Relationship between microstructure andacoustic emission in Mn-Mo-Ni A533B steelC. B. Scruby, C. Jones, J. M. Titchmarsh, and H. N. G. Wadley

The sources of acoustic emission in a steel of theA533B composition with low sulphur content have .been investigated using two different recordingtechniques. First, systematic narrow band (0·1-1'OMHz) measurements were made of acousticemission activity during uniaxial tensile tests ofspecimens with a range of controlled microstructures.Mechanical properties and fracto graphic data wereanalysed, and the most probable deformation andfracture mechanisms deduced for these micro-structures. It appears that only a few processesgenerate detectable emission in this material. Forinstance, high-velocity dislocation motion can bedetected during the yield of slow-cooled or long-tempered specimens, and cleavage fracture in rapidlyquenched specimens. A model is presented to relatethe magnitude of the emission to the dynamicoperation of dislocation and fracture sources. Second,very wide band (25 MHz) calibrated measurementswere made of the most energetic emission sources inrapidly quenched material. These indicate the sourcesto be consistent with cleavage microcracking. Atypical micro crack propagated over a distance of"" 50 Jlm with an average velocity of 450 m s -1 .

MSj0627© 1981 The Metals Society. Manuscript received 8 August 1980;in final form 10 February 1981. C. B. Scruby, J. M. Titchmarsh,and H. N. G. Wadley are with AERE Harwell, Oxon, and C. Jonesis in the Department of Metallurgy, University of Bath.

When a sensitive piezoelectric transducer is attached to ametal undergoing deformation, motion of the surface canoften be detected. This motion is very small in amplitude,with displacements of 10-13 m or less, and can also be veryrapid, with frequencies extending to 50 MHz or more.1 Thissurface motion is called acoustic emission and is believed tobe generated by localized relaxations in the elastic straindistribution within the metal. 2 These relaxations are usuallyshort lived and occur, for instance, when dislocations orcracks rapidly propagate. It is probable that every mobiledislocation and every growing crack will cause thegeneration of elastic waves that propagate to the surface ofmetal. However, due to the limited sensitivity and frequencyresponse of piezoelectric transducers (typically "" 10-14 mand ;:;;;1 MHz) only some of the deformation or fractureprocesses that occur are detected1 in most acoustic emissiontests.

Thus, the acoustic emission technique can only be used to

investigate deformation and fracture processes whichgenerate surface displacements greater than the detectionthreshold, and whose frequency content is at least partiallycovered by the range of the detection instrumentation.These two criteria for detectability are partly responsible forthe strong dependence of acoustic emission activity upontesting variables (sample geometry, strain, strain rate,temperature) and metallurgical variables (crystal structure,composition, grain 'size, precipitate structures, etc.). Whilethere is now a reasonable empirical understanding of someof the effects of test variables such as strain rate or gaugevolume, the effects of metallurgical variables are much lessclearly understood, especially in ferritic steels where theoccurrence of the austenite to ferrite transformation oncooling from above "" 850°C results in complex micro-structures. 2

From the practical viewpoint, when acoustic emission isused to detect growing flaws, it is most important that oneshould have a good understanding of the origins of thedetectable acoustic emissions from ferritic steels. Previousstudies have indicated that some important types of crack-growth process, e.g. ductile fracture by void growth andcoalescence, may not be detectable in all ferritic steels.3 Inthese materials one must look for acoustic emission fromplastic zone deformation processes in order to detec.t thepresence of growing flaws.

Detailed metallurgical studies of the sources of acousticemission during plastic deformation have indicated that inferritic steels the fracturejdecohesion of inclusions4,5 (andpossibly carbide precipitates) together with the suddenmotion of groups of dislocations in the vicinity of generalyield 5 are detectable under some conditions. There isalso good evidence that acoustic emission is generatedduring brittle fracture below the ductile to brittletransition temperature or following temper or hydrogenembrittlement.6 Clark and Knote have suggested that atleast one ductile mechanism of fracture, that of alternatingshear fracture, may also be a source of energetic emission.Both the occurrence of these emission sources and theproperties of the emission they generate may change in asteel of fixed composition but which undergoes differentthermo mechanical treatments before testing.5

In this work tensile specimens were used to studysystematically the effect of changes in microstructure of anexperimental ferritic steel of fixed composition uponacoustic emission activity. The composition chosen forstudy was 0'2%C, 1·3%Mn, 0'5%Mo, 0·5%Ni and wassimilar to that of A533B. The sulphur content was reducedto a minimum and the forging treatments adjusted to avoidlarge numbers of elongated MnS inclusion stringers, aprominent acoustic emission source in some experiments.Thus, the experiments were designed to determine theacoustic emission from metal deformation processes in theabsence of emission from elongated MnS inclusions.

Metal Science June 1981 241

242 Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel

Table 1 Bulk chemical composition of Mn-Mo-Ni aHoy, wt-%

Element C Si Mn S Mo Ni Cu N o

Content 0·24 0·27 1·35 0·0045 0·5 0·55 0·02 0·001 0·005

Heat treatmentsSpecimens were given one of two types of heat treatment. Inone, dumb bell and Yobell specimens were austenitized I hat 1000°C in a dynamic argon atmosphere and cooled atvarying rates. The differing cooling rates were achieved byquenching different specimens into either iced brine at-10°C, 10%NaOH solution at room temperature,unstirred water, oil, or by air cooling or furnace cooling at'" 50 K h -1. These treatments resulted in a wide range of

molybdenum is shown in Fig. 2. The spacings between theMo-rich areas varied with orientation in part as aconsequence of the non-uniform forging deformations; onaverage, the spacing was ",0·5 mm. Silicon and manganesecosegregated with the molybdenum. Attempts were made toreduce the segregation by prolonged austenitizingtreatments at temperatures above I10aoC. These weresuccessful at reduCing the segregation to insignificant levelsbut resulted in an unacceptably large prior austenite grainsize. Consequently, for the studies reported here, somealloy-element segregation was present.

direction.--forging--.

Specimen geometriesTwo specimen geometries were used. For the majority oftests where narrow band measurements of relative acousticemission activity were made, a modified dumb bell specimenwas used (Fig. 3a). The tensile axis of the specimen wasparallel to the solidification direction of the original ingot.The gauge was 2 mm long x 3 mm dia. and this short gaugesection was used in order both to simplify interpretation bylocalizing the emission source and to facilitate comparisonwith measurements of the Yobell specimen (Fig. 3b) whichhad been developed for wide band measurements of theacoustic emission wave form.8 It was used here tocharacterize the more energetic emission sources found fromnarrow band experiments.

The majority of tests in the study were performed usingconventional 'narrow band width' measurements withpiezoelectric transducers. These measurements areinherently limited and provide only qualitative informationabout the acoustic emission source. To overcome thelimitations of the narrow band technique, additional testswere performed with calibrated measurements up to25 MHz, using capacitance transducers, in order tocharacterize the most energetic emission sources.

EXPERIMENTALMaterialA 100 kg vacuum-degassed ingot of a Mn-Mo-Ni steel ofA533B nominal composition was produced by Ross andCatherall Ltd using a low-sulphur Japanese Electrolyticiron. The bulk chemical composition of the ingot wasdetermined at three positions across the cross-section foreach of five equally spaced locations along the ingot length.These 15 analyses showed there was no significantmacrosegregation. The mean concentration· values fordetected elements are given in Table 1. Chemical analysisindicated the phosphorus concentration was less than200 ppm, the trace impurities Sn, Sb, and As were less than100 ppm, while Ti, V, Nb, and Cr were all less than 40 ppmin concentration.

The ingot, originally '" 60 cm long and '" 16x 16 cm incross-section, was cut'into four shorter billets each", 13 cmlong. These were heated at 1050°C for 2 h and press forgedto 5 cm thick plate so that the plane of the final platecontained the ingot solidification direction.

Small samples of the forged plate were examined forevidence of microsegregation. The metal samples werepolished by standard metallographic techniques, lightlyetched in 2% nital, and the surface composition measuredusing a Cameca X-ray microprobe analyser. Using anelectron beam with a spot size of '" I Jlm, significantvariations in Mo, Mn, and Ni concentrations were detected(Fig. 1). Molybdenum showed the severest segregation;some regions contained up to four times the mean Moconcentration. The segregation in two dimensions of the

1 Relative concentration profiles for Si, Ni, Mn, andMo in hot-forged material; major alloying elementsshow differing degrees of microsegregation

2 Segregation of Mo in Mn-Ni-Mo steel after hotforging; darkest shading corresponds to highestconcentration of Mo

Metal Science June 1'981

CD==-I" \([J:Jomm2mm 5mm dia.

(a)

45mm

16mm

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 243

cooling rates with a correspondingly large range ofmicrostructures, mechanical properties, and acousticemission activities. There was no evidence of quenchcracking in any of the specimens.

The second set of treatments were performed on dumbbell specimens only. Specimens were water quenched from1000°C and pairs subjected to isothermal temperingtreatments at 650± 5°C in argon for times of 6-10020 min,and water quenched. It was recognized that those specimenstempered for the shorter times had probably not reached650°C for more than a few minutes before quenching. Allspecimens were then electropolished in a 10% solution ofperchloric acid in acetic acid to remove the surface oxidelayer as the cracking of this layer has been found to be acopious source of extraneous acoustic emission.

Acoustic emission measurementsAcoustic emission measurements were made with twodifferent recording systems. First, a more conventionalnarrow band technique incorporating a piezoelectrictransducer was used to determine systematically the relativeemission activity of the dumb bell specimens. Second, a wideband calibrated technique incorporating a capacitancetransducer and digital recording was used in an attempt tocharacterize the emission sources in Yobell specimens.

2mm(b)

a dumb bell specimen used for narrow band acoustic emiSSionmeasurements and for measuring reference mechanical properties;b Yobell specimen used for measuring acoustic emission wave formsover wide frequency range

3 Specimen geometries for mechanical testing

N arrow band techniqueDumb bell specimens were tested in tension using a model1195 Instron screw-driven machine. Specimens weregripped at the 45° shoulders at either end of the specimen.All tests were performed at a constant crossheaddisplacement rate of 0·2 mm min -1 (""' 1·7x 10- 3 S-1 strainrate). The grips contained special insulators which were

pow~r

chart r~cordar+ machanical

~n~rgy proparti~s data

oscilloscop~

RF pow~r m~t~r+ intagrator

58d

mainamplifi~r

o0'1 MHz 1·0band-pass

'filt~r

2·5 X 105y A-1

at 1MHz

piazo~l~ctrictransduc~r

dumbb~"sp~cim~n

(a)

pr~amplifi~r

charg~amplifi~r

3-3 x1011 Y C-1

capacitancatransducar

o0.Q3MHz45

band-passfiltar

+10

mainampllfi~r

wav~ forms

digitalcomputar

Yoballspaciman

(b)

4 Schematic diagrams of a narrow band and b quantitative broad band acoustic emission detection systems; forboth systems, specimen, transducer, and preamplifier were enclosed in screened room

Metal Science June\t981'

244 Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel

used to isolate electrically the specimen from the testingmachine in order to reduce ground-loop coupled noise.However, these insulators were mechanically soft and thus acomponent of the crosshead displacement was accommo-dated in the grips rather than the specimen gauge. Duringeach test the load was continuously monitored togetherwith the acoustic emission activity.

The acoustic emission system is shown schematically inFig. 4a. Care was taken to standardize all experimentalprocedures so that valid comparisons could be madebetween different tests. The acoustic emission was detectedwith a damped piezoelectric transducer attached by a thinlayer of grease, under a constant small pressure, to thepolished end of each specimen. The transducer had anapproximately flat frequency response below 1 MHz andwas connected to a wide band preamplifier which bothamplified and differentiated the transducer output. Thus,the output voltage of the head amplifier was expected to beproportional to the surface velocity of the specimen. Thesensitivity of the system to small-amplitude emissions wasimproved by reducing the ambient noise with acoustic andradio frequency shielding. A 0·1-1·0 MHz bandpass filterwas used to increase signal/noise ratio. The signal wasfurther amplified and monitored with a Hewlett-Packard435A power meter which permitted the electrical power dueto an emission to be recorded. This was approximatelyproportional to the acoustic power impinging on thetransducer in the pass band 0·1-1·0 MHz. The power-meteroutput was also integrated so that a measure of the energy.of an emission could be obtained. Both the emission powerand energy together with the load were plotted on a chartrecorder as a function of crosshead displacement during atensile test. The electrical signal was also continuouslymonitored on an oscilloscope.

Wide band techniqueThe objective here was to measure accurately the surfacedisplacement wave forms from acoustic emission eventsduring the deformation and fracture of Yobell specimens.While such wave forms can be measured for any specimengeometry, a meaningful analysis can be performed only ifthe effects of specimen boundaries are removed, or at leastminimized. The Yobell geometry permits this while stillallowing a simple mechanical test to be performed. All theYo bell specimens were tested to failure in an Instron 1195screw-driven machine at a crosshead speed of0·1 mm min -1. The transient displacement at the centre ofthe polished surface of the Yobell specimen (designed to bethe epicentre), due to the arrival of an acoustic emissionsignal, was measured with a capacitance transducer(Fig. 4b). The transducer air gap was maintained at 2·7 Jlmfor every test, giving a constant transducer sensitivity of1·72x10-3Cm-1.

The transducer was coupled to a charge-sensitivepreamplifier with as short a length of cable as possible(about 20 cm in practice). This was followed by two filters: ahigh-pass filter at 30 kHz to eliminate low-frequencymachine noise and a low-pass filter at 45 MHz to eliminatesignal aliasing errors .during later analogue to digitalconversion. The filtered signal was further amplified beforedigitization in a Biomation transient recorder and storage ina PDP8/E minicomputer. The recording system had a flattransfer function of 5·35 x 109 vni-1 from 80kHz to25 MHz, the upper band width limit of the Biomationrecorder. The recorder threshold was set just above theambient noise and corresponded to a surface displacementof 1 pm.

Metal Science June 1981

Mechanical properties dataThe low compliance of the Instron prevented accuratestress-strain measurements of the specimens tested foracoustic emission response. To overcome this, additionalspecimens of identical geometry and heat treatment weretested in a servohydraulic machine at a constant strain rateof '" 2·1 x 10 - 4 S - 1. Knife-edge extensometers weremounted at the ends of the gauge length so that the strainwithin the gauge could be measured as a function of loadthroughout the test. After testing, the load-extension datawere converted to give nominal stress-nominal strain andtrue strain-true plastic strain data, so that proof andultimate stresses, strains to fracture, and work-hardeningexponents could be measured.

Microstructural characterization andfractographyAfter testing, the un deformed ends of dumb bell specimenswere ground and polished to allow the microstructure to becharacterized by optical metallography. In addition,extraction replicas were taken from the polished and etchedsurfaces and the precipitate type and composition for someof the long-tempered specimens determined using micro-diffraction techniques and energy-dispersive X-ray analysis.The precipitate size distribution was determined fromscanning electron micrographs of nital-etched specimens.

The fracture surface of one specimen in each conditionwas examined by scanning electron microscopy todetermine the fracture mode. In addition, fractured ends ofspecimens were also mounted, ground, and polished so thatthe longitudinal section behind the fracture face could beexamined by optical and scanning electron microscopy.

RESULTSCharacterization of microstructureAll specimens were austenitized for 1 h at 1000°C whichresulted in a prior austenite grain size of '" 70 Jlm. Variationof cooling rate from 1000°C led to specimens of widelydiffering microstructures (Fig. 5). Both the iced-brine andNaOH quenches produced acicular martensite containingno visible carbides when examined by optical microscopy.(Fig. 5a and b). Oil quenching (Fig. 5c), and to a lesserextent water quenching, also resulted in a martensiticstructure which now contained barely resolvable carbides.Air cooling resulted in a bainitic microstructure (Fig. 5d),the visible carbides being < 0·5 Jlm in size. The much slowerfurnace-cooling treatment gave rise to a complex micro-structure with randomly distributed areas of upper bainite,fine pearlite, and proeutectoid ferrite containing isolatedcarbides (Fig. 5e).

Isothermal tempering at 650°C of water-quenchedspecimens also resulted in major changes of microstructure(Fig. 6). Tempering for as little as 6 min (Fig.· 6a) resultedin the appearance of small carbides on lath and prioraustenite boundaries. Increasing the tempering time causedthe growth of carbides (Fig. 6b-e). After tempering for about4-8 h the largest carbides were '" 1 Jlm in size and there wasevidence of the formation of subgrains of ferrite, formed byrecovery of the quenched-in dislocation structure. Furtherprolonged tempering increased the carbide size and led tothe complete disappearance of martensite laths. Thesecarbides are shown in Fig. 7. Carbides less than '" 0·2 Jlmwere unresolvable, which accounts for the fall indistribution at small carbide size. After testing, largecarbides located close to the fracture surface had themselvescracked (Fig. 8).

Specimens tempered for more than 4 h were also

a

c

e

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 245

b

. d

a 1000°C, 1 h, iced-brine quench; b 1000°C, 1 h, water quench; c1000°C, 1 h, oil quench; d 1000°C, 1 h, normalized (air cool); e1000°C,1 h, furnace cool "'50 K h-1

5 Optical micrographs showing effect of cooling rateupon microstructure; 20/0 nital etch

examined using convergent beam microdiffractiontechniques. The majority of carbides present in thesespecimens tempered ~ 8 h were found to be of the M3Ctype. However, prolonged tempering resulted in theappearance of M23C6 precipitates, especially at prioraustenite grain boundaries (Fig. 9). Energy-dispersive X-rayanalysis revealed that the M3C carbides had a Mn/Fe ratioof 0·15± 0·03 and a Mo/Fe ratio of 0·04. The Mn/Fe ratio

for M23C6 was 0·13± 0·03, similar to that of M3C, whereasthe Mo/Fe ratio was significantly greater at 0·08.- It was noted in the specimens tempered for the .longest

times that the apparent volume fraction of carbides variedover the specimen surface. This was thought to be aconsequence of the diffusion of carbon to the micro-segregated alloying elements resulting in a non-uniformcarbide distribution.

Metal Science June 1981

246 Scruby et al.

a

c

e

Microstructure and acoustic. emission in Mn-Mo-Ni steel

b

d

a 6 min; b 36 min; c 4 h; d 24 h; e 167 h

6 Optical micrographs showing effect of isothermaltempering at 650°C upon microstructure; 20/0 nitaletch

Mechanical propertiesThe effect of varying the quench rate upon the relationshipbetween nominal stress and nominal strain is shown inFig. 10. True stress (J'T was also calculated as a function ofthe true plastic strain eTP using the relations

(J'T = (J'n(l +en) . (1)

eTP = In (l +enp) (2)

Metal Science June 1981

where (J'n is the nominal stress, en the nominal strain, and enpthe nominal plastic strain defined as

enp = en -(J'T/E . (~)

where E is Young's modulus. The true stress v. true plasticstrain data were plotted logarithmically to calculate thework-hardining exponent n for each material condition

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 247

21lm

8 Example of cracked carbide precipitate close tofracture surface of specimen tempered 167 h

(Fig. 11). A summary of mechanical properties data and n-value estimates is given,in Table 2 for each condition. As thecooling rate decreased on going from an iced-brine quenchto a furnace cool, the 0'1% proof and maximum stressesdecreased while the plastic strain to fracture and theapparent n value showed a small increase. The true plasticstrain at maximum nominal stress did not varysystematically with cooling rate.

The effects of isothermal tempering at 650°C for periodsof 6-10 020 min upon mechanical properties are shown inFigs. 12 and 13 and summarized in Table 3. Temperingreduced the 0·1% proof and maximum stresses while causingan increase in plastic strains to maximum stress and tofracture and an increase in n to 0·22.

axis---.tensile4--

Fracture appearanceVariation of cooling rate had a major effect upon the modeoffracture at room temperature. The iced-brine and NaOH-quenched specimens underwent a predominantly brittlefracture (Fig. 14a) even though preceded by significantplastic deformation (Table 2). The fracture surface could bedivided into two regions: a central region that had fracturedapproximately transverse to the tensile axis and an annularouter region where a shear lip had formed. The central areafractured by transgranular cleavage (Fig. 14b), theindividual microcracks appearing to have changedorientation at lath boundaries. Water-quenched specimens(Fig. 14c) exhibited only a few isolated areas ofquasicleavage, the majority of the central region of thefracture had occurred by alternating shear (Fig. 14d).Optical metallography of material adjacent to the fractureindicated that at the intersection of alternating shear facets a

from the relation

aT= ke¥p (4)where k is a constant. While values of n could be measuredfor some specimens, others had n varying with strain

~30

a 025 075 1-25 175 2·25 2·75050 lOa 1-50 2·00 2·50c LENGTH,~m

7 Scanning electron micrographs of nital-etchedspecimen tempered 167 h at 650°C showing aprecipitates at grain boundaries, b precipitates incentre of grain, and c size distribution for graininterior precipitates

a

b

Table 2 Mechanical properties of specimens cooled at varying rates

Nominal stress, MN m-2 True maximum True plastic Nominal plasticstress, strain at strain to

Heat treatment 0"0'1"" O"max MNm-2O"max fracture n value

1000oe, 1 h, iced-brine quench 1100 1790 2000 0·068 0·30 0·051000oe, 1 h, NaOH quench 1020 1690 1880 0·065 0·32 0·081000oe, 1 h, water quench 1110 1750 1945 0·068 0·32 0'01-0'19*1000oe, 1 h, oil quench 905 1460 1580 0·057 0·33 0'04-0'15*1000oe, 1 h, air cool 535 960 1035 0·061 0·51 0'08-0'32*10OO°C, 1 h, furnace cool 50 K h-1 540 750 850 0·074 0·46 0·10

* n-value variation with strain.

Metal Science June 1981

248 Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel

a

A

B

MoLa,

<2~1> zona

b

9 Use of convergent beam microdiffraction and X-ray energy analysis to characterize carbide types andcompositions in Mn-Mo-Ni steel of A533B composition

Metal Science June 1981

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 249

a iced-brine quench; b water quench; c oil quench; d air cool; efurnace cool - 50 K h-1

10 Representative nominal stress v. nominal straincurves showing effect of quench rate uponmechanical properties

0·6

2000

o

tf)

~ 10000:::r-tf)

-.J<[Z~oz

a water quench; b 6 min; c 36 min; d 4 h; e 24 h; f 167 h

12 Nominal stress as function of nominal strain,showing effect of isothermal tempering at 650°Cupon mechanical properties

0·6o

•....tf)tf)

~ 1000r-tf)

-.J

~ 500~oz

C'\J 2000IEz21500

crack occasionally propagated parallel to the tensile axis.Slower cooling rates resulted in the occurrence of a classicalcup and cone fracture mode (Fig. 14e). The fracture surfacewas covered in fine dimples which were deepest in theslowest cooled specimen. In the furnace-cooled condition,spacing and size of dimples varied over the fracture. Someparticles, comparable in size with the carbides observed byoptical metallography, were observed within the dimples.This suggests that the large carbides acted as void-nucleation sites.

Tempering specimens, even for only 6 min, changed thefracture mode of water-quenched samples (Fig. 15a). Theareas of alternating shear had almost fully disappeared andthe fracture surface was covered in a fine array of barelyresolvable shallow dimples (Fig. 15b). Increasing thetempering time had two effects. First, as can be seen bycomparing Fig. 15a, c, and e, it led to the formation of a

. classical cup and cone fracture. Second, the size and depthof the fine dimples increased (Fig. 15b, d, andf). Few oftbedimples in any tempered specimen contained resolvableparticles. Some of those observed were spherical inclusions(Fig. 15f).

Acoustic emission measurementsN arrow band testsFigure 16 shows the effect of quenching rate upon the stressand strain dependence of the acoustic emission fromquenched specimens. The two most rapidly quenchedspecimens (iced brine and NaOH) behaved similarly: littleor no emission was observed during elastic loading and onlya few emissions were detected during post-yield deformationbefore attainment of maximum stress (Fig. 16a). As thespecimens were deformed past maximum stress large-energy

n=0·19

°maxI~ \~ ----n=0·32 n=0·08

10-3 10-2 10-1

TRUE PLASTIC STRAIN

a 1000°C, 1 h, iced-brine quench; b 1000°C, 1 h, water quench; c1000°C, 1 h, oil quench; d 1000°C, 1 h, normalized (air cool); e1000°C,1 h, furnace cool -50 K h-1

True stress as function of true plastic strain(logarithmic axes) showing effect of quench rate onwork-hardening exponent n (n < 0 at high strainserroneous due to plastic instability and necking)

11

(b)

1o'maxI

a'maxI

n=0·04

(fmaxI

c:==:::n= 0·15

1CT3 10-2 10~1TRUE PLASTIC STRAIN

2 «(I)10

10-4

V)V)

lIJ1030::l-V) (e)

lIJ:J~ 0=0·10

103

17 Metal Science June 1981

a 6 min; b 36 min; c 4 h; d 24 h; e 167 h

13 True stress as function of true plastic strain(logarithmic axes) showing effect of tempering at650°C on work-hardening exponent n (n < 0 at highstrains erroneous due to plastic instability andnecking)

10-2 10-1

PLASTIC STRAIN

250 Scruby et at. Microstructure and acoustic emission in Mn-Mo-Ni steel

n=0·07 n=0·08

ermax dmax103 I

I\

\N (a) (b)IE 102

z~ n= 0·10 n= 0·16

~103 ermaxen \

"W ~(r~en

~ 102 (c) Cd)(r

10-4 10-3 10-2 10-1~n=0·22 TRUE PLASTIC STRAIN

1030max

burst emissions were generated and these increased both inmagnitude and rate of occurrence as the specimen wasstrained towards final fracture. Similar behaviour wasobserved with water-quenched specimens (Fig. 16b)although in this case the onset of burst emissions wasdelayed until the final stages of deformation just before finalfailure. The oil-quenched (Fig. 16c) air- and furnace-cooledspecimens generated little emission at any stage ofdeformation. The data for these tests have been summarizedby measuring the acoustic emission energy for each test andare given in Table 4.

Tempering at 650°C also had a marked effect upon thestress and strain dependence of the acoustic emission power.Tempering water-quenched material for 6 min resulted in

the complete disappearance of detectable emission exceptfor a single emission at fracture (Fig. 17a). Tempering forlonger times up to ",,4 h gave similar results. However,specimens tempered for 8 h and more began to exhibit anacoustic emission peak in the vicinity of yield (Fig. 17b andc). This peak consisted of many small-amplitudeoverlapping signals (usually called continuous emission)with occasional larger amplitude signals superimposed. Theacoustic emission energy data for these specimens aresummarized in Table 5.

In all the tempered specimens the only emission generatedafter O'max occurred at the instant of final fracture and thisdecreased with specimen tensile strength. The trends inemission activity are presented graphically in Fig. 18.

Table 3 Mechanical properties of isothermally tempered specimens

Nominal stress, MN m - 2 True maximum True plastic Nominal plasticTempering time, stress, strain to strain tomin 0"0'1"" O"max MNm-2 maximum stress fracture n value

0 1110 1750 1945 0-047 0-32 0-01-0-19*6 830 960 1066 0-061 0-48 0-07

13 635 940 1053 0-070 0-48 0-0836 680 930 1045 0·067 0·53 0-0760 680 870 971 0-070 0-54 0-06

120 700 850 957 0-072 0·50 0-08240 695 795 906 0-083 0-51 0-10480 680 785 870 0-028 0-49 0-05

1440 405 600 721 0-116 0-55 0-166000 270 510 642 0-160 0-61 0·15

10020 330 545 690 0-157 0-68 0-22

* n-value variation with strain.

Metal Science June 1981

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 251

a b

c d

e "'·iWk;l,i·",'(;,m'"'.:. fa and b cleavage fracture mode of iced-brine quenched sample; c and dfracture mode of water-quenched sample; e and {fracture mode of furnace-cooled sample

14 Effect of quench rate upon fracture mode

Wide band testsFive Yobell specimens were tested at an exten'sion rate of0·1 mm min -1 using the screw-driven Instron machine. Thespecimens had been austenitized 1 h at 1OOO°C; two werequenched in iced-brine solution at - 8°C and the rest inwater at room temperature. A representative stress v. straincurve with superimposed acoustic emission activity is givenfor each condition in Fig. 19.

Acoustic emISSIons were detected by the capacitancetransducer from all the specimens. These emissions, in themain, were generated after the attainment of maximumstress. The strain dependence of the emission activity of theYobell specimens was qualitatively similar to that of dumbbells of similar quench method. More emissions weredetected from the more rapidly cooled iced-brine quenchedsamples (see Table 6).

Metal Science June 1981

252 Scruby et at. Microstructure and acoustic emission in Mn-Mo-Ni steel

a

c

e

b

d

fa and b 6 min; c and d 18 min; e and f 100 h tempering time

15 Effect of tempering on fracture mode of water-quenched samples

The calibrated detection system enabled each emission tobe recorded as a time-varying surface displacement and oneexample from each condition is shown in Fig. 20. All ofthese surface-displacement wave forms had a characteristicpulse at the longitudinal (L) wave arrival time followed by aless distinct step at the arrival of the shear (5) wave. Whenaccount is taken of the low-frequency filtering of themeasured wave forms, good agreement exists with the waveforms calculated for a modelled crack-growth event. The

Metal Science June 1981

model did not predict the oscillatory form of the so-called'wash' following the L arrival and preceding the 5 arrival.This may have been due to interactions of the generatedwave with nearby free surfaces such as subcriticalmicrocracks or the boundaries of the gauge.

Because of the generally good agreement of measuredwave forms with theory, it was possible to apply the sameanalytical procedure as reported elsewhere9 to obtain thesource function from each measured wave form. First, the

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 253

0::W~

0·8 0Q..

0·7 z0·6 00·5 Vi

CJ)0·4 ~0·3 w0·20·1a

0·80·70·50·50·40·3

0·2 ~E0·1o

u•...CJ)::>o~

0·80·70·50·50·40·30·20·1o

0·8

(b)

(a)

(c)total anargy = 2·1 mJ

total enargy = 0·82 mJ

total anargy = 0·61 mJ

o

500

1000

1000

1500

.J<{

~ 0~ 1500z

C\I 0IE 1500

z:If"o1000

CJ)en1&.10::tn 500

0'2 0·4 0-6NOMINAL STRAIN

a6min;b8h;c167h

17 Strain dependence of stress and acoustic emissionpower for A533B composition steel waterquenched and tempered at 650°C for various times

transfer function from source to detector was calculated.Second, the inverse transfer function was calculated whichwas convolved with the measured wave form to obtain theactual source function. This source function is in the form ofa time-varying crack volume when the epicentre surfacedisplacement is measured. The time dependence of thesource function for the two representative wave formsshown earlier is also presented in Fig. 20.

In every case the volume rose to a positive maximum andthen started to fall. The fall was artificial and was due tolow-frequency filtering and for short-duration events thiscould be ignored. Previous work has shown that the twomost useful parameters to measure are the maximumvolume attained and the time taken. In order to reducepossible errors from gauge-boundary reflections themaximum volume was calculated as twice the volumereached when the surface displacement of the L arrival pulsereached its maximum value.9 This is equivalent to assumingthat the true L wave pulse was symmetric about a verticalline drawn through its maximum. The time for which thesource operated, the source lifetime, was likewise taken astwice that for the surface displacement at the L wave to gofrom zero to its maximum. The mean values of thismaximum volume and lifetime are summarized for each

0·80·70·60·50·40·30·20·1o

0·8

~ (Jmax Total

5·0 5·14'5 5·12·5 2·80·7 0·71·1 1·60·3 0·6

~ (Jmax

Acoustic emission energy, m]

0·2 0·4 0·6NOMINAL STRAIN

o

500

1000

(a)

1500 ••.•10 amissions\

0·8

1000 0·7

0·60·5

0·4500 0·3

0·2

0·10 0

(b) ~E

C\I 1500 0::IE wz 0·8 ~

0~ 0·7 Q..1000 0.6 z~

en 0CJ) 0·5 Viw en0:: 0·4 ~.-CJ) 0.3 w..J 0·2 U<{ 0·1 ~z~ 0 :J

00 Uz total anargy = 0·7 mJ (c) <

1500

Iced-brine quench 0·1NaOH quench 0·6Water quench* 0·3Oil quench 0·0Air cooled 0·5Furnace cooled 50 K h - 1 0·3

Quench method

* Average of four tests.

Table 4 Effect of quench rate on acoustic emissionenergy

a iced-brine quench; b water quench; c oil quench

16 Strain dependence of stress and acoustic emissionpower for A533B steel austenitized 1 h at 1000°C

Metal Science June 1981

254 Scruby et a/. Microstructure and acoustic emission in Mn-Mo-Ni steel

3·0 3-5

6-0 6-5

Sway(/,...,

2·5 3·0 3·5

S wava,Il

5·0 5·5

2·0 2·51·0 1·5

0·5 1·0 1·5 2·0TIME,J.lS

0'5

L wava,II

L wava,r:-1

(b)o

o

a

50

100

5

o

10

500

-502'5 3·0 3·5 4·0 4·5

-500

o-5

2·5

1000

-1000 (a)o

Ea.1--ZW2wu«..J0..(/)

a

Ea.1-•...ZW2wu«..J0..(/)

C(Y)

E::l;. 10000w~ 5000::>..Jo>~ -50000::5-10000(/)

20 Examples of acoustic emission displacement waveforms for a iced-brine quenched specimen and bwater-quenched specimen; correspondingsource-volume-time histories were calculated byconvolution with inverse transfer function forspecimen; apparent fall in volume at long times isdue to high-pass filtering

6000

4000

2000(")E::t

W2:o :3~wu

6000~oIf)

....,E>-~~2wz

\W\

Z \

0 \ bli) aIf)

~W1u ,i=If) , ,:)0u«

1000

If)(/)

~ 0I-If) (b)

<1.1500z~oz

~ 500Ez2:

(a)

1500

• •o 10' 102 103

quenchedTEMPERING TIME) min

18 Acoustic emission generated a before and b aftermaximum stress as function of tempering time

1000 4000

Table 5 Effect of isothermal tempering on acousticemission energy

Acoustic emission energy, m]

Tempering time, min ~(Jrnax ~ (Jrnax Total

0* 0·3 2·3 2·66 0·0 0·9 0·9

18 0·3 0·9 1·236 0·2 0·7 0·960 0·1 0·7 0·8

120 0·1 0·6 0·7240 0·5 0·5 1·0480 0·3 0·7 1·0

1440 1·9 0·4 2·36000 1·6 0·3 1·9

10020 1·1 0·3 1·4

* Average of four tests.

2000500

oo 0·5 1·0 1·5CROSSHEADDISPLACEMENT, mm

19 Nominal stress as function of displacement forYobell specimen tested in a iced-brine quenchedcondition Y1 and b water-quenched condition Y5;superimposed on each is acoustic emission activity;see text for discussion of parameter representingthis activity, the source volume

Metal Science June 1981

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 255

0·2 0-4(a) (b)

0-3

0'1 0·2

en~ 0'1zw>w~ 0 00zO'2 0'40 (e) Cd)~u~ 0·3Q:~

0·1 0-2

0,1

0103 104 105 106 0 100 200 300 400 500CRACK VOLUME, J,lm3 CRACK LIFE TIME,ns

a and b iced-brine quench; c and d water quench

21 Maximum source volume and source lifetime for all recorded wave forms from specimens in each quenchedcondition

specimen in Table 6. Histograms of both source parametersare shown for both quench rates (Fig. 21). The emissionsource volumes from water-quenched samples were slightlylarger on average than those from iced-brine quenchedsamples. There was little significant difference in source'lifetimes. The emission activity parameter shown in Fig. 19was the calculated maximum source volume.

DISCUSSIONVariation of specimen heat treatment caused changes inacoustic emission activity during narrow band tests. This,the authors believe, is a consequence of changes in thedetailed deformation and fracture mechanisms broughtabout by microstructural variation. 2 The authors will brieflyexamine the variation in deformation behaviour and thendeduce which of these processes was the origin of detectablesignals. Finally, the wide band results will be used to obtainquantitative data about the most energetic emission in thesematerials.

Deformation and fracture mechanismsIt is convenient to separate the processes occurring beforemaximum stress (lmax from those leading to ,final fracture.

Premaximum stressGeneral yield is characterized by the widespreadpropagation of dislocations. The transition in yield-pointbehaviour from quenched to either slow-cooled or heavilytempered conditions (Figs. 10 and 12) is indicative ofchanges in the mechanism of motion. In the quenched state,yield is accommodated by movement of the quenched-indislocation structure, while in tempered/slow-cooledsamples, where the dislocation density is initially very low,deformation is accommodated by the nucleation and rapidpropagation of fresh dislocations.lO

In martensitic microstructures the yield stress is expressedas a sum of strengthening components associated with thematrix friction stress (10' the lath packet size D, the lath sized, and the forest dislocation interactions.11-13 Norstrom13

Table 7 Crack length and velocity results: see Fig. 22

Metal Science June 1981

256 Scruby et at. Microstructure and acoustic emission iri Mn-Mo-Ni steel

0·3

0·2

~ 0·1Zw>w~ 0

0·3zoI-U~0·2LL

0·1

(a)

(e)

(b)

(d)

o 100 200 300 0 500 1000 1500CRACK LENGTH, J.1m CRACK VELOCITY) m s-1

22 Crack length and crack velocity for a and b iced-brine, C and d water-quenched A533B steel

has assumed that a Hall-Petch relation exists for both lathand lath packet sizes so that the expression for the yieldstrength of low-carbon marten sites is

(Jy = (Jo+(Ji+kyD-l/2+ksd-l/2+r:J.j1bpl/2 . (5)

where (Ji is the solution hardening (here due to both carbonand metallic alloying elements), ky and ks the grain-sizecoefficients for the lath packet and lath grain sizes, p thedislocation density, r:J. the dislocation strengtheningcoefficient, and j1 the shear modulus. Expressions such asequation (5) for martensitic steels assume independentadditive contributions to strength from the components ofthe microstructure and the existence of a Hall-Petchrelation for the grain-size parameters. Both assumptions aredifficult to test rigorously but this is of limited relevancehere: their importance is that ky and ks are non-zero,indicative of dislocation interactions both with the lath andlath packet boundaries. Thus, a few dislocations, at least,can propagate a distance comparable to the width of a lath,i.e. several micrometres.

An additional strengthening mechanism in moderatelycooled or tempered samples arises from the presence of adispersion of hard precipitates12 (M23C6, M3C). Themagnitude of this strengthening depends upon the size and

separation of the precipitates. Gladman et al.14 have shownit to be

(J oc Jl; log (~) . (6)

where x is the mean planar intercept diameter of theprecipitates and A. is the surface to surface precipitatespacing given by

(7)

where ns is the number of precipitates per unit area of slipplane. Assuming a fixed volume fraction of precipitates, thegreatest strengthening due to this process occurs for a finedispersion of small particles, e.g. as found in the air-cooledsample. In the absence of cross slip, the presence of adispersion of impenetrable precipitates will limit thedistance dislocations can freely propagate to ""A.. Thepresence of a dispersion of precipitates results in anenhanced work-hardening rate since the dislocation loopleft around a bypassed precipitate reduces the effectiveinterparticle spacing. This in part is the origin of theincrease in n value with decrease in cooling rate.

The furnace-cooled material contained fine pearlite inaddition to dispersion-strengthening carbides. The

Table 8 Comparison of mean crack length and velocity data

Material

Pure A533BPure A533BMild steel9

Quench method

Iced-brine quenchWater quenchWater quench

Test temperature, Number of Crack length, Lifetime, Velocity,K crack emissions Jlm ns ms-1

290 181 52 117 450290 53 57 124 46077 106 57 110 490

Metal Science June 1981

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 257

0·3(a) (b)

500 1000 1500CRACK VELOCITY, m 5-1

3000100 200CRACK LENGTH, J.lrn

o

0·2

V) 0·1....zw>W

l.L. 000·3z (d)0i=u: 0·2l.L.

0·1

23 Comparison of crack length and velocity for a and b cleavage of A533B steel at room temperature, C and dcleavage of mild steel at 77 K

dispersion-strengthening contribution to the flow strengthshould be much less in this case since there is a lower volumefraction of carbides within ferrite grains (most carbon is inthe pearlite). Within the pearlite, deformation of both

,cementite and ferrite will occur. Porter et al.15 have shownthat the cementite layers within fine pearlite fracture by aductile necking process.

Post-maximum stressThis was the region where the load was falling since theeffect of reduction in specimen cross-section due to neckingexceeded the increase in strength due to work hardening. Inall specimens it was in this region that subcriticalmicrocracks and void growth and coalescence occurred,ultimately causing failure when the remaining uncrackedligaments were no longer capable of supporting the appliedload.

In the most rapidly quenched specimens the true stressreached very high levels, '" 2 GN m - 2 (Fig. 11), andcleavage microcracking occurred (Fig. 14b). Fractographicobservations indicated that cracks made small orientationchanges at martensitic lath boundaries. Other studiesl6, 17have shown that the major change of crack orientation atlath packet and prior austenite grain boundaries effectivelystops a propagating microcrack. The authors thus expectindividual microcracks to be approximately a lath packet insize and, on average, to be transverse to the tensil~,axis. It islikely that some plasticity accompanies the growth of thesemicrocracks which may cause blunting after crack arrest ata boundary. Water quenching produced a slightly lowerstrength martensite and much less cleavage was observed(Fig. 14c). Here, the subcrltical cracking occurred mostlyby shear which resulted in a zigzag fracture topography asshear separation occurred on alternate planes. This

mechanism of fracture has been observed in other steels ofhigh strength and limited work-hardening capacity.7 It isthought to occur when the deforming ligament between apair of large voids (perhaps formed at inclusions) exhaustsits work-hardening capacity, resulting in a local intenseshear band and the creation and coalescence of a microvoidsheet. Slower cooling rates resulted in the formation of cupand cone fracture (Fig. 14e). The period of falling load nowcharacterized a time when voids previously nucleated atinclusions grew and new voids were nucleated after highelastic strain at carbides. The voids formed at carbides mayhave been nucleated by decohesion of the carbide/matrixinterface or by cracking of the precipitates themselves. Asdeformation continued the ligaments between voidsinternally necked down and void coalescence occurred.

Tempering of water-quenched specimens, even for only6 min, resulted in the disappearance of alternating shear andthe initiation of cup and cone fracture (Fig. 15). Extendingthe tempering time increased the size of the ductile dimpleson the fracture surface as the number of carbides presentdecreased although their size increased. There wasfracto graphic evidence that carbide fracture did occur butwas mostly confined to regions which had experiencedintense plastic strain.

Origin of detectable acoustic emissionIt is likely that all mechanisms of deformation and fracturedissipate some energy in the form of elastic waves. However,many studies in the past have clearly shown that the elasticwaves from some mechanisms are more easily detected thanthose from others.2 This study attempts to determine thefactors controlling the generation of detectable acousticemission in a steel within the A533B composition range. Ifone ig~ores the effects of ultrasonic attenuation there are

Metal Science June 1981

258 Scruby et at. Microstructure and acoustic emission in Mn-Mo-Ni steel

. (10)

· (15)

· (16)

· (12)

· (11)

hCIXmin

amin = bnvc~ .

Crack-growth sourcesIt is assumed that the source is the creation of a smallhorizontal microcrack of cross-sectional area A, crack faceseparation 2b, and volume V, under mode I loading. Thenat a distance h, much greater than the crack dimension, thestrain field approximates to that of an edge dislocation loopof area A and effective Burgers vector 2b. The area under thelongitudinal wave pulse is given by9

VSc = -- . . (13)

2nc1h

Assuming a parabolic increase in crack area with time, thepeak amplitude of the longitudinal component is

VX=-- . . (14)

nC1Th

This enables the minimum detectable crack volume to beestimated, provided T is known. However, it is more usualto work in terms of crack length. Equation (14) can beexpressed in terms of length if it can be assumed that thecrack behaves in a linear elastic manner during its growth.In this case, the crack will have an ellipsoidal shape ofvolume

4nazbV=--

3

where a is the radius of an assumed circular crack ofmidplane area naz. Under elastic loading the crack facedisplacement, at the centre of the crack b is a function ofcrack radius and applied stress

b = 2(1- vZ)aa.E

If Xmin is the displacement detection threshold, the smallestdetectable dislocation loop must have grown to a radius

hCI Xmin

amin = bvc~ .

For the detection system used here Xmin ~ 10-14 m. For steel,C1 = 5960 m S-1, Cz = 3200 m S-1, and an average sourcedepth h was 0·03 m. Using a Burgers vector of",3 x 10-10 m and assuming the dislocation velocity was200 m s -1, amin '" 100 Jlm. Thus, if a single dislocation loopwere to expand to a radius of '" 100 Jlm at a radial velocityof 200 m s -1 it would be just detectable. Loops of smallerradius could be detected according to the model if theyeither grew at greater velocity or if n loops situated close toeach other expanded together over the same time interval sothat their combined r~dii exceeded 100 Jlm. In the latter case

where E is Young's modulus and v is Poisson's ratio.Substituting for V and b, equation (14) becomes

8a3(1- vZ)ax=---- . (17)

3C1 ThE

Assuming a constant radial crack velocity v so that a = VT

and rearranging equation (17) to give

amin = 3c 1Ehxmin . (18)8(1- v2)va

From equation (18) it is seen that the smallest crackadvances are most easily detected when they are close to the

T = a/v . (9)

then T will correspond to the width of the longitudinal wavepulse. If the height of the pulse is x, then for a loop area thatgrows parabolically in time to its final value of naz, thesurface displacement first arrival commences with a ramp ofslope X/To Thus

bavc~x = hCI .

two practical factors which are likely to have a majorinfluence on the detectability of an acoustic emission. One isthe sensitivity of the detection system to the surfacedisplacements caused by arrival of elastic waves from anemission source. Narrow band systems with their bettersignal-to-noise characteristics have a much greatersensitivity compared with wide band techniques. The otheris the frequency response of the detection system comparedwith the spectrum of the source.

Sensitivity limitations·All detection systems have some kind of detection thresholdand only acoustic emission signals whose amplitudes exceed·this are recorded. The nature of the threshold depends onthe type of acoustic emission instrumentation in use. Insystems where ring-down counts are made, the threshold isa voltage level which the emission signals must exceed andmay correspond to a minimum detectable surfacedisplacement or velocity of the specimen surface dependingon the mode of operation of the transducer. The wide bandsystem used in this study had a voltage threshold in thetransient recorder (set to be just above the noise level) whichcorresponded to a minimum detectable surface displace-ment of 10-1z m. The narrow band system, however, hadno such clearly defined threshold. Here, the transduceroutput was monitored continuously and the effectiveminimum detectable signal was that producing the smallestpower fluctuation which could be distinguished from noise.This corresponded to a frequency-band limited powerrather than a displacement, so that the minimum detectabledisplacement was frequency dependent. However, in orderto simplify the following discussion the authors assume it tobe constant at '" 10-14 m. We can then proceed to comparethis value with the displacements from dislocation motionand crack-growth sources.

Dislocation sourcesThe authors consider the expansion of a glissile dislocationloop of Burgers vector b from an initial radius of zero to a. Itis assumed the loop is inclined at 45° to the tensile axis of thespecimen and that the final loop radius is much less than thesource to transducer distance and the shortest wave lengthdetected, so that the loop may be considered a point source.The operation of such a source will cause the simultaneousgeneration of longitudinal (compressional) and transverse(shear) elastic waves with velocities C1 and Cz respectively.The time dependence of the surface displacement at theepicentre of a half-space produced by such a source hasalready been evaluated.1 The early part of the wave form isdominated by a pulse due to the arrival of the longitudinalwave; the area under this pulse is given by

c~bASd = -2 h 3 • (8)

n C1

where h is the depth of the loop below the transducer and Ais the final loop area.

If the loop grows at constant radial velocity v so that theduration over which it grows

Metal Science June 1981

Scruby et at. Microstructure and acoustic emission in Mn-Mo-Ni steel 259

Thus, for this model, it would be expected that as adetection system was made more sensitive, it wouldpreferentially detect small dislocation loop sources. A lesssensitive detection system such as the broad band sys'tem(Xmin = 10-12 m) should more readily detect crack-growthsources than dislocation loop sources.

transducer, of short duration (high velocity), and occur in ahigh-strength matrix.

Using E=2x1011Nm-2, v=500ms-1, v=0·3,(J= 109Nm-2 and h, c1, and Xmin as before, the authorsfind that the smallest detectable crack with the narrow bandsystem has a diameter 2amin = 10- 6 m. amin has a differentdependence on Xmin for dislocation loop and crack-growth

, sources (equations (12) and (18)). The relative sensitivity

The detected displacement is reduced by a factor equal tothe ratio of the ,filter frequency to the band width of theincoming signal.

Now consider the same incoming pulse filtered by a high-pass filter which can be represented as the complement of asecond Gaussian. of unit area and width r 2' where r 2

corresponds to the corner frequency of the filter, i.e. thefilter function is

[b(t)- T2~ exp (- t2/d)]

When convolved with the input pulse this gives thedifference between two Gaussian functions, the first equal tothe input and the second of equal--' area, but of width

Band width limitationsAll acoustic emission detection systems have a limitedfrequency response. The most serious consequence is thatthe wave form is distorted so that dynamic informationabout the source is lost. Thus, the wide band system wasdeveloped to have the widest possible band width to enablemeasurement of dynamic source properties. There is,however, a second effect of a limited frequency response: itleads to a reduction in effective sensitivity of a recordingsystem, dependent upon the band width of the detectionsystem relative to the spectrum of the source. The narrowband system was band-limited to the frequency range 0·1-1·0 MHz (and many other systems in current use are evenworse). It is therefore importan t to examine the effect of thisupon sensitivity.

For simplicity, we assume that the emission pulse isGaussian of width r and height x. Its area is krx where k is aconstant. First examine the effect of low-pass filtering.Suppose the pulse is filtered by a filter whose impulseresponse is the Gaussian pulse of width r 1 and unit area, i.e.(1/r 1~) exp ( - t2 /rf). The effect of filtering, expressedmathematically by convolution in the time domain, is togenerate a further Gaussian pulse as the output, of widthJ r2 +rr and height X'.

The areas of the Gaussian pulses multiply so that the areaof the input and output pulses are equal, i.e.

kxr = kx'Jr2+rr . (20)

Th\ls, by rearranging equation (20) and noting that r ~ r 1

· (23)

· (22)

· (24)xr2

I 2x-x =-.2r2

J r2 + r~ and height x', i.e.

kxr = kx'p+r~

The height of the output pulse is therefore

X-X'=X[l- ~]

For r2 ~ r

Detectable acoustic emission sources in A533BsteelThe above considerations have indicated limitations ofacoustic emission techniques. Because of these, not alldeformation and fracture processes that occur in a metal arelikely to be detectable. In the material studied here, a widerange of deformation and fracture processes were inducedand the authors will now attempt to determine which ofthese were responsible for the observed emission. This isgreatly simplified by the exclusion of the MnS stringers fromthis material since these may be a source of acousticemission that could mask the emission from other processes.

Premaximum stressThe four most rapid quench rates resulted in no significantemission during this period of deformation. The few isolatedand random emissions that were observed could have beenthe decohesion-cracking of spheroidal inclusions. Itappears that deformation processes in material in theseconditions generate no detectable emission. This is probablybecause the high dislocation density and the small width ofmartensite laths result in very short distances for dislocationpropagation, well below the detection limit of '" 100 Jlm fora single dislocation. Air and furnace cooling gave increasedlevels of continuous emission around yield. The occurrenceof these correlates with an increase in the distance of

Due to the high-pass filter the height of the pulse is reducedby a factor 0·5 (r 2/r)2 , proportional to the square of the ratioof the band width of the signal to that of the filter.

A pulse whose width was 100 ns (a typical value for anemission from a crack) subjected to a 1 MHz low-pass filter(rl = 0·35 Jls) will have its amplitude attenuated by a factorof 0·27 (from equation (21)). Furthermore, the degree of thisattenuation will vary with the band width of each emission.Similarly, a broad pulse of 10 JlSwidth will suffer significantamplitude attenuation when subjected to a 100 kHz high-pass filter (r2 = 3·5 Jls). From equation (24) the amplitude isred,uced by a factor 0·06.

The narrow band detection system thus tends tounderestimate, and may even fail to detect, signals whoseband width extends beyond the pass band of the detectionsystem.

Further complications also exist due to reflections fromboundaries of the specimen. These are difficult to analyseand are ignored here. Their main effect will be to increase,somewhat, the probability of detecting a source. In thequalitative discussion to follow the authors assume thateffects due to reflections are approximately cancelled bythose due to limited band width.

In summary, while narrow band techniques -have highsensitivity and are thus able to detect weak signals, theydistort the acoustic wave form so severely that little usefulinformation can be deduced about the source. Conversely,wide band measurements record only energetic events butwith sufficient fidelity to facilitate source characterization.

. (19)

. (21)

amin (dislocation) 1/2------- ex Xminamin (crack)

I xrX"'-

r1

Metal Science June 1981

260 Scruby et at. Microstructure and acoustic emission in Mn-Mo-Ni steel

dislocation propagation, due to the disappearance of thelath structure and a reduction in initial dislocation density.

The tempering of water-quenched specimens had agreater effect. Lightly tempered martensites generated littleor no emission which was consistent with dislocationpinning by other dislocations, carbides, and lathboundaries. Prolonged tempering, however, resulted in astrong increase in emission as the quenched-in dislocationstructure and the lath boundaries were annealed-out, andthe size and carbide spacing increased. The calculationsabove estimate that a single detectable dislocation loopwould need to expand to a radius of at least '" 100 Jlm if itsvelocity were '" 200 m s -1. In the quenched condition thelath width was ~ 10 Jlm, so even at a velocity of 200 m S-1 avery large number of dislocations would have to movecooperatively in order to generate a detectable signal. Shorttempering treatments, by introducing small impenetrablecarbides while not greatly changing dislocation density orlath structure, would tend to reduce still further the effectivedistance of propagation between pinning events. In thespecimens tempered for long times many of the barriers todislocation motion are removed or their effect reduced,resulting in the detection of many very small amplitudeemissions for a short interval of strain around yield. Forestinteractions during work hardening presumably limitdislocation propagation distances during post-yieldstraining and lead to a progressive disappearance of theemission.

Post-maximum stressVarying the cooling rate of quenched specimens markedlyaffected the emission activity· during post-maximum stressdeformation (Fig. 16). The two fastest quenched specimensgenerated very large amplitude emissions as final fractureapproached. This period of deformation was characterizedby the formation of subcritical cleavage cracks. If the crackswere a lath packet in length ('" 30 Jlm), then for these steelsemissions of amplitude about 30 times the threshold wouldbe expected. It thus appears that the most likely source ofemission was the formation of cleavage microcracks andthat the strain dependence of the detected emission was anindication of the rate of cleavage microcrack formation.

When the cooling rate was reduced slightly (by quenchingin water) the emission activity was both reduced in energyand delayed, occurring closer to final failure (Fig. 16c). Herethe fracture mode was alternating shear with only a fewisolated areas of cleavage. Both fracture processes couldgenerate detectable emission. 5,9 The areas of cracking fromeach were sufficient to generate signals whose amplitudewould greatly exceed the threshold value. The narrow bandresults provide insufficient evidence to determine whetherone or other or both were the sources of emission, althoughthe reduction in emission may be an indication that in thesematerials the alternating shear processes generate weakersignals than cleavage.

Neither the slow cooled nor any tempered specimengenerated any emission other than at final failure. Thesespecimens fractured essentially by ductile void coalescence.However, the wide range of heat treatments resulted in awide range of void densities and sizes. These results areconsistent with those of earlier studies of a 4Ni-Cr steel6and indicate that for a wide range of microstructures voidgrowth and coalescence (which involve no long-distancepropagation of individual dislocations or sudden fractureover several micrometres) generate no detectable acousticemission under uniaxial loading.

Carbide cracking occurred in the samples containing

Metal Science June 1981

large carbides (Fig. 8). These were M23C6 and M3C incomposition (Fig. 9) and ~ 2 Jlm in size. Signals from thefracture of these particles are just on the limit ofdetectability if the fracture occurs in a brittle manner. Theabsence of acoustic emission at the large plastic strains whenfracture is expected to have occurred indicates that crackingof these composition carbides is not a major acousticemission source.

Acoustic emission characterization ofmicrocrack growthThe acoustic emission results discussed above have shownthat only a few of the wide range of deformation andfracture events are detectable, even when using sensitiverecording techniques. Because of the severe distortionexperienced by the signals during the· sensitive narrow bandtests, these can provide only a qualitative description ofemission source activity.

However, selecting only those specimens. that emittedlarge-amplitude emission signals, i.e. the iced-brine andwater-quenched specimens, so that the less sensitive broadband system could be used, it is possible to measureemission wave forms sufficiently free from distortion toobtain a single parameter description of the source in termsof a time-varying crack volume (Table 6). The resultsshowed distributions of source volume and lifetime for bothheat treatments (Fig. 21). It is noted that typical lifetimesare '" 100 ns, so that the wave forms have significantfrequency content outside the band width of all except wideband detection systems. Thus, with narrow band systemsthere are likely to be the errors in amplitude estimationoutlined above.

It is more usual to deal with crack length rather thanvolume for crack-growth events. The authors calculatecrack length for each recorded wave form making the sameassumptions, principally that the crack is horizontal,circular, and under mode I elastic loading, and using thesame equations ((13)-(18)) as above. A second usefulparameter, the crack velocity, is calculated by dividing thecrack diameter by its lifetime, assuming uniform radialvelocity of the crack front. The results are plotted ashistograms in Fig. 22 and the mean values (with standarddeviations also given as a guide to the widths of thedistributions) tabulated in Table 7 which shows consistencybetween the results of individual specimens, but withsomewhat greater deviations for the specimens with thepoorest statistics.

By comparing the histograms for the two different coolingrates and the mean crack lengths and velocities (Table 8) itcan be seen that while the slightly slower cooling rate givesmarkedly fewer emissions, it gives virtually indistinguish-able distributions of crack length, lifetime, and velocity, Itwas concluded above that the source of emission duringpost-maximum stress deformation in iced-brine quenchedmaterial was cleavage fracture. In the case of water-quenched material there was some doubt about whether thesource was again cleavage or alternating shear fracture. Thequantitative emission data show no difference withinstatistical error between the two materials, suggesting thatthe emitting crack-growth process is the same in bothmaterials. Furthermore, the estimated total detected crackareas (calculated by summing the individual microcrackareas) is '" 4 and '" 1% of the gauge section for iced-brineand water-quenched specimens respectively. The lowdetected area in the latter case would be consistent withfracto graphic evidence of a lower incidence of cleavagemicrocracking. This is an example of the extra power of

Scruby et al. Microstructure and acoustic emission in Mn-Mo-Ni steel 261

calibrated broad band techniques over qualitative narrowband techniques for source characterization.

The typical emission source in both materials appears tobe a cleavage 'microcrack about 50-60 Jlm -long,propagating at a speed of 450 m s-1. The length issomewhat greater than the lath packet size which is likely tobe the most significant dimension and is somewhat less thanthe mean prior austenite grain size ('" 70 Jlm). However, itmust be borne in mind that the broad band detection systemhas a threshold of '" 10 J.1mcrack diameter, so that loss oflow-amplitude signals tends to give too high a mean cracklength. With this proviso, the agreement with fractographyis again good. It is not possible to measure crack speedindependently, but the mean value (0'07cd, deduced byacoustic emission analysis, is consistent with theoreticalcalculations for crack growth with limited plasticdeformation at the crack tip.18

The results are also in excellent agreement with dataobtained from cleavage microcrack growth in mild steeltested at 77 K9 (Table 8 and Fig. 23: the histogram for'pure' A533B inc1udesthe data for both quench rates). Someagreement should be expected since the grain and lathpacket sizes are comparable in the two steels; however, theagreement is made more remarkable when it is rememberedthat the mild steel specimens were tested at 77 K to reducecrack-tip plasticity. It might, for instance, have beenexpected that there would be a greater difference betweenthe crack speeds. The effects of assuming elastic loadingwhen in reality there is some plasticity at the crack tip in thepure A533B samples are to underestimate the crack opening{)and thereby overestimate the crack length and velocity.Thus the true crack lengths and speeds are likely to besomewhat less than those calculated above.

CONCLUSIONS1. The origin of yield-region acoustic emission from Mn-

Mo-Ni A533B steel containing a low concentration ofsulphur appears to be the motion of groups of dislocationswhose combined slip distance exceeds about 100 J.1m.Suchevents are restricted to slowly cooled or heavily temperedmicrostructures as these have the most widely spaceddislocation barriers. Carbide cracking does not generatedetectable signals in this alloy system in the heat-treatedconditions investigated.

2. Acoustic emission before fracture appears only ifcleavage or possibly alternating shear micro fracture occurs.The wide range of ductile dimple fracture modes produced,in thiss.teel generate no detectable emission over _thefrequency range 0'1-1'0 MHz.

3. Quantitative broad band recording techniquescoupled with time domain inversion of measured waveforms enabled the emission source during fracture to becharacterized. This source corresponds typically to acleavage microcrack propagating '" 50 J.1mat an averagespeed of ",450 m S-1.

ACKNOWLEDG M ENTSThe authors are grateful to Dr S. G. Druce for helpfulcomments during the preparation of the manuscript, toA. Forno of NPL for the forging of the plate material, andto R. Hawes for help with the microsegregation studies.

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Metal Science June 1981