FWHM measurements can beused to indicate the impending failure caused by fatiguedamage, which would help in preventing the catastrophicfailure in components during service.

Acknowledgements

Authors are thankful to Dr. Placid Rodriguez, Director,Indira Gandhi Centre for Atomic Research, Kalpakkam forhis constant encouragement and support. Authors are alsothankful to Dr. S.L. Mannan, Head, Materials DevelopmentGroup and Mr. P. Kalyanasundaram, Head, Division for PIEand NDT Development, for many useful discussions.

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