High resolution transmission electron microscopy study of interfaces

17

Materials Chemistry and Physics, 32 (1992) 77-85

High resolution

transmission

electron microscopy study of interfaces

Joel Douin Laboratoire d’Etude des Microstncctures, CNRSKINERA (France)

Thierry

(OM), 29 u~w~~~ede la Divkion Leclerc, RI’ 72, 92322 Chatillon Cedex

Epicier

Cetztre d ‘Etudes et de Caractirisation Microstructurules, GEMPPM, (France)

Jean-Michel

u.a. CNRS 341, INSA, B&t. 502, 69621 Villeurbanne Cedty

Penisson

D~~a~e??l~nt de Recherche Forldame~tale, Service de Fhysiqae, CENT,

and Alain

85X, 38401 Grewble

Cedex (France)

Thorel

Centre des Matkiau,

Ecole des Mines de Paris, BP 87, 91003 Evry Cedex (France)

Presenred at thr 3rd Meeting on the Science of Materials, 27th-29th October 1991, Oran (Algeria) (Accepted

February

28, 1992)

Abstract is discussed The study of interfaces by means of High Resolution Transmission Electron Microscopy (HRTEM) through selected observations conducted on the Atomic Resolution Microscope (NCEM, Berkeley, USA). The precipitation of germanium into aluminium, the study of inter-facial non-crystalline films in silicon nitride, the determination of the chemical nature of twin planes in CuAlO,, and the structure of the X3 grain-boundary in aluminium are the examples that serve to illustrate the important problems of relevance in the atomistic study of interface with HRTEM.

1 Introduction Most real solids contain interfaces, grain-boundaries or phase-boundaries. The transmission of a given property within the material may depend on the physical and chemical structure of the interfaces; for example, one can refer to the propagation of a dislocation from one grain to another, which is greatly influenced by the relative misorientation between both crystals [I]; also, the limitation or modification of transport properties, as the conductivity in superconducting materials can be cited [2]. The microstructural characterization of such interfaces is then an obligatory path. Among the various investigation methods, Transmission Electron Microscopy (TEM) is an advantageous technique which offers the possibility of getting detailed information at a very local scale; moreover, the increase of the resolving power of modern instruments allows, in the high resolution mode (i.e. HRTEM}, the condensed matter to be imaged in terms of atomic columns at a 0.2 nm or slightly

0254.(584/92/$5.00

better resolution (see [3] for an excellent presentation of HRTEM). The goal of the present paper is to illustrate the benefits of HRTEM for the structural characterization of interfaces through dedicated examples; for further information on this topic, the reader is referred to the following specific reviews on TEM and/or HRTEM studied of interfaces [4-81. Basically, microstructura~ features of any interface can be described according to the following simple classifications (from the ‘macroscopic’ to the ‘microscopic’ scale): orientation relationships - since the interface generally corresponds to a misorientation behveen two phases, the geometrical relations between both compounds have to be determined. HRTEM will allow the question of variants to be examined (section 2); infergranular structure - in numerous real cases, interfaces act as seeds for the appearance of new wetting phases, such as interfacial films (section 3), precipitates or local segregation;

@ IV92 - Elsevier Sequoia. All rights reserved

78 atomic structure of the interface - the ultimate characterization scale concerns the nature (section 4) and the position of atoms - or atomic columns (section 5) at or near the interface. In order to tackle all these aspects, two kinds of approach can be followed: either the sample is a ‘model’ (in such a case, the characteristic of interest is generally well selected), or the ‘real’ material is directly investigated because of its specific properties. The easiest approach for the HRTEM or conventional TEM work is the first one: appropriate ‘bi-materials’ (such as bi-crystals, or planar interfaces prepared through cross-sectional methods, see [9, lo] for detailed information on the thin foil preparation techniques) are simple to handle from the point of view of the geometrical constraints imposed by the instrument itself (e.g. tilting performances, and their influence on the optical quality of the objective lens). In many cases however, this academic approach cannot be developed, and one is forced to deal with the ‘real’ material: at this point, the success of the study is strongly dependent on the luck of the investigator, who has to find adequately oriented interfaces among all the interfaces which are most frequently randomly oriented and distributed. A typical example is the examination of the grain-boundary structure within fine-grain ceramics, where edgeon interfaces have to be found. The HRTEM observations presented in this work have been conducted on the Atomic Resolution Microscope at NCEM (National Center for Electron Microscopy, Berkeley, USA); the performances of this instrument are summarized in Table 1. From these data, it is important to note the unique association of an excellent resolution (in a routine way, details at a 0.17 nm level can be obtained in axial illumination) with unsurpassed

TABLE

1. Characteristics

of the JEOL Atomic Resolution

accelerating voltage (kV) electron waveiength h (nm) measured coefficient of spherical aberration Scherzer defocus: Afs= -m (nm) theoretical resofution” 0.67 Cu4h3” (nm) semi-angle of beam convergenceb (mrad) defocus spread” (nm)

characteristics

of the goniometer

C, (mm)

tilting capabilities on a HR machine; such a combination makes this instrument a well-adapted one for the study of interfaces within both types of specimen evoked previously.

2 Orientation relationships: germanium in aluminium

precipitation

of

The aluminium-geranium system, although metallurgically simple, is of a special interest because of the importance of a knowledge of metallsemiconducting material interfaces in the electronics industry; moreover, it can be noted that the precipitation of germanium into aluminium leads to a great variety of shapes and orientations, which is of intrinsic interest (for a review, see [12]). HRTEM allows direct and detailed information to be obtained on the orientation relationships (OR) between Ge precipitates and the Al matrix, on the nature of interfaces and the role played by finning during the nucleation and growing of the precipitates. Attention will be more specifically focused on the OR. 2.2 Experimental procedure An Al - 1.1% Ge alloy has been annealed at 550 “C, and quenched at - 60 “C in order to create the vacancy supersaturation required for the precipitation of Ge in Al (the generalized precipitation being indeed obtained after further annealing at 250 “C for 4 hours). Because of the polyc~stalline nature of the samples, and due to the large variety of the orientations adopted by the precipitates, the large tilting possibilities of the ARM are required for the observation of edge-on particles.

Microscope

(from [ll]) 1000 0.000872 2.27 * 0.22 - 54.3 0.131 0.55

800 0.001027 1.93_to.12 - 54.4 0.143 0.76 12-15

sample diameter ~3 mm, adjustable ‘z’, ‘double-tilting’ + 45”

“l/g”, with h: first cut-off in the contrast transfer function at the Scherzer defocus ‘usual observation conditions (with a condenser aperture of 100 and 150 grn at 800 and 1000 kV respectively) ‘due to the chromatic aberration (typical estimated value)

79

2.3 Precipitates - matrix OR Among all the configurations which have been encountered for the Ge precipitates in Al [13], the most frequently observed are those where < llO> Ge needles are parallel to < lOO> (see Fig. la) or < llO> (Fig. lb) axis of the Al matrix. Such relationships are favored by the good match between the Ge and Al lattices in < 100 > A, and < llO> CiC, and < llO> Al and < llO> C;e directions. Faceting of < 110 > precipitates generally occurs along the {lll}G, planes, frequently parallel to one of the (100) set of Al-planes (arrows in Fig. la). However, the Ge needles are mainly twinned, and twinning strongly influences the final morphology of the precipitates: the intersection of one or more twin plane(s) with the surface of the particle modifies the matrix-precipitate OR, as well as the faceting. Such a ‘pseudo-rotation’ of

the precipitate with respect to the matrix leads to new OR, such as {lll}G, ll{310}A, in the case presented in Fig. la. A further twinning leads to a configuration where the (001) planes from both compounds are almost parallel: this can be directly observed when the twinning occurs in two different (111) planes in zone with the growth direction of the precipitate (Fig. lc). The ‘pseudo’ five-fold symmetry arises from the existence of a 5-atom Ge-ring at the intersection of the twin planes, which is directly related to the favourable atomic arrangement of the Ge-seed [13]. A detailed examination of this intersection (enlargement of Fig. Id reveals that the (111) plane of part (1) of the precipitate is parallel to (OlO),,, and that the (OOl),, plane is parallel to (OlO),, in part (3). Since five (111) twins within a cubic structure lead to a total rotation of 352.65”, a residual stress is produced by the elastic distorsion due to the ‘closure failure’ (360”-352.65” = 7.35”). Such a stress is clearly accomodated in both (1) and (2) twinned crystals, the growth of which is consequently inhibited. In the case of needles parallel to < llO>,,, both relationships {111},,11{001},1 can be considered as correlated and {OO1}G, II {OOl},, products of a unique growth mechanism [13]. Indeed, it is clear from Fig. Id that, starting from ‘cube-cube’ orientation relationship (i.e. ;ool),, II (001) *,, one, two and three twin operations lead to the successive parts (2) (3) and (4) of the precipitate. 2.4 Conclusion The direct observation of the internal structure of Ge precipitates in Al by means of HRTEM allows a better understanding of the different orientation relationships between germanium and aluminium. Indeed, it is shown that this ‘macroscopic’ feature (i.e. the orientation relationships) is directly governed by atomic events (nucleation processes e.g. atomic configuration of the nucleus, multiple twinning etc) which can only be isolated in very small precipitates, making the HRTEM study a required step for the understanding of such precipitation mechanisms.

Fig. 1. Precipitation of Ge in Al (ARM, 800 kV). (a) High resolution image of a < llO> oc twinned needle parallel to ,,,. Faceting occurs parallely to {lll}oC (notice the orthorhombic shape). One of the Ge planes is perfectly aligned with a (OOl),, (arrow). The top part of the precipitate shows the (111)o,11(310),, OR; (b) HR micrograph of a < llO> oc twinned needle parallel to < 110 > Al. Faceting occurs parallel to (11 l}oe; (c) Another example of a < llO> oc multiple-twinned needle parallel to ,,; (d) enlargement of the multipletwinned region of the needle in c); note the five-atom ring at the center.

3 Atomic structure nitride

of interfacial

films in silicon

3.1 Introduction The question, of great importance for the practical use of such materials in high temperature environments, of the presence of non-crystalline films at grain-boundaries (GB) in polycrystalline ceramics has been widely discussed in the last 20

80

years (for a review, see for example [14]). Although the presence of such GB films has been ascertained in a lot of cases, and mostly by TEM work, the structure of these phases is much less documented. In the lack of direct experimental evidence, the literature suggests that the nature and composition of such interfacial films in ceramics are identical to those of a bulk vitreous silicate phase, such as exists at grain triple junctions under the form of large glassy pockets. However, in silicon nitride or sialon-based ceramics, the thickness of the GB films varies from 1 to 2 nm: it is then highly improbable that a true vitreous structure i.e. a random distribution of SiO, tetrahedra linked by their corners, can exist under such conditions. From such considerations, a model has been proposed [14], which suggests a kind of epitaxy between the crystal (i.e. grain) and the non-crystalline film (see Fig. 2a). At the high temperatures involved during the elaboration process, the glassy phase, based on the chemical impurities incorporated in the starting powders as sintering aids, is a liquid and wets the grains: at this stage, it is argued by Clarke that a preferential orientation of the SiO, molecules develops under the action of the short distance van der Waals forces, expected to be relatively large near the surface of the covalent grains of nitride. It is the purpose of the present work to check experimentally such an hypothesis: a HRTEM study has been undertaken in order to confirm (or invalidate) this type of organization within the interfacial films in silicon nitride-based ceramics. 3.2 Experimental determination of the atomic structure of an interfacial film in &Si,N, It is indeed shown that a HRTEM work, coupled with numerical simulations of the micrographs, is the only way to win the proposed challenge [15]. The observation azimuth has to be chosen in order to visualize the film edge-on, while conserving resolvable distances within the adjacent grain(s). In p-silicon nitride, interfacial films frequently lie along (1010) surfaces of the grains (see Fig, 2b). Figure 2c shows that the convenient [OOOl] projection leads to easily resolvable inter-reticular distances (e.g. dlloio)p = 0.66 nm), although distances between atomic columns are smaller (distance between ‘(6h)’ bi-columns are equal to 0.26 nm, and the bond Si-N is ~0.165 nm, Fig. 2b). Since an accurate description of the atomic structure is desired, it is necessary to predict the contrast of the experimental HRTEM images, and thorough simulations have been run for that purpose ([15] and Fig. 2~).

Under these conditions, the p-/3 interfaces can be of 2 kinds: - both grains are in the [OOOl] orientation (Fig. 2d). This configuration is most probably of low energy, and no GB film is present. A local reconstruction of the lattices occurs in order to accomodate the misorientation between both grains, which may lead, as in Fig. 2d2), to a dislocation network; - only one grain has a [OOOl] orientation (Fig. 2e), a GB film is present and its organized structure, considered as ‘frozen in’ during cooling at the end of the elaboration process, is partially but positively visible. 3.3 Discussion and conclusion The experimental observations have shown that the condition ‘both c-axis aligned’ is sufficient in order to get a clean grain-boundary between two adjacent P-Si,N, grains: indeed, the GB is of low energy and the interfacial film is not stable during cooling down. Whatever the misorientation between {lO?O} planes from each grain, the boundary is a near-coincidence GB [17]; however, for such a covalent material, it appears that the question of ‘speciality’ strongly depends on the nature of the interface plane. When the c-axis of both grains are not parallel, their relative misorientation is generally random and the boundary possesses a higher energy. The presence of an interfacial film can decrease this energy. The present HRTEM observations bring evidence for a topotaxial region within the film near the surface of the grains; however the accuracy of the [OOOl] orientation must be better than 3” nominally in order to allow the ‘organized’ zone to be comfortably discerned (thus, it is only identified beside the grain perfectly oriented along [OOOl] [15]). Another important feature is that the contrast of the topotaxial zone exhibits a partial periodicity as observed at interfaces between pS&N4 grains and bulk glass, modelled on the basis of a silicate compound [16]. It is emphasized that such ‘organized’ structures in a thin interfacial film might exist in other systems for which there is a similarity between the structural units from the crystal and the glass, as is the case with S&N, and SiO+ 4 Chemical nature of interface planes: boundaries in CuA102

twin

4.1 Introduction The binary oxide CuAlO* is an occasional interfacial product of the (pre-oxidized) cop-

Xl

[OOOil t

(6

60 )

Fig. 2. Atomic structure of interfacial films in S&N, (ARM, 1 MV). (a) Preferential orientation of Sit& tetrahedra at the immediate vicinity of the silicon nitride grains; (b) visu~i~ation of the P-S&N4 structure; (c) HR [OOOl] images for two defocusing conditions: (c,)-60 nm, the electronic potential is black; (~2) +30 nm, the electronic potential is white; (d) GB between grains with parallel c-axes; (e) GB where only one grain has the (0001) orientation: HR images for two defocusing conditions (e,)-90 nm; (er)- 60 nm; (ez) schematic interpretation of the contrast; (e.,) enlargement and contrast simulation based on a silicate structure (see details in 115, 161).

82

per-alumina bonding; in the general frame of the technological development of metal-ceramic junctions, the system Cu/Al,O, is of particular interest (on the one hand, the metal is ductile and its oxide can easily be decomposed; on the other hand, alumina is a very stable ceramic [18]). Previous conventional TEM observations have shown that CuAIOz (with a trigonal structure*) exhibits a particular twinned microstructure, with defects parallel to the basal plane, and corresponding to enantiomorphic configurations (i.e. ‘left-handed’ and ‘right-handed’ crystal configurations, as described later) [20]. The aim of the present work is to determine the chemical nature of the twin planes, e.g. Al or Cu. HRTEM is, in such a case, the only available method. Adequate thin foils have been prepared by ion thinning of massive polycrystalline CuAlO, compounds [20]. 4.2 Required conditions for HRTEM imaging Due to the polycrystalline nature of the samples, the large tilting capabilities of the ARM appear to be a great advantage for the observation of edge-on (0001) twin planes. Taking into account the crystallography of t_he oxide, the useful viewing direction is of the < 1120 > type: in this orientation, well-separated (~0.25 nm, see Fig. 3a) ‘Al’ and ‘Cu’ atomic columns are projected. However, the access to a resolution better than 0.2 nm allows the direct transfer of structural information arising from the presence of oxygen atoms at 0.19 nm away from the Al columns: as a consequence, the Al columns are imaged as distorted dots (due to a dumbbell effect between close Al and 0 columns), easily distinguishable from the undistorted Cu dots. Figure 3b illustrates this feature which is confirmed by the numerical simulation obtained with the NCEMSS programs

Pll. It might be precised here that the separation of both Al and Cu columns can easily be obtained on a current high resolution instrument; however, the distinction between both types of cations might be tedious, depending upon the crystal thickness, and/or the defocusing conditions (see, for example, approaches by van Tendeloo et al. [22] in the case of metallic alloys, and Ourmazd et al. [23] or Glaisher et al. [24] in the case of III-V semiconductors, for more detailed discussions of the in-

*Space group R3m (delafossite structure, type CuFe,); parameters: a = 0.2858, c = 1.6958 nm (in hexagonal axis); atomic positions: A1:3(b), Cu:3(a), 0:6(c) with ~=0.1098 [19].

Fig. 3. Study of enantiomorphism twin planes in the mixed oxide CuAIOz (ARM, 800 kV). (a) Projected structure along [1210] (the hexagonal cell is underlined). The configuration shown is of the ‘left-handed’ type; the enantiomorphic crystal (‘righthanded’ type) is symmetrical (relatively to the horizontal basal plane); (b) HR micrograph of a thin ‘left-handed’ crystal (thickness = 2 nm); the superimposed simulated image has been calculated at a defocus value near the Scherzer value (Af= -45 nm); (c) diffractrogram from region (b), Af= -50+5 nm; (d) low magnification HR image showing the presence of enantiomorphism twins (arrows); (e) enlargement of a detail from the upper twinned band in (d), the matching image (t=2.5 nm, Af= 10 nm) corresponds to an elementary ‘right-handed’ cell inserted in a ‘left-handed’ crystal; twin planes in Cu planes are indicated by arrows; (f) same as in (e) for a defocus value Af= -40 nm).

fluence of such experimental conditions on the ‘selective imaging’ of sublattices by HRTEM). 4.3 Results

Figures 3d to f summarize the study of a twin band corresponding to the insertion of a ‘righthanded’ crystal of elementary height (i.e. a onecell thick material in the c direction) into the ‘lefthanded’ variant. The examination of the HRTEM images e and f shows that the twin planes are unambiguously copper planes. Academic simulations from a twin model where twin planes are

83

aluminium planes confirm this conclusion. A further discussion of the reasons for which the ‘Cu’ model is more favourable than the ‘AI’ model is reported in 1251: briefly, the twinning on the copper planes leads to a smaller perturbation of the chemical bonds in the vicinity of the interface.

5 Atomic positions near an interface: (112) twin in aluminium

the C3

5.1 Introduction Grain boundaries in crystalline solids can be regarded as interfaces of a specific type. The grains which are separated by the boundary have the same crystallographic structure, and simple geometrical operations, such as a rotation and a translation, transform one grain in the other. Among all the possible boundaries, tilt boundaries possess their rotation axis in the plane of the frontier; they are made of a set of dislocations parallel to this axis. This relative simplicity, as well as the possibility of making dedicated samples where such GI3 are present fe.g_ ‘bi-crystals’) justify the abundant literature in this field and on various metals: MO [26], Au [27-291, Al f30-311, Ni [32]. In most cases, the experimental HRTEM images are compared with images simulated from models obtained from a numerical relaxation technique. At the present time, such a comparison is limited to an accuracy of about 0.02 nm. This limitation comes from different factors, such as the noise present in the experimental image and the presence of optical defects: slight misalignment of the incident beam, astigmatism or distorsions. If differences between two competitive models are less than this limit, they cannot be distinguished [30]. However, a rigid body translation between both grains remains measurab1e. 5.2 ~)~se~ation and disc~s~o~ An incoherent X= 3 (112) twin in aluminium has been observed by HRTEM at 800 kV. Figure 4 shows a typical image of this GB; atomic columns are white for the defocusing conditions, used, i.e. Af= -80 nm. The thickness of the sample, which has not been ascertained, is estimated to about 10 nm. It is seen on this micrograph that the (111) planes perpendicular to the bounda~ run continuously across the interface, i.e. no rigid translation can be detected. Starting from a geometrical model of coincidence, a relaxation has been made with the computing ‘Embedded Atom Method’ 1331 under conditions similar to those reported in [3OJ. This relaxation leads to two configurations of

Fig. 4. Atomic positions close to an interface (i.e. X3 (112) twin in Al). (a) Incoherent 83 (112) twin in aluminium observed at 800 kV with a defocus equal to -80 nm. The atomic positions are white; no rigid body translation is visible across the boundary; (b) superimposition of the calculated atomic positions onto the experimental image (confi~ration with no translation).

comparable energy, which differ by the presence or not of a rigid translation parallel to the boundary. Figure 4 b shows the superimposition of the theoretical atomic positions, in the case without any translation, on the experimental image. The agreement is correct for all atomic columns, even in the core of the boundary. A similar type of boundary has been studied in [34], where a translation of amplitude 0.03 nm was found by the technique of the cy fringes. The same twin has also been observed in gold [35]; from a pure geometrical interpretation, these authors found a translation of the same order of magnitude, The portion of twin which is presently examined belongs to a micro~inned domain in a < 011 > matrix, and the longer portions bounding this domain are indeed C = 3111 l> twins. Their presence near the (112) portion explains the absence of rigid translation. Such an influence of the environment on the magnitude of the rigid body translation was also noted in the case of gold [35f.

5 Discussion

and general conclusion

Through the previous exampies, the interest of High Resolution Transmission Electron Microscopy for the (micro)structurai characterization of interfaces has been demonstrated.

84

For an accurate description of a given interface (or any structural detect), the problem of the resolution is a crucial one. Because of the lattice reconstruction near the defect, distances that are to be resolved in its vicinity are smaller than those present in the perfect crystal: thus, the HRTEM study of a defect generally imposes a resolution better than that sufficient for studying the perfect crystal where it occurs. Moreover, defects are not always aligned or parallel to low-index directions, and it is not always easy to prepare perfectly oriented thin foils for their study; in the case of real materials, such ideal foils do not exist, since interfaces frequently lie in different planes. Hence, all these constraints require a microscope with an excellent resolution and, simultaneously, a large tilting latitude. In most cases, the correct description of interfaces requires to overcome the difficulty of the physical interpretation of their images as they are produced by the microscope; hence, image simulations as an aid to test structural models is frequently necessary. Although this part was only briefly discussed in the present test, this is a fundamental step of any specific HRTEM work. It is fortunately becoming easier and easier to undertake, since simulation packages are now available for personal computers (see, for example,

long-term stays at the NCEM (National Center for Electron Microscopy, Lawrence Berkeley Laboratory, ~niversi~ of California, Berkeley, USA); it is a pleasure to thank the staff of the Center for their help and permission to conduct HRTEM observations on the ARM.

WI).

10

A major dimple remains, regarding the fact that an isolated HRTEM observation cannot have a wide statistical meaning: the technique is a local one, and can still be considered as doubtful when applied for describing macroscopic properties of a given material. Nevertheless, HR information is unique, owing to the power of such a direct and detailed imaging, and HRTEM remains an unsurpassed chara~ter~ation method for a lot of applications. Because of the tendency of investigators to study more and more complex interfaces, i.e. interfaces closer to those existing in real materials, the search for a compromise between the good resolution and the large tilting possibilities should be a guide for further technological progress. In this sense, the Atomic Resolution Microscope was the instrument of the eighties, well-adapted to this type of observations; projects for new equipment do exist in different countries, concerning improved versions of such a microscope, which will be available in the very near future.

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