Measurement of the grain boundary diffusion of In in Ni bicrystals by the SIMS technique

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MEASUREMENT OF THE GRAIN BOUNDARY DIFFUSION OF In IN Ni BICRYSTALS BY THE SIMS TECHNIQUE W. GUST’, M. B. HINTZ’,

A. LODDING3,

H. ODELILY? and B. PREDEL’

‘Institut fur Metallkunde der Universitlt Stuttgart and Max-Planck-Institut fiir Metallforschung. Institut fur Werkstoffwissenschaften. D-7000 Stuttgart 1. Seestrasse 92. West Germany “Michigan Technological University. Department of Metaliur~ical Engineering, Houghton. Michigan 49931. U.S.A. and ‘Materialcentrum. Chalmers Tekniska Hogskola. S-41296 Gothenburg. Sweden (Receiced S June 1981)

Abstract-The diffusivity of indium as an impurity along grain boundaries in nickel bicrystals of known orientation has been measured using secondary ion mass spectrometry ISIMS). The values obtained are very large and comparable in magmtude to those observed by Hillert and Purdy. and by Smidoda. Gottschalk and Gleiter for migrating interfaces. These results along with other data from the literature lead to a different interpretation of the migrating interface results than that presented by original authors. It is suggested that the high observed interfacial diffusivities may be the cause rather than the result of the observed interracial migration. R&utn&--Nous avons mesure par spectrometrie de masse d’ions secondaires. la diffusivite intergranufaire de I’impurete In dam des bicristaux de nickel d’orientation connue. Les valeurs obtenues sont tres grandes: elles sont comparables a celles observks par Hillert et Purdy. et par Smidoda, Gottschalk et Gleiter pour des interfaces en tours de migration. Ces resultats et d’autres don&es de la litterature conduisent B une interpretation des rtsultats concernant les interfaces en cours de migration. differente de celle avancee par les auteurs. Nous pensons que les diffusivites interfaciaies Clevees qui ont CtCobservees pourraient itre la cause plutot que I’effet de la migration de f’interface. Zuaammenfaastmg--Mittels Sekundlrionen-Massenspektrometrie (SIMS) wird das Korngrenzendiffusionsvermogen von In in definierten Ni-Zweikristallen gemessen. Die ermittelten Werte sind sehr grolj. Sie sind in der GroBenordnung vergleichbar mit den von Hillert und Purdy. und Smidoda. Gottschalk und Gleiter fir wandernde Korn-bzw. Phasengrenzen ermittehen Werten. Diese Ergebnisse ftihren zusammen mit anderen Literaturangaben zu einer neuen Interpretation des Diffusionsverm~gens wandernder Kern- und Phasengrenzen. die von der anderer Autoren abweicht. Es wird dargelegt. daB das hohe. experimentell ermittehe Diffusionsverm~gen der Grund und nicht das Ergebnis der Korn- bzw. Phasengrenzenwanderung ist,

1. IlQRODUCTION Recent experimental work [l. 21 has indicated that the properties of moving grain boundaries may differ significantly from those of stationary grain boundariest. Hillert and Purdy’s [l] study of chemically induced grain boundary motion during the zincification of poly~ystaIline iron. and Smidoda er 01,‘s[2] study of the recrystallization process in deformed. supersaturated Ni-Cr-Al solid solutions resulted in values for the grain boundary diffusivity that were two to four orders of magnitude larger than those previously measured for stationary grain boundaries in similar systems. This was interpreted in both papers as indieating that the properties of moving grain boundaries differ from those of stationary boundaries.

In the present work the diffusivity of indium along stationary grain boundaries in nickel bicrystais of defined orientation has been determined and is found to be comparable in magnitude to the moving boundary diffusivities referred to previously. In what follows, arguments are developed which relate the present results, the recent moving boundary results and other expe~menta1 evidence in a plausible manner, 2. EXPERIMBNTAL PROCEDURE Techniques used in the preparation of and subsequent diffusion measurements on defined bicrystais have been described in some detail elsewhere [4]. Discussion of these topics in the present paper is therefore relatively brief. 2.1 Specimen preparation

+ Borrowing

Gleiter’s terminology [3], ‘grain boundaries’ in the present context will be taken to mean highangle grain boundaries unless otherwise specified.

Symmetric tilt bicrystals were prepared from single crystal nickel rods of better than 99.99*, purity.

75

76

GUST rr al.: ON GRAIN BOUNDARY DIFFUSION

Their nominal orientation was 39’ about (1 IO) as measured between j 110; type planes (39’ (110) :llO;). The reasons underlying the choice of this particular orientation are discussed in section 4 of this paper. Briefly, specimen preparation technique consists ot cutting suitably oriented discs of 2 mm nominal thickness from the approximately 12 mm diameter single crystal rods using spark erosive methods; wet grinding, lapping and electropolishing the faces of the discs to produce a flat, essentially strain-free surface, and then sintering the prepared surfaces together. For the present specimens the sintering was performed in an ultra-high vacuum environment (pressure p < low6 Pa) for 16-18 h at about 1450 K. After sintering the bicrystals were cut in half using a low speed diamond wafering saw (Buehler ‘Isomet’) such that the surfaces formed during cutting were per~ndicular to the (110) tilt axis. These surfaces were wet ground on 600 grit silicon carbide paper and then lapped on cast iron surface piates using 6jtrn diamond paste to improve surface flatness. The deformed and contaminated surface layers were removed by electropolishing for 1.5s at a 20 volts potential in an electrolyte consisting of 3OOml of 909; HNOs and 600 ml of absolute methanol at 260 K. After optical microscopic inspection. the as-polished bicrystals high vacuum were thermally etched under (p z 2 x 10e5 Pa) at approximately 1030 K for 1.5 h to reveal the grain boundary, and again examined using optical and. in some cases, scanning electron microscopy. Specimens exhibiting any evidence of recrystallization, grain boundary porosity, or other imperfections were discarded. The remaining bicrystals were again wet ground, lapped, and electropolished to prepare the surfaces for acceptance of the indium diffusant film. The indium was deposited on these surfaces by vacuum evaporation. The amount of indium applied was measured during the deposition process with a quartz crystal film thickness monitor. 70~g/cm2 f 209,; was applied to all of the specimens: based on the density of crystalline indium, this results in films of about 1OOnm thickness. performed in a rarified Annealing was (p 2 10-l Pa) flowing atmosphere of high purity argon. The temperature of the furnace was held constant to within i: 1 K. Absolute accuracy of the temperature measurements was + 2 K or better. Since the surfaces upon which the In was deposited were perpendicular to the (IlO} tilt axis- the diffusional flux was parallel to this axis. Previous work by others [5,6] on oriented tilt boundaries of various orientations has shown that the diffusivity is generally higher along a direction parallel to the tilt axis than a direction perpendicular to it. 2.2 Meusurement The impurity concentration distributions in the specimens after annealing were determined using

secondary ion mass spectrometry (SIMS). The SIMS technique in general is reviewed in Ref. [7,8]; some recent examples of its application to impurity diffusion measurements can be found in Ref. [9-l 1-J. The measurable penetration depth of the indium along the grain boundaries was of the order of hundreds of pm for the specimens examined here. This is an order of magnitude or more larger than the distances over which info~ation can be conveniently obtained by direct in-depth profiling with SIMS. Therefore. with the help of a precision goniometer, a small-angle bevel (typically between 2 and 4 (+_O.l) degrees) was ground, lapped and electropolished into the surfaces upon which the diffusant had been applied. This exposed the edge of the grain boundary plane at depths suitable for analysis. The measurements were performed by a Cameca IMS-300 instrument equipped with a double focusing electrostatic sector. 5.5 keV 0; primary ions were employed. Measurements were taken in a stepwise fashion along the exposed edge of the grain boundary on the bevelled surface. This is shown schematically in Fig. 9 of Ref. [4]. As will be seen, the data obtained at each location can be reduced to the equivalent of one section in a serial sectioning measurement. The primary ion beam was scanned over an approximately 3OOpm on side square area at each location. The area of analysis, however, was restricted to a 1OOpm diameter circle by an aperture in the image plane of the secondary ion extraction lens system, minimizing the possibility of anomalous contributions to the signal from contaminants that may be present near the crater rim where the sputtering rate is low [ll]. The mirror-symmetric nature of the bicrystals with respect to the boundary plane made it possible to place the specimens in the instrument such that the primary ion beam incidence angle was crystallo~aphically equivalent f$r both crystailites. Orientation induced preferential sputtering [E] of the crystallites was thus minimized. By using the Cameca instrument in the ion microscope mode, it is possible to ascertain that the grain boundary is centered in the analyzed area at each location. At the end of each measurement the scan generator was stopped for a few seconds, allowing the focused primary beam to make a small deep hole which marked the center of the analyzed area. The secondary ion signals of four ionic species were recorded at each location. These were Ni;,; In:%,; and (Ni2);ib_ About 30 data points were In;,,; obtained for each ionic species at each location on the bevelled surface. The total depth of the sputtered craters at each location was of the order of 1 pm. Since this distance is very small in comparison with the total measurable penetration depth, the impurity concentration remains effectively constant for any given location. Possible near surface enhancement and contamination effects are therefore easily separated from the steady state ionic signal. Following SIMS analysis, grain and boundary

GUST er (II.: ON GRAIN BOUNDARY

orientation relationships were determined via the back reflection Laue method. These relationships are schematically described in Fig. 1. The single crystal from which the bicrystal is to be constructed is positioned in an orthogonal coordinate system such that its [OOl]. [llO] and [ilO] directions are initially parallel to the X. J’ and : axes, respectively. A stereographic triangle is shown. as it would appear if projected upon the x: plane (i.e. a t 110) standard projec-

DIFFUSION

77

Table I. Measured angular reiationshtps between the crystallites and grain boundary plane. & i unless specified.

tion. cf. Fig. I(a)). The crystal is then cut in half along the ST plane. Rotation of the crystal halves in opposing directions about the -_axis will result in the orientation relationships illustrated schematically by the standard triangles in Fig. l(b). Symmetric [ilO] (110) tilt misorientation corresponds to 0, = n2 = Y

’

11103

8,‘2. Subsequent rotations of @r and damabout the x axis and rl about the ,Vaxis (cf. Fig. I(c)) result in additional undesired components of inter~ystalline misorientation. During these orientation measurements. it became apparent that one half of specimen 2 had recrystallized during sintering: its measured orientation differs greatly from what was desired. Its orientation relationships may be visualized in an essentially identical manner, however. The only additional operation required is a 90’ rotation of the left crystal half about [ilO] immediately after cutting. The subsequem rotation of 0, about [ilO] will then be measured between the X; plane and the (001) planes of this crystal haIf. The measured 0, @and 1 values are listed in Table 1. Specimens 1 and 3 correspond reasonably well to the target value of 39 /ISO) - ;I 10;. with the preparation technique assuring that 0, = 82 and @, = @z to within 1 3. RESULTS Figure 2 is an optical micrograph (LM) showing the bevelled surface after SIMS analysis. Data typical of those obtained at each location are plotted in

fb)

X

i

f

4-y

Fig. 1. Schematic ihustratron of intercrystalline orientation relationships. la) An uncut single_crystal and corresponding stereographic triangle. tbt @ [l lO]-(l IO) symmetric tilt misorientation obtained by cutting and rotation about ;. tct Other undesired misorientation components caused by subsequent rotation about x and ,r.

Fig. 2. A portion of the bevelled surface on specimen I. The edge of the step produced by electropohshing is seen at the far left. The first 7 craters formed during SIMS analysis are plainly visible. Their depth in the center where ‘marked’ by the focused ion beam is of the order of 3 Mm.

GUST er al.: ON GRAIN BOUNDARY DIFFUSION

78

Fig. 3. In order to cancel small differences in instrument focusing between one location and the next, each of the background corrected In;ls data points was normalized to its corresponding (Ni2):ib matrix signal. For the present case of a low solute concentration in an otherwise nominally pure matrix the mean of these normalized points, I/1M, may be taken to be proportional to the average In concentration in the sputtered volume [9]. The corresponding distance normal to the original surface, X, for each rlr, value can be easily calculated from measurements of the electropolishing step height, the distance from the step to the center of the sputtered crater, and the known angle of the bevel. A semi-logarithmic plot of m vs x615for a given specimen (Fig. 4) should he linear if the x values are sufficiently large so that contributions to the impurity concentration by volume diffusion from the original surface can be neglected [ 131. In general, the rjr, values obtained at the analysis location nearest the original surface were somewhat larger than extrapolation of the otherwise fairly good straight line fit of the remaining data points indicated. Since the electropolishing step heights were of the order of 7-10 pm, it is possible that, at least at the higher annealing temperatures, this may be due to volume diffusion from the surface [IO]. Additionally, the step introduces a perturbation into the otherwise uniform electric field in front of the secondary ion extraction lens which may affect its focusing. Hence. data obtained near x = 0 were deleted from the data analysis.

. . . . . . . . . . . . ..*...............*. (Nizt;,

SDccimcn

3

10

20

30

N- 1, Fig. 3. Typical brhaviour of the measured matrix and Impurity secondary ion signal intensities. I, vs data acquisition cycle number IV. which is proportional to the sputtering time r, for one location on the bevelled surface.

Specimen 3

I

0

200

LOO 600 X6’5(wnP5)

600

1 1003

Fig. 4. A semi-logarithmic plot of n vs x6 ’ for specimen 3 containing a symmetric 39. (1 IO> - : I IO: tilt boundary.

Some comment is warranted regarding the background correction procedure. For the case of volume impurity diffusion measurements in single crystals with SIMS [9, IO] the usual procedure is to sputter through the impurity concentration profile to a depth where the impurity signal approaches a steady-state plateau and obtain several data points on this plateau. The mean of this plateau value is then subtracted from the impurity ion signal data to correct for contributions to the impurity signal from the matrix itself. The present study, the small number of m data points obtainable for each specimen (limited by the dimensions of the bicrystal and the size of the area measured at each location) and their scatter at the lower values of Iii, introduced some ambiguity into this procedure. A ‘plateau’ appeared to be approached, but was always a factor of 1.5-2 larger than the I/I, values obtained in the bulk of either crystallite away from the boundary plane. Although there is no evidence suggesting that either crystallite sputtered faster than the other, Fig. 2 indicites that the region in the immediate vicinity of the boundary sputtered somewhat more rapidly than the rest of the material, as may be expected. The effect appears stable with respect to time for the measurement times employed here since the impurity and matrix signals remain stable for the duration of each measurement after the near surface transients. It was therefore felt that background levels obtained from measurements including the boundary plane were more representative of the background contribution to the measured m values. Using these values for background correction resulted in a better fit of the least squares line (as indicated by the least squares correlation parameter) and a slope with an absolute value which was about 15 to 209, larger than obtained using the bulk background levels. There is, therefore. additional uncertainty associated with the calculated grain boundary diffusivities. but it should be noted that this error, if present, tends to make the calculated diffusivities smaller than their actual values.

79

GUST PI al.: ON GRAIN BOUNDARY DIFFUSION Table 2. Annealing time r. annealing temperature the determined sSDB values.

7: and

Specimen

r(h)

T(K)

s6DB (cm3is)

1

42.0 16.0 17.5

967 1007 1081

1.3 x lo-l3 1.8 x lo-I3 1.0 x lo-‘2

2 3

The product of the grain boundary diffusion coefficient, Dg. the grain boundary thickness, 6. and the segregation factor [143, s. can be calculated from the slope of the least squares lines using the relationship presented by Le Claire [13]:

where (2) DV represents the volume diffusion coefficient which is taken from Ref. [IO], r is the diffusion annealing time, and i is the average diffusant concentration, which for the present measurements can be replaced by m. For serial sectioning measurements, equation (I) is applicable within experimental error to either Whipple’s [15] constant source concentration or Suzuoka’s [l6] instantaneous source solution of the diffusion equations developed for Fisher’s [ 173 grain boundary slab model [ 13.18-J. The last term in equation (1) is a weak function of /I and may be iteratively determined using equation (1) and the following relationship [16]: ? In E -12

?I(@

05

1

=

-

0.72 ~“~“““.

diffusion in polycrystalline nickel at comparable temperatures and 2 to 4 orders of magnitude larger than for ~if-diffusion along grain boundaries in defined nickel bicrystals of various other orientations [24,253. The lack of suitable grain boundary impurity diffusion data necessitates comparison of the present measurements with nickel grain boundary selfdiffusion data. This probably contributes to the observed differences, but it seems unlikely that a complete explanation of the present observations can be formulated in terms of this alone. Until such data are available, however. the possibility cannot be completely excluded. Although the present data represent only a small number of measurements, there appears to be no reason to doubt their validity. The measurements at 967 and 1081 K were performed on alternate halves of the same bicrystal. The measurement at 1007 K was performed on a different bicrystal of the same nominal starting orientation. although as mentioned previously the actual orientation changed during sintering. Unfortunately, to the knowledge of the present authors there is no available information which would allow one to reasonably estimate how the diffusivity of a 39’ (1 lo?_: 110; interface should compare with the asymmetric boundary in specimen 2. The data from the two different bicrystals are in reasonably good agreement, however, as the Arrhenius plot in Fig. 5 indicates. The high diffusivity is therefore not a property of a particular specimen. Also, assuming that an equation in the form sSD, =

b%). exp(-QIR T)

adequately describes the temperature dependence of the grain boundary diffusivity s6D,, where Q is the activation energy, (sfiD,j, the temperature indepen-

The values of s8DH calculated from equations (1) and (3) using the experimentally determined values of r’ ln (~)~~(~’ ‘f are presented in Table 2.

1081

1007

967

K

I

In in NI

4. DtSCUSSION The 39’ (110) - I1101 tilt orientation was chosen for this study as previous measurements [5,19] on grain boundary diffusion in bicrystals of approximately the same orientation indicated unusually high diffusivities. Independent studies of aqueous corrosion [20] and liquid metal induced grain boundary embrittlement 1211. both performed using oriented aluminium bicrystals. also indicate that grain boundaries of this orientation may exhibit additional ‘special’ properties. The results of the present study tend to substantiate the above observations, although the comparable magnitude of diffusivity measured for the asymmetric boundary in specimen 2 indicates that high diffusivity is not exclusively a property of the 39’ (1 IO) - 1110: orientation. The values presented in Table 2 are roughly 2 orders of magnitude larger than those obtained by others [22.23] for grain boundary self-

(4)

390 CllO> - {llOl Tilt

-

Boundary

ul

lo-13, 9 lo‘/

, 10 T f K-’ 1

11

Fig. 5. An Arrhenius plot of the experimentally determined s6D, products.

80

GUST et ul.: ON GRAIN BOUNDARY DIFFUSION

dent pre-exponential factor, R the universal gas constant, and T the absolute temperature, the values Q = 161 kJ/mole

(5)

(s&D& =i 5.4 x 10m5cm’/s

(6)

ratios for the relative energies of the three boundaries: c3:c2:g1 = 1.W:O.80:0.64

(8)

Kinetic measurements on the same specimen showed that the growth rate observed for the discontinuous precipitation reaction varied strongly from one can be calculated. The ratio of this activation energy boundary to the next and that the boundaries of to that for the diffusion of indium as an impurity in higher energy exhibited higher growth rates. As can the nickel lattice [lo] is about 0.64. This is a reasonbe seen in Fig. 6. the differenps in growth rates able value in light of data available for grain boundbetween the different boundaries are quite large, with ary and volume diffusion in other alloy systhe grain boundary of highest energy exhibiting a rate tems [ 19,261. It is therefore believed that the present of growth between 30 and 100 times greater than that results are not attributable to experimental error. observed for the lowest energy boundary. Grain boundary difTusivities comparable in magniIt should be emphasized that caution must be exertude to the present results have been recently obcised when comparing these previously mentioned served by others for the case of migrating interfaces. discontinuous precipitation results with experiments In particular, Hillert and Purdy (HP) Cl] have studied performed on grain boundaries since one is dealing the solute redistribution resutting from chemically with an interphase boundary as soon as the disconinduced grain boundary migration during the zincifitinuous precipitation reaction front begins to migrate. cation of iron, and Smidoda, Gottschalk and Gleiter However, for the case of discontinuous precipitation (SSG) [2] have examined precipitate formation during in the Ag-Cu system, the lamellae of solute depleted recrystallization of deformed supersaturated Nimatrix take on the orientation of the grain away from Cr-Al solid solutions. In both cases calculated grain which the reaction front migrates [30]. The orientaboundary ~ffusivities were 2 to 4 orders of ma~itu& tion relationships present at the original grain boundlarger than those previously measured for stationary ary are therefore preserved for the Ag-rich phases on grain boundar/es in similar systems. This was intereither side of the reaction front. Relating the reaction preted in both cases as an indication that the strucfront migration rates to the original grain boundary ture and properties of moving grain boundaries may energies therefore seems justifiable as a first approxidiffer considerably from those of stationary grain mation since better data are currently unavailable. boundaries. Since it appears to be generally accepted that growth The present results clearly demonstrate that high kinetics of the discontinuous precipitation reaction grain boundary diffusivities (i.e., shDB products of the are limited by the rate of diffusion along the interface order of lo- ’ 3 cm3/s and higher) are not exclusively a at the reaction front [31-333, the tricrystal results property of migrating interfaces. On the other hand, [29] indicate that high energy interfaces tend to also the articles by (HP), (SGG), the discussion of have high diffusivities. Speich’s [27] data for discontinuous precipitation in The preceeding is supported by the work of Hasson the Fe-Zn system by (HP), and examination of Gust, et al. [S]. They compare the results of theoretical Leininger and Predel’s (281 discontinuous precipigrain boundary energy calculations for symmetric tation data for the Ni-In system for which a s&‘& (110)-11101 tilt grain boundaries with experimental pr@duct of 4.2 x IO- I3 cm3is (at T = 1049 K) results, results for intergranular diffusion of zinc into aluminclearly indicate that a high diffusivity is likely to be ium bicrystals containing grain boundaries of the associated with moving interfaces. The (HP) and (SGG) interpretation therefore seems reasonable. In the following paragraphs, however, additional information will be presented that suggests a different interpretation. Gust or ~1.[29] have measured discontinuous precipitation reaction kinetics in Ag-6.2 at.:, Cu tricrystals that were grown from the melt. The relative grain boundary energies were determined by contact angle measurements at the junction of the three crystallites in the center of the specimen using the familiar relationship: 01 -J------csT--. sin HZ.3

=2 sin @,,,

=3

sin &.2

(7)

where CTiis the specific grain boundary energy of the ith boundary and Oj.k is the angle between grain boundaries j and k. The use of equation (7) and the measured contact angles resulted in the following

id/T

tK-‘1

Fig. 6. A comparison of the growth velocity, c. temperature dependence for the discontinuous precipitation in a melt grown Ag-Cu tricrystal starting at three different grain boundaries (1,2,3). After Ref. [ZS].

GUST cr ai.: ON GRAIN BOUNDARY DfFFUSiON

same type. In general. it can be seen that the grain boundaries of highest energy are also those exhibiting the highest diffusivities. although the authors are careful to point out that there is no one-to-one correspondence between these parameters and that the calculated energies alone are not sufficient to account for all the observed orientation dependent property changes. It is also interesting that their energy calculations (which neglect entropy effects and are therefore most applicable as T approaches 0 K) show a small cusp at approximately 39. (1 lo>-; 110; which corresponds to a (122) twin boundary. One may speculate that this cusp, if retained at the temperatures over which experimental measurements are performed, may tend to stabilize the grain boundary in a high energy configuration. There is experimental evidence which indicates that for the case of aluminium. such a cusp does indeed exist at temperatures approaching the meking point [34]. This may help account for experimental observations [5.19.21] which indicate that this particular orientation exhibits unusual properties. Using a total free energy minimum criterion, however, one would not expect these high energy. high diffusivity (HEHD) grain boundaries to be frequently observed in polycrystalline materials. In fact. work done by Warrington and Boon [353 indicates that twin boundaries (many of which would be expected to exhibit a low energy and low di~usiv~ty [S]) occur much more frequently in many polycrystalline metals than calculations based on a random probability of occurrence wbuld predict. Our interpretation of the data presented in this paper from the present study and the work of others is therefore the following: (1) The high diffusivities measured by (HP) and (SGG) in their migrating interface experiments are not exclusively properties of migrating interfaces since comparably large values have been measured for stationary grain boundaries. (2) Grain boundaries possessing a structure which results in a high diffusivity may be frequently expected to also have a high specific interfacial energy. (HEHD) boundaries are therefore not expected to make up a large fraction of the interfaces found in a polycrystalline material. When (HEHD) boundaries do occur. however, they wili dominate in reactions involving the redistribution of a solute aiong a migrating interface. (HEHD) boundaries are therefore the ones that are frequently observed in measurements of such reactions. (3) Consequently, it is not surprising that large discrepancies are observed when comparing diffusion measurements made on polycrystals or low-angle bicrystal orientations with values calculated from moving boundary experiments. This interpretation does not directly refute the assumption of (HP) and (SGG) that a change in the structure of the migrating interfaces is the cause of the

81

observed enhanced interfacial diffusivity. Recent speculative work [36] does suggest that under certain conditions the vacancy concentration near a migrating interface may be appreciably larger than for the case of one that is stationary: this could indeed lead to enhanced interfacial diffusivity. However. it is suggested that the observations of (HP) and {SGG) may also be explained without the assumption of an interfacial structural change being necessary if the present evidence is considered. The present results indicate that interfaces possessing particular structures will exhibit high diffusivities irrespective of whether they are migrating or not. and that these interfaces are likely to play a larger role in reactions involving solute redistribution by interfacial diffusion than considerations of their frequency of occurrence would suggest. The present interpretation is also compatible with recent observations which indicate that the ‘breakaway’ behaviour observed during grain boundar! migration in otherwise pure materials containing controlled amounts of impurities may result from a transformation of the grain boundary structure at a particular temperature [37]. Experimental evidence has been presented by Gleiter [3] which indicates that such structural transformations occur in stationary grain boundaries as well. It is obvious. however. that additional experimental and thzorerical work is necessary if the phenomena discussed in the preceding paragraphs are to be well understood.

Acknosledgemenrs-This work was financially supported by the Deutsche Forschungsgemeinschaft and the Swedish Board of Technical Development. The authors are grateful to Mr F. Eckstein for the production of the single crystals. REFERENCES 1. M. Hillert and G. R. Purdy. Acra mrtall. 26, 333 11978). 2. K. Smidoda. W. Gottschalk and H. Gleiter. Arta metall. 26. 1833 (1978). 3. H. Gleiter. Z. Met&, 61, 282 (1970). 4. W. Gust, M. B. Hintz. B. Predel and U. Roll. Am metal/. 28, 1235 (1980). 5. G. Hasson. J.-Y. Boos. 1. Herbeuval. M. Biscondi and C. Goux. Su$ Sci. 31, 115 (1972). 6. I. Herbeuval and M. Biscondi. C.r. hehd. SPanc Acad. Sci. Paris. C 273, 1416 (1971). 7. H. W. Werner. Vacuum 24, 493 (1974). 8. H. W. Werner. in Appiird Scwfucr Analjws (Edited b) T. L. BARR and L. E. DAVIS) p. 81. Am. Sot. Test. Mater.. Spec. Tech. Publ. 699 f 19801. 9. P. Dorner. W. Gust. M. B. Hintz. A. Lodding. H. Odelius and B. Predel. Actu tneraif. 28, 291 (1980). 10. W. Gust. M. B. Hintz. A. Loddine and H. Odelius. Phil. May.. A43, 1205 (1981). 11. A. Lodding. Adr,. Moss Specrron~. 8. 471 (1980). 12. M. Bernheim and G. Slodzian, Inr. J. Mu.v Spwtrom. Ion Phys. 12. 93 (1973). 13. A. D. Le Claire. Br. J. uppl. Phys. 14, 351 (1963). 14. C. B. Gibbs. Ph?sica status soiidi 16, K?7 (19661. 15. R. T. P. Whipple, Pl~ii. May. 45, 1125 (1954). 16. T. Suzuoka, Trans. Ju,~an Inst. .Merclls 2, 25 (1961). 17. J. C. Fisher, J. uppl. Phys. 22, 74 f 1951 I. 18. T. Suzuoka, f. Ph~s. Sot. Japan 19. 839 (1964).

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GUST et a/.: ON GRAIN BOUNDARY DIFFUSJON

19. H. Gleiter and B. Chalmers, Progress in Materials Science. 16, 1 (1972). 20. J.-Y. Boos and C. Goux, C.r. hebd. Seanc. Acad. Sri. Paris,

Metal/k.

C 271, 978 (1971).

21. J. A. Kargol and D. L. Albright. Merall. Truns. A 8. 27 (1977). 22. W. Lange. A. Hassner and G. Mischer, Physica sfutus sohdi 5, 63 (1964). Phvs. 36. 3596 (1965). 23. A. R. Wazzan. J. uool. M: J. Sinnott. ‘Am. Sot. Merals 24. W. R. Upthegroveand SO, 1031 (1958). 25. R. F. Canon and J. P. Stark. J. uppl. Phys. 40, 4366 (1969). N. A.

Gjostein, in Diffusion, Proceedings of ASM Seminar, coordinated by H. 1. Aaronson, ASM Press. Metals Park, Ohio, 241 (1974). 242, 1359 27. G. R. Speich, Trans. Mefall. Sot. A.I.M.E. 26.

29. W. Gust. B. Predel and K. Diekstall, Z. Metal/k. 68, 619 (1977); 69. 75 (1978): 69, 445 (1978). 30. W. Scharfenberger. G. Schmitt and H. Borchers, Z.

(1968).

28. W. Gust, U. Leininger and B. Predel, Proc. fnt. ConJ: So/id-So/id Phuse Transformations (Edited by H. I. Aaronson, D. E. Laughlin.-R. F. Sekerka and C. M. Wavman) August 10-14. 1981, Pittsburgh, Am. Sot. Metals, Metals Park, Ohio (1982).

63, 553 (1972).

H. I. Aaronson and Y. C. Liu, Scripru mecull. 2, 1 (1968). 32. M. Hillert, Metull. Trans. 3, 2729 (1972). 33. W. Gust. in Phase Transformations. Sorina Residential Con/:, April 4-7, 1959. -University of York:England. Series 3. No. 11. The Chameleon Press. London. 1. 31.

II-27 (19791. 34. G. C. Hasson

and C. Goux, Scripta metal/. 5, 889 (1971). 35. D. H. Warrington and M. Boon. Acta mefall. 23, 599 (1975). 36. Y. Estrin and K. Liicke, Acru metull. 29, 791 (1981). 37. C. J. Simpson, W. C. Winegard and K. T. Aust, in Grain Boundary Structure and Properties (Edited by G. A. Chadwick and D. A. Smith) p. 201, Academic Press, New York (1976).