High Resolution Transmission Electron Microscopystudy of nanoscale Ni-rich Ni-Al films evaporated on to

NaCl and KCl

M. Yandouzi°, L. Toth* and D. Schryvers°

° EMAT, University of Antwerp, RUCA, Groenenborgerlaan 171, B-2020 ANTWERP, Belgium

* MFKI, Hungarian Academy of Sciences, P.O. Box 76, H-1325 BUDAPEST, Hungary

(Accepted February 11, 1998)

Abstract - The atomic and microstructure of Ni-rich B2 NixAl100-x thin films with

nanoscale dimensions, for thickness as well as lateral grain size, and produced by evaporation

on to NaCl and KCl (001) surfaces are investigated by high resolution electron microscopy.

Different textures and orientation relationships with the substrate are found exhibiting grains

with [001], [011] and [111] zones parallel with the film normal. Local distortions including

precursor-like domains and surface shearing as well as dislocations are observed in the

grains. No micromodulations as in bulk martensite precursors are visible. No martensitic

transformation was observed when cooling the free standing films to -170 °C.

INTRODUCTION

In the literature only a few reports of thin films of NixAl100-x with nanoscale dimensions

have been published so far. Most of these studies refer to the mechanical characteristics of Ni-Al

or Ni-Al based superalloys (1,2) or to potential applications, e.g. as metal-semiconductor

contacts (3,4,5,6) . The long term goal of the present study is to obtain Ni-Al nanoscale films in

which the martensitic transformation, known to exist in bulk material with compositions between

x = 62 and 69 (7,8) , and the resulting microstructures can be investigated. Under bulk

conditions, this transformation, including its shape memory behaviour, has been extensively

documented by several experimental as well as theoretical studies (for a recent review see (9) ).

Until now most studies of martensitic transformations in thin metallic films only considered

films with a thickness of the order of microns and focused on materials already in use as bulk

material in commercial applications such as Ni-Ti (10,11) . In this case the microscopic

behaviour of the film still follows the normal bulk conditions. However, when the film thickness

is decreased to the nanometer scale the anisotropic nature of the film dimensions can play an

important role in the martensitic transformation. Indeed, several investigators have already noted

difficulties for transforming a thinned bulk sample in-situ or, on the contrary, the existence of

relaxation-induced transformation leading to surface martensite.

As the overall dimensions of the material are decreased also the internal dimensions, e.g.

the grain size, are potentially affected. The latter, of course, will strongly depend on the

deposition technique and subsequent treatments of the film. If no special efforts for epitaxial

growth are made most film preparation techniques will yield a film consisting of small particles,

sometimes in the nanoscale region. In fact, some investigations on the martensitic transformation

have already been performed in other systems for nanoscale particles as such, i.e. not in a

continuous and free-standing film, essentially concluding on a reduction of the transformation

temperatures with particle size (12,13,14,15). Parameters which have been suggested for

affecting the martensitic transformation in such particles include the surface-volume ratio, lattice

defects such as dislocations and vacancies and imperfect ordering (16,17). When the particles

are combined to form a continuous film the effects of the type and number of grain boundaries,

relative particle orientations and the free surface plane can also be expected to play a role (18). In

this respect the epitaxy on the substrate and the texture of the growing particles is essential. The

influence of the latter parameters will be enforced when monocrystalline thin films are used.

Moreover, in such films the lattice orientations are usually well known so that the correlation

with theoretical predictions on the transformation behaviour and resulting microstructures of the

martensite films can be investigated (19). Of course, in all of the above cases, care should be

taken to properly measure the composition of the films since in many systems the martensitic

transformation temperatures are strongly dependent on composition (for Ni-Al see, e.g., ref. (8).

The emphasis of the present paper is on the atomic and microstructure of nanoscale and

polycrystalline Ni-rich Ni-Al films crystallised in the parent or austenite phase and investigated

by conventional and high resolution transmission electron microscopy (CTEM and HRTEM).

This phase has an ordered B2 (CsCl, bcc based) structure in which, for bulk samples of this

composition range, the off-stoichiometry is accommodated by randomly distributing Ni atoms

on the Al sublattice (20). Depending on the actual goal in mind, the most ideal configuration

could be a single crystal thin film of austenite with a simple normal direction, e.g. [001], or a

polygrain film with controllable grain size. In the latter case different degrees of epitaxy or

texture could be envisaged. In any case, the current aim is for monophase films of B2 austenite

with different possible microstructures. As received arc-melted bulk ingots with x ≤ 62.5 reveal

large single crystals of the austenite phase at room temperature and with very few dislocations

and virtually no anti-phase boundaries. In such samples the martensitic transformation can be

obtained by applying an external stress (21) or at crack tips in thin foils (22) or by cooling the

material to below -100 °C (23). Also bombardment with electrons of 400 keV of a thinned foil in

the electron microscope often induces the transformation in small regions (24).

EXPERIMENTAL PROCEDURES

The results presented here all originate from samples prepared by a conventional

evaporation technique of Ni-Al on a (001) NaCl or KCl surface under a vacuum of

1.5 10-5 mbar. The substrate crystals are air-cleaved and could be heated up to 350 °C for NaCl

and 550 °C for KCl during evaporation. The present samples are all prepared using a single

alloy source with the desired composition. The starting compositions are x = 62.5 and 65,

further denoted as 62.5Ni and 65Ni, typical values encountered in investigations of bulk Ni-Al

martensite. The deposition rate is monitored by a quartz crystal and is 10 Å/sec. as determined

from calibrations using cross-sectioned films deposited on Si. Films of 70 to 100 nm are

produced and floated off in water followed by capturing on a Cu grid yielding plan view

samples. The internal cohesion of the films is sufficiently high so that no extra support film is

needed. By cutting the films parallel with the vertical (100) sides of the substrate crystal before

dissolving the latter, the in-plane orientation of the films with respect to the substrate could be

determined. The advantage of this procedure is that such samples are directly suitable for

CTEM, selected area electron diffraction (SAED) and HRTEM studies without the need for

further thinning which could severely influence the atomic and microstructure of these nanoscale

films. The samples were investigated in a CM20 Philips and 4000EX JEOL microscope, both

equipped with a LaB6 filament. The presented SAED patterns and TEM images were all

obtained with the tilting angles of the sample holders placed at zero so that the orientation

relation with the substrate could directly be visualised.

The composition of the films was measured by energy dispersive X-ray analysis (EDX)

using a Si(Li) detector in a Link QX-2000 system with an atmospheric thin window attached to

the CM20 instrument and using factory provided standards. This configuration was calibrated

before using known bulk material of the source compositions. Averaging over several

measurements (at least ten locations) indicates an enrichment of both films by 4 ± 2.5 at.% Ni

which, for bulk material, would increase Ms by at least 400° bringing the martensitic

transformation temperature above room temperature for both samples.

RESULTS

Crystalline and monophase B2 films were obtained in all depositions as long as the

temperature of the substrate was kept at or below 400 °C, which partially confirms earlier reports

(1) . The largest grains were found for evaporation of the 62.5Ni alloy on to NaCl at 300 °C. A

representative example of such a nearly contiguous film with B2 grains between 10 and 20 nm

in lateral diameter together with the corresponding SAED pattern, obtained using an aperture

with an image size of 600 nm in diameter, is shown in the bright field (BF) image of fig. 1a. The

thickness of the film was measured during deposition as 70 nm. As a result of the use of the BF

condition the well oriented grains which yield the strongest diffraction peaks appear darker than

the other ones. Depending on the actual region of the film, 40 to 70 % of the grains are in this

condition. Figs. 1b and 1c are dark field (DF) images of the central region of fig. 1a and

obtained by selecting one strong Bragg reflection (1b) and a part of the first diffraction ring in

between two strong Bragg reflections (1c), as indicated on the SAED inset in fig. 1a by a circle

and square, respectively. These two selections will thus highlight grains in different orientations

(see also below). The locations of the grains contributing to the Bragg reflection are indicated as

“+” signs whereas those belonging to the ring are given as “x” signs. From the

correspondence of the “+” and “x” signs with the highlighting or non-highlighting grains it is

seen that no overlapping exists between the highlighted grains in both images which is a first

indication of a preferably columnar growth of the grains.

Figure 1. (a) SAED and BF image revealing the microstructure of a B2 62.5Ni thin filmevaporated on to (001) NaCl at 300 °C. (b) and (c): DF images of the central region in (a)obtained using a strong reflection (circle in the SAED) and part of the first diffraction ring

(square in the SAED), respectively, as indicated on the SAED.

From the SAED pattern, which is repeated in fig. 2a, it is clear that the grains in the film

are not randomly oriented, but that there is a certain tendency for textured or epitaxial growth.

Indeed, although a ring pattern is clearly eminent, several reflections can readily be discriminated

forming two-dimensional patterns superposed on to the rings. On the other hand, some freedom

of orientation must still exist even for the well oriented grains since each reflection does reveal

some spread of approximately 10° (± 5°) along its ring. The observation of superstructure or

ordering reflections of the B2 structure is in accordance with the measured grain size which is

larger than the critical size for ordered structures in Ni-Al films on (001) NaCl as obtained by

Kizuka et al. (1) . The square pattern in fig. 2a clearly reveals a preference for the [001] zone of

the B2. The reflections corresponding with this orientation are indicated by filled circles in the

left schematic of fig. 2b, their sizes being a relative measure for the corresponding intensity.

Moreover, the primary cubic directions in this pattern are found to lie parallel with the [100]

directions of the substrate, also indicated in the figure. Still some other sharp reflections are

observed in the experimental pattern, especially in the first ring (100) and at angles of 45° with

respect to reflections of the former configuration. These, together with related reflections, are

indicated in the right schematic of fig. 2b as triangles of different types (see also below). It

should be remarked that the latter ones are less sharp than those of the former set. From this

diffraction pattern alone no unambiguous choice of zone for the latter reflections could be made

(the indicated rectangles and indices are based on HRTEM images as shown below in fig. 5).

From the current ring pattern a lattice parameter for the B2 of (0.28 ± 0.01) nm is measured

using an internal calibration by reflections of undissolved NaCl crystallites on a film rapidly

recovered from the dissolving water. This value corresponds well with bulk lattice parameters

around the present composition (20) .

The SAED pattern in fig. 2a is observed in the entire film except for a small portion,

approximately 10 % of the total film area, which yields the SAED pattern in fig. 2c. In this

region the grains are apparently preferentially oriented along a sixfold [111] zone although the

most intense 110 reflections are slightly more spread along the circles than in fig. 2a, which is

correlated with a less well defined texture. Moreover, the intensity of the rest of the 110 ring is

also stronger than in fig. 2a and even locally concentrated at other points which indicates that

[111] grains with different in-plane orientations will also be present. The relative orientation

between both SAED patterns in fig. 2 is retained in the entire figure and it is seen that the <110>

direction from bottom-right to upper-left coincides when the pattern in (a) is compared with the

primary one in (c).

Figure 2. SAED patterns of a B2 62.5Ni thin film evaporated on to (001) NaCl at

300 °C revealing preferential orientation for (a) [001] and (c) [111] zones. (b) and (d) are

schematics of the most prominent configurations of reflections in (a) and (c), respectively.

Essentially the same results are obtained when using a Ni65Al35 alloy as evaporation

source or a KCl (001) crystal surface as substrate but the preferred epitaxy is less well

established, especially for the 65Ni films, and the grain size is some 20 % smaller than in the

previous case. In the following all images stem from the 62.5Ni source films but some are

evaporated on to NaCl while others on the KCl substrate. The basic conclusions are the same for

all cases mentioned above.

The above results are confirmed when viewing the films under high resolution conditions.

An example corresponding with the SAED pattern of fig. 2a is shown in fig. 3. Here the

HRTEM image in a single grain clearly reveals expected white dot patterns for the [001] zone of

the B2 structure, confirming the above interpretation of the ring patterns and lattice parameter

measurement. Depending on the actual location in the grain the square white dot pattern reveals

the [001] projected superstructure (A) or basic structure (B). Indeed, for column structures such

as the B2 viewed along a cubic direction, the geometry of the white dots in the HRTEM image

corresponds well with that of the projected atomic columns (25,26) . However, in some regions

(C) the square symmetry of the dot pattern is distorted yielding a nearly one-dimensional pattern

with fringes parallel with the <110> directions in the plane, i.e. the elastically soft directions of

this material and visible in the present orientation (27) . Also, when viewing the image along a

grazing incidence parallel with the lattice planes, especially {110} ones, numerous local

distortions of the lattice fringes are observed. On the edges of the grain (D) a few of these

planes are clearly seen to shear in the transverse <110> direction. However, no micromodulated

domains revealing the precursor phenomena well-known for bulk Ni-Al austenite are observed

(26) . Although the edges of the grain are certainly not perfectly planar, some preference for the

{100} and {110} planes is observed, which confirms earlier images of even smaller grains (1) .

A lower magnification but still revealing the atomic resolution is shown in fig. 4 in which

three adjacent but separate grains with a [001] zone orientation are visible. The cubic directions

in all three grains are parallel to within 5° which again indicates that all [001] grains have

essentially the same epitaxial relation with the substrate. This is emphasised by the power

spectra shown as insets from which the average directions can easily be measured and compared

as done in the centre for the [100] direction. Again preferential {100} and {110} edges are

found. The two grains in the left part seem to be separated by a thin amorphous or less

crystalline layer while a grain with a different orientation exists between the two lower ones,

indicating that the three grains were nucleated and grown separately. In practise, all observed

[001] grains in a given film are found to have their lateral directions parallel to within a few

degrees.

Figure 3. [001] HRTEM image of a B2 grain in a 62.5Ni thin film evaporated on to(001) NaCl at 300 °C. The square pattern is clearly visible and the grain is bounded by

{100} and {110} planes. Lettering A, B and C indicate locations with superstructure, basicstructure and distorted HRTEM image. At D transverse shearing of the {110} planes at the

edge is observed.

Figure 4. HRTEM image of three adjacent [001] B2 grains in a 62.5Ni thin filmevaporated on to (001) KCl at 300 °C. The close to parallel orientations in the plane areobvious from the included power spectra and compared [100] directions in the centre.

In fig. 5 a coalescence of three well oriented grains but with different zone axes is shown.

In this case the zone orientation of the lower-left grain is clearly sixfold [111] but that of the

other two is less perfect. Nevertheless, the image of the central grain still reveals the square

symmetry and this one is thus close to a [001] orientation. The upper-right grain, however,

reveals a rectangular symmetry with a ratio of 0.7 which indicates the [011] zone. These

orientations are confirmed by the power spectra shown as insets. The interface between the

[001] and [011] grains is parallel to a (11- 0) plane for the former and to a (100) plane of the

latter. As a result one family of {110} planes continues over both grains, although the match

between them is not perfect and they make an angle of approximately 5°. The interface between

the central and left grains is parallel with a (100) plane of the [001] grain and with a (101- ) plane

of the [111] grain. This match yields a close to parallel configuration for the [100] and [101- ] or

[010] and [12- 1] directions in the former and the latter, respectively. In other words, this

correlation is thus different from the one suggested by the SAED patterns in fig. 2 where two

<110> directions coincide. The edges of the [001] and [011] grains are again parallel with {100}

and {110} planes.

In all of the above HRTEM images the observed symmetry of the pattern belongs to a

single zone of the B2 structure, regardless the chosen defocus of the objective lens. This

confirms the earlier conclusion of columnar growth based on selective DF imaging. Moreover,

image simulations indeed confirm the observed HRTEM patterns for the current film

thicknesses and thus column height. Ni-Al films deposited on Si (001) and investigated in

cross-section for calibration of the thickness monitor also revealed a columnar shape for single

grains. As the films deposited on NaCl or KCl are only weakly attached to their substrate, cross-

sections of these samples could not be made.

Viewing the central grain in fig. 5 along the [11- 0] direction a few dislocation cores with a

projected Burgers vector of 1/2[110] and roughly aligned along the [010] direction are observed.

This part is enlarged in fig. 6 where one dislocation is indicated by a Burgers circuit and the

other ones by the conventional sign at the dislocation centre revealing opposite signs for the

observed projected Burgers vectors. The partial 1/2[110] displacement vector can be measured

from the circuit in the first case. The distances between two adjacent dislocations of opposite

sign are measured as 2 nm (left couple) and 0.5 nm (right couple).

On one occasion a large region revealing the sixfold symmetry already shown in fig. 2c

was observed in a film which was for the most part [001] oriented. A typical example of a grain

with a [111] normal is seen in the right part of the HRTEM image of fig. 7. Although the actual

HRTEM pattern differs quite strongly within the grain, the sixfold symmetry is still apparent as

seen from the power spectrum. The observed deviations from a perfect sixfold image resemble

Moiré patterns which is a typical HRTEM imaging feature indicating small differences in

orientation of subgrains within a given column with an average [111] zone axis.

Figure 5. HRTEM image revealing the orientation relations between grains withdifferent zone orientations in a 62.5Ni thin film evaporated on to (001) KCl at 300 °C.

Figure 6. HRTEM image of the central part of the [001] grain in fig. 5 revealingdifferent dislocation cores roughly aligned along the cubic directions.

This will of course result in the elongation of the respective reflections along their circles as

already mentioned above. In the same image an [011] oriented grain, completely separated from

the former, is seen in the lower-left corner. Again the relative orientation between both grains can

be determined from their respective power spectra and it is concluded that the [011] direction of

the [011] grain is parallel with one of the <110> directions in the [111] grain.

When the above described monophase films are cooled to in-situ -170 °C no martensitic

transformation is observed. In view of the measured film compositions which in bulk have an

Ms above room temperature, the origin of a suppressed martensitic transformation should be

found in the nanoscale features of the films.

Figure 7. HRTEM image of a large [111] grain revealing Moiré-like contrast due tooverlapping of slightly rotated subgrains within a given column. An [011] grain is seen in

the vicinity and their relative orientations are clear from the inset power spectra.

DISCUSSION

As mentioned in the introduction, the long range aim of the present study is to investigate

the martensitic transformation in nanoscale thin films of Ni-Al. The nanoscale features primarily

refer to the thickness of the film but for the present samples they also correspond with the lateral

dimensions of the grains. A first step is to properly characterise the atomic and microstructure of

the as-deposited films which should be monophase B2, i.e. only contain the ordered austenite

phase with respect to the martensitic transformation. The results presented above indeed reveal

the B2 structure and the films selected are those which reveal the largest grains with the

strongest texture or epitaxial relation with the (001) NaCl (or KCl) substrate surface.

The HRTEM image of a typical [001] grain in fig. 3 reveals some specific imaging

features on the atomic scale. First of all, the lattice image changes quite rapidly over very short

distances. Indeed, the imaging focus of the instrument is chosen such as to clearly visualise the

ordered structure in most of the grain, i.e. one works at Scherzer defocus which is approximately

-40 nm in the used instrument. However, in regions denoted by B an image of the basic lattice is

prominent. Image simulations indicate that for films of a thickness between 50 and 100 nm a

thickness change of minimum 10 nm is needed to change the image accordingly at Scherzer

defocus. Changing the defocus by this amount reveals the superstructure image at regions B but

decreases the contrast at regions A, proving the overall ordered character of the grains.

The distorted lattice images at C and D can not be explained by a change in defocus. Here

the symmetry of the lattice must be affected in some way in order to produce the distorted

patterns. At some locations the lattice strains around dislocation cores, which are indeed

omnipresent in the grains as also seen from figs. 5 and 6, will be at the origin of this. Another

explanation could be local distortions of the austenite lattice following the elastically soft

directions of the lattice. In bulk Ni-Al material such distortions give rise to the well-known

precursor tweed pattern in conventional TEM images (28) which is due to an underlying

contiguous patchwork of inhomogeneously strained domains (ISD’s) in which the transverse

shear mechanism of the ensuing martensitic transformation is anticipated (26) . Usually such

domains are correlated with the existence of potent defect sites (vacancies, interstitials, antisite

atoms) in their centres and with a strain field to which the surrounding lattice adapts its local

structure. The gradient of the amount of distortion when moving away from the centre of the

ISD yields imperfect atomic columns along the viewing direction which results in line instead of

dot HRTEM patterns. In 62.5Ni bulk samples the ISD’s include a micromodulation of a

wavelength close to that of the stacking sequence of close packed planes in the martensite. No

such micromodulations are seen in the present grains which could indicate that the room

temperature films are much further away from Ms than bulk material with the same composition.

The slight but massive transverse shear at the {110} grain edges (D) differs from the previous

mentioned distortions since here the dot pattern is retained. This indicates that the atomic

columns are still well aligned along the electron beam which means that the entire plane has

sheared as such. This could be interpreted as the first signs of the actual martensitic

transformation, i.e. not only a precursor, happening on the edges of the grains where the

retaining function of a surrounding lattice has fallen away such as in the case of surface

martensite in thinned bulk samples, e.g. of Ni-Al (29) , and which is usually referred to as

relaxation-induced martensite (30) .

From the SAED pattern in fig. 2a and the HRTEM images of figs. 3 and 4 the most

common orientation relationship found between substrate (s) and grain (g) is [001]s//[001]g

indicating the parallelism of both zone axes. Since the lateral directions of these grains are also

oriented parallel among each other and since the cubic ones were found to be parallel with those

of the substrate a strong crystallographic relation with the directions in the substrate (001) plane

is clear. When the lattice parameters of the B2 structure are compared with those in the substrate

(001) plane good correspondence is indeed found as seen from table I and depicted

schematically in fig. 8a. All drawings in fig. 8 only show the atoms in the contact planes and

should be interpreted in view of the lattice directions, not the actual atomic positions. The valuesin the table are a measure for the lattice mismatch δ between the simple directions in the contact

planes. They are obtained from the conventional formula

δ = aB2 - a’NaCl

a’NaCl

with a’NaCl being the actual distance between two successive holes in the [100] direction on the

NaCl surface, i.e. the basic lattice parameter, rather than that of the real ordered lattice.

TABLE-I Lattice Mismatch between Thin Film B2 Ni62.5Al37.5 and NaCl

B2 \ NaCl [100] = 0.56402 nm [110] = 0.79764 nm

[100] = 0.28 nm -0.7 % -29.8 %

[110] = 0.40 nm -29.8 % -0.7 %

From table I it is seen that the relative deviation between the [100]B2 distance and half of

the [100]NaCl distance, i.e. the distance between two successive holes, is very small. In other

words, the (001)B2 plane can easily fit on to the (001)NaCl plane by placing the respective [100]

directions parallel. Thus the atoms of the (001)B2 plane, which is a pure Ni or Al plane, the latter

with randomly distributed Ni atoms, will probably be placed into the holes of the (001)NaCl

plane. For these grains an orientation relation of (001)s[100]s//(001)g[100]g can thus be

concluded. The smaller dimensions of the grains as found on KCl (a = 0.6293 nm) substrates

then correlates with the fact that the lattice mismatch here becomes 11.0 % instead of -0.7 % for

NaCl. As the films are finally detached from the substrates the relevance of these mismatches for

the martensitic transformation lies in their influence on the epitaxy and texture of the grains.

The good match between the NaCl (001) surface and the Ni-Al (001) plane can be

expected to yield a coherent interface with small coherency strains or, for large grains, a semi-

coherent interface with interfacial dislocations. In the latter case the expected distance D betweenthe dislocations can be calculated from the lattice mismatch δ by

D = b|δ| =

aB2 + a’NaCl

2|δ| = 40 nm

which is larger than the lateral dimensions of the grains and thus certainly larger than the

distance between the partial dislocations measured in fig. 6. Moreover, when the substrate is

removed potential misfit dislocations, lying in the (001) plane, are expected to disappear at the

new surface. The observed dislocations in the plan views of the deposited films thus have to

originate by a different mechanism such as the difference in lattice contraction between film and

substrate when cooling down from the deposition temperature to room temperature. Owing to

the very small grain size in our films, these dislocations are very likely to be trapped at the grain

boundaries, and thus are also removed from the inside of the grains. The dipole configuration as

observed in fig. 6, however, may prevent the dislocations from being swept to the next grain

boundary by the driving force of the thermal stress during cooling. Unfortunately, the nanoscale

dimensions and the local lattice distortions of the grains inhibit a three-dimensional investigation

of the observed dislocations. Whether the observed 1/2[110] displacement vectors stem from

partial 1/2[110] or 1/2[111] dislocations could not be determined as in stoichiometric bulk

samples in which the latter type has been found to appear from the dissociation of a [110]

dislocation (31). On the other hand, it was calculated that 1/2[111] dislocations are not likely to

result from the splitting of <111> dislocations in stoichiometric bulk Ni-Al (32).

The existence of strong strain centres such as dislocations can play a major role in the

nucleation of the martensitic transformation. On the other hand, dislocations can also stabilise

the austenite (16) . A description of the influences of a large density of dislocations, certainly

when compared with homogenised bulk Ni-Al, should thus certainly be treated with care.

The second type of preferentially oriented grain has its [011] zone parallel with the [001]

normal to the NaCl surface. From the HRTEM image of fig. 5 it is seen that the [100] direction

in the plane of this grain makes an angle of approximately 45° with the <100> directions in the

adjacent [001] grain. The [011- ] direction of the former does remain parallel with the [110]

direction of the latter. So it is the correspondence between the <110> distances rather than that of

the <100> ones in the contact planes which yields the preferred orientation. From table I it is

seen that the simple calculation used cannot discriminate between these two options. Again the

atoms of the B2 contact plane, which is now a stoichiometrically mixed plane, can be placed in

the holes along the [110] direction on the (001) NaCl surface as seen in fig. 8b. For these grains

an epitaxial relation of (001)s[110]s//(011)g[011- ]g can thus be concluded. With this information

the SAED pattern of fig. 2a can now better be understood. Indeed, the four reflections in the 100

ring and placed at 45° of those belonging to the [001] type grains can now be interpreted as

originating from two sets of [011] grains rotated 90° with respect to one another. Both sets are

indicated by up- and downward pointing triangles in the right part of fig. 2b, open and closed

ones indicating the relative intensity of the reflection. The second independent direction in the

plane for these grains is the [011- ] of which the reflections coincide with those of the [110]

direction already present in the pattern of fig. 2a (or the left part of fig. 2b) from the [001]

grains.

Figure 8. Schematics of epitaxial orientation relations between (001) NaCl and differentNiAl B2 orientations: (a) (001)s[100]s//(001)g[100]g, (b) (001)s[110]s//(011)g[011

- ]g, (c)

(001)s[110]s//(111)g[11- 0]g and (d) (001)s[100]s//(111)g[11

- 0]g. Only atoms in the contact

planes are shown.

The third type of grain with a well defined texture has a [111] zone, as seen in figs. 5 and

7. The most prominent correlation with the [001] texture can be found from the SAED patterns

in fig. 2 which combine the diffraction effects from several grains within the selective aperture.

As mentioned above, a coincidence between one set of <110> directions from the [001] pattern

and the [111] pattern is observed. As this <110> direction in fig. 2a also belongs to one type of

the [011] oriented grains (upward pointing triangles), this correspondence is the same as the one

observed in fig. 7 and implies a (001)s[110]s//(111)g[11- 0]g relationship with the substrate in

which at least one direction has a very low misfit of -0.7 % and the contact plane for the B2 is

pure Ni or Al, the latter with random Ni atoms. On the contrary, the correlation between the

[111] and [001] grains in fig. 5 implies a (001)s[100]s//(111)g[101- ]g relation which is thus also

possible but occurs less often. Such a situation will yield some extra intensity in between the

prominent reflections as seen in fig. 2c. From table I it is seen that for this case the geometrical

mismatch is clearly worse - more complex directions do not yield any better fit - explaining the

fewer cases observed. Both possibilities are also shown schematically in figs. 8c and 8d from

which again the preference for the first type can be argued on the basis of contact plane

geometry.

The actual reason why a mosaic of textures is observed correlates with other observations

of metal and alloy depositions on NaCl crystals (33,34,35) . Generally it might be expected that

when the experimental conditions improve, a single texture with good epitaxy will prevail. The

current set-up using air-cleaved samples and a simple evaporation unit leaves a small amorphous

film as contamination between the substrate and the deposit. Films obtained after shorter

deposition times reveal worm-like conglomerates of small grains and separated by a few tens of

nanometers of amorphous material indicating an island growth or Volmer-Weber mode

mechanism at the early stages of the deposition followed by a combination of this with

coagulation or continued growth into chains of grains. The texture and orientation of the

nucleating particles can also be influenced by surface steps resulting from the cleavage of the

substrate crystals (36) . The growth mechanism leading to the worm-like configuration could

then explain possible intergrain relations thus yielding large regions of the same epitaxy. When

more material is deposited the growth continues as column growth which, under the present

conditions, indeed corresponds with the diagram of structured zones as proposed by Movchan

and Demchishin (37,38) . Generally one could suggest that the present observations together

with the lattice mismatches imply the [001] grains to be the normal mode with respect to the

perfect (001) substrate surface. Surface defects or decreased substrate-film interactions due to a

contamination layer could then possibly alter this preferential mode yielding [011] and [111]

grains. On the other hand, well oriented metal films have since long been produced on NaClsurfaces even with very large mismatches indicating the limited influence of δ [27]. The use of

an alloy, however, further complicates the interactions with the substrate due to different

contacting atoms.

The existence of a mosaic of different epitaxial grains in these films can again disturb any

potential martensitic transformation. Indeed, the strains occurring during a martensitic

transformation will not be aligned over a long distance in such a polycrystalline film which could

ultimately stop a passing transformation front (18) .

When the appearance of a monophase B2 film at these compositions and lower

temperatures is compared with the phase diagram of the bulk material it is seen that in the latter

case one would expect a decomposition with the Ni3Al (L12) or Ni5Al3 phases. Ni-Al films of

the present composition but deposited on a polycrystalline Fe alloy yield a multiphase structure

including pure Ni and Ni3Al grains (39) . Rings from these phases, however, are clearly missing

from the diffraction patterns in fig. 2. As for the observed phases, our results confirm those of

earlier reports for the case of an evaporating Ni-Al alloy on NaCl (1) . No disordered phases

were observed which is understood from the fact that the grains are indeed substantially larger

than the critical grain size of 5 nm (1) . In view of the observed epitaxial relations it can thus be

concluded that the formation of the B2 grains is possibly aided by the underlying substrate. Still,

the influence of the geometry alone of the substrate should not be overestimated. Indeed, several

reports have shown that metal grains on NaCl not necessarily orient along the directions as

expected from simple geometrical rules (40) and that also much larger misfits can produce good

epitaxy (41) . When the substrate is held at higher temperatures the influence of the epitaxy is

decreased and the expected decomposition takes place.

Recent results in Ni-Al as well as other materials indicate that the austenite can strongly be

stabilised by the size of the grains or the introduction of defects such as dislocations, grain

boundaries and point defects (15,16,17,42,43) . The large number of grain boundaries and

dislocations is obvious in the present case while the relatively low temperature of the substrate

can be assumed not to provide enough diffusion for perfect ordering kinetics and thus leaving

the austenite with a high density of vacancies and antisite defects. On the other hand, strain

centres generally act as nucleation sites so a competition between these different forces will

exist. Also the two-dimensional nature of the film could allow for some extra strain relief

without the need for a martensitic transformation.

CONCLUSIONS

Monophase B2 Ni-Al films with a thickness below 100 nm and a grain size between 10

and 20 nm were produced by evaporation of Ni62.5Al37.5 and Ni65Al35 alloys on heated NaCl

and KCl (001) substrates. Different epitaxial relations were observed including

(001)s[100]s//(001)g[100]g, (001)s[110]s//(011)g[011- ]g and (001)s[110]s//(111)g[11

- 0]g.

Different observations indicate that the present grains are actually columns formed following the

structural zone diagram of Movchan and Demchishin (37,38). HRTEM images indicate that the

grain thickness varies by as much as 10 nm and that internal distortions common to Ni-Al B2

material are present although no micromodulations are observed. Most grains exhibit several

dislocations with a density much higher than in as-received bulk material.

No martensitic transformation was observed upon cooling to -170 °C. At this point it is

unclear which of the observed features such as the nanoscale size, dislocation density or mosaic

texture is primarily responsible for this drastic decrease in Ms.

ACKNOWLEDGEMENTS

We like to thank Ludo Rossou from the University of Antwerp for his extensive help with

the thin film preparations and Dr. V. Teodorescu from the Institute of Atomic Physics,

Bucharest-Magurele, Romania, for many stimulating discussions on the thin films. Part of this

work was sponsored by a collaboration program between the Academies of Science from

Belgium and Hungary and by a RAFO project from the University of Antwerp. We also like to

acknowledge support from the IUAP program from the Belgian Government on Reduced

Dimensionality Systems.

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