Some aspects of interfacial structures from intergranular and surface experiments in the same material

Surface Science 162 (1985) 519-529 North-Holland. Amsterdam

519

S O M E A S P E C T S OF INTERFACIAL S T R U C T U R E S F R O M I N T E R G R A N U L A R AND SURFACE E X P E R I M E N T S IN T H E S A M E MATERIAL Jean B E R N A R D I N I , Franqoise CABANI~ and Jean CABANa. l.ahoratoire de MOtallurgie. UA 443, Facultb des Sctences et Techniques de Saint JOrbme, Rue ttenri Poincar~. F-13307 Marseille (.¥,dex 13. France

Received 1 April 1985; accepted for publication 4 April 1985

This paper deals with grain boundary structural m~lifications deduced on the one hand from both grain boundary diffusion and segregation measurements and on the other hand from surface segregation experiments performed on the same alloys. Similar behaviour is found for surfaces and grain boundaries in the various binary and ternary solid solutions studied.

I. Introduction The direct observation of grain boundaries (GB) is still difficult to perform in spite of recent improvement in electronic microscopy and localized analysis. Nevertheless, structural information is available from experiments related to atom movement in the grain boundaries. In this paper we propose to discuss some experimental results on GB diffusion and segregation in comparison with observations performed on the surface for the same material under quite similar conditions.

2. Some typical aspects of surfaces and grain boundaries The present state of our knowledge on interfaces allows us to draw up a framework for a comparison of surfaces with grain boundaries: - Interfaces are extended defects producing a local excess of energy in comparison with the bulk; but a specific regular array for the interfacial atoms tends to minimize this excess of energy. From this point of view, GB sites are more similar to bulk sites than surface sites when either energy or number of nearest neighbours is examined. - The notion of surface or of GB includes a great deal of widely disparate surfaces and grain boundaries; it is worth noting that some surfaces (e.g. (111)cfc) can be described in terms of only one kind of site while several kinds 0039-6028/85/$03.30 ~ Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

520

J. B e r n a r d m t et al. / Interracial gtructure~" m s a m e rnatertal

of site are always necessary to describe even the simplest GB (e.g. a twin GB). The crystallographic structure of a surface can be deduced from observations of the surface itself (e.g. LEED); for a GB, only the relative orientation of the crystals is usually well known and it is almost impossible, except in special cases, to check that the actual GB on an atomic scale is in fact the predicted interface. Simple surfaces (e.g. (111) fcc) have given the greatest part and the most complete sets of experimental data [1]; conversely, "general" random grain boundaries have been the more often examined as they are generally more sensitive to segregation and to embrittlement than simple grain boundaries [2]. - As the thickness of an interface is not more than a few atomic rows, the number of atoms concerned in an interface is always limited; hence the maximum amount of segregated matter is of the same order of magnitude as a monolayer, atoms of the solvent or of other solutes can be rejected from the interface when one solute accumulates as the interface, a change in the equilibrium structure of the interface must be expected when the segregated amount is large. - It is generally accepted that the tendency for one solute to segregate in a binary solid solution is the same when either the surfaces or the graiq boundaries are considered. However, in the presence of more than one solute, the surface may not behave in the same way as the grain boundaries [3]. Given this situation, it appears that simple grain boundaries resemble complex surfaces rather than simple ones: it is then difficult to look for precise comparison of surface and GB atomic structure from experiments: this is especially true under segregation conditions when simple surfaces and general grain boundaries are involved. Nevertheless it is possible to obtain some less precise but realistic relationships between GB and surfaces when interfaces are studied by means of surface techniques (AES. LEED) associated with the radiotracer technique. Thus GB diffusion which mainly involves atoms on the sites of highest energy is a precious tool to localize the segregated atoms among the GB sites of various energy levels.

3. Main characteristics of materials and techniques 3.1. Materials

We have studied binary A(B) and ternary A(B, C) solid solutions in which B (or C) solubility is very low. Thus the correlation between the inverse of solid solubility and segregation [4] makes possible a prediction of a large solute segregation in the interfaces. The nature of the phases in the phase diagrams [5] reveals a general tendency to form either A - A and B -B bonds; their surface tension ), may be taken as a significant quantity when interracial phenomena

J. Bernardini et al. / I n t e r f a c t a l structures in same matertal

521

are involved. So, according to these criteria, our materials can be listed in four classes. Alloys of the two first classes are similar by the absence of intermediate phases in the phase diagram which shows a tendency to form A - A and B- B bonds rather than A - B bonds. They are different by the value of the ratio " / , o l u t c/'~,.o1~, e n t :

(a) Ag(Pb) solid solution: Y~,,lutJY~,,l,t.nt< 1. A solute segregation can be predicted in any way. (b) Ag(Ni) and Cu(Fe) solid solutions: Y,,,lut~./Y~,,i,~,,> 1. This case is not favourable to solute segregation according to the usual thermodynamical prediction [6]. The presence of intermediate phases characterizes the two last classes: it shows that bonds between different atoms are stronger than bonds between like atoms. (c) Fe(Sn), Fe(Sb), Ag(S), Cu(S), Ni(S) solid solutions: The y values are always smaller for the intermediate phases than for the pure solvent. When both elements are metals, y is smaller for the pure solute than for the pure solvent: this should be favourable to solute segregation. (d) Fe( Ni, Sb). Cu( Fe, S), Cu( Ni, S), Ag( Ni, S)solidsolutions: One solute is a transition metal and the other one is a non-metallic element. Moreover. strong B C attraction may be deduced from the existence of B,,,C,, compounds. It worth noting that the pre-eminence of A - A and B B bonds in alloys (a) and (b) and of A - B bonds in alloys (c) is very marked since it gives rise to a very narrow solubility range, even at high temperature, when solid and liquid phases are in equilibrium. From this point of view, it is quite different from cases such as Au(Ni) or Au(Cu) alloys in which either clustering or order only occurs at relatively low temperature within a solid solution range. As a consequence every model which implies ideal or regular solutions [7] in the bulk and in the interfaces and interactions in the interfaces similar to those in the 3D solid solutions are unlikely to provide an adequate mean of understanding segregation in the above mentioned alloys (a)-(d).

3.2. Experirnental procedure The experimental procedure used for preparation of materials and for segregation and diffusion studies have been described in detail elsewhere, for each alloys [8]; only some major points are summarized below. The initial materials were always ultra-pure metals (99.999%). The heat treatments were carried out in pure gases (N60 hydrogen or argon) in order to avoid evaporation and oxidation. In all binary or ternary alloys, the solute content was within the solubility range at the temperature selected for diffusion and segregation measurements. The diffusion anneal was performed in a previously equilibrated solid solution. The alloys were generally in the form of poly-

522

J. Bernurdmt et al. / Interracial structures m s a m e material

crystals; some bicrystals were also used. Diffusion data were obtained by means of radiotracers; segregation data were obtained either by means of scanning AES on samples fractured "in situ" (Fe alloys) or by radiochemical techniques (metal-sulfur and A g - N i alloys). Notice that radiotracers and A E S - L E E D techniques, which were both used for observation on the surfaces (metal-sulfur, Ag(Ni), Ag(Ni, S)) complement each other: the whole a m o u n t of segregated atoms is determined with radiotracer without any assumption for calibration while the strictly superficial atoms play a main role in the intensity of an AES peak, especially when its energy is low.

4. Experimental results on grain boundary diffusion and segregation We were able to determine two quantities in the grain boundaries of one sample: P is the parameter characteristic of mass transport and n the n u m b e r of atoms segregated per unit area. Indeed P is related to the diffusion coefficient D(;B: P = 6D(mCcm/('b, where 6 is the GB width, Ccm and (7 h (expressed in atoms per unit volume) are the concentrations of the diffusion species in GB, and in the bulk near to the interface; n can be connected to a mean GB concentration ~ ( ~ = n/6 and to a segregation ratio a = ~(~/C,. It is then possible to calculate D(;~ from P and n data to the extent that ~(;H = (~;f~. The main results are presented here in reference to the classification of our materials: (a) Ag(Pb): With the gradual addition of Pb to silver, the silver diffusivity decreases at low concentration and then reaches a constant value (Table 1 ). (b) Ag( Ni): The presence of Ni in silver does not affect the silver diffusivity at 462 and 517°C; yet a large GB segregation coefficient ( a = 3 0 0 0 ) was measured at 517°C (table 1). (c) Fe(Sb): P in relation to Fe and Sb was measured at 550°C together

"['able I Silver grain b o u n d a r y diffusivity in Ag(Pb) and Ag(Ni) alloys Ag(Ni), T = 5 1 7 ° ( `

Ag(Pb), T = 2 8 1 ° ( ` Bulk lead concentration (at%)

101"~ p a 8 (cm ~ s ~)

Bulk nickel concentration (10 4 a t % )

101,s p a ~ (cm 3 s ~)

0 0.04 0.09 0.23 I).5

10 4.3 3.2 3.5 3.4

0 41 82

1.7 1.5 1.9

J. Bernardmt et al. / lnterfacial structures in s a m e material

0.02

o.o5

0.1

02

05

523

1 tat%)

F i g l. Segregation isotherms of Sb in solid solutions Fe(Sh) (a), Fe(Ni 1 ate,, Sh) (h) and Fe(Ni 2 a t e , Sb) (c) at 550°C.

with the segregated amount [9], fig. 1. It is thus possible to draw P variations with GB antimony concentration (figs. 2a-2c). The decrease in P, as the Sb concentration increases, is quite the same for the solvent as for the solute Sb. The same behaviour has been observed for Fe(Sn) alloys [10]. Ag(S) and Cu(S): It is possible to obtain DGB coefficients for sulfur in Ag(S) as a function of temperature without any variation of the GB concentration and to determine the activation energy QC;B [11]. The results are reported in fig. 3 together with various bulk and GB diffusion coefficients in silver. It is noticeable that sulfur diffusion in the grain boundaries is slower than silver self-diffusion while the reverse is true for bulk diffusion. The same results were obtained in copper [12]. Ni(S): Some experiments have been performed under the same conditions in two Ni grain boundaries of different (100) tilt angle but with the same energy level. The amount of segregated sulfur was almost equal to the maximum amount adsorbed on a (100) surface for one bicrystal; it was about 3 times lower in the other one. Simultaneously, the D(~, coefficient, always relatively small, decreases as the sulfur segregation increases [13]. (d)Fe(Ni, Sb): The addition of Ni of one Fe(Sb) alloys produced an increase in the amount of GB antimony (fig. 1) which implies N i - S b cosegregation. At the same time the decrease in pr~ as the GB antimony concentration increases, is identical to the evolution of pV~ and pS~ (fig. 2). However, for C sb > 5 at%, P ~ values are lower in ternary alloys than in binary alloys whereas the opposite is observed for pNi and pSb. This difference disappears when dealing with calculated De; B as a function of c(S~ [14].

J. B e r n a r d m l et al. / Interracial structure~ in s a m e m a t e r i a l

524

o. ~t~,

z

i

~ 8

j~.O~ ~1 .c

I

i (t..'F, £wo9t.Ol. ),N d

..c L"

..z 0

.z

,"5 [.a. ...z

rOt) q~d o

~n

e-

~a

v: z

^

> ,7

- 0

,--i

J. B e r n a r d i m el al.

/lnterfactal

structures tn s a m e m a t e r i a l

525

1

g ,v /T Fig. 3. Temperature dependence of bulk and GB diffusion coefficients in silver.

Cu(Ni, S); Cu(Fe, S): An addition of Ni or Fe in one Cu(S) alloys produces an increase in the amount of segregated sulfur and a decrease in the GB diffusion coefficient (fig. 4). However the variations induced by Ni

CM

Cu(Ni),.,

80

60

T.- 6 0 0

i:

" C

40

20 -"

2°

-'1"4T- . . . .

6'

S amount inG3.

(at x.) Fig. 4. GB sulfur diffusion coefficients as a function of the sulfur G B amount in Cu(Ni, S) and Cu( Fe, S).

526

J. Bernardim et aL / I n t e r f a c t a l structure~ in .~ame material

a d d i t i o n are very small in c o m p a r i s o n with those induced by Fe a d d i t i o n ; this is consistent with a F e - S cosegregation [ 15,16].

5.

Discussion

Some salient features emerge from all the above results: - the mobility is generally very low in G B when segregation occurs: the diffusion p a r a m e t e r s and the diffusion coefficients rapidly decrease for low c o n c e n t r a t i o n and then remain constant as the c o n c e n t r a t i o n increases: the effect of segregation on diffusion is more m a r k e d in alloys (c) and (d) than in alloys (a) and (b). It was shown in a previous p a p e r that no clear correlation a p p e a r e d between the solvent diffusion b e h a v i o u r and the s o l v e n t - s o l u t e b i n d i n g energy [17]. Conversely the following a s s u m p t i o n s were consistent with the e x p e r i m e n t s

[181: G B behaves in effect as a 2D phase where the nature of the main b o n d s is consistent with the bulk tendency to f o r m a t i o n of either clusters or c o m pounds: - Solute diffusion can be s u b m i t t e d on the one h a n d to an " a c c u m u l a t i o n effect" which tends to increase the transfer of m a t t e r as the G B c o n c e n t r a t i o n C(;)) increases and on the other hand to a " t r a p p i n g effect" which tends to decrease the p r o b a b i l i t y of a segregated a t o m t o j u m p out of one site since this site is favourable for segregation. F o r the solvent, a " c o n c e n t r a t i o n effect" a n d a "'trapping effect" due to s o l u t e - s o l v e n t attraction can also occur; but they both tend to decrease the solvent diffusion. W e have also to take into account well known observations on the surfaces: Segration on metals leads to 2 D phases with m e t a l - s u l f u r b o n d s stronger than those in the c o r r e s p o n d i n g 3D sulfide [19]. A c o m p a c t 2D phase from which the solvent is c o m p l e t e l y rejected can form on Ag(Pb) simple surfaces [20]. - A " s a n d w i c h " structure forms on Ag(Ni) surfaces where the solvent a t o m s form the strictly superficial layer while the segregated solute a t o m s are mainly in the atomic layer i m m e d i a t e l y below [21]. In the first case ( m e t a l - s u l f u r ) , the m a x i m u m a m o u n t of segregated sulfur is almost i n d e p e n d e n t of the orientation and c o r r e s p o n d s to a b o u t I sulfur a t o m for 2 metal atoms; in the other cases, the m a x i m u m a m o u n t can c o r r e s p o n d to a very close c o m p a c t a r r a y of solute atoms, a n d even a b o u t 1.5 m o n o l a y e r in Ag(Ni), on simple surfaces but it remains quite less than one m o n o l a y e r on c o m p l e x surfaces and polycrystals [22]. Given that we c a n n o t reach from e x p e r i m e n t s a precise description for the structure of GB, we represent a G B as a schematic d i s t r i b u t i o n of a t o m s on

527

J. Bernardmt et aL / Interracial structures tn same material

I

III s.,..

.

.

~

- :::::::::::::::::::::

~c'N~&Igdd ~''~'

II , ~

~I' .

I' .

.

IV

)4)00( .

.

-

t()tL)i~.)--Sur|ice .

.

.

.

__,£1£~11

• ~I'I(X~(IJ(2~J_LLt.,

[IIIXIXXXXIIIXXI

.~.x.~.

......

555;V/4

Fig. 5. Schematic representation of a pure grain boundary (1), and "'segregated" interfaces (11, II1, IV).

two types of interracial sites (high energy and low energy sites) between two compact arrays of bulk atoms (fig. 51). In alh)ys of class (c), the solute diffusion is largely slowed down when c~'~;B increases. The solvent diffusion undergoes a similar change. The " t r a p p i n g effect" is then very strong in comparison with the "accumulation effect". This is consistent with the shape of the GB segregation isotherms which implies for Fe(Sb) and Fe(Ni, Sb), as for Ag(S) [11] strong attractions in the segregated phase and a phase transformation in the interface. Such a segregated GB can then be represented by a regular array of solute and solvent atoms located in the high energy sites (fig. 5II). Another similarity between surfaces and GB is that equivalent maximum amounts of segregated sulfur are found on these interfaces. It is very likely that a F e - S b 2D compound also forms on the surface, although F e - S b attractions are weaker than metal-sulfur attractions. In ternary Fe( Ni, Sb) alloys (class d) there is a particularly large decrease in pF¢ in comparison with its decrease in binary alloys (fig. 2a): this leads to the conclusion that cosegregation produces a regular arrays of Ni and Sb atoms in the high energy sites, the solvent atoms being rejected from these sites. This interpretation is consistent with all the above mentioned results concerning sulfur segregation. It is also in fair agreement with observations of cosegregation performed on the surfaces for the same alloys. We may thus conclude that for all these alloys, as for Cu(Fe, S) [16], cosegregation can be expressed by a shift of the non-metallic solute segregation isotherm towards low bulk concentrations. The degree of shift must be related to both the strength of the bond between the metallic and the non-metallic solute and to the chemical activity of the metallic solute in the ternary solid solution. The small effect of Ni on sulfur diffusion and segregation in Cu(Ni, S) alloys, despite N i - S cosegregation, can be related to the complete miscibility of Ni and Cu and their immediate vicinity in the periodic chart. Thus partial or

528

J. B e r n a r d m i et al. / Interfa('ial structure~" m s a m e m a t e r t a l

complete rejection of Cu from the grain boundaries should not greatly affect the segregated amount and the GB diffusion of sulfur. In alloys of class (a) the decrease in P for the solvent cannot be due to solute-solvent attraction since the main tendency in these systems is to form clusters; it should be attributed to the presence of solute atoms in high energy sites in the grain boundaries (fig. 5III.). This would explain why Ag diffusivity in Ag(Pb) is only decreased by a factor of 3 whereas Fe diffusivity in Fe(Sb) is decreased by a factor of 10 for the same ratio T / T , , = 0.45. The maximum Pb concentration can be estimated at 0.6 Pb monolayer with the assumption that the Pag decrease is only a consequence of the decrease of (C(m/Ch)A~. The same maximum amount of lead has becn found on a complex surface [20]. In alloys of class (b). solute segregation does not change the solvent diffusion; this implies that silver atoms occupy the high energy sites in GB. the segregated Ni atoms being most probably in the low energy sites (fig. 5IV). This confirms the "sandwich" model proposed for surface segregation in these alloys [23,24].

6. Conclusion

Even with bicrystals, it was not possible to obtain for GB the same crystallographic data obtained for surfaces. However, a discussion of experimental results on GB diffusion and segregation, associated with observations performed on the surfaces for the same solid solutions, allows us to have an overview (summarized in fig. 5) of the main structural effects of segregation in the interfaces. In all above-mentioned solid solutions (a)-(d) the sparingly soluble element always accumulates in the interfaces; the segregation of one solute produces a similar structural change on surfaces and in grain boundaries in ternary alloys as well as in binary alloys. Moreover the maximum segregated amount is generally of the same order for grain boundaries as for a complex or polycrystalline surface. Cosegregation, if any, goes hand in hand with a noticeable change in the GB structure, from the binary to the ternary solid solution. This is particularly true for Ag(Ni) and Ag(Ni, S) or Cu(S) and Cu(Fe, S) solid solutions: the solvent atoms which are in the high energy sites in binary interfaces are rejected toward GB low energy sites in ternary alloys, while the reverse is observed for the segregated metallic solute. It is worth noting that this conclusion may have very important practical consequences: cosegregation of a residual impurity with an alloying element can completely modify the structures and thus the properties of the interfaces.

J. Bernardim et al. / Inter[acial structures tn same matersal

529

References [1] J. PAnard, Ed., Adsorption on Metal Surfaces - An Integrated Approach (Elsevier, Amsterdam, 1983). [2] E.D. Hondros, Sci. Progr. (Oxford) 68 (1982) 35. [3] C. Molinari and J.C. Joud, in: Physical Chemistry of Solid States, Ed. P. Lacombe (Elsevier, Amsterdam, 1984) p. 151. [4] E.D. }tondros and M.P. Seah, Intern. Rev. 222 (1977) 262. [5] M. Hansen, Constitution of Binary Alloys (McGraw-Hill, New York, 1958). [6] P. Wynblatt and R.C. Ku. Surface Sci. 65 (1977) 511. [7] F.F. Abraham and C.R. Brundle, J. Vacuum Sci. Tcchnol. 18 (1981) 506. [8] F. Caban~ and J. Bernardini, J. Physique 43 (1982) C6-163. [9] P. Gas, M. G u t t m a n and J. Bernardini, Acta Met. 30 (1982) 1309. [lO] J. Bernardini, P. Gas, E.D. ttondros and M. Seah, Proc. Roy. Soc. (London) A379 (1982) 159. [11] B. Aufray, F. Caban~-Brouty and J. Caban6, Acta Met. 27 (1979) 1849. [12] F. Moya and G.E. Moya-Gontier. J. Physique 36 (1975) C4-157. [13] M. Pierantoni, B. Aufray and F. Cabane. in: Proc. Intern. Conf. on the Structure and Properties of Internal Surfaces, Irsce, Fed. Rep. of Germany, 1984 (J. Phys., in press). [14] P. Gas, S. Poize and J. Bernardini, Acta Met., in press. [15] B. Aufray, Tht~se, Marseille (1982), [16] A. Pinea, B. Aufray, F. Cabane-Brouty and J. Cabanc, Acta Met. 31 (1983) 1047. [17] J. Cabana, in: Physical Chemistry of Solid States, Ed. P. Lacombe (Elsevier, Amsterdam, 1984) p. 201. [18] J. Bernardini and E. Caban/:, in: Proc. Intern. Conf. on the Structure and Propcrties of Internal Interfaces. Ir~,ee. Fed. Rep. of Germany, 1984 (J. Phys., in press). [19] J. Oudar, Mater. So. Eng. 42 (1980) 101. [20] A. Rolland, J. Bernardini and M.G. Barthes, Surface Sci. 143 (1984)579. [21] A. Rolland and F. Caban& Nouveau J. ('him. 8 (1984) 485. [22] A. Rolland and J. Bernardini. Scripta Met., in press. [23] A. Rolland, B. Aufray, F. Cabane-Brouty and J. Cabana, Compt. Rend. (Paris) II. 292 (1981)

1477. [24] A. Rolland and B. Aufray, Surface Sci. 162 (1985) 530.