The effect of carbides on the high strain fatigue resistance of an austenitic steel

THE

EFFECT OF CARBIDES ON THE HIGI-I STRAIN RESISTANCE OF AN AUSTE~ITIC STEEL* J. T.

and

BARNBY?

F. M.

FATIGUE

PEACEt

High strain fatigue experiments have been conducted at strain amplitudes from f 1 per cent plastic strain to &4 per cent plastic strain on an A.I.S.I. 316 austenitic stainless steel. The effect of the presence of 2-3 oA by volume of a chromium carbide is to shorten high strain fatigue lives down to around ) of that of the same steel in a solution treated condition. Metallographic evidence shows that internal fracture of carbides precedes the formation and propagation of fatigue cracks. Short life tests correspond to rapid fatigue crack propagation by ductile tearing between the main crack front and voids produced by carbide fracture. INFLUENCE

DES CARBURES SUR PAR DEFORMATION

LA RESISTANCE A UNE FATIGUE DANS UN ACTER AUSTENITIQUE

~~PORTANT~

Deux experiences de fatigue importante par deformation ont Bte effect&es SUP un acier inoxydablr austenitiqne AISI 316, avec des amplitudes de deformation allant d’une contrainte plastique de f 1% it une contrainte plastique de &4%. La presence de 2-3 ‘A en volume d’un carbure de chrome diminue la duree de vie sous fatigue Blevee # de jusqu’it environ sa valeur relative au meme acier dans des conmontre de fac;on Bvidente ditions correspondant a une solution solid%. L’etude metallographique que la rupture interne du oarbure precede la formation et la propagation des fissures de fatigue. Les experiences dormant une vie courts correspondent B une propagation rapide des fissures de fatigue par dechirure ductile entre le front de lafissure principale et les trous produits par la rupture du carbure. DER EINFLUB VON KARBIDEN IN EINEM WIDERSTAND GEGEN ERMUDUNG BE1

AUSTENITISCHEN STAHL AUF HOHEN DEHNUNGSAMPLITUDEN

DEN

An A.LS.I.-316-Edelstahl wurden Ermiidungsexperimsnte bei hohen Dehnungsamplituden von *I % plastische Dehnung und +4 % plastische Dehmmg durohgefuhrt. Die Gegenwart von 2-3 Vol. % Chromkarbid verkiirzt die Lebensdauer bei hohen Dehnungsamplituden auf l/3 der Lebensdauer desMetallographisohe Untersuchungen zeigen, da13 der innere selben Stahls nach Liisungsbehandlung. Bruch der Karbide der Bildung und Ausbreitung von Ermtidungsrissen vorausgeht. Kurze Lebensdauern entspreehen der schnellen Ausbreitung von Ermtidungsrissen durch duktile Schenmg zwischen der Hauptfront des Erm~dungsrisses und den Hohlraumen, die durch den Bruch der Karbida emstehen.

INTRODUCTION

The presence of second phase particles is known to affect the fatigue resistance of metallic materials.“) It seems possible that the mechanics of internal fracture of second phases are similar to those operative in a tensile failure,t2*3) though it is surprising that these mechanisms can operate at such low stress levels in the fatigue situation. Evidence of large Bauschinger effects in two phase materiaM4) shows that slip bands are severely blocked by the presence of high strength brittle particles, and the question naturally arises as to whether the large stresses borne by these particles can rise in a cumulative manner, during fatigue cycling, to a level suiKcient to break open the particles. In order to test whether this is so it seemed appropriate to conduct experiments where the overall plastic strains and stresses during a fatigue cycle were less than t#hosenecessary to initiate particle cracking in monotonic tensile loading. Internal damage to particles may be viewed as the initiation of the large scale failure mecha~sms of ductile fracture or fatigue fracture, and approaches at t,his size scale are productive in forwarding an * Received October 20, 1970;

t Department of Metallurgy, ingham 4, England. ACTA

METALT~URGICA,

VOL.

revised April 10, 1971. University of Aston, Birm19, DECEMBER

1971

understanding of the relative resistance of materials to fracture failure mechanisms.(4-6) In the work described the effect of the presence of carbides on resistance to high strain fatigue is illustrated by comparison between results on a solution treated A.I.S.I. 316 sta,inless steel and the same steel heat treated to contain Z-3% by volume of the carbide (CrFe),,C,. This dispersion of carbides does not affect the yield stress of the steel but does enhance the workhardening rate over the initial 10 per cent of plastic extension. This work forms part of a general programme on the effects of two phase structure on internal damage and fracture failure mechanics. MATERIALS

AND

EXPERIMENTAL

TECHNIQUES

The A.I.S.I. 316 type stainless steel used corresponded to the analysis of that in Ref. 3. and surface electropolishing techniques were also as described in Ref. 3. As in Ref. 3, careful checks were carried out to make sure that the fracture of carbides on specimen surfaces corresponded with similar carbide fracture in the interior of the specimens. This was always found to be so. Specimens were tested in a screw driven machine programmed to cycle at a few- minutes per cycle about

1351

ACT-4

1rid:!

ME:TALLURGICA,

Extensometer grooves

VOL.

19,

1971

Cyclic stress strain loops were recorded during test-

/\

ing in order to follow the Bauschinger seemed

possible

gressive deterioration

in the ability

raphic slip planes. plastic

The Bauschinger

in a reversed

culties

in

defining

measured.

for the onset

direction

plastic flow in one direction. precisely

It is not clearly

subsequent

There are many how

this

a zero mean stress giving &l, plastic

strain.

Strains

t)randucers mbunted grooves

f2,

13 or -&4 per cent

were monitored

by electrical

ou extensometers

on the specimen

mens were located,

illustrated

and locked

locating

test

complete

cyclic

into,

stress

already

into

tapered

strain

dif% be

significant, to measure, at the

strain (reversed) t,o the yield strain in the

forward direction. has

In most cases reversed plastic flow

commenced Load IO’ lb

in Fig. 1. Speci-

which allowed no backlash in the mountings. each

equivalent

of to

should

for instance, the stress in the reverse direction Fm. 1. Standard Mend-type high strain fatigue specimen.

phase

on crystallog-

effect is basically

of the stress necessary

flow

show a pro-

of second

particles to obstruct plastic deformation

a lowering

effect since it

that this effect would

grips

at, t,his st,rain.

WiIson’s

t

During

loops

were

recorded, from the transducer and load cell outputs, on an X/Y

recorder.

The stress and strain 0utput.s were

also fed to a data logging could

be used

calculates

with

device

a computer

the work hardening

so that the data programme

exponent’,

which

n, for each

half cycle of deformation. Tests

were periodically

interrupted

and formvar

replicas taken from the elect,ro-polished

gauge length

of a specimen.

Second st8age carbon replicas were then

produced

viewing

Fracture

for

in the

surfaces, gaugelength

sections were directly

electron

microscope.

Solutmtreokd Compteswx Load

surfaces and polished

examined

moter~ot

by optical and scan-

ning electron microscopes.

LW3d IO3 lb

t

Tension

The cyclic stress strain loops of Fig. 2 illustrate the work hardening behaviour precipitate

containing

is summarised

of the solution treated and

steels.

Data on this behaviour

in Fig. 3 by plotting

stress against cumulative

the maximum

plastic strain.

The mono-

tonic stress strain curves for the two heat treatments are superimposed

for

comparison

of

stress

levels.

~~omparison of Figs. 3(a) and (b) shows that the precipitate containing steel undergoes more rapid cyclic hardening over the initiai 5-20 per cent cumulative plastic strain, but by 50 per cent the two materials reach equivalent stress levels which depend only on the strain amplitude.

These saturation

levels of true

stress are always iess than that produced by monotonic straining to 50 per cent plastic strain.

Compression

Load I

(b) FIG. 2. The initial stages of cyclic strain hardening. Cycles 1-5: (a) solution treated material; (b) precipitated material.

BARNRY

ANI) PE:ACE:

HIGH

STRSIS

c

‘20,

140,

I Precrpitated

material

120 -

P.ATIGUE

OF

AUBTENITIC

1363

RTEET,

Using this method an n value wits ca.lculated for each half cycle of the test and plotted against the number of strain reversals, 2N, where N is the cycle number. Whereas this does not directly measure a Bausohinger effect it is considered that the change in ‘1~with reversal number should indicate any large changes in the size of the Bauschinger effect, and should also indicate any large change arising from progressive breakdown of the ability of second phases to hinder plastic flow. The results of these tests are shown in Figs. 5(a)-(d) where comparisons are made between the precipitated and solution t,reated materials. Clearly the large change in n arises from the first stress reversal. Subsequently t.he n values are relat,ively constant to the end of the test, which is a surprising result,taken in relation to the metallogrSraphicobservations of partiole damage described below. Fracture of carbides

!

2’3 !

I

The initial electro-polished surfaces of specimens showed no fractures in the lightly et~ched carbides, but fractured carbides were observed in the early stages of life at all four plastic strain amplitudes, A general tendency wits that the smaller, rod-shaped carbides fractured first. Broken carbides are illustrated in Figs. 6(a) and (b). At the smallest strain amplitude slip bands were of small step height in

FIG. 3. True stress vs. cumulative strain (irrespective of sign): (a) solution treated material; (b) precipitated material.

elegant experiments measured the ratio of stress in the reverse direction to that for continued plastic flow in the forward direction for a, range of strains. Thus the smaller is the ratio, the greater the B&uschinger effect at that strain. The absence of Bauschinger effect would result in a stress ratio of unity. This type of result is shown in Fig. 4. In order to follow this behaviour through many cycles it is necessary to test one specimen for the continued forward straining curve on cycle i in order to compare it with the reversed stresses on a specimen which continues to cycle on the cycle i. This is an expensive and time consuming method and so an alternative measurement was made of the work hardening exponent n using the equation : crT = K&r/’ Here 5~ is the true stress, K a constant and +, the true strain.

”

FIG. 4. The

uvz Reversedstrain Bauschinger &feet material.

for

0.03

the

004

precipitated

ACTA

5

IO

15 Reversal

20

25

number, (a,)

METALLURGICA.

30

35

VOL.

19.

1971

40

2N

Solution treated 0’ 0

5

IO

15 Reversal

20

25

number,

30

35

40

2 N

(b)

Solution

5

IO

15 Reversal

20

25

number,

treated

30

35

40

2N

I

5

IO

15 Reversal

20 number,

25

30

35

40

2N

(d) Fm. 5. Work hardening exponent vs. reversal number for t,he solution treat,ed and precipitated material (a) +I per cent plast,ic strain; (b) &2 per cvrlt: plastic strain; (c) +3 per cent plastic strain; (d) -&a per

FIG. 6. (a) Precipitated material. Carbon replica. Diffuse slip hands. 550 cycles &I per cent plastic stxain. i( 6000. (b) Precipitat,ed matarial. Gaugelength surface. Fractured grain boundary carbide. 14 per per cent plasbic strain. \’ 1.760.

W.4RNRY

Fractured carbides

not observed

ANI)

PEACE:

HIGH

STRAIN

FATIGUE

OF

AUSTENITI(’

STEEL

1355

’ \ \ \ \ 20

60

40 Cycle number, N

J?“ra.7. The rclatkmship be+wccn maximum true stress in t,ension itnd oarbide frwture for the precipitdetl mat~rrial.

comparison

with

amplitude.

At a plastic

dominant,

slip bands

process

was

replicas

taken

at the

highest

strain of 14 carbide

strain

per cent the

fracture

at

grain

boundaries. Surface

at

stages

of

the

cyclic

straining allowed an estimate of the onset of particle fracture. the

Figure 7 shows plotted lines w-h&h bracket,

onset

material.

of

particle

damage

in the

~recipitatecl

This figure also shows the cyclic hardening

curves for each plastic strain range and so indicates the applied

stresrJ levels at which

carbide

fractures

were observed. The stress level for onset of carbide fracture in a t,ensile t,esG3) IS . included for comparis(?n. Fra. 8(h).

C~rackiwitiath

and p?Jropagation in the

transgranular crack propagation, at &I per cent, plastic strain, to intergranular crack propagation at

solution-treated material was transgranular at all strain amplitudes. I~tiation occurred from surface rumples and propagation was by stage II type ripples.

&4 per cent plastic strain as shown in F’igs. S(a)-(c). Crack surfaces showed ripple formation at &l per cent plastic strain, Fig. 9 and at *t2 per cent plastic

Ripples

strain a kind of ripple formation

outlined

was

of

1. Xolution-treated material.

Crack initiation

were finer at low strain amplitudes.

fracture occurred

Final

by ductile dimple formation.

2. Precipitated material. In the precipitated material there was a progressive change from predominantly

observed

on

some

areas

grain

fracture, Fig. 10. No striation formation

by cavities boundary

was observed

during crack propagat,ion at, &4 per cent plastic strain,

ACTh

1 356

METALLURGIC.%,

FIG. 8(c). FIG. 8. Precipitated material. Gaugelengt,h surface: * 1 per cent. (a) tr~n~gran~ll~r crack propagation. plastic strain. ~492; (b) mixed transgranular and intergranular crack propagation. *2 per cent plastic .&rain. X ‘WO. &4 > (c) intergranular crack propagation. x 260. per cent pla&ic strain. rather

the process was one of void linkage

as in ductile

fracture. Summarising material,

crack propagation

in the precipitated

it appeared that at low strain amplitudes

number of cracked carbides

was relatively

the

small at the

VOL.

stage

19,

of macroscopic

~agat.ion

occurred

poration

of cracked

As

crack

damage ductile In

increased, the

initiation.

ripple

inbo the

proceeded

producing

at macroscopic

crack

cycled

t,o &4 per cent

plastic

immediately

dimple

forInat~ion without

trates

the

onset

of

front. particle

t)ransition

to

than ripple production.

of damaged

greater

then

proincor-

crack

general

a gradual

rather

number

Crack

f(~rInat,ioll and

carbides

dimple tearing

appearance

&age II rippIes on fracx 1400.

crack

by

propagation

co&fast

gation

FJG. 9. Precipitated material. ture crurface.

197 1

particles

initiation strain.

Crack

commenced ripples.

carbide

was

for material by

Figure

fractures

propaductile 11 illus-

and

the

of surface cracking.

Frc. Il. Total plastic strain range vs. cycle number showing init,iation of carbide fracture and surface cracks, and final fracture for the precipit.nt.ed material.

BARNBY

PEACE:

AND

HIGH

STRAIN

FATIGUE

OF AUSTENITIC

solution-treated

STEEL

and aged conditions.

ance of this enhanced direct correlation

hardening

The disappear-

in Fig. 3 shows no

with the onset of carbide fracture

shown in Fig. 7. Figure 4 shows the persence of a large Rauschinger Solutwn

treated malerlol

effect on the first reversed steel as would be expected This corresponds work hardening

cycle in the precipitated from the work of Wilson.c4)

to a very considerable

drop in the

in the second half cycle in Fig. 5(a),

but the work hardening

exponent

n remains roughly

constant from then onwards to the end of the test in all cases.

It is noteworthy

in n is much

that the sudden initial drop

Iarger for the precipitated

though this drop does not correspond carbide

fracture

evidence

as bracketed

in Fig. 7.

general build-up

by the metallographic

It seems that the Bauschinger

effect and work-hardening c-3

material,

to the onset of

characteristics

of internal

relate to the

stresses and dislocation

structure and are not sensitive to the isolated regions FIG. 12. Total plastic strain range vs. life (cycles to failure) for the solution treated and precipitated material.

in which the local stresses are relaxed by precipitate fracture. The metellographic mechanism second

Lives of precipitated and solution-treated materials Figwe

12 shows the life data for solution

and precipitated Coffin plot.

material

in terms of the Manson-

Data for the solution

falls very close to that of published represented

treated

treated

material

work(7*sJ and is

by :

Here N, is the number

of cycles to failure and A.Q~

the total plastic strain range. particle

rupture

The effect of internal

in the precipitated

material is to shorten lives by at least a factor of 4. This changes the constant in the Manson-Coffin

phase particles

law so t,hat this data is a

good fit to: A7FJ5 Ae,r = 0.39

shows the change in

fracture

are present.

when damaged There is a clear

transition from plastic gr0wt.h of the crack by dimple formation

to ductile

t*earing which links more preIt is

existing voids with the crack front on each cycle. obviously

important

to pursue the changes in life, or

fatigue crack propagation size and

No.46 AEPF = 0.65 f

evidence

of the fatigue

dispersion

of

r&es, with changes in the hard,

brit,tle second

phase

particles. A key feature which

of these results is shown in Fig. 7

demonstrates

a remarkable

applied stress level at which mences. The procedure of

variation

particle

in the

fracture

bracketing

the

oomcycle

number at which particle fraet,ure commences is relatively inaccurate, but the informat.ion given on the stress level for onset of particle fracture is nevertheless

DISCUSSION

accurate

strain amplitude

and significant.

Thus

of 1 per cent particles

at a plastic are observed

The life data of Fig. 12 shows clearly the significant reduction in life of precipitate-contai~ng steel at all

t,o break at less than 60 k,s.i.,

the plastic strain amplitudes

contrast,, particles

do not break until a stress of 90

solely from the presence of the hard brittle carbides,

k.s.i.

strain

which

There is little doubt

undergo

internal

tested.

fracture,

This must arise since

the

yield

properties of solution-treated and precipitated steels are identical. The only difference in their stress strain curves is the enhanced work-hardening in the precipitated steel, and this must arise directly from the presence of the precipitates. The cyclic work-hardening

of Fig. 2 gives rise to an

enhancement of hardening at low cumulative strains in the cyclic hardening curves of Fig. 3 for both t*he

tSest, no

particle

at a plastic

governed

whereas in a tensile

fraet,nres occur

until

amplitude

that particle

by the achievement

75 k.s.i.

In

of 4 per cent. fracture

of a critical

local stress which must be characteristic material.

must be internal

of the particle

Imernal stress levels may be differently related t,o to the applied stresses because of a change in the effective shnrpness(lO~ll) of a slip band with strain amplitude, or because of the increased level of cyclic work hardening with increasing strain amplitude.

ACTA

1358

METALLURGICA,

CONCLUSIONS

The major conclusion

VOL.

19, 1971

and for encouragement

drawn from this work is that

second phase particles, in the form of around 2-3%

by

in this work.

sponsored by a research Research Council.

grant

from

The work is the

Science

volume of hard, brittle carbides bonded to the matrix, very significantly

shorten the life of material under In the case investigated

REFERENCES

high strain fatigue conditions.

the carbides were only around 1~ in thickness. further concluded significantly

that particle fracture

can occur at

lower applied stresses, under high strain

fatigue conditions, It is intended

than in tensile straining.

to pursue this work by investigating

the effects of carbide propagation

It is

distributions

rates using fracture

on fatigue mechanics

crack testing

techniques. ACKNOWLEDGEMENTS

The authors Alexander

would like to thank Professor

for the provision

of laboratory

W. 0. facilities

R. M. N. PELLOUX, T~ans. Am. Sot. Metals 57,511 (1964). J. T. BARNBY, Quantitatizx Relation Between Propertiecr and Microstructure, edited by D. G. BRANDON and A. RESEN, p. 381. Israel Universities Press (1969). J. T. BARNBY, Acta Met. 15, 903 (1967). D. V. WILSON, Acta Met. 13, 807 (1965). J. M. KRAFI+Y,App.?. Mater. Res. 3, 88 (1964). A. J. BIRKLE, R. P. WEI snd G. E. PELLISSIER, Trans. Am. Sot. Metals 59,981 (1966). S. S. MANSON, N. A. C. A., TN2933 (1954). L. F. COFFIN and J. F. TAVERNELLI, Trans. Am. Inst. Min. Engrs. 215, 794 (1959). J. T. BARNBY and M. R. JOHNSON, Metal Sci. J. 3, 155. E. SMITH, Acta Met. 16, 313 (1968). P. M. HAZZLEDINE and P. B. HIRSCH, Phil. Mag. 15, 121 (1967).

h: 3. 4. 5. 6.

8. 9. 10. 11.