Solid metal-induced embrittlement of steel

51

Materials Science and Engineering, 18 (1975) 51--62 © Elsevier Sequoia S.A., Lausanne -- Printed in The Netherlands

Solid Metal-Induced Embrittlement of Steel*

J.C. L Y N N * * , W.R. W A R K E and P. G O R D O N

Department of Metallurgical and Materials Engineering, Illinois Institut~ of Technology, Illinois (U.S.A.) (Received J u n e 12, 1974)

SUMMARY

The embrittlement of steel by surface coatings of solid Zn, Pb, Cd, Sn and In has been studied. This embrittlement manifests itself as a reduction in tensile ductility over a temperature range extending from about three quarters of the embrittler absolute melting temperature up to this melting temperature. It is shown that the embrittlement is caused by the growth of stable, subcritical intergranular cracks. The growth rate and activation energy for growth of these cracks were estimated. The data suggest that the severity of embrittlement is limited by the rate of supply of embrittler atoms to the crack tips. Various transport mechanisms are considered; it is concluded that multilayer surface self-diffusion of the embrittler over the coated crack surfaces is the most likely mechanism.

INTRODUCTION

The loss in ductility and fracture strength experienced b y a metal when stressed in tension while in the presence of a liquid metal environment has come to be called liquid metal embrittlement (LME). Generally, it has been assumed that this embrittlement would n o t be observed below the melting point of the embrittler metal. For example, in a recent review article, it was stated [1], " O f course, if the * This paper is based on the thesis research o f J.C. L y n n s u b m i t t e d in partial fulfillment of the req u i r e m e n t s for the Ph. D. degree. ** N o w at A m s t e d Research Laboratories, Bensenville, Illinois (U.S.A.).

temperature falls below the solidus of the wetting agent, no embrittlement is observed". There is some evidence in the literature, however, which casts d o u b t on the validity of this statement. A case in point is the embrittlement of steel by cadmium. In 1960 Rostoker, McCaughey and Markus [2] d o c u m e n t e d the LME of steel by this metal. At a b o u t the same time, Kennedy [3,4] observed delayed hrittle failure of cadmium plated high strength steel at temperatures t o o high for hydrogen embrittlement b u t definitely below the cadmium melting point of 610°F. This problem has also been recognized by the high strength steel fastener industry, where it is termed stress alloying [5], and use of cadmium plated bolts above 450°F is forbidden. Iwata, Asayama and Sakamoto [6] and Fager and Spurr [7,8] have studied solid cadmium embrittlement of steel and of titanium, and Hildebrand [9] observed cracks on cadmium plated notch tensile specimens of various steels below the melting point of cadmium. Tiner [10] has reported that zinc is embrittled by solid mercury. In our own laboratories we have found that the Cu - Hg system also exhibits this p h e n o m e n o n [ 11 ]. In an extensive study, Mostovoy and Breyer [12] demonstrated that leaded high strength steel is susceptible to a severe embrittlement in the range of the melting point of lead and this embrittlement was clearly shown to be a manifestation of LME. However, it was found that the embrittlement begins to set in around 400°F, some 200 degrees below the lead melting point (622°F), and that it is continuous through this melting point; there were no discontinuities or anomalies in the ductility versus temperature curves at or near 622°F. These observations would indicate that similar

52

fracture processes and embrittlement phenomena are active above and below the melting temperature of the embrittler. The purpose of the present study was to investigate whether the extension of liquid metal embrittlement of steel to temperatures below the embrittler melting points is a general phenomenon in steel, and to gain some insight into the mechanism by which the embrittler is transported to the tip of growing cracks. Those low melting point elements known to embrittle steel include Sb, Te, Zn, Cd, Pb, Bi, Sn and In. Of these, all but Sb and Te were employed in the present work.

EXPERIMENTAL PROCEDURE

AISI 4140 vacuum degassed steel was chosen for this investigation; the steel composition is shown below:

Composition, wt. % Material

C

Cr

Mn

Si

4140 Steel

0.41

0.96

0.87

0.23

Composition, wt. % Material

Mo

S

P

Fe

4140 Steel

0.18

0.016

0.014

Balance

Specimen blanks were austenitized at 1525°F for one hour, oil quenched and tempered at 875°F for one hour. This heat treatment was known to produce a room temperature ultimate tensile strength of 200,000 p.s.i. The specimens were then finish machined to standard 0.252 inch diameter tensile specimens. A globule of the embrittling element was soldered to the surface of each tensile bar at the point of minimum diameter along the gage length. This soldering was usually accomplished by positioning a chip of the low melting metal and some=zinc chloride flux on a bar surface and placing the assembly into a furnace held at an appropriate temperature. The temperatures employed were 675°F, 665°F, 575°F, 500°F

and 450°F for lead, cadmium, bismuth, tin and indium respectively. A modified technique was employed for zinc in order to avoid softening the bars by heating above the temperature at which they had previously been tempered. In this case, the zinc and some flux were melted in a small graphite container and the steel bar laid across the molten metal until wetting occurred. Prior to soldering, the embrittling metals were o f at least five nines purity; in the cases of lead and cadmium, six nines pure material was used. It is not implied that these high purity levels were retained through t h e soldering operation, however; the metal in place on the bar surface may have become somewhat less pure. Tensile tests were carried out at various temperatures from room temperature up to the respective melting points of the embrittlers. A hydraulic universal test machine equipped with a split furnace and temperature controller, and a high temperature extensometer, was employed for testing. The elastic strain rate of the tensile specimen was of the order of 10 -3 in./in./min while that in the plastic range after yielding was of the order of 10 -2 in./in./min. The specimen temperature was monitored by attaching thermocouples at three locations along the gage length; the given temperature was that at the embrittler globule at the onset of straining. During plastic straining, a small b u t significant temperature rise, due to thermomechanical effects, was noted. This temperature increase amounted to from 5 to 30 deg F depending on the amount o f strain prior to fracture. In some cases, for samples tested just below the melting point of the embrittler, the heat evolved was enough to melt the embrittler globule resulting in true liquid metal embrittlement. These cases will be noted in the results section. The properties obtained from the test results were the yield strength (0.2% offset), ultimate strength, true fracture strength, percent elongation (in 2 in.) and percent reduction of area. The fractured bars were studied macroscopicaily up to 30X. Selected specimens were sectioned longitudinally and prepared for metallographic examination using (mainly) dark field illumination to delineate the fracture paths.

53

/\

360 O Z

320

~------.-~/ zf

ZINC 320

L~X~_L~

TRUEFRACTURE STRENGTH

280

TRUEFRACTURESTRENGTH

/X_

0 Z

280

\ 240

240

z~

ULTIMATE TENSILE STRENGTH 200

:'-Z

YIELD

~

YIELDSTRENGTH O.O~,-~/

\

STRENGTH - - e

-O.13"O"O E)~? " O

a..

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120 120

.< Z

Ot-~..z.

"%'0

ELONGATION

[] 80

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o°

160 160

\

ULTIMATETENSILE STRENGTH

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~-U

l

~---B -B~E~43r~m43

Z O

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ELONGATION

[]

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]

80

o

20

z O

0

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o

REDUCTION OF AREA ~ % . ~ - ~

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I

I

I

200

4oo

6oo

800

z ~ o,, =, r__=u E3uJ Luc¢

REDUCTIONOF AREA

0

<

TESTTEMPERATURE,OF

Fig. 1. Elevated temperature tensile properties of 4140 steel at 200 k.s.i, nominal strength level.

40

I

I

I

200

400

600

800

TESTTEMPERATURE,°F

Fig. 2. Elevated temperature tensile properties of 4140 steel at 200 k.s.i, nominal strength level, and surface wetted with pure zinc (99.999%).

EXPERIMENTAL RESULTS

Tensile test

The base line (unembrittled) elevated temperature tensile properties of the 4 1 4 0 steel used throughout this study were determined and are shown in Fig. 1. A peak, due to dynamic strain aging, was observed in true fracture strength around 550°F. The reduction of area increased gradually from 55% at room temperature to 75% at 800°F. Elevated temperature tensile properties in the presence of zinc, lead, cadmium, tin or indium are shown in Figs. 2, 3, 4, 5 and 6, respectively. The solid points in each Figure represent those cases in which the thermomechanical effect (heat evolved during plastic strain) raised the temperature above the melting point of the embrittling metal. These cases, then, constitute true liquid metal embrittlement wherein the embrittler is present in the liquid state and is able to be transported into a growing crack by fluid flow. For these samples, a layer of solidified embrittler could later be seen on the fracture surfaces; this was n o t true for the other samples. In Figs. 2 - 6 it will be noted that for every embrittler

A O Z

280

Z :ff u,l~_

240

/

A

~

A

TRUE FRACTURE STRENGTH

l'

U LTIMATE ~. ~ L~ u

N<

/k\

LEAD

M.P. J I

I I

200

120

I I

ELON GATION

[ 3 - -

Z O

I

r : l ~

20

1

I I

80

O Z O~_

~z ~u

40

REDUCTION OF AREA

2OO

4OO

6OO

800

Fig. 3. Elevated temperature tensile properties of 4140 steel at 200 k.s.i, nominal strength level, and surface wetted with pure lead (99.9999%).

~

54 CADMIUM A'----'---- A~A

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MJJ.P.

280 TRUE FRACTURE \ STREN GTH ~'/~

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ELONGATION

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%

REDUCTION OF AREA

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TRUE

JFRACTURESTRENGTH

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lal \l

200 " 160

TENSILESTRENGTH YIELDSTRENGTH

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8O

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REDUCTION OF AREA

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I 200

YIELDSTRENGTH

"

12(:

ELONGATION

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-

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I

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6OO

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J

REDUCT,ON OF A R , A \

0

O Z

J

40

a. Z O

I 200

I

300

I 400

TESTTEMPERATURE,,OF

M.P.

280

~U O< ZL u~

160

Fig. 6. Elevated t e m p e r a t u r e tensile properties o f 4 1 4 0 steel at 2 0 0 k . s . i , n o m i n a l strength level, and surface w e t t e d w i t h pure i n d i u m ( 9 9 . 9 9 9 % ) .

TIN

~

O-.~... i

100

Fig. 4. Elevated temperature tensile properties of 4140 steel at 200 k.s.i, nominal strength level, and surface wetted with pure cadmium (99.9999%).

0

0--0-...O~.

o.

3U 0

•

80

0

I

4O

J

B--D--B--ra~ Z

8O

I

ULTIMATE TENSILE STRENGTH

20

I3E]I~E]~ I

Z ~.

uc

A .....~ TRUE FRACTURESTRENGTH ~']:k

;=

I

a;

~

~.~

,',Z ._1 u.i u.Ja~

STRENGT.

2~ ~z

28(:

ZOz <~.=

~IEI-~.-'U

Z

_j~ ~ Z~ "~

,T__u~

~-~200

,.<,

INDIUM

:I:

TESTTEMPERATURE,°F

Fig. 5. Elevated temperature tensile properties of 4140 steel at 200 k.s.i, nominal strength level, and surface wetted with pure tin (99.999%).

the true fracture stress and reduction of area were observed to decrease drastically even though the temperatures were well below the melting points and the embrittling species remained solid. Little if any change was noted in the other mechanical properties (yield and ultimate tensile strengths); as in the case of true LME, the flow curves were shortened but their shapes up to the point of fracture were not changed, indicating the similarity in basic embrittlement process between LME and the present solid state embrittlement. With the exception of zinc, where the embrittlement occurred at a higher temperature, each of the embrittlers prevented the dynamic strain aging peak from being observed. When submelting point tests were attempted for the Bi - steel couple, the brittle Bi globule invariably spalled off of the specimen surface soon after plastic deformation began. This removal of the embrittler from the bar surface precluded embrittlement and therefore no further tests were carried out on this couple. The embrittlement is most clearly revealed in terms of the normalized true fracture stress and normalized reduction of area. The values are norrealized by dividing them by the corresponding

55 o

value for an unembrittled specimen tested at the same temperature. Thus, any negative deviation from a value of unity indicates embrittlement and the extent of such deviation is an indication of the severity. The resulting curves are plotted in Figs. 7 - 11. Again it is evident that there is a progressive drop in ductility and strength extending over a considerable range of temperatures below the respective melting points. It is not clear from the present results whether the mechanical property changes are continuous up to the melting point since relatively few tests were carried o u t with the embrittlers molten. This is a point which needs further checking. Our evidence, however, does seem to agree with that previously reported for steel - Pb by Mostovoy and Breyer [12] and for steel - Cd by Iwata, Asayama and Sakamoto [6]. This continuity up to the melting point again suggests a mechanistic similarity between LME and solid metal embrittlement. In order to make comparisons most conveniently among the five elements employed in this study, the data for all five were plotted on a c o m m o n temperature scale, making use of the so-called homologous temperature -ZINC M.P.

TFSRATIO

LEAD ~

~

~

ACX

O Z

I TFS RATIO

0.8

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o.6 1.0

<

~

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0.8

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I RA RATIO I

O.6

Ol Z O U a

0.4

<~

0.2

I

I

~1

% 0

I

200

I

IJ

I

400

600

800

TEST TEMPERATURE,OF

Fig. 8. N o r m a l i z e d t r u e f r a c t u r e strength and reduct i o n o f area f o r 4 1 4 0 steel surface w e t t e d w i t h p u r e lead.

CADMIUM M.P.

O 1.0 Z

1.0

0.8

t.L

0.6 1.0

o

\

?

0.6

0.4

2

°

I

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o 0.6

RA RATIO

t

RA RATIO

0.4

z o 0.2

0.2

I

I 0 200

400

600

800

TEST TEMPERATURE,°E

Fig. 7. N o r m a l i z e d t r u e fracture strength and reduct i o n o f area for 4 1 4 0 steel surface wetted with p u r e

zinc.

Z

0.6 ~

~

~

0.8

o -r

RATI 0

~o

J

~

I

\

~A

1.0

M .P.I

2O0

4O0

6O0

8O0

TEST TEMPERATURE, OF

Fig. 9. Normalized true fracture strength and reduction of area for 4 1 4 0 steel surface wetted with pure cadmium.

56 TIN ].o 9

M.P.

A ~ A

I II

\&A

1.0 0

TFS RATIO

I

TFSRATIO

0.8 Z

~1~

'~0~

I

0.8 NO D70.6

if__

0.6 ~

g

i

1.0

u.

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Ill

~ Q ~

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0.8

1'

0

i-

0.6

X

0

I

I

RA RATIO

I I

0.4

A Pb

~_0.2

[] z. ~ s.

O

Q

,

/x

O ,°

I I

I-

0.4

0

0.4

I 0.5

I 0.6

I 0.7

I 0.8

' 0.9

i 1.0

HOMOLOGOUSTEMPERATURE, Tn :T/'r M

0.2

Fig. 12. C o m p a r i s o n o f n o r m a l i z e d true f r a c t u r e s t r e n g t h a n d r e d u c t i o n o f area as a f u n c t i o n o f hom o l o g o u s t e m p e r a t u r e for 4 1 4 0 steel surface w e t t e d as indicated.

I I

I

200

400

I

,

I

600

800

TESTTEMI~RATURE,OF Fig. 10. N o r m a l i z e d t r u e f r a c t u r e s t r e n g t h and reduction o f area f o r 4 1 4 0 steel surface w e t t e d w i t h p u r e tin.

INDIUM A ~A

1.0 O

M.P.

z

I

0.8

TFS RATIO ~ ' ~

i,

~0 u_ u,i i--

1.0

I

0.8

0

\\

the test absolute temperature divided by the respective embrittler absolute melting temperature; the results are shown in Fig. 12. It may be seen that all of these embrittlers begin to be effective when the test temperature reaches approximately three quarters of the respective melting points (TH = 0.75). Beyond that point, the curves diverge and the five elements vary in the severity of the embrittlement which they cause. The elements, listed in order of the severity of embrittlement with respect to normalized reduction of area, fall in the decreasing order Cd, Pb, Sn, Zn and In. A slightly different list can be obtained based on normalized fracture strength (Cd, Pb, Zn, Sn and In).

0.6

u. 0.4

0 Z 0

RA RATIO

~t

i-

0.2

e

$

I

I

II

100

200

300

!

TESTTEMPERATURE,OF Fig. 11. N o r m a l i z e d true f r a c t u r e s t r e n g t h a n d reduct i o n o f area f o r 4 1 4 0 steel surface w e t t e d w i t h p u r e indium.

Frac tograp h y The fractures of the base line specimens were cup-and-cone type ductile fractures typical of those in high strength steel [13]. The fractures were initiated fibrously in the center of the specimen and propagated outward. When a metal embrittler was present on the specimen surface, on the other hand, fracture initiated at the metal - steel interface and propagated into the bar. A set of specimens surface wetted with zinc and tested over a temperature range including the melting point of zinc (787°F) are shown in Fig. 13. Stable secon-

57

499°F

590°F

695°F

760°F

791 OF

Fig. 13. Elevated temperature tensile specimens of 4140 steel at 200 k.s.i. UTS, surface wetted with zinc and tested at indicated temperatures. (Melting point of zinc is 787°F.) 2.2X. dary microcracks were observed when the test temperature was below the melting point of zinc and these became larger and fewer as the test temperature approached the melting point of zinc. The size of the environmentally-induced brittle fracture zone on the fracture surface increased with increasing temperature and was always oriented per-

pendicular to the loading axis. The remainder of each fracture surface was dominated by ductile fracture, inclined at 45 ° to the loading axis. When melting of the zinc occurred (791°F in Fig. 13) a single fracture was seen; there were no secondary cracks. The first crack to form was able to grow and cause catastrophic failure before another crack could initiate.

Zn

I.C. LAYEr,

STEEL,

Fig. 14. SEM Photomicrograph of a secondary crack in a zinc coated tensile specimen tested at 758°F. 500X.

58 The same general fracture characteristics were noted for samples tested in the presence of Pb, Cd, Sn and In.

Metallograph y Selected specimens fractured in the presence of each of the five embrittler metals were sectioned longitudinally, m o u n t e d and polished for metallographic examination. The following observations were made: (1) Secondary microcracks were in evidence and seemed to follow prior austenite grain boundaries, Fig. 14. (2) Microcracks were found only in those areas of the interface where intimate metal-tometal contact had occurred. {3) Layers of intermetallic c o m p o u n d were observed at the embrittler - steel interface in zinc and tin coated samples, e.g., Fig. 14. This c o m p o u n d layer did n o t interfere with the embrittlement process. At one time it was thought that c o m p o u n d formation would preclude liquid metal embrittlement [2], b u t more recent analysis has modified this viewpoint somewhat [14].

DISCUSSION The above results represent a specific example of the general class of p h e n o m e n a which may be termed environmentally induced embrittlement. Other examples of this type of behavior include liquid metal embrittlement, hydrogen embrittlement and stress corrosion cracking. It is suggested that a more uniform terminology would result and the similarities among these various embrittlements would be emphasized if they were described as specific cases of environmentally induced embrittlement (EIE) with the particular environment duly inserted. Thus, we would speak of hydrogen induced embrittlement (HIE), liquid metal induced embrittlement (LMIE), or, as in the present case, solid metal induced embrittlement (SMIE). Such terminology avoids ambiguity such as is found in the term "solid metal embrittlement" and also avoids premature association of the description of the phenomenon with some mechanism, such as in "adsorption induced embrittlement". The present research has shown that Zn, Pb, Cd, Sn and In, each of which produces LMIE of steel, also embrittle steel at tempera-

tures well below the respective melting points of the embrittlers. These results suggest that all metals which in the liquid state embrittle steel also will embrittle the steel below their melting points, provided that the solid embrittler possesses sufficient ductility to plastically deform along with the steel up to the point of fracture without spalling off. Since this p h e n o m e n o n has also been observed for Cd on Ti [8], for Hg on Zn [10] and for Hg on Cu [11], the phenomenon may well be found in many, if not all, other metal - metal systems which are subject to LMIE. If such a correlation exists, the conclusions drawn regarding mechanisms of SMIE may well apply to LMIE and, in fact, to other forms of environmentally induced embrittlement as well. The commercial significance of the present results in terms of permissible service temperatures of structural metal parts in contact with low melting metals either internally, as surface coatings, or as adjacent contacting components, is immediately apparent. The scientific significance lies in the inferences which can be drawn regarding possible fracture mechanisms. As with many other phenomena in alloys, the process of 8MIE can be considered in terms of nucleation and propagation -- in this case, of the embrittlement crack. In view of the large a m o u n t of secondary cracking we observed in our test samples below the embrittler melting points, it is clear that at these temperatures crack extension rather than initial crack formation is the rate-controlling factor. In each case observed, many cracks nucleated before catastrophic fracture was produced by the propagation of these cracks until some critical situation was attained. The controlling mechanism is then presumably a crack extension mechanism. Since it is known that propagation of the crack can occur only so long as the embrittler can follow the crack tip, we assume that~crack growth is controlled by the rate of embrittler transport along with the advancing crack. Several possible mechanisms have been proposed for embrittlet transport in LMIE and can be considered as possibilities for SMIE. Chief among these mechanisms are gross fluid flow of the liquid embrittler [2], evaporation of the embrittler into the crack and transport in the vapor state [12], diffusive flow by surface diffusion of the embrittler along the substrate metal

59

surface (heterogeneous surface diffusion) [2] and transport by second m o n o l a y e r heterogeneous surface diffusion of the embrittler [ 15]. Still other possibilities are transport by volume diffusion through the substrate metal or along its grain boundaries, or by multilayer surface self-diffusion of the embrittler on the substrate metal surface. In the present case, the first of the above hypotheses can be eliminated immediately, for the embrittlers are not in the liquid state, nor in any of these cases does alloying take place to lower the embrittler melting point sufficiently, as indicated by the appropriate phase diagrams. In the second -- vapor -hypothesis it is assumed [12] that the vapor of the embrittler is present at equilibrium at the crack tip and that the increasing embrittlement with increasing temperature is due to the temperature dependence of the vapor pressure. It seems clear then that the actual severity of embrittlement should be given by an expression of the form:

I= g P fl ~7 where ! is the severity of embrittlement, K is a proportionality constant specific to a given base metal (the steel here), P is the vapor pressure of the embrittler, fl expresses the efficiency of deposition of the vapor at the crack tip (the "sticking coefficient") and 77 is the specific embrittling effect of the particular embrittling species. Our data show that I for all five embrittlers is approximately the same at a TH (homologous temperature) = 0.75, i.e., the embrittling effect first becomes measurable at this temperature. Assuming that fl~ unity at these relatively low temperatures, it follows that the product of P ~? is the same for all five embrittlers at TH = 0.75. The vapor pressures of these five embrittlers are listed below:

vapor hypothesis is inconsistent with our experimental data, unless the improbable (though, of course, possible) assumption is made that variations in 77 (or in fiT?if/~ ¢ constant) just balance these variations in P. Furthermore, the above elements are listed in the order of decreasing severity of embrittlement at TH = 0.95 as based on normalized reduction of area. Again, it seems that unless there is an improbable balancing of P and ~?, the vapor pressures at TH = 0.95 are not in the right order. It is also possible to estimate the number of metal vapor atoms present in a typical crack based on the ideal gas law, the equilibrium vapor pressure, an estimated crack volume and the test temperature. Such estimates yield unreasonably low values for Sn and In (less than one atom) and marginal values for Pb. We are thus left with the various diffusion hypotheses as most likely mechanisms for transport of the embrittler. As in the case of vapor transport, these may be tested on the basis that the embrittlement effect first becomes evident at the same homologous temperature for all five embrittlers. In addition, it is possible to calculate -- albeit very roughly -- the diffusion coefficients, Dexp, indicated by the experimental data and compare these with the expected values based on the various hypotheses. The "experimental" D can be estimated on the basis of the equation: x 2

Dexp ~ 2~where x = distance of crack propagation to failure, and t = the propagation time. Assuming that the crack propagated concurrently with plastic deformation beyond yielding, it was found that t was of the order of one minute. The propagation distances (crack depths) were of the order of one mm, so that:

Vapor pressure (atm.)

Cd Pb Sn Zn In

(10-1) 2 2 × 60 ~ 1 0 - 4 c m 2 / s e c "

at T H = 0 . 7 5

at T H = 0 . 9 5

Dexp ---

6X 1 X 4 X 1 X 5 X

4X 5 X 1 X 5X 7 X

The predicted value of Dvolume may be estimated by noting t h a t diffusion literature indicates that the diffusion rates of solutes in a given metal are n o t more than a few orders of magnitude different from that of the solvent self-diffusion. Data for volume self-diffusion and surface self-diffusion are given in Fig. 15, adapted from Gjostein [16]. It may

10 -8 10 -17 10 -36 10 -7 10 -34

10 -5 10 -i3 10 -27 10 -5 10 -26

The very wide differences in the listed P's at T H = 0.75 would seem to indicate t h a t the

60 TH = T / T M 1.0

-3

0.8

0.6 o

0.5 w

0.4 o

0.35 ,

-5

-10

\

- - ..

%o

-15

q

o.N -20

\ \ \ \ \

M.P. -25

J 1.0

la.4

1~8

21.2

21,6

310

TM/T Fig. 15. Rates o f self-diffusion in solid metals. Adapted from Gjostein [16].

potheses, we may consider the processes to be represented as shown in Fig. 16. By analogy with the volume diffusion case, and by consideration of atomic bonds which must be broken in forming and moving the surface adatoms, we should expect that the first monolayer of embrittler atoms should diffuse at rates more closely related to the surface selfdiffusion rate of the steel than to t h a t of the embrittler. In the second monolayer this should be less true, and by the 3rd, 4th or 5th layer -- that is, in multilayer surface diffusion -- the rate should be essentially that of surface selfdiffusion of the embrittler. Though to date the a m o u n t of experimental data on heterogenous surface diffusion is not large and not y e t completely consistent, the data available do support the above deductions; an excellent summary and discussion of these data are given by Bonzel [17]. This r a t e - the surface self-diffusion rate of the embrittler -as shown in Fig. 15, is related to Tm of the embrittler rather than the much higher T Fe , and the rate is in turn much higher than that of the atoms in the first layer. Thus, the prevailing rate is the fastest one -- t h a t of the embrittler surface self-diffusion -- and the layers with lower rates, including the first em-

be noted that at the temperatures of our experiments -- TFme/T ~- 3, where TFme = the melting temperature of Fe -- the volume selfdiffusion coefficient of Fe is ~ 10 -24 cm2/sec so that our: Dvolum e _-__ 1 0 - 2 4

+ 3

cm2/sec.

Comparison of this with Dexp eliminates the volume diffusion hypothesis for the embrittler transport mechanism. Some data given by Gjostein [16] for grain boundary diffusion in body-centered cubic metals indicate that for iron at TFme / T = 3 this value would be about 10 -14 which is also too low to support the concept of transport along the grain boundaries near the crack surface. Further corroboration may be noted in t h a t these diffusion hypotheses allow no way of rationalizing the constant TH at which embrittlement sets in for all five embrittlers, for these diffusion rates should be more closely related to the melting point of the steel rather than that of the embrittler (i.e., the melting point of the solvent rather than the pure solute). Turning then to the surface diffusion hy-

ADATOM

(o)

ADATOM

(J) Fig. 16. Two possible mechanisms of embrittler surface diffusion. (a) Diffusion o f embrittler atoms (dark) in a monolayer along the steel crack surface (light). Embrittler atoms move by heterogeneous surface diffusion over steel substrate atoms. (b) Multilayer movement o f embrittler on substrate by surface self-diffusion o f embrittler over its own surface. Moving embrittler atoms (dark) in top layer effectively "see" only other embrittler atoms. Lower atomic layers o f embrittler atoms formed by "waterfall" effect rather than diffusion on substrate (light).

61

brittler layer, are formed by a kind of "waterfall" effect in which atoms from the upper layers "spill over" and are "captured b y " ledges in the lower layers. On this basis, we see that the constant TH for onset of embrittlement with the various embrittlers is readily rationalized using this embrittler surface selfdiffusion hypothesis b u t n o t with single-layer surface diffusion, nor for that matter with second monolayer heterogeneous surface diffusion. Estimation of the diffusion rates confirm selection of embrittler surface self-diffusion as the probable active embrittler transport mechanism. Referring to Fig. 15, noting that for single-layer heterogeneous surface diffusion the appropriate temperature is T Fe/T ~- 3, and using the curve for body-centered cubic metals (Fe = b.c.c.), we find that D Fe ~- 10-12; if the rate of single-layer heterogeneous surface diffusion, D~ mb , is n o t much different from this, then: Dsemb ~_ 10-12 ~+3 cm2/sec.

This is 8 + 3 orders of magnitude lower than D~xp. For embrittler surface self-diffusion, however, T~ n o w refers to the melting point of the embrittler, and the appropriate T ~ / T for the temperatures at which D~xp was calculated is approximately 1.05, giving: Demb ___ 1 0 - 4 c m 2 / s e c , ss

that is, the same order of magnitude as D~xp (the agreement is u n d o u b t e d l y fortuitously good). We conclude that our experiments reported here support the concept that embrittler surface self-diffusion is the active mechanism of embrittler transport during the metal induced embrittlement of steel in the solid metal embrittlement range. It should be noted that our conclusion as to the transport mechanism predicts that the activation energies for embrittlement in this temperature regime [ T ~ / T (= 1/TH)] from 1.0 to 1.33) may well be very high, even though the diffusion rates are themselves very high, corresponding to the region in Fig. 15 where the surface diffusion curves turn up rapidly just below Tin. There are at present two sets of values in the literature which bear on this point. Iwata, Asayama and S a k a m o t o [6] presented data for Cd embrittlement of steel in the temperature range TCd/T from 1.14 to 0.94, giving an activation energy for time to failure in tensile delayed failure tests of

39,000 cals/mol and data which give a D of 10 -6 cm2/sec at TCd/T = 1.04. Fager and Spurr found for Cd embrittlement of steel [7] and for Cd embrittlement of titanium [8] an activation energy of 13,500 cals/mol in the range TCmd/T from 1.10 to 1.91 and a D of 2 × 10 -s cm2/sec at TCd/T = 1.49. These D values are within two orders of magnitude of those deduced for Cd surface self-diffusion from Fig. 15. The activation energy for such self-diffusion is, from Fig. 15, only 7700 cals/mol, suggesting that the embrittlement experiments do indeed fall in the regime where the surface diffusion curves turn rapidly upward. We are n o w carrying o u t tensile delayed failure tests on the five steel embrittlement systems of present concern in the appropriate temperature range to test the embrittler surface selfdiffusion hypothesis further; these results will be reported shortly. In addition, experiments on fracture mechanics type specimens are also n o w being started in order to measure the propagation rates more directly.

ACKNOWLEDGEMENTS

The authors are pleased to acknowledge the financial support for this work of Project THEMIS under Contract # DAAA-25-69C0608 and of the National Science Foundation under Grants # GH-34664 and GH41636.

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