Astm stp 559 1974

Contents Fracture Toughness Test Methods for Abrasion-Resistant White Cast Acoustic Emission from 4340 Steel During Stress Corrosion Cracking-- H.. Diesburg 1 Fracture Toughness Test M

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FRACTURE TOUGHNESS AND SLOW-STABLE CRACKING

Proceedings of the 1973 National Symposium on Fracture Mechanics, Part I

A symposium sponsored by Committee E-24 on

Fracture Testing of Metals, AMERICAN SOCIETY FOR TESTING AND MATERIALS University of Maryland, College Park, Md., 27-29 Aug 1973 ASTM SPECIAL TECHNICAL PUBLICATION 559

P C Paris, chairman of symposium committee

G R Irwin, general chairman of symposium List price $25.25

04-559000-30

~(~l~ AMERICAN SOCIETY FOR TESTING AND MATERIALS

1916 Race Street, Philadelphia, Pa 19103

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(~) A M E R I C A N SOCIETY FOR T E S T I N G AND MATERIALS 1974 Library of Congress Catalog Card N u m b e r : 7 4 - 8 1 1 5 4

N O T E Thc Society is not responsible, as a body,

for the statements and opinions advanced in this publication

Printed in Baltimore, Md

August 1974

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Foreword

The 1973 National Symposium on Fracture Mechanics was held at the University of Maryland Conference Center, College Park, Md., 27-29 Aug 1973 The symposium was sponsored by the American Society for Testing and Materials through Committee E-24 on Fracture Testing of Metals Members of the Symposium Subcommittee of Committee E-24 selected papers for the program Organizational assistance from Don Wisdom and Jane Wheeler at ASTM Headquarters was most helpful

G R Irwin, Dept of Mechanical Engineering, University of Maryland, served as general chairman Those who served as session chairmen were

H T Corten, Dept of Theoretical and Applied Mechanics, University of Illinois; C M Carman, Frankford Arsenal; J R Rice, Div of Engineer- ing, Brown University; D E McCabe, Research Dept., ARMCO Steel;

J E Srawley, Fracture Section, Lewis Research Center, NASA; E T Wessel, Research and Development Center, Westinghouse Electric Corp.; and E K Walker, Lockheed-California Co

The Proceedings have been divided into two volumes: Part 1 Fracture Toughness and Slow-Stable Cracking and Part II Fracture Analysis

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Related ASTM Publications

Stress Analysis and Growth of Cracks, STP 513 (1972), $27.50

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Contents

Fracture Toughness Test Methods for Abrasion-Resistant White Cast

Acoustic Emission from 4340 Steel During Stress Corrosion Cracking

H H CHASKELIS, W H CULLEN, AND J M KRAFFT 31

Comparison of Acoustic Emission to the Fracturing Process 36

Effects of Shot-Peening Residual Stresses on the Fracture and Crack-

Fracture Properties of a Cold-Worked Mild Steel E, J RIFLING

Materials and Procedure

Results and Discussion

More on Specimen Size Effects in Fracture Toughness Testing~J 6

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Dynamic Compact Tension Testing for Fracture Toughness -P c pARIS,

Further Studies of Crack Propagation Using the Controlled Crack Propa-

Double Torsion Technique as a Universal Fracture Toughness Test

Method J o OUTWATER, M C MURPHY, R G KUMBLE, AND

Theoretical Basis for the Double Torsion Technique 130

Measurement of KIe on Small Specimens Using Critical Crack Tip Open- ing Displacement J N ROBINSON AND A S TETELMAN 139

Correlation Between Fatigue Crack Propagation and Low Cycle Fatigue

Properties -SAURINDRANATH MAJUMDAR AND JODEAN MORROW 159

Description of the Fatigue Crack Propagation Model 163 Mechanics and Fatigue Analysis of the Crack Tip Region 164 Influence of Material Properties on the Coefficient of Eq 13 168

Comparison of Eq 14 with Barsom's Data on Steel 172

Effect of Stress Concentration on Fatigue-Crack Initiation in HY-130

Steel J M BARSOM AND R C M e NICOL 183

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Evaluation of the Fatigue Crack Initiation Properties of Type 403 Stain- less Steel in A i r and Steam Environments -w G CLARK, JR 205

Subcritical Crack Growth Under Single and Multiple Periodic Overloads

in Cold-Rolled S t e e l ~ F H GARDNER AND R I STEPHENS 225

Effects of R-Factor and Crack Closure on Fatigue Crack Growth for

A l u m i n u m and Titanium A l l o y s - - M KATCHER AND M KAPLAN 264

Application of the Linear Superposition Method to the Fastener

Achieving a Fatigue Stress Intensity Threshold 290

Rapid Calculation of Fatigue Crack Growth by lntegration T R

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STP559-EB/Aug 1974

Introduction

Readers of this volume will not be disappointed with regard to novelties

of current and practical interest in fracture toughness and slow-stable cracking These range from unusual test methods to a puzzling effect

of lateral specimen dimensions on Kic values for an aluminum alloy Observational techniques include acoustic emission, both in relation to onset of rapid fracture and stress corrosion cracking, tape recordings as

an assist for rapid load testing, and use of rubber castings to verify measurements of crack opening stretch Toughness measurements are reported for white cast irons and cold-rolled steel The papers dealing with fatigue cracking include a low cycle fatigue viewpoint on fatigue crack growth, effects of shot peening, initiation of fatigue cracking as a function

of notch root radius, as well as effects of overloads, mean K, and mechanical fastener pressure

The development of technology in this field has prospered over the years so that often novel approaches soon become routine techniques to solving problems This volume is another contribution to the engineer and metallurgist faced with fracture problems

With two exceptions, all of the papers in this volume were presented

at the 1973 National Symposium on Fracture Mechanics held at the College Park campus of the University of Maryland, 2 7 - 2 9 Aug 1973 The two exceptions were a paper offered for this symposium but not presented and a late submission of a paper from the 1972 symposium The companion volume, STP 560, covers fracture analysis

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D E Diesburg 1

Fracture Toughness Test Methods for

Abrasion-Resistant White Cast Irons

Using Compact Specimens

REFERENCE: Diesburg, D E., "Fracture Toughness Test Methods for

Abrasion-Resistant White Cast Irons Using Compact Specimens," Frac-

Society for Testing and Materials, 1974, pp 3-14

ABSTRACT- The fracture toughness of abrasion-resistant white cast irons has been measured, using precracked compact specimens Some procedures used for precracking the brittle cast irons were outside the ASTM Test for Plane-Strain Fracture Toughness of Metallic Materials (E 399-72) requirements but still gave valid results The excellent reproducibility, combined with a range in toughness values of 17.5 to 28.5 ksiVTff (19.2 to 31.4 M N / m 8/~) for abrasion-resistant white cast irons, provided the sensitivity necessary to distinguish differences in the toughness of white cast irons resulting from variations in composition or microstructure The fracture toughness of three commonly used irons, 27Cr, 9Cr-6Ni, and 20Cr-2Mo- 1Cu, was compared in the as-cast (and stress-relieved) condition Heat treating the 20Cr-2Mo-lCu iron substantially increased the hardness and reduced the fracture toughness slightly

KEY WORDS: abrasion-resistant iron, white cast iron, fracture properties, toughness, evaluation, mechanical tests, fatigue (materials), mechanical properties

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also tend to be brittle Not only must the processing equipment withstand abrasion, but it must do so without fracturing during service Therefore,

in every application where abrasion is a factor, the material used must provide abrasive wear resistance and adequate toughness

Research laboratories have been able to determine the wear rates of many materials under various types of abrasive environments, but the fracture resistance of these materials has been difficult to evaluate before the material is placed into service Although abrasion-resistant materials are all relatively brittle, lack of toughness may shorten the service life

by an amount depending on the application Extensive investigations of fractured components has led to the conclusion that there can be a large difference in fracture resistance in materials, which, when tested in a Charpy impact testing machine, may exhibit less than 2 ft-lb (0.3 kgfm/cm2) Before a research laboratory can investigate the metallurgical parameters providing the best fracture resistance, it is first necessary to devise a test method that can reproducibly provide enough sensitivity to distinguish the various levels of toughness that can be present in materials that fail at low levels of absorbed impact energy

Plane-strain fracture toughness is a measure of fracture resistance in the early stages of crack propagation and has been successfully measured for a cast steel 2 and gray and ductile cast irons?, ~ Fracture toughness measurements are sensitive to changes in microstructure, which is exactly the type of measurement needed to investigate the factors controlling the toughness of white cast irons

The method of testing for the plane-strain fracture toughness of metallic materials ( A S T M Test for Plane-Strain Fracture Toughness of Metallic Materials ( E 3 9 9 - 7 2 ) ) provided the basis for the development of a test that can reproducibly provide enough sensitivity to measure the toughness

of various abrasion-resistant white cast irons The attempt was made to follow exactly the method outlined in A S T M E 399-72, but it was soon realized that a few of the specifications could be relaxed and still provide

a valid measurement of plane-strain fracture toughness for these brittle irons This paper (1) describes the development of the test procedure used to measure K~e in white cast irons and (2) cites a few examples

to illustrate the spread in values that can be expected

2 Greenberg, H D and Clark, W G., Jr., Metals Engineering Quarterly, American Society for Metals, Aug 1969, pp 30-39

3 Glover, A G and Pollard, G., lournal of the Iron and Steel Institute, Feb 1971,

pp 138-141

"Lazaridis, A., Worzala, F J., Loper, C R., and Heine, R W., Transactions,

American Foundrymen's Society, 1971, Vol 79, pp 351-360

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DIESBURG ON ABRASION-RESISTANT WHITE CAST IRONS 5 Experimental Procedure

The irons used to establish the test procedure were available from previous laboratory investigations and, in many cases, were not in the recommended heat-treated conditions These available irons will be referred

to as Irons I, II, III, IV, V, and VI and are further described in the appendix

The white cast irons (compositions given in Table 1 ), used to illustrate typical values of fracture toughness expected of white cast irons in the properly heat-treated condition, had been cast into baked-sand molds as 1-in ( 2 5 - m m ) thick plates from 125-1b (57-kg) induction-melted heats The irons were tested in the as-cast plus stress-relieved condition The

2 0 C r - 2 M o - l C u iron was also tested after the matrix microstructure had been changed from austenite to predominantly martensite through heat treatment

Specimen Preparation

Compact test ( C T ) specimens with dimensions as shown in Fig 1 were prepared for each iron The outer dimensions were obtained by grinding, while electrical discharge machining ( E D M ) was used to prepare the pin loading holes and the crack initiating notches The notch was machined

in two steps, the final step producing a 0.01-in (0.3-mm) wide by 0.09-in (2.3-mm) deep slot A fatigue crack was grown at the base of the 0.01-in (0.3-mm) wide slot using an SF-1U Sonntag fatigue testing machine with

a loading cycle that always kept the specimen loaded in tension The R-values (ratio of minimum to maximum load) were always less than 0.1 The crack length was measured on both broad surfaces of the C T specimens (polished through 600-grit paper)

Testing

Once the specimens had been precracked, the load required to extend the crack was determined by pulling the specimen in a tension testing machine The same clevis fixtures used on the fatigue machine were used

to load the specimens in the tension testing machine

TABLE 1 Chemical analysis of the irons

Element, weight percent

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FIG 1 Compact specimen

A cantilever extensometer, calibrated to meet the requirements of ASTM, was used to monitor crack extension The extensometer was clipped between two parallel knife edges (clip gage support, Fig 1), which had been machined separately and attached to the specimen with LocTite 310 metal bonding adhesive A special fixture was used to hold the knife edges in a parallel position until the adhesive cured The strength

of the, adhesive bond was proven to be sufficient by clipping the extensometer between the knife edges of one of the specimens for a duration of

24 h No extensometer movement, and therefore no slippage, was observed under the load applied by the extensometer

The precracked specimens were installed in the tension testing machine and the extensometer was attached A crosshead rate of 0.005 in./min (0.13) ram/rain) provided a loading rate of approximately 700 lb/min (320 kgf/min) A typical load versus crack displacement curve is shown

in Fig 2 Oftentimes Po was found to be Pr x (Fig 2)

Results and Discussion

An experimental procedure was developed for measuring the plane- strain fracture toughness of brittle white cast irons An attempt was made

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DIESBURG ON ABRASION-RESISTANT WHITE CAST IRONS 7 EXTENSION ACROSS KNIFE EDGES (MM)

I j~.-EXTENSION OF ELASTIC REGION -~800

l /

I / / ~ L I N E HAVING S~'o LESS SLOPE /

Test R e q u i r e m e n t s

The requirements outlined by ASTM E 399-72 can be divided into three main categories: (1) specimen geometric requirements, (2) precracking requirements, and (3) testing requirements All specimens did meet the dimensional requirements: B_> 2.5 (K~e/~ys) 2, fatigue crack

c was at least 0.05 in (1.3 mm) in length and greater than 5 percent of

L, and B was at least 0.25 W but less than W ~ The testing requirements were also met The requirements that were not always met involved the growth of the fatigue crack

The difficulty in controlling the growth of the fatigue crack often resulted in cracks that did not quite meet specifications The lowest stress

5 B, c, a, L, and W are defined in Fig 1

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intensity required to propagate a fatigue crack was usually very close

to the maximum limit of 0.6 of the stress intensity, KQ, routinely deter-

mined in fracture toughness testing Often Kf(m,~.~, the maximum stress

intensity used to grow the last portion of the fatigue crack, exceeded 0.6

KQ and reached as high as 0.9 K~ A comparison of fracture toughness

values (KQ) with the values obtained from specimens precracked with a

stress intensity less than 0.6 KQ (Table 2), indicated that Kf( ~ could

be as large as 0.84 KQ (Iron VI) without altering the measured fracture

toughness of brittle white cast irons

The fatigue crack path in white cast irons was very dependent on

dendrite orientation Nonrandom orientation of dendrites could cause the

paths to deviate from the plane of symmetry for the specimen Fortunately,

the resulting fracture toughness determinations were not affected, even

when the angle from symmetry reached 20 deg (Iron VI in Table 2,

Specimen 4)

ASTM recommends determining the length of the fatigue crack by

taking measurements from the fractured surface However, it was very

difficult to make precise measurements of the fatigue crack on the frac-

tured surface because there was usually no distinct boundary between the

fatigued surface and the fractured surface It was observed, however, that

the crack front usually formed a linear boundary between the surface

traces This observation was made by exposing five precracked specimens

to moisture to form a slight layer of rust, which clearly outlined the loca-

tion of the crack tip prior to testing The crack lengths on subsequent

specimens were measured along the two surface traces of the fatigue crack

and then averaged The nonrandom cast orientation sometimes resulted in

nonuniform crack lengths About one sample in six could be expected to

have one side precracked to a length less than 90 percent of the average

crack length Again, the resulting fracture toughness determinations were

not affected (Iron III in Table 2)

The shortest fatigue crack length was 0.18 in (4.6 mm), which was

substantially greater than the required 5 percent of L The distance a

from the crack tip to the loading plane was greater than 0.45 W for all

specimens However, there were two specimens (Iron VI, Specimens 2

and 3, in Table 2) for which this distance was greater than the maximum

limit of 0.55 W, as set by ASTM E 399-72 The effect of this extra

length crack could not be detected The difference between the KQ obtained

from the specimen having the correct notch depth a and the specimen

having a greater than 0.55 W was 0.9 and 8.6 percent The fact that

one specimen with a long crack length gave a consistent KQ indicates

that the difference of 8.6 percent was not caused by a exceeding 0.55 W

Downloaded/printed by

University of Washington (University of Washington) pursuant to License Agreement No further reproductions authorized.

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10 FRACTURE TOUGHNESS AND SLOW-STABLE CRACKING

However, since the seemingly low value of KQ for one specimen cannot be

explained, it was decided that the 0.55 W maximum limit for crack

length a should be maintained as a requirement for a valid Kit test of

white cast irons

Branching of the fatigue crack was observed in a few specimens (Irons

II, IV, and V) During testing one branch stopped while the other propa-

gated to complete fracture The measured fracture toughness values were

always higher in the branched specimens compared to the values obtained

from the specimens that did not branch It was concluded that valid

fracture toughness measurements cannot be obtained with specimens in

which cracks branched during precracking

According to ASTM E 399-72, about half of the fracture toughness

determinations in Table 2 can be labeled at Kic However, it was pointed

out that those which did meet all of the requirements did not differ sig-

nificantly from the "so-called" invalid determinations, with the exception

of the specimens that had a branched precrack It was concluded that

although all ASTM E 399-72 requirements must be met for ductile

materials, they may be relaxed for brittle white cast irons The only

requirements that must be met are that the crack length not exceed 0.55 W

and the fatigue crack not be branched prior to testing

Toughness of Common Abrasion-Resistant White Cast lrons

The range of fracture toughness values (Table 2) obtained for the

irons Used in establishing the testing procedure, combined with the ex-

cellent reproducibility, permitted further investigation of the effect of

microstructure and heat treatment on the resistance of white cast irons to

fracture

Three commonly used abrasion-resistant white cast irons are 27Cr,

9Cr-6Ni, and 20Cr-2Mo-lCu The compositions of the irons tested are

given in Table 1 All three irons are used in certain applications in the

as-cast (and stress-relieved) and appropriately heat-treated conditions

These irons were all tested in the as-cast condition The 20Cr-2Mo-lCu

iron was also tested in two heat-treated conditions The heat treatments

and resulting matrix microstructure and hardness are given in Table 3

Representative microstructures of the irons in the conditions tested are

shown in Figs 3 and 4

The fracture toughness values given in Table 3 are each an average of

three determinations The spread in values obtained for any given iron

was always within 3 percent of the average value Heat treating the

stantially increased the hardness and slightly reduced the fracture tough-

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FIG 3 As-cast white irons: (a) 27Cr iron, eutectic carbides in a predominantly

austenitic matrix; (b) 9Cr-6Ni iron, eutectic carbides in a matrix o] austenite and

large plates of martens#e; (c) 20Cr-2Mo-lCu iron, eutectic carbides in a matrix of

austenite containing a few patches of fine carbide particles Etched with 1 percent

picral and 5 percent HCI ( •

11

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] 2 FRACTURE TOUGHNESS AND SLOW-STABLE CRACKING

FIG 4 Microstructures in 20Cr-2Mo-lCu white cast iron after heat treating at

5 percent HCI ( X 5 0 0 )

TABLE 3 Fracture toughness of common white cast irons

Fracture

A=austenite, M=martensite (listed in the order of decreasing volume percent)

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DIESBURG ON ABRASION-RESISTANT WHITE CAST IRONS 13 ness from the value the iron h a d when in the as-cast (and stress-relieved)

condition

T h e choice of iron for a given application will depend on the com-

bination of hardness and fracture toughness desired The iron with the

highest fracture toughness would be the 2 0 C r - 2 M o - l C u or 27Cr iron in

the as-cast condition However, if a hardness of 63 Rc was desired, the

2 0 C r - 2 M o - l C u iron could be heat treated and very little fracture tough-

ness would be lost

Conclusions

1 A valid fracture toughness of abrasion-resistant white cast irons

can be measured even if there is a slight relaxation of the requirements

outlined by A S T M E 399-72

2 An average fracture toughness, as determined from three compact

tension specimens, c a n be expected to be accurate to within at least

3 percent

3 T h e range of fracture toughness values, obtained f r o m the irons

tested, provides m o r e than enough sensitivity to permit the testing tech-

nique to be used in the evaluation of the effect of various microstructures

on toughness

TABLE 4 Description of Irons 1 through VI

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APPENDIX

The white cast irons used to establish plane-strain testing techniques

were not necessarily in a recommended heat-treated condition Irons II,

IV, V, and VI had not been stress relieved, and Iron III had been furnace

cooled Table 4 gives the heat treatment and resulting matrix microstruc-

ture and hardness of each iron

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J C R a d o n 1 a n d A A P o l l o c k 2

Development of Fast Fracture

in a Low Alloy Steel

REFERENCE: Radon, J C and Pollock, A A., "Development of Fast

Fracture in a ' Low Alloy Steel," Fracture Toughness and Slow-Stable

1974, pp 15-30

ABSTRACT: Fracture mechanics, acoustic emission, and fractography

were used to study the process of crack growth in 2-in thick double canti-

lever beam specimens of a low-alloy steel over a wide temperature range

Kzc values ranged from 33 ksiVin, at 200~ to 240 ksiVTff, at -t-18~

The acoustic emission rate rose steadily with load up to fast fracture

Changes in the emission amplitude distribution were observed shortly before

fracture at one of the warmer temperatures Emission during load holds was

measured and emission waveforms were photographed Fractography

showed a small region of ductile tearing prior to fast fracture at warmer

temperatures A model of the interaction of ductile tearing and cleavage

fracture is proposed Ductile tearing is seen as a process taking time and

strongly dependent on temperature Cleavage is seen as a rapid process

averted by ductile flow and tearing Emissions are believed to be produced

by plastic zone growth and by cleavage The experimental facts can be con-

sistently interpreted through this small set of assumptions

KEY WORDS: fracture properties, acoustic emission, crack propagation,

fractography, brittle fracture, mechanical properties

T h e o b j e c t i v e of this p r o j e c t was to c o m b i n e t h r e e t e c h n i q u e s : f r a c t u r e

m e c h a n i c s , a c o u s t i c emission, a n d f r a c t o g r a p h y - - i n o r d e r to g a i n a

d e t a i l e d p i c t u r e of t h e p r o c e s s of c r a c k g r o w t h in a l o w - a l l o y steel F r a c -

t u r e m e c h a n i c s p r o v i d e s the b a s i c t o o l for r e l a t i n g the p h e n o m e n o n of

c r a c k g r o w t h to the a p p l i e d stresses a n d is u s e d in e n g i n e e r i n g design

15

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TABLE 1 Material characteristics

Strength, Tensile Strength, Charpy

Steel tonf/in ~ tonf/in 2 50% FATT Remarks

tempered

mation events that precede fracture and is used in engineering inspection Fractography provides insight into the microstructural and metallurgical factors that steer the course of crack growth and is used in failure investigations

The project shows that this combination of techniques is productive

In considering the results of experimental work in the three areas, we were able to develop a model of interacting deformation mechanisms, providing a simple basis for understanding the development of fast fracture in this material

Material, Specimen Geometry, and Fracture Mechanics Results

The material used for these tests was a low alloy steel whose chemical composition and mechanical properties are shown in Table 1

The specimens used for these tests were of the double cantilever beam ( D C B ) parallel edge type, which had previously been found particularly suitable for the study of fast fracture in materials of this type T h e development of this geometry is described in detail in Ref 1 3 Due to the fact that the D C B specimen could b e designed in any convenient size, there was no need for an excessively large testing machine A specimen length of 30 in and a width of 2 in was used, as shown in Fig 1 The other dimensions were carefully balanced to avoid yielding of the aims and to give satisfactory performance over a wide temperature range ( - 2 0 0 ~ to + 1 8 ~

After sharpening the starter saw-cut by fatiguing in a three-point bending rig at one-fourth nominal yield and 6 0 / 1 0 0 kc, the specimen was m o u n t e d

in a standard liquid nitrogen/petroleum ether cooling bath and acoustic

a The italic numbers in brackets refer to the list of references appended to this paper

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RADON AND POLLOCK ON FAST FRACTURE 17

SAW 0.)1' STAR'rER

/ "

I ' 4

I

P L O P A G ~ " r l N ~ ~ RACI~

~ m O , ~ 6 t tr

I I -~

4 ~ ~ ~o'ro, J (.oo~' ~ o o ' r ~A~

FIG 1 DCB specimen (straight edge type)

emission transducers were attached The specimens were tested in a Tinius-

Olsen 120 000 Ib capacity testing machine at a crosshead speed of 0.02

in./min On reaching the critical load the crack propagated rapidly along

the median plane until a new equilibrium condition had been attained

T h e crosshead speed was increased to 0.05 in./min and further crack

jumps followed, typically 2 to 4 per specimen according to temperature

At the lowest temperatures ( - 135~ and below) a type of slow tear was

observed, with the crack advancing in a large number of small unequal

steps A typical record of load versus extension is shown in Fig 2 The

exact length of the crack jumps was measured directly from the fracture

surface after the test

In calculating Kic values, the critical loads for crack initiation and arrest

were read directly from the load record and values for crack opening

were corrected for the amount of machine extension The slope of the

elastic loading lines, which may be extrapolated to the origin, was used in

the calculation of specimen compliance ~ Alternatively, d~/da was

measured on another test piece by introducing sharpened slots of increas-

ing length The derivative of the compliance, which therefore includes

the effect of the side grooves, can be used in the calculation of strain

energy release rate from which K~c values were derived Calculated values

G~

for initiation and arrest are shown as a function of temperature in Fig 3

T h e curves show some similarity with the behavior of medium strength

steels [2], but the arrest curve is much closer to the initiation curve, indi-

cating that the crack does not propagate so freely as might be expected

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I~a, 5 ~ o o o I bliwl sii

o.t o.~ o , 4 0.5 O.G 0.7 0 8 O.~

valid by ASTM E39g criteria

Initiation) 2 in plate 1:1 Arrest )

q -200 -~50

'I~ II

n -I00 Temperature ~

F I G 3 Fracture toughness versus temperature: Ducol steel D C B specimens (2-in

thick), crack propagation at 90 deg to rolling direction and plate surface

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Acoustic Emission

The equipment u s e d to monitor acoustic emission in these tests was a

conventional Dunegan system, counting threshold crossings, and a Cam-

bridge Consultants amplitude sorter which allows emissions to be batched

according to peak amplitude, the ranges being 20 dB apart Dunegan

140 Series transducers were used with both instruments

As a simple cheek on some possible sources of spurious emission (for

example, pin rotation), several of the tests were interrupted midway and

the load was reduced to a low level then reapplied A good Kaiser effect

was observed in all these cases, that is, emission was negligible during

unloading and during reloading until the previous maximum load was

reached However, there was some other evidence that sources of irre-

versible spurious emission (possibly fixture noise) were acting at low

stress levels, so the data acquired below 30 kips may be somewhat

unreliable

Amplitude sorter results for five tests are shown in Fig 4 The parameter

plotted is the number of events per second at the first level of the ampli-

tude sorter (50 to 500 ~V at the transducer) This emission rate data is

averaged over appropriate intervals up to the first crack jump, and plotted

as a function of load Since the specimens are behaving in an essentially

elastic manner, the plots are virtually the same as would be obtained by

plotting emission rate against time Emission rate increases with rising

Emission rate

(level I events/see)

-~ 135~ ~ -55~ ~ -32~ g8 -19~ a -8~

F I G 4 Emission ]rom first set of five specimens

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load These experiments did not show any systematic dependence of low-

load emission on temperature The cumulative emission to failure was

much higher for the warmer-temperature tests, because they proceeded to

higher loads with a rising emission rate

Analysis of the amplitude distribution yielded some interesting results

The tests at warmer temperatures tended to yield a larger proportion of

high-amplitude emissions A simple empirical parameter for describing

the amplitude distribution is

nz=number of emissions with amplitude exceeding 500 ~V

nl = number of emissions with amplitude exceeding 50 ~V

This parameter, averaged over the load rise to the first crack jump, is

plotted as a function of test temperature in Fig 5 One test gave an

anomalous result which is believed to be due to transducer mounting

problems encountered in that test Apart from this anomaly, there is good

FIG 5 Variation of amplitude distribution with temperature

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TABLE 2 Variation of amplitude distribution parameter during test at IO~

The drop in the 50 to 500 ~V emission rate prior to failure at - 8 ~

(Fig 4.) was confirmed in a second test at - 1 0 ~ The second test

showed, however, that this effect was outweighed by an increase in the

rate of high amplitude emissions The amplitude distribution changed

markedly at 85 to 90 percent of the failure load Data from this test is

shown in Table 2 The change in the parameter n 2 / n l above 65 kips is

statistically significant at the 0.1 percent fiducial level The approach to

failure was difficult to recognize by other means and it appears that in

this material at least, amplitude distribution analysis[3,4] offers the

most promising diagnostic of incipient failure At the lower temperatures

this effect was not observed; the amplitude distribution parameter was

constant, apart from statistical variations, up to the point of instability

Figure 6 shows a graph of cumulative emission count (smoothed)

against load for the test at - 1 0 ~ This data was taken from the Dune-

gan system at a gain of 86 dB The concave form of the graph is typical

for flawed specimens There is no feature in this graph that would serve

as a precursor to failure, but it is possible that by operating at a lower

gain, a recognizable feature would be obtained as a result of the pre-

viously described change in amplitude distribution

Figure 6 also illustrates the effect of "hold" periods in the test at

- 1 0 ~ The machine crosshead was stopped for several minutes at 30

and 60 kips Large quantities of emission were produced in both these

periods On resumption of loading, the emission rate was low, as if to

compensate for the "extra" emission produced during hold Figure 6

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FIG 6 Emission from Specimen D21 ( IO~

shows that the effect of an extended "hold" period is not entirely relieved

until the load has risen nearly 10 kips above the "hold" load

The implication is that we were conducting the tests at quite a high

strain rate for this material, in that the material was lagging well behind

the equilibrium state even at relatively low loads

Emission during hold is of interest because in some situations it may

serve as a warning of incipient failure[5,6] Parameters which may be

used to characterize "relaxation emission" are (1) counts and ( 2 ) time

distribution

In the test on Specimen D21, the 60 kips "hold" yielded nearly twice as

much emission as the 30 kips "hold." The time distribution of the relaxa-

tion emission is shown in Figure 7 The statistical quality of the data is

not very good, but the emission rate appears to fall off with a time con-

stant of the order of 1 min There are no obvious differences in this

respect between the two "hold" periods

With only two "hold" periods to work with, it would be premature to

say whether relaxation emission could be used for incipient failure diag-

nosis in this material The theory of relaxation has still to be developed;

these tests provide interesting grounds for thought but no obvious con-

clusions can be drawn

A number of emission waveforms were photographed as the tests pro-

Trang 30

F I G 7 Emission during hold at 30 kips (lower curve) and 60 kips (upper curve)

during the test of Specimen D21 ( 10~

ceeded A two-transducer technique allowed the rising edges of emission

waveforms to be photographed in real time Commonly, the natural

emission waveforms had a slower risetime than artificial waveforms pro-

duced by pulse injection from another transducer This rather surprising

result suggests that many of the deformation events had a fine structure

extending over 100 to 300 ~s A case was also recorded where the initial

event was followed by other short-rise events after 1.3 and 2.3 ms

Figure 8a shows a typical emission waveform Figure 8b shows a

more unusual waveform, which was produced by the second crack jump in

a test at - 1 0 8 ~ Instrumentation systems always saturate for the largest

emissions, so to capture this event we connected an accelerometer (nominal

resonance 50 kHz, sensitivity 60 m V / g 3 to 10 kHz) directly to an oscillo-

scope at 10 V / c m without intermediate amplification The resulting wave-

form is a remarkably perfect exponential decay Peak-to-peak voltage is

plotted against time in Fig 9 Extrapolating back to the time of origin, the

original signal level may be estimated at 140 V peak-to-peak, an unprece-

dentedly high figure

Emission waveform photography of rapid fracture events might welt be

used as a method of determining crack velocity We would expect the

transducer response to follow a perfect exponential decay only after the

crack has arrested; by examining the earlier part of the waveform it

Trang 31

24 FRACTURE TOUGHNESS AND SLOW-STABLE CRACKING

FIG 8 - - ( a ) Typical emission -0.5 ms/cm; (b) last [racture lO V/cm, 0.5 ms/cm

Trang 32

FIG 9 Peak-to-peak amplitude of decaying waveform as a function of time

should be possible to determine the length of time over which the crack is

in motion

Fractography

A number of stereoscan micrographs of selected regions of the fracture

surface were prepared Interest concentrated on slow growth bands which

were visible at the points of initiation and arrest These bands are wider

and more distinct at higher temperatures (width of the order of 0.015 in

at - 1 0 ~ and are almost imperceptible at temperatures below

- 1 0 0 ~ Figure 10 shows micrographs of the slow growth bands and of

the fast fracture region at two temperatures, - 1 1 5 and - 1 0 ~ The slow

growth bands are in clear contrast: ductile tearing is evident at - 1 0 ~

but at - 1 1 5 ~ quasi-cleavage is seen with only traces of ductile tearing

The fast fracture regions are essentially similar: failure is by transgranular

cleavage However, there are local traces of ductile tearing even in the

fast fracture region at - 1 0 ~ Above - 1 0 0 ~ the amount of ductile

tearing prior to failure increased strongly with temperature

Trang 33

FIG lO -(a) Slow growth band at 115~ (>(1000), (b) slow growth band at

- - 1 0 ~ ( • (c) [ast #acture at 115~ ( X 5 0 0 ) , and (d) last fracture at

Trang 34

Discussion

The main findings of these tests can be summarized as follows:

(a) K~e and K~rr increase with rising temperature, as is common with

steels

(b) The acoustic emission rate rises with K and the cumulative emission

prior to fracture increases rapidly with rising temperature

(c) In a test at - - 1 0 ~ the emission amplitude distribution changed

shortly before fast fracture In low-temperature tests this did not happen

(d) Fast fracture was by transgranular cleavage, with small amounts of

ductile tearing at warmer temperatures

(e) A slow growth zone of ductile tearing was well formed at warmer

temperatures At lower temperatures, this ductile tearing at the edge of

the fatigue precrack was progressively replaced by quasi-cleavage and the

slow growth zone became imperceptible

Before turning our attention to events at the crack tip itself, we should

consider the plastic zone which plays such an important part in fracture

mechanics In these tests it was impractical to examine the plastic zone

directly, either before or after the tests We can, however, make a rough

estimate of its size from the formula:

where

d _ m K 2

37ray 2

, = 33.6 t o n / i n ? , ignoring the variation of yield stress with temperature,

K = 34 ksi~/in., corresponding to fast fracture initiation at - 135~

we find d = 0 0 2 2 in., and

K = 228 ksi~/in., corresponding to fast fracture initiation at - 7 ~

we find d = 0 9 8 in

This latter figure is probably so large as to invalidate the simple formula

just used However, there can be no doubt that yielding extended well into

the specimen and well ahead of the region of slow crack growth, before

rapid fracture took place

Plastic zone considerations are important because in many materials of

this type, acoustic emission is associated with yielding and the growing

edge of the plastic zone is considered to be the most likely source of

emissions This is the basis of a well-known model which gives a power-law

relationship between cumulative emission count and stress intensity

Ductile tearing is a relatively quiet process, in low-strength steels at least

[6] Cleavage, on the other hand, is a credible source of relatively high-

amplitude emissions

The wide variations in supposed plastic zone size are broadly consistent

Trang 35

with the wide variations in cumulative emission count prior to fast frac-

ture Comparing results of tests run at different temperatures, the cumu-

lative emission count to failure varies approximately as the square of K~,,

Also, the curve shown in Figure 6 approximates closely to a parabola

Turning to the crack tip itself, we have an interplay between two con-

flicting mechanisms: cleavage and ductile fracture A simple model is

proposed in the following as a description of the interaction on micro-

scopic and submicroscopic levels

(a) Ductile rupture is a relatively slow process, taking time and involv-

ing substantial flow of the material on the fracture surface This process is

strongly dependent on temperature, probably because some of the micro-

scopic flow mechanisms are thermally activated

(b) Cleavage is a relatively rapid process For activation it requires a

stress concentration that is very high on the microscopic level; this stress

concentration will be relieved by local ductile flow Cleavage produces

acoustic emissions of relatively large amplitude There is relatively little

flow of material on the fracture surface

The different microscopic characteristics of the two mechanisms, as just

described, give an understanding of the macroscopic effects that were

measured in the course of these experiments As the temperature rises,

cleavage is increasingly inhibited by plastic flow which relieves stresses in

the microscopic regions of highest stress concentration As the temperature

rises, the macroscopic stress intensity can rise to much higher levels

before cleavage finally supervenes Cleavage events at the crack tip have a

tendency to lead on to fast fracture, since relatively little energy is

required to sustain crack propagation in the cleavage mode

At a temperature of - 1 0 ~ some local cleavage is taking place before

fast fracture; the material is ductile enough to arrest these incipient fail-

ures locally, until the stress intensity becomes too high This local cleavage

leads to a change in emission amplitude distribution to 85 to 90 per-

cent of the failure load

At low temperatures, local cleavage tends to lead straight on to brittle

fracture since there is little chance that an incipient failure will be arrested

by local plasticity Thus, the amplitude distribution does not change before

failure, but remains characteristic of the activity at the edge of the grow-

ing plastic zone that is the dominant source of emissions

Fast fracture is predominantly by cleavage even when the preceding

slow growth zone shows ductile tearing On the basis of our model, this

happens because ductile tearing takes time, and time is no longer available

once cleavage initiates at a high enough stress intensity In the first stage

of fast crack growth, stored elastic energy is made available faster than

Trang 36

it is absorbed[7] Ample energy will be available to sustain cleavage, the

crack will accelerate and ductile flow will become less and less likely

Nevertheless, some traces of ductility were observed even in the fast

crack region at warm temperatures This residual ductility is of some

importance An area of slow ductile tearing may absorb nearly 20 times

more energy than an equal area of fast fracture[7] Ductile tearing over

small regions of the fast fracture surface may make a disproportionately

large contribution to energy absorption This could be one of the causes

of the large variation in dynamic surface energy with temperature (22,

52, and 117 • 105 ergs/cm 2 at - 197, - 135, and - 60~ respectively,

in U X W steel[2]) Our model a l s o predicts that the amount of ductile

tearing would be greater at the beginning and end of the crack jumps,

where the crack was running more slowly; we do not have enough infor-

mation at present to say whether this effect would be readily observable

The model is consistent with two other well-known effects in the frac-

ture behavior of steel The lowering of critical stress intensity with increas-

ing specimen thickness is generally recognized to be a result of material

constraints; increasing the thickness reduces plastic flow and thereby in-

creases the likelihood of cleavage Second, increases in strain rate lead to

reductions in K~c[2]; on our model, this would happen because plastic

flow takes time At high strain rates, flow is delayed and the likelihood of

cleavage is higher The persistence of emission for several minutes during

hold periods is an indication that this is a real effect at the strain rates

used for these tests

Conclusions

In summary, a large number of experimental factors can be brought

into line on the basis of a very limited set of suppositions, namely

(a) Ductile rupture takes time and the process is strongly dependent

on temperature

(b) Cleavage is rapid on the microscopic scale, and is averted by

ductile flow and rupture at the points of highest local stress

Hopefully, in this model we have isolated some of the most salient

factors involved in the ductile-brittle transition The model lends itself

naturally to making predictions and affords an outline picture whose

details are yet to be filled in

Acknowledgments

Fractography for this project was conducted under the direction of

C H Jones All DCB tests were performed by F A Johnson The work

was carried out under contract to the Ministry of Defence ( N a v y )

C o p y r i g h t b y A S T M I n t ' l ( a l l r i g h t s r e s e r v e d ) ; M o n D e c 7 1 3 : 1 0 : 5 5 E S T 2 0 1 5

Trang 37

3 0 FRACTURE TOUGHNESS AND SLOW-STABLE CRACKING

References

[1] Turner, C E and Radon, J C., "Fracture Toughness Measurements in Low

Strength Structural Steels," Paper 14, Proceedings, 2rid International Conference

on Fracture, Brighton, April 1969

[2] Pollock, A A and Radon, J C., "Acoustic Emission in the Fracture Toughness

Test of a Mild Steel," Document IIW-X-595-70, Lausanne, Switz., 1970

[3] Nakamura, Y., Veach, C L., and McCauley, B O in Acoustic Emission, A S T M

STP 505, American Society for Testing and Materials, 1972, pp 164-186

[4] Pollock, A A., Nondestructive Testing, Vol 6, No 5, Oct 1973, pp 264-269

[5] Harris, D O., Dunegan, I4 L., and Tetelman, A S., "Prediction of Fatigue

Lifetime by Combined Fracture Mechanics and Acoustic Emission Techniques,"

Technical Bulletin DRC-105, Dunegan/Endevco Corp., 1970

[6] Pollock, A A and Smith, B., Nondestructive Testing, Vol 5, No 6, Dec 1972

[7] Radon, J C and Pollock, A A., Engineering Fracture Mechanics, Vol 4, 1972,

Trang 38

H H C h a s k e l i s , 1 W H Cullen, 1 a n d J M K r a f f t 1

Acoustic Emission from 4340 Steel

During Stress Corrosion Cracking

REFERENCE: Chaskelis, H H., Cullen, W H., and Krafft, J M "Acoustic

Emission from 4340 Steel During Stress Corrosion Cracking," Fracture

Toughness and Slow-Stable Cracking, A S T M STP 559, American Society for Testing and Materials, 1974, pp 31-44

ABSTRACT: A commercially available system is employed to detect and count acoustic emissions emanating from the aqueous stress corrosion crack propagation in 4340 steel Standard ASTM E 399 compact tension specimens with short notches, a/W=0.25 versus 0.50 standard, are prepared in four tempering temperatures: 204, 316, 427, and 538~ (400, 600, 800, and 1000~ Crack length is monitored with a notch opening clip gage Com- parisons show the time rate of emission events to increase with stress intensity in rough correspondence to, but much more rapidly than, the crack velocity Tempering back tends to suppress the count rate as a warning of fast fracture instability Some similarity is observed between the areal rate

of acoustic emission and factors which influence degree and intensity of plastic flow instability in the fracture process zone

KEY WORDS: acoustic properties, crack propagation, fracture properties, fatigue (materials), acoustic detection

That fracturing is a source of sound should have been common knowl- edge in any age What is new, in view of Kaiser[I] 2 and others[2], is that even extremely slow, ostensibly stable fracturing is not silent Hart-

b o w e r [ 3 ] h a s o b s e r v e d t h a t the p r o g r e s s of a s t a b l e s u b c r i t i c a l c r a c k is

s i g n a l e d b y a s t e a d y s u c c e s s i o n of d i s c r e t e e m i s s i o n events, o b s e r v a b l e b y

" l i s t e n i n g " t h r o u g h sensitive p i e z o e l e c t r i c t r a n s d u c e r s a n d amplifiers

P r i o r to his o b s e r v a t i o n , t h e r e h a d b e e n o u r o w n [ 4 ] t h a t the event of

1 Physical science technician, metallurgist, and head, Ocean Materials Criteria Branch, respectively, Ocean Technology Division, Naval Research Laboratory, Wash- ington, D C 20375

The italic numbers in brackets refer to the list of references appended to this paper

31

Trang 39

criticality Kie eould often be associated with an audible "pop in." Brown

and Srawley's application of a phonograph type transducer in such tests

[5] did indeed show that the sound was emitted not as a single but as a

succession of events as the point of criticality K~e was reached and sur-

passed in the path toward general fracture instability K~

An annunciator for K~, while useful in testing procedures, is rather

too late in the life of a structure to serve as a warning of destruction

On the other hand, a warning of the early stages of subcritical cracking,

as proposed by Hartbower, could be of great value in proof testing, even

perhaps for in-service surveillance of fracture-sensitive structures[6] It

would be helpful in such endeavors to understand the sources of the

sound in subcritical propagation; ideally a correct predictive model of the

sound producing event But falling short of this, even a discriminating

characterization of the emissions to be expected in various materials and

corroding environments should be helpful This paper reports an attempt

to supply such information for 4340 steel, quenched, then tempered back

to levels of hardness typical of its utilization in structural and machine

components

Experimental Procedures

Briefly, the cracks, serving as sources of acoustic emission, are grown in

short notch compact-tension specimens (ASTM Test for Plane-Strain

Fracture Toughness of Metallic Materials (E 3 9 9 - 7 2 ) ) , immersed in

distilled water, and subjected to constant load The opening of the notch

was used as a measure of compliance, thus providing a measure of crack

growth The accumulation of emissions is recorded, along with the load,

versus the notch opening Tensile stress-strain curves, needed to assign a

degree of fracturing instability within the crack tip region, were measured

on small tensile specimens The crack growth data, compared here to the

stress wave emission, has been reported elsewhere[7], and some of the

procedural description is repeated here

The 4340 specimens were cut from a 13-mm-thick plate, Interlake Co

heat ~ B 0 - 5 3 5 - 5 0 0 Compact tension specimens were cut, in full plate

thickness, so as to be subject to crack propagation in the long transverse

direction of the plate, TL Of the 2 in of width W between the back of

the specimen and the load line of the holes, about 1 in is useful for

crack propagation a measurements, 0.25<_a/W~_0.75 Four batches of

about 20 each were austenitized at 843~ (1550~ oil quenched, then

tempered for 1 h at temperatures of 204~ ( 4 0 0 ~ 316~ ( 6 0 0 ~

427~ ( 8 0 0 ~ and 538~ (1000~ followed by air cooling to room

temperature Small 4.3 by 12-mm tension-compression (T-C) specimens

were cut from the same material, axes normal to the crack plane (T-direc-

Trang 40

CHASKELIS ET AL ON ACOUSTIC EMISSION FROM 4340 STEEL 33 tion) and heat-treated along with the fracture specimens All fracture

specimens were prepared with low-stress fatigue cracks

The compact tension specimens were tensioned in an N R L dynamic

loader[8] (Fig 1) set up to maintain constant load through a selective

gas pressurization Some tests were run in stages in which the head was

locked, allowing a "load shedding" which held KI to a rate of increase

about one third that for the constant load condition The crack growth was

detected with a displacement clip gage inserted in the notch, with out-

put, time modulated, plotted versus the load on a Hewlett Packard Model 2

F R A X-Y1-Yz recorder

A standard P Z T transducer was positioned atop one arm of the speci-

ment (Fig 1) using an interface of viscous petroleum grease and a dead

weight of some 89 to 1 kg The counting instrument, Dunegan 301 total-

izer, was set to provide 94 dB gain, 40 dB from the 801 P preamplifier and

54 dB from 301 totalizer to the received electrical signal, then cutting off

all signals below 1.0 V Of the several filters available in this instrument,

that passing the Band 0.3 to 1.0 M H z was employed The natural fre-

quency of the P Z T transducer was about 0.16 MHz, with a second peak in

the 0.3 M H z range, 10 dB down from the primary peak This secondary

peak falls at the lower end of the band pass window which was used, a

FIG 1 Specimen and grips with load cells, A E transducer and notch opening

clip gage Immersion of the back half of the specimen in water reduces by one fourth

the A E count rate

Tiêu đề	Fracture Toughness And Slow-Stable Cracking
Tác giả	P. C. Paris
Người hướng dẫn	G. R. Irwin, General Chairman
Trường học	University of Maryland
Thể loại	Báo cáo chuyên đề
Năm xuất bản	1974
Thành phố	Baltimore

Định dạng
Số trang	319
Dung lượng	6,41 MB