The Role of Hydrogen in the DuctileFracture of Plain Carbon Steels

H. C I A L O N E A N D R. J. A S A R O

In this report we consider the problem of hydrogen induced ductility losses in a plaincarbon spheroidized steel. Specifically, the effect of internal hydrogen on the formationof voids from second phase (cementite) particles and t h e i r subsequent growth and coales-cence was studied by careful microscopic inspection of uniaxially strained b a r s , both initi-ally cylindrical and circumferentially notched, with and without hydrogen. Void initiationo c c u r r e d with lower strains and s t r e s s e s with hydrogen, although an equally importantcontribution t o the ductility loss was from hydrogen accelerated void growth and coalescence.This latter process takes p l a c e by the propagation of voids a l o n g the g ra in , and possibly sub-g ra in , boundaries which interlink the cementite spheroids. The results indicate that hydrogenfacilitates interface separation, possibly by accumulating at the boundaries during hydrogena-tion of the specimen and lowering the cohesive strength, thereby m a k i n g void initiation andgrowth along them e a s i e r .

1. I N T R O D U C T I O N

O F the many ways in which hydrogen can degrade themechanical properties of materials two, in particular,seem most prevalent in steels: hydrogen e i t h e r causesa change in the mode of fracture, usually from a " n o r -mal" ductile process involving void initiation andgrowth to less ductile transgranular or intergranularcleavage, or hydrogen enhances, without changing thecharacter of the " n o r m a l " mode of failure. 1'2 Thef i r s t instance is common in high strength steels whereductile fractures are often found t o give way to cleav-age presumably brought on by a relaxation of loca lfracture s t r e s s (or strains) due t o hydrogen segrega-tion and accumulation at internal interfaces.3 In highstrength alloys that would normally display cleavagefractures, possibly due to the p r i o r embrittling ef-fects of impurities, for example, hydrogen can havethe second effect and enhance a n o r m a l intergranularor transgranular cleavage by lowering even further thestresses required locally to cause failure. 4 On theother hand, for low strength steels such as the plaincarbon spheroidized steels discussed herein, conditionsleading t o a change from what is normally a ductilefailure mode t o brittle cleavage with hydrogen are notreadily attained. Yet when these steels contain r e l a -tively large amounts of hydrogen, introduced for in-stance by electrolytic charging, they are subject t osignificant losses in ductility. This is aga in a casewhere hydrogen seems to accelerate, without changing,the n o r m a l mode of fracture, which in this case is duc-t i l e .

This paper contains a report on some experimentsconducted on hydrogen charged spheroidized plain car-bon steels, the purpose of which was to m e a s u r e theinfluence of hydrogen on a process of ductile fracture.The ductile fracture process we studied was quite s i m i -lar t o others discussed in the literature, but we willindicate how certain microscopic events (viz interracialseparations d u r i n g void growth) which lead t o an ac-

H. C1ALONE and R. J: ASARO are GraduateStudent and AssistantProfessor of Engineering, respectively, Division of Engineering,BrownUniversity, Providence, RI 02912.

Manuscript submittedMay 1, 1978.

METALLURGICAL TRANSACTIONS A

celerated void development have perhaps been giventoo little attention in r e c e n t work. In fact, the resultsof the present study suggest an explanation for thecommonly observed ductility losses in hydrogen con-taining environments that is more in the spirit ofmodels which have until now been exclusively associa-ted with brittle intergranular fracture of high-strengths t e e l s - t h a t is the accumulation of hydrogen at inter-f aces with a resulting loss in the resistance t o inter-f ac i a l separation,s

Our experiments were conducted on spheroidizedplain carbon steels containing between 0.30 and 1.05wt pct carbon. Cylindrical specimens were tested intension w h e r e they were observed to undergo uniformdeformation followed by necking; ductile fracture tookp l a c e within the neck of the specimen by a process ofvoid initiation, void growth and coalescence. The onesignificant macroscopic effect of subjecting theses t e e l s to hydrogen environments was a loss in the finalreduction in area and in this sense t h e r e was a loss inductility. The microscopic events leading up to com-plete fracture w e r e , both in kind and sequence, un-changed. T o understand the role played by hydrogen itis f i r s t necessary t o develop a certain perspective onthe detailed process of ductile fracture in t h e s e steels.This we do briefly in the next section. Some relevantexperimental details follow in Section 3 with discussionin Section 4.

2. DUCTILE TENSILE FRACTURE IN PLAINCARBON STEELS

It is generally accepted that ductile fracture in plaincarbon spheroidized steels b e g i n s with the formationof voids at second phase particles, notably carbides, sVoids initiate by particle cracking, separation of con-terminous particles, or by particle-matrix separation,although in our experiments we observed essentiallyonly the latter. Voids grow by plastic deformation,which like void initiation, is greatly augmented withinthe t r l a x i a l s t r e s s field imposed by the neck. When asignificant void volume fraction has developed, voidcoalescence by direct impingement is frequent andmay be further accelerated by localized deformation

ISSN 0360-2133 /79/ 0312-0367500.75 /0© 1979 AMERICAN SOCIETY FOR METALS AND VOLUME 10A, MARCH 1979-367

THE METALLURGICAL SOCIETY OF AIME

Fig. 1--Example of a centralcrack propagating towardsthe lateral surface of a hy-drogen charged specimen.The outlined area on the leftis magnified on the righthand side of the photograph.The micrograph shows howcrack formation and propa-gation is clearly influencedby existing voids. The tensileaxis orientation is vertical.

over extended regions in the m a t r i x6 or by a processof localized necking in the ligaments lying betweennearly contiguous holes. Another mechanism that wefound to be quite important in these steels and thatwill be emphasized in the discussion of our resultswas the separation of g ra in boundaries lying betweenparticles that have initiated voids. Void coalescenceresults in a large central c r a c k7 within what, at thiss t age in the deformation, is usually a r a t h e r deep local-ized n e c k . An example of a central c r a c k which hasjust begun to propagate and coalesce with existingvoids is shown in Fig. 1. The c r a c k then propagatestoward the specimen surface resulting in f i n a l sepa-ration. C r a c k propagation is also ductile and often in-volves shearing off a l o n g bands of concentrated strain.This, then , is the genera l process that is affected bythe presence of hydrogen.

Although all of the individual s teps just describedcould be affected by hydrogen, the prevailing viewsare that it is either the initial process of particle-m a t r i x separation, or particle-particle separation,or void growth that is facilitated. ~,~ Our procedurewas t o make a detailed comparative study of the f r a c -ture in steels with and without hydrogen.

3. EXPERIMENTAL

Experiments were c a r r i e d out on s e v e r a l plain car-bon steels containing 0.30, 0.45 and 1.05 wt pct carbon,although in this paper we discuss only our results fora 1045 s t e e l (see Table I). The material, obtained as0.75 in. (1.91 cm) diam rods, was austenitized at825°C for 2 h, quenched in a 25 pct solution of water-soluble oil, tempered for 15 h at 700°C, and then aircooled. The resulting microstructure containedspheroidal carbides with an average d i a m e t e r of 1~tm in a f e r r i t e m a t r i x with an average gra in size of4.7 ~tm. This represents an average planar density ofFeaC particles of about 7 × 107 cm-2. T h e r e was a sec-

ond significant distribution of particles identified bymicroprobe analysis as MnS. They were from 3 t o 5~tm in size, of ellipsoidal shape, appeared to be randomlydistributed, and represented a density of about 5.5× 104 em-~. T h e i r participation in the fracture processwas duly noted as described later. It was common tofind la rge r carbide particles situated preferentially onthe gra in boundaries with s m a l l e r particles located inthe gra in matrices. Furthermore, for steels spheroid-ized by this simple quench and temper process it isknown that the carbides are generally interlinked bya network of subgrain boundaries. 9 Axisymmetric ten-sile specimens were then machined with a 1 in. (2.54era) long and 0.25 in. (0.63 em) diam gage. Specimenswith premaehined deep n e c k s also had a 0.25 in. (0.63era) diam minimum section with a 0.50 in. (1.26 era)diam away from the neck. All specimens were deformedin tension at room temperature until visible necksformed and were then unloaded. The extent of neckingin the initially smooth specimens was varied, withsome of the specimens completely fractured. A portionof each specimen, containing the necked down region,was longitudinally sectioned, polished t o the centerand then etched in a 4 pct P i c r a l and 3 pct Nital solu-tion. T h e s e sections, as well as the fracture surfaces,were then examined by optical and scanning electronmicroscopy. Seve ra l important quantities were deter-mined by quantitative metallography and are plotted v s

the distance from the plane containing the minimumsection of the neck in F i g s . 5 through 8. Determinationof the critical conditions leading to void initiation wasfacilitated by u s i n g the circumferentially deeply

Table I. Composition of the Steel in Weight Percent

C Mn P S Si Fe

0.44 0.69 0.007 0.024 0.15 Balance

368-VOLUME 10A, MARCH 1979 METALLURGICAL TRANSACTIONS A

F i g . 2--Stereoscopic m i c r o g r a p h of the fracture surface of ahydrogen charged specimen.

4 . R E S U L T S A N D D I S C U S S I O N

Hydrogen charging, as just described, typica l lycaused a los s in ductility of about 22 pct for the 1045s t e e l . Again, ductility is be ing defined as the net r e -duction in a r e a in the minimum sect ion of the neck.The elongation to fracture was only m i l d l y af fected( l e s s than 1 pct for the initially smooth spec imens)and, as mentioned ear l i er , the fracture mode was un-changed from a normal ductile rupture (cf Figs . 2 and3). Furthermore , there was no noticeable change inthe stress -s tra in behavior of this s t e e l , as displayedin Fig. 4. To determine why failure occurred ear l i er ,that i s with less necking, in the hydrogen chargedspec imens several microscopic features of the f r a c -ture process were quantified. T h e s e are described be -low.

Void Initiation

Voids generally initiated by the separation of par-t i c l e -matr ix interfaces; Fig. 5 shows how this oc-curred within the necks of the initially smooth spec i -mens . The figure sugges t s that at a given stage in the

Fig. 3--Stereoscopic micrograph of the f r a c t u r e s u r f a c e ofa n uncharged specimen.

grooved spec imens ske tched in Fig. 6. Both theinitially smooth and deep ly grooved spec imens of the1045 s t e e l conformed accurate ly in des ign and m a t e -rial properties to those used in e a r l i e r studies, of voidinitiation by Argon and coworkers zone

Hydrogen was introduced by a more or l e s s stand-ard electrochemical charging technique, n with caretaken to minimize the extent of irrevers ible damage.Charged blanks were inspected with optical and scan-ning e lectron microscopy to determine the extent ofsurface or subsurface damage. When the currentdensities were be low about 6 m A / c m2, for chargingt imes up toi 24 h, no damage was evident . However, itshould be noted that at current dens i t i e s larger thanthis considerable surface damage was generally pro-duced, ranging from smal l subsurface microcracks tomassive surface bl i s ter ing with increas ing currentdens i ty . The charging conditions employed for our ex-per iments were a current dens i ty of 4 m A / c m2 and acharging time of 24 h. Under t h e s e conditions the hy-drogen ef fects are n e a r l y revers ible in that most ofthe ductility (70 pct as measured by the reduction incross-sect ional a r e a at fracture) was recovered bythermal outgassing for 400 h at 150°C.

~_ 1450

1305

1160 I x H Y D R O G E N C H A R G E D¢J %• U N C H A R G E D ; F R A C T U R E

Z 1015 G R A I N S I Z E ~ 4 . 7 # m ,

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f 1~1¢ NECKING B E G I N S

~" 29o

I ~ : 145

0 I I I I I i I I I I I I I I I [ I•0 2 .04 .06 ,08 I0 12 14 .16 .18 .20 .22 .24 .26 .28 3 0 5 2 3 4

A X I A L . S T R A I N ~

F i g . 4- -Stress - s tra in c u r v e s for f ine-grained spec imens , hy-d r o g e n charged and uncharged, both taken to f r a c t u r e . Flows t r e s s , a s d e f i n e d by B r i d g m a n ,22 plotted on the ordinate .

ffl

> .4aILl .3b-"

~ . ~m 0

l-F- .2

W 0_J

~ . ~

Z .20I'-

~ 0

U-

;<xX o ) R E D U C T I O N I N A R E A ( R A ) = 0 . 4 5x 0

× ×• • • X X

b ) R . A . = 0 . 4 7K X x

x xx

c ) R . A . = 0 . 5 5 x H Y D R O G E N C H A R G E D• U N C H A R G E D

~ . / , U N C H A R G E D , F R A C T U R E D( R . A . = 0 . 6 7 )

~ ~ x'

0 .05 .10 .15 .20 2 5 3 0 3 5 . 4 0 . 4 5 . 5 0 . 5 5 . 6 0 6 5 .70 .75 . 8 0 . 8 5 . 9 0 . 9 5 1.00V E R T I C A L D I S T A N C E FROM M I N I M U M S E C T I O N OF N E C K ,c m , (Z ) --~-

F i g . 5--Fraction of partic les wi th assoc ia ted voids v s v e r t i -c a l dis tance from m i n i m u m s e c t i o n of neck (Z) for s e v e r a ld e g r e e s of necking. Uncharged f r a c t u r e d s p e c i m e n i n c l u d e dfor c o m p a r i s o n to c h a r g e d f r a c t u r e d specimen.

METALLURGICAL TRANSACTIONS A VOLUME 10A, MARCH 1979-369

deformation process, at comparable locations in thet r i a x i a l f i e l d of the neck, t h e r e is a distinct t r e n d formore of the particles t o have associated voids in thecharged specimens. The effect is not l a rge , but n e v e r -theless suggests that void initiation may be influencedby hydrogen.

Argon et a l~2 have developed a technique for d e t e r -mining the c r i t i c a l conditions for void initiation atparticles and have given n u m e r i c a l details for spheroid-ized 1045 steels. The procedure involves determiningthe position within the neck of a tensile specimen, andthereby a l o n g a gradient of t r i a x i a l s t r e s s and plasticstrain, w h e r e void initiation appears t o just begin.Then if the (local) stresses and strains are known bynumerical analysis, conditions required, or at leastsufficient, for void initiation can be determined. How-e v e r , determination of the local conditions requiredfor particle cavitation is generally complicated by the

10x I 0 s

T 8Ec~" 7

3500(510)

3100(459)

2800(408)

2450(357)

O> 6r (o /R) i = 1.0

x HYDROGEN CHARGED, o/oi =0.855 L._..~.l-% • UNCHARGED, (a/(3i = 0.85

N 3 1 ~ . X ~ O ' r r FOR 0/o i=0.85

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L4

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~3

I

00

2100(306)

1750(255)

1400(204)1050(153)

700( i02)

550(51)

00.1 0.2 0.3 0.4 0 5 0.6 0.7 0,8 0.9 1.0

z / o 0(u)

3500(5 i0)

x HYDROGEN CHARGED, o / a t = 0.78531OO

• UNCHARGED, o / oi = 0.774 (459)

2800(408)2450

~\ (357)x~x / O'rr for 0/o i =0.774 2 I O0

/ (3o6)

\ ~ ~ ,4oo\ " x [ " ~ 1(204)

~ • ~ ~ O-rr f o r 0 / 0 i = 0 . 7 8 5 -J 1 0 5 0\'N \~ ~ I (153)

",. \ ~ -I 7oo." . "~- ~ / (io2)

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÷

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?,%

+

%

F i g . 6 - - ( a ) A r e a l densi ty of voids v s vertical distance frommin imum section of neck normalized with the ini t ial r a d i u saway from the neck ( z / a g ) for circumferentially notchedspecimens s t r a i n e d to a / a i = 0 .85 . The s t r e s s crrv acting ona part ic le i n t e r f a c e i s computed us ing A r g o n ' s model , l°'t2(b) A r e a l densi ty of voids v s z / a o for an uncharged c i r c u m -ferentially notched specimen strained to a / a i = 0.774 and ahydrogen c h a r g e d specimen w h i c h f r a c t u r e d at a / a i = 0 .785 .

nonuniform size and spatial distribution of particles.For example, as Fig. 5 shows, nowhere along the gagesection (until the grip portion) of the initially smoothspecimens does the void density vanish. Uniform plas-tic strains, in this c a s e , on the o r d e r of 0.2, were evi-dently sufficient to initiate voids at particularly sus-ceptable points typically within particle clusters andaround particles with large aspect ratios. T o utilizeArgon's method the above measurements were re-peated on specimens prepared with deep circumferen-tial grooves and with correspondingly much steepergradients of s t r e s s and s t r a i n - t h e details and resultsare illustrated in Fig. 6. The figure shows that thedensity of voids aga in does not smoothly vanish butapproaches a well defined low leve l of about 0.5 × 106voids-cm -2. This v a l u e is well within a factor of 2 ofthe lower values for the fraction of separated parti-cles shown in Fig. 5 and we note equally comparable tomost of the last data points listed by A r g o n and Im~z int h e i r Fig. 9. Interestingly enough it was found that ap-proximately 15 pct of these voids were associated withMnS particles which would suggest that the sulfidescavitate more easily than the carbides and that a f t e r auniform axia l strain of 0.2, most of the sulfides haveformed voids. Taken together with the e a r l i e r resultsof Gurland,13 which showed that voids initiate from thevery beginning of plastic strain, and the recent re-view of Goods and Brown~4 these results suggest thatthe initiation of voids in these ductile steels cannot beviewed as simply a " c r i t i c a l s t r e s s " phenomenon butr a t h e r as a combined stress-plastic s t r a i n event. In-deed to the extent that large interface stresses whichcause cavitation a r i s e from incomparable plasticstrains around particles stress and s t r a i n c r i t e r i aare nearly equivalent x4'15 whereas s t r e s s increasesdue to constraints also contribute significantly t o voidinitiation, provided plastic strains sufficient to sepa-rate m a t e r i a l accumulate.

With the above perspective we have interpreted theresults on void initiation in Fig. 6 by picking out thepoints where the line of void density v s axia l distancecrosses the base line of 0.5 x 106 voids-cm -z. The datathen yield consistent results computed by Argon'smethod. For the uncharged specimens the critical in-terfacial s t r e s s is found to be 965 (140) and 937 (136)MN//m2 (ksi) as compared to the 1400 MN/m2 (203 ksi)reported by A r g o n and Im.12 This lower estimate isdue in part t o the la rge r average particle size em-ployed here and t o the ambiguity which we note is in-herent t o this procedure, of choosing a base line voidcount. A base line count of about 1.0 x 106 for examplewould yield critical stresses in the uncharged speci-mens of 1275 and 1160 MN/m2, For the hydrogencharged specimens the critical s t r e s s e s are found tobe 593 (86) and 517 (75) MN/m2 (ksi), indicating that asignificant effect does exist. Such an effect would al-ways be found so long as the base line count wasla rge r than the 0.5 × 106 voids-cm-e used in Fig. 6;lower values would appear to diminish the dependence.Finally, we point out that the data and conclusionsd r a w n from Fig. 6 should not be influenced by voidgrowth s i n c e the specimens involved were deliberatelystrained so as t o only beg in the void initiation proc-ess, although in the case of one of the hydrogen chargedspecimens which fractured, some void growth was evi-dent near the fracture surface.

370-VOLUME IDA, MARCH 1979 METALLURGICAL TRANSACTIONS A

Development of Voids

The manner in which voids grew and coalesced wasquite interesting and was made apparent when longi-tudinal sec t ions were inspected at high magnification.Voids initiated above and below the p~trticles a longthe tensile axis and then grew into the matr ix with sub-sequent plas t ic deformation. Coalescence of the voids,though, very commonly occurred by the propagation ofvoids in an almost crack like fashion along interfaces(Figs . 8 through 10), notably grain or poss ib ly evensubgrain boundaries. We reca l l that the grain bound-ar ies in t h e s e s t e e l s are preferentially populated withthe larger part ic les and that even the grain matr ixcarbides are interlinked with subgrain boundaries. In-terface separations such as the examples shown inFig. 10 have not been given much attention in recentwork but we believe are observable in the photomicro-

-x

I

. 0 3

. 0 2

.01

0~'~0 . 0 4 <

. 0 5( ~ . 0 2

~- .010 0

~ .09~ .08,¢( . 0 7

"~ .05.04.03.02.01

0

~x x a ) R.A.:O.45t xx

.04

b) R.A. = 0.47

~ : xx x

•03,a..~ .02h

.01ahi 0

0 3 'n,,"0 20..

" ' .01>... 0

.09

Z .08: 30 0 7m

.06b.0 05~) ,04

.03~ .02U- .01

C) R . A . : 0 . 5 5

x HYDROGEN CHARGED• UNCHARGED

UNCHARGED,FRACTURED(R.A,=O,67}

x

xx, , ~ ` , , , , , ~., r@,, ~ , ,~, ,

0 .05 .10 .15 .20 .25 .30 .55 .40 ,45 .50 .55 . 6 0 . 6 5 . 7 0 , 7 5 . 8 0 . 8 5 . 9 0 . 9 5 LO0

Z,cmF i g . 7 - - A r e a f r a c t i o n of voids v s Z for severa l d e g r e e s ofnecking , including data for an uncharged fractured specimen.

- x

a ) R.A.=0.45

.04

x b ) R . A . = 0 . 4 7 ~ ,

• • x ~ f l

c) R . A . = 0 . 5 5

x HYDROGEN CHARGED F i ~ ' / i• UNCHARGED bs: (~)

X A UNCHARGED,FRACTURED-- (R.A.=O.6?) " N=NUMBER OF GRAINS

P• AVERAGE PERIMETEROF EACH GRAIN

x

I I I I I 1 41]~ I I ~ I III~I I III I l~lll I• 5 .10 .15 2 0 . 2 5 .30 3 5 . 4 0 . 4 5 . 5 0 . 5 5 . 6 0 .65.70.75.80.85~90.95 1.00

Z,cm

F i g . 8--Fraction of boundary separated v s Z for severa l de-g r e e s of necking , including data for a n uncharged fracturedspecimen. The average g r a i n perimeter i s determined bymeasuring the g r a i n s ize before deformation, and then con-sidering the la tera l and long i tud ina l strains at each pointalong the n e c k .

graphs of o ther recent reports on fracture in spheroid-ized plain carbon s t e e l s (Figs . 3 and 4 in Liu and Gut-]andz6 and Figs . 8, 15, and 16 in Argon and ImZ2).Rogers, on the o ther hand, has emphasized his ob-servat ions of void growth along the grain boundariesof OFHC copper heat treated in a hydrogen atmo-sphere .6

In Fig. 7 the a r e a fraction of voids along the centerline of the initially cylindrical spec imen i s plotted v sthe vertical distance from the minimum sect ion of theneck. Although the resu l t s do not typica l ly form asmooth curve, they show that at a comparable stagein neck development and at comparable locationswithin the triaxial s tress and plas t ic s tra in f i e l d ofthe neck, the hydrogen charged s t e e l s develop largervoid volume fractions. In the neck's center this repre-sen t s an almost three fold increase as compared to thetwo fold increase in the number of voids in a unit a r e a(or roughly the fraction of part ic les with assoc iatedvoids) suggesting that void growth as well as void initi-ation is accelerated. Our study of the micrographssugges ted that the explanation for this involved an ac-ce lerated coalescence of voids by interface separa-t ions . Figure 8 displays how this separat ion of bound-ary ligament ly ingbetween decohered part ic les de -veloped within the neck. As the data indicate, bound-ary separation was great ly acce lerated in the center-most portion of deep necks• Thus a void volume frac-tion sufficient to induce the f inal s tages of fracture wasattained with a reduced neck. The data indicate that thelarge transverse s tres ses assoc iated with the neckspromote void propagation along grain boundaries (withor without hydrogen charging) but that , s ince boundaryseparat ion i s apparently much e a s i e r ( i . e • occurs withreduced s tres ses on the boundaries) in the hydrogen

t t

" ~ - ~ --TS-TRESSES"-i--.-~TO NECK ~

A B C D E

tor

A

4

O" O"

F R A C T U R E

EF i g . 9- -Schemat ic representat ion of void f o r m a t i o n and de-velopment by growth and coalescence along a boundary drawnwith approximate correspondence to the s t a g e i n deformation.

METALLURGICAL TRANSACTIONS A VOLUME 10A, MARCH 1979-371

(a) (b)

2/.Lm 2 iu,rn

(c) (d) (e)Fig. 10--(a) and (b) M i c r o g r a p h s of void formation and development in a hydrogen c h a r g e d specimen. Examples of v o i d s p r o p a -g a t i n g a l o n g i n t e r f a c e s a r e indicated by a r r o w s . The t e n s i l e a x i s orientation i s v e r t i c a l . (c) through (e) Micrographs of voidformation and development in an u n c h a r g e d s p e c i m e n .

charged steels, interface fracture, void development,and hence specimen failure o c c u r r e d with less s e v e r enecks. The end result is a loss in the reduction inarea at failure with an otherwise i d e n t i c a l mode offracture. The area fraction of voids v s the longitudi-nal position from the minimum section of the neck foran uncharged but completely fractured specimen isalso shown in Fig. 7(c). It is readily seen that , withinthe fracture zone, void densities are comparable t othose developed at fracture in a hydrogen cha rged ma-t e r i a l . We see a s i m i l a r correspondence in Fig. 8(c),in which the fraction of boundary separated is plottedv s the v e r t i c a l distance from the min imum section ofthe neck. This is aga in consistent with the hypothesisthat hydrogen does not a l t e r the fracture process butonly accelerates one or more of the individual steps.Our data suggest that void initiation is promoted atlower s t r e s s e s and strains but that indeed a m a j o rinfluence of hydrogen is enhancing void growth andthat this enhancement comes about through an ac-celerated interracial separation. This acceleration ofvoid growth by hydrogen is consistent with the ob-servations of Garber e t a l1 who reported a very m i n o reffect on void initiation. However, since they make no

mention of profuse boundary separation with or with-out hydrogen, we i n f e r that the mechanistics were dif-ferent.

Further evidence that interfacial fracture accountsfor a significant part of void development was obtainedby a comparative study of fine v s c o a r s e grainedsteels. Seve ra l cylindrical specimens were made fromthe same 1045 s t e e l austenitized at 1200°C for 2 h,furnace cooled t o 825°C, then quenched in a 25 pct solu-tion of water-soluble oil. Tempering was c a r r i e d outfor 15 h at 700°C so as t o insure a carbide size al-most identical t o that of the fine grained specimens,but with an average gra in size of approximately 50/zm. Typical results are presented in Table II whichdisplays data for fractured hydrogen cha rged speci-mens in tabular form due t o the wide range of datapoints. An i n c r e a s e in gra in size lead t o an evenla rge r loss in ductility, that is, the neck in the c o a r s eg ra ined s t e e l was less deep than in the fine grainedm a t e r i a l by a factor of approximately 15 pct. Fracturesurfaces were comparable to those of the fine grainedspecimens, with approximately the same average dim-ple s i z e . Within the fracture zone the area fraction ofvoids in the c o a r s e grained s t e e l was comparable

372-VOLUME 10A, MARCH 1979 METALLURGICAL TRANSACTIONS A

Table II. Tabular Comparison of the Quantities in Fig. 5, 7, and 8 for Coarse and Fine Grained Charged and Fractured Specimens

Z, cm ~ 0.000 0.013 0.025 0.038 0.125 0.200 0.250 0.300

Fract!on of boundary separatedCoarse grained (RA = 0.47) 0.929 0.286 0.271 0.156 0.216 0.441 0.076 0.110Fine grained (RA = 0.55) 0.084 0.046 0.056 0.037 0.049 0.011 0.017 0.010

Area fraction of voidsCoarsegrained 0.122 0.033 0.048 0.014 0.022 0.026 0.007 0.013Fine grained 0.085 0.052 0.059 0.051 0.046 0.016 0.028 0.016

Fraction of particles with associatedvoidsCoarsegrained 0.490 0.239 0.190 0.155 0.122 0.285 0.082 0.094Fine grained 0.580 0.247 0.420 0.225 0.243 0.103 0.150 0.135

Both specimens charged and fractured.

with, and even la rge r than in the fine g ra ined steel,yet the fraction of particles with associated voids was,everywhere a l o n g the specimen's ax i s , less in thec o a r s e grained steels. The comparable void volumefraction at the fracture surface can be accounted forby the noticeable increase in the fraction of boundaryseparated. With an increase by a factor of 10 in thegra in size ( i . e . with one-tenth the n u m b e r of bound-a r i e s per unit a r e a ) a roughly tenfold i n c r e a s e in thefraction of boundaries separated is r e q u i r e d t o fullyaccount for a comparable void volume fraction by in-terface splitting. Our data indicate that this is thec a s e . The increase in separated boundary is slightlyl a rge r than tenfold and the micrographs indicate thatthe extent to which boundaries separate is also la rge rin the c o a r s e grained material. Thus, the fact thatthe area fraction of voids is approximately 30 pctl a rge r in the minimum section of the neck of thec o a r s e g ra ined material, coupled with the more thantenfold increase in the fraction of boundary separated,indicates that the void development is indeed achievedby interracial separation. Since this interfacial " f r a c -t u r e " apparently plays a much more important r o l e ,that is, a much la rge r fraction of the boundary mustbe separated in the c o a r s e grained m a t e r i a l than inthe fine grained material, the hydrogen charging hasa more m a r k e d effect.

F r a c t u r e Surfaces

The fracture surfaces exhibited r a t h e r equiaxeddimples (Figs. 2 and 3), both shallow and very deep.It is indeed curious that the deep holes which were ar e s u l t of g ra in boundary separation, as discussede a r l i e r , were in genera l quite axisymmetric. Sincegra in boundaries are two dimensional, we might ex-pect many lenticular dimples to be f o r m e d by thisg ra in boundary separation process. We also note thatthe deep holes such as the example shown in Fig. 10(e)were not associated with the few sulfides which wereeasily identified by microprobe analysis. T h e i r axi-symmetric shape indicates that growth was plasticitycontrolled, but as suggested by our observations, ac-celerated along the gra in boundaries. F u r t h e r evi-dence of axisymmetric void growth is shown in Fig.11 which is a cross-sectional view of a charged speci-men polished just below the fracture surface.

An increase of 10 pct in the average dimple size wasachieved by hydrogen charging. This was attributableto a reduced n u m b e r of s m a l l dimples, which arecaused by the very s m a l l voids formed just p r i o r t o

fracture. Failure with a less s e v e r e neck, and thuslower transverse stresses, results in the formation off e w e r of the very s m a l l voids and thus in a la rge raverage dimple s i z e .

It is not c l ea r that our results and conclusions on ahydrogen charged spheroidized 1045 s t e e l can readilybe generalized so as to provide insight into hydrogeninduced ductility losses in other materials, including" c l e a n e r " plain carbon steels. The fairly high phos-phorous level in our s t e e l (Table I) may have contri-buted t o a g r e a t e r amount of boundary and perhapsparticle-matrix separation than might otherwise haveoccurred. However, it would indeed seem worthwhilet o attempt t o do so since hydrogen induced interfacialfractures are common. 2'17 The results for our plaincarbon steels, on the other hand, do seem c l e a r . Hy-drogen reduces ductility, as m e a s u r e d by a loss in thenet reduction in a r e a , by accelerating void develop-ment t o a c r i t i c a l s t age at which complete fracturesets in. This acceleration is brought about by facili-tating void initiation and by a reduction in the t r a n s -v e r s e stress, and therefore in the degree of neckingrequired to grow voids preferentially a l o n g internalboundaries.

From a mechanistic viewpoint, it is difficult t omake firm conclusions concerning the role of hydro-gen in inducing premature fracture. Void growth byhydrogen pressure buildup within the voids has notbeen clearly established, either in the literature or in

Fig. ll--Cross-sectional view of a c h a r g e d s p e c i m e n pol ishedjus t below the f r a c t u r e s u r f a c e .

METALLURGICAL TRANSACTIONS A VOLUME 10A, MARCH 1 979-373

our own preliminary calculations. However, if hydro-gen migration to the voids, e i t h e r by diffusion or dis-location transport with a resulting development of highpressures within them, plays an important role in voidgrowth then it must do so when the void volume f r a c -tion is quite s m a l l and hence e a r l y in the deformationas perhaps in e a r l y void initiation. This must be sobecause the specimen in the charged condition is notan infinite source of hydrogen, a fact which is all toooften overlooked, and thus when the m a t r i x concentra-tion of hydrogen begins t o drop the driving f o r c e for,and therefore the rate of, hydrogen transport to thevoids also decreases. Furthermore, as the void vol-ume fraction increases, so does the n u m b e r of sitesfor hydrogen deposition, and thus the rate of m a t r i xdepletion increases.

From our circumferentially notched specimens, wecan make some useful observations concerning voidgrowth. An indication of the average void size can beobtained by dividing the area fraction of voids by thea rea l density of voids. This procedure yields the re-sult that in the specimens strained t o ac/ai = 0.85, foridentical l oca l s t r e s s e s and strains, the average voida r ea is 36 pct l a rge r in the charged specimen than inthe uncharged specimen. Since the void density is s t i l lquite low and the voids have not yet begun t o propa-gate a l o n g internal boundaries, a hydrogen pressureeffect, if it exists, must be l a r g e s t at this stage forthe reasons outlined above. However, most of this in-c r e a s e d void growth can be accounted for by the sim-ple fact that some 92 pct of t h e s e voids ( i . e . the num-ber in e x c e s s of the base line count of 0.5 × 106 cm"¢)having in genera l formed at lower strains (andstresses), have had more time t o grow. Rice andTracey 18 have developed a formalism for describingthe growth of noninteracting initially spherical voidsin ideally plastic and l i n e a r hardening materials withhigh trlaxiality. For the l i n e a r hardening m a t e r i a l ther a t e s of increase in the principal r a d i i are given by

'~L! 4 T~ j R ° K , L =I, II, III.

w h e r e a~ is the trlaxiality, labelled aT in the work ofA r g o n e l a l l°,z2 and in our Fig. 6. r ~ is the equivalentflow stress in s h e a r . Because of the exponential natureof the resulting expressions for the R K , and since themacroscopic stress and s t r a i n history is identical forboth the charged and uncharged specimens, the r a t i oof the average void size in a charged specimen to thatin an uncharged specimen should not change from thetime (or actually the strain) which a void is initiatedin the uncharged specimen. S ince void initiation oc-curs e a r l y in the deformation (as we quantified t o someextent e a r l i e r ) for both the charged and the un-charged specimen, we used the results of Brownand McMeeking19 to determine Cr~°/T ~. Our calcula-tion predicts the average area of a void in the hydro-gen cha rged specimen t o be approximately 28 pctg r e a t e r than that of the uncharged specimen. Thevoid volume fraction in the cha rged specimens atthis s t age is m e a s u r e d t o be 0.005. B a s e d upon thestudies of Smialowski, 2° the hydrogen content of oursteels was estimated 2~ and a maximum internal hy-drogen p r e s s u r e calculated by the a d hoc procedureof assuming that an equilibrium amount of hydrogen

374-VOLUME 10A, MARCH 1979

accumulated in the voids as a gas. This yields apressure of approximately 20 MN/m2 which wouldhave a minimal effect on growth r a t e s . Thus whilehydrogen gas pressure c o u l d have influenced p r i o rvoid growth, that is at much s m a l l e r volume fractionsand assuming suitable kinetics prevailed, there shouldbe little or no pressure effect by this point or cer-tainly afterwards. The numerical results using theRice and T r a c e y growth law for a nonhardening ma-t e r i a l fitted to the yield s t r e s s of our m a t e r i a l arewithin 2 pct of the 28 pct quoted above.

Further, plasticity controlled hole growth cannot suf-ficiently account for the difference in average void sizea f t e r further straining. Using the procedure just dis-cussed on specimens strained t o a /a i = 0.774 (thecharged specimen fractured at a / a i = 0.785) the a v e r -age void area for a hydrogen charged specimen was400 pct l a rge r than for an uncharged specimen. As wenoted before, the hydrogen charged specimen exhibitedsome boundary separation near the fracture surface.Furthermore, many of the voids appeared t o growa l o n g particle-matrix interfaces. Hydrogen acce l e -r a t e d growth a l o n g interfaces can be accounted forwithout the necessity of an internal pressure mecha-n i sm. Hydrogen at a g ra in or subgrain boundary (orfor that matter a particle-matrix interface) could en-hance void growth a l o n g the boundary by decreasinginterracial cohesion,z Unfortunately, the magnitude ofthis effect is also not c l ea r due to the lack of a suit-able absorption isotherm for hydrogen at high fugaci-t i e s . Thus, at the present time we can only concludethat the m a j o r role of hydrogen in accelerating theductile fracture of our spheroidized 1045 s t e e l was t oenhance void development by reducing the stresses(or strains) necessary to form voids and to propagatethem along gra in and possibly subgrain boundaries,without making a definitive statement as to the mecha-nistic details of this process. Continuing r e s e a r c h isb e i n g c a r r i e d out in o r d e r t o understand more fullythe mechanisms involved.

5. CONCLUSIONS

Internal hydrogen facilitates void initiation at FeaCparticles causing voids to form at lower stresses andstrains. Voids initially grow into the matrix with plas-tic deformation but subsequentlypropagate, in aseemingly ductile fashion, along grain and perhapssubgrain boundaries. Hydrogendoes not have a majoreffect on initial void growth, as the larger averagevoid size in the hydrogen charged specimens at smalldeformations can in large part be accountedfor byplasticity controlled growth in the triaxial field im-posed by the specimen's neck. Hydrogendoes have animportant effect on the intermediate and latter stagesof void growth and coalescence w h e r e the voids propa-gate along internal interfaces.

ACKNOWLEDGMENTS

The authors gratefully acknowledge the support ofthe Energy R e s e a r c h and Development Administrationthrough Contract EY-76-02-3084.

METALLURGICAL TRANSACTIONS A

R E F E R E N C E S

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2. CI D. Beachem:Met. Trans., 1972, vol. 3, p. 437.3. J. R. Rice: Effect o fHydrogen on Behavior o f Materials, A. W.Thompson

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METALLURGICAL TRANSACTIONS A VOLUME 10A, MARCH 1979-375