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The characterization of NbAl2O3 and NbMgO interfaces in MBE grown NbMgONbAl2O3 multilayers
Acta matall, mater. Vol. 40, Suppl.,pp. $237-$247, 1992 Printed in Great Britain. All fightsreserved
0956-7151/92 $5.00+ 0.00 Copyright © 1992Pergamon Press Ltd
THE CHARACTERIZATION OF Nb-A1203 AND Nb-MgO INTERFACES IN MBE GROWN Nb-MgO-Nb--A1203 MULTILAYERS D. X. LI 1, P. PIROUZ 1 and A. H. HEUER l, S. YADAVALLI2 and C. P. FLYNN 2 IDepartment of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106 and 2Dfpartlnent of Physics and Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana, IL 61801, U.S.A. /dmtrlctwNb-MgO--Nb-A1203 multilayer stacks were prepared by sequential deposition of Nb and MgO films on the (0]'12) (R) plane of sapphire, using Molecular Beam Epitaxy (MBE). Cross-sectional specimens of this multilayer were prepared and various interfaces investigated by high resolution transmission electron microscopy (HREM). Orientation relationships between Nb and AI203 and between Nb and MgO were determined and misfit dislocations, steps and facets at the Nb-AI203 and Nb-MgO interfaces were analyzed. The HREM observations show that, under electron beam irradiation, a modified layer forms at the Nb-AJ203 interface by a solid state reaction. Electron beam-induced grain growth in the Nb film also takes place. g/mmnt---Des empilements de multicouches Nb/MgO/Nb/A1203 sont pr6par~s par d6p6t S&luentiel de films de Nb et de MgO sur le plan (01].2) (R) du saphir, par 6pitaxie par jet mol~daire (EJM). Des 6chantillons de sections droites de ees milticouches sont pr6par~s et diff~rentes interfaces sont 6tudi/~s en micrmcopie 61ectronique par transmission en haute r~tsohition. Les relations d'orientation entre Nb et AI203 et entre Nb et MgO sont d6termin6es et les dislocations de d~,saccord, les marches et les facettes montrent que, sous irradiation 61ectronique, une couche modifi6e se forme ~i l'interface Nb/A1203 par r6action ~ l'6tat solide. I1 se produit aussi une croissasnce des grains indulte par le faisceau d'61ectrons daus le film de niobium. Zammmenfamung---Durch aufeinander folgende Abscheidung yon Nb/MgO-Filmen mittels Molekularstrahlepitaxie werden Nb/MgO/Nb/AI203-Schichtefolgen auf der (0112) (R)-Ebene yon Saphir hergestellt. An Querschnittspriparaten dieser Vielfachschicht werden die verschiedenen Grenz~chen mit hochaufl6sender Elektronenmikroskopie untersucht. Die Orientierungsbeziehungen zwischen Nb und A1203 und zwischen Nb und MgO werden bestimmt; die Feldpassungsvereetzungen, Stufen und Facetten an den beiden Grenzfl~chen werden analysiert. Aulkrdem ergibt sich, dab sich w~thrend der starken Elektronenbestrahlung bei der hochaufl6senden Abbildung eine modifizierte Schicht an der Nb/A]203-C_vrenztiiche dutch eine Festk6rperreaktion bildet. Elektronenstrahi-induziertes Kornwachstum lauft zusatzlich in den Nb-Filmen ab.
1. INTRODUCTION The combination of two or more materials as a way of making a composite has resulted in many new materials with novel properties. One class of such composites is the structural type whereby the purpose of making such a material is to achieve better mechanical properties than the constituent materials from which the composite is made. Such composities are generally made from a combination of metals, ceramics or polymers and the important issues are mechanical properties such as strength, fracture toughness, creep, etc. A second class is the electronic composites in which the constituent materials are either different semiconductors, or combinations of a semiconductor and an insulator (or a metal). The properties of interest in the latter type of composites are the electronic properties where they may have unique properties such as two-dimensional conductivity.
The methods employed in the production of different types of composites are often very different. In the case of electronic composites, techniques such as Chemical Vapor Deposition (CVD), Liquid Phase Epitaxy (LPE), or Molecular Beam Epitazy (MBE) are used where a film of one material is deposited from the vapor or liquid phase onto a different substrate of another material. F o r structural composites, the method of production is often a technique such as diffusion bonding, eutectic bonding, liquid infiltration, internal oxidation/reduction, etc. Because of the cleanliness of the systems and the control of growth parameters such as temperature and pressure, production of electronic composites is usually better controlled than is production of structural composites. Such strict control is essential in electronic composites because small changes in the purity of the crystals, or small variations in the growth parameters, can have drastic effects on the electronic properties of the composite, while in structural composites, small changes
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in the crystal purity and manufacturing parameters can generally be tolerated. In addition, because of the dimensional scale of the different types of composites, economical factors can play a determining role in the choice of the technique of composite production. An important part of any composite structure is the interface between its different constituents. This problem is well known in electronic composites where considerable effort has been invested in the investigation of interface structure. X-ray diffraction, electron diffraction, and transmission electron microscopy (TEM) are the usual methods employed for investigating the interface structure. An interface may have its own unique properties, which often will have a controlling influence on the properties of the composite as a whole. Thus, in a ceramic-matrix (structural) composite, the interface between two strong ceramics may be weak because of weak chemical bonding between the two constituent ceramics. In this case, the fracture toughness of the composite is controlled primarily by the weak interface rather than the toughness of the constituents. A different case is the presence of a conducting interface in a semiconductor-on-insulator electronic composite (such as silicon-on-sapphire, SOS) which dramatically changes the electronic properties of the material. The interface properties often depend on the atomic structure of the interface, For example, it is widely believed that the Scbottky or ohmic behavior of a semiconductor junction is determined by surface states arising from strained or broken bonds at the interface. Another example is the mechanical strength of an interface in a structural composite, which may be determined by the arrangement of the different atomic species on the two sides of the interface and the way that they bond to each other. The investigation of a composite interface is much easier when it is made in a controlled fashion, so that effects of temperature and purity on important properties can be studied. Partly because of this, investigation of interfaces in electronic composites has progressed more than for structural composites, resulting in a richer literature. Recently, however, interfaces in structural composites have received increasing attention. One of the more useful techniques for studying bimaterial interfaces is high resolution transmission electron microscopy (HREM) which, in favorable cases, can determine the detailed atomic arrangements at interfaces. This technique has been applied to various metal-ceramic and ceramic-ceramic systems (for a review see, e.g. [1, 2]). Nb-Al203 is an extensively investigated metalceramic "model" system. Nb and Al203 possess nearly the same thermal expansion coefficients, so that interfacial bonding can be studied without the added complication of gross thermal stress. The pioneering work was carried out by Durbin et aL who grew artificial superlattices of b.c.c, metals on different sapphire substrates by MBE [3]. McWhan first stud-
ied the growth of Nb, and other metals, on the (01"[2)Aa2o3 (R) plane of sapphire [4]. Both Durbin et al. and McWhan used RHEED and X-ray diffraction for characterizing the interfaces of these systems [3,4]. HREM study of the Nb-AI203 system was pioneered by Rfihle, Mader, and colleagues [5-11]. The first Nb-AI203 interface studied by this group was made by diffusion bonding [5]. Subsequently Nb-AI203 interfaces were obtained by precipitation of AI203 particles during internal oxidation of a ]~a--Al alloy [6-9]. Later, HREM studies were made on other Nb-AI203 interfaces obtained by MBE growth of single crystal Nb on different sapphire substrates including (0001)~a2o3, (1T00)~a2o3 and (1120)~a2o~[10, 11]. Other HREM studies include that of Knowles et al. who investigated the atomic structure of the (001)Nb/(01T2)~a203 interface [12]. The aim of the present work was to characterize the orientation relationships, atomic structures, and interface reactions of several interfaces in a metalceramic multilayer grown by MBE. This is important because, as McWhan has suggested, MBE growth appears to reflect the three-dimensional symmetry of the substrate rather than the symmetry of the surface [4]. The (001)Nb/(01T2)~a203 interface studied in the present work is similar to the interface studied by Knowles et al. [12]. This orientation is of special interest in connection with the three-dimensional relationship which links the Nb and A1203 lattices when Nb is grown on principal sapphire planes at an optimal temperature. In particular, (001)Nb grOWStilted by 2.5-3 ° from (01"[2)m2o3 in order to conform with this relationship [3,4]. In addition to the (001)Nb/ (01T2)m2o3 interface, a previously unstudied metalceramic interface, Nb-MgO, was also investigated here. For this purpose, a multilayer of Nb-MgONb-A1203 was grown by sequential deposition of Nb and MgO films on a sapphire substrate. Subsequently, the multilayer composite was studied by selected area electron diffraction (SADP), conventional transmission electron microscopy (CTEM) and HREM. 2. EXPERIMENTAL Nb and MgO were sequentially deposited on the (01"[2) (R) plane of an epitaxial grade sapphire substrate by MBE to produce Nb-MgO-Nb-A1203 multilayer stacks. The thicknesses of Nb and MgO films were, respectively, ~100 and ~60nm. The substrate was preheated and cleaned by annealing at 1150°C for one hour in a vacuum of about 10- ~0torr. The Nb and MgO films were deposited at about 0.05 nm/s in the same vacuum onto the sapphire substrate which was kept at a temperature of 920°C for Nb and 650°C for MgO [13]. The cross-sectional HREM specimens were prepared by the following procedure. First the multilayer stack was oriented by the X-ray Laue back-reflection technique from the sapphire substrate side, and
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LI et aL: CHARACTERIZATION OF Nb-Al203 AND Nb-MgO INTERFACES parallel slices were cut by a diamond saw in the proper orientation.Then, two pieces of the sapphire substrate carrying the N b and M g O films were glued together, with the film surfaces facing each other, using M-bond 610 adhesive.t These "sandwiches" were cut in strips 1.5 mm long and embedded in conducting epoxy contained in a copper ring. The cross-sectional specimens were prepared from these couples using standard techniques, involving mechanical grinding to a thickness of 0.13 mm, dimpling to a thickness of 20 #m and ion beam milling at 6 kV [14]. The cross-sectional TEM specimens were prepared with the foil normal parallel to the [220]]~,o3 direction of the sapphire substrate. The samples were subsequently examined in a top-entry JEOL 4000EX high resolution electron microscope with a point-to-point resolution of ~0.18nm. Image simulations were performed using the SHRLI programs developed by O'Keefe [15] for the operating parameters of the JEOL 4000EX electron microscope, viz. coefficient of spherical aberration Cs = 1.0mm; defocus spread 6f= 8 rim; divergence angle half-width =--0.75mrad. The Scherzer defocus at 400 kV for this microscope is ~ - 50 nm. In order to overcome specimen charging which causes image drift and, sometimes, microscope instabilities resulting in discharges, the top surface of the specimen was covered with a 400 mesh copper grid after it was placed in the microscope sample holder.
Fig. 1. Bright-field image of Nb-MgO-Nb-AI203 multilayer.
electron beam simultaneously parallel (or nearly parallel) to the zone axes of Nbt, Nby, M g O and Al203. A schematic diagram of this SADP, indexed with respect to the four crystals,is shown in Fig. 2(b). According to this composite diffraction pattern, deposition on the (0112) (R) plane of sapphire
3. RESULTS AND DISCUSSION 3. I. Growth morphology and orientation relationship
Figure 1 is a low magnification cross-sectional TEM micrograph of the multilayer showing the growth morphology of the Nb and MgO films. The thicknesses of the Nb and MgO layers determined from bright-field TEM micrographs were 106 and 63 nm, respectively. The niobium film deposited on the sapphire substrate is denoted as Nb I while the niobium film deposited on MgO is designated as Nb2. As Fig. 1 shows, three interfaces are present in the multilayer: one between the Nb deposited on the sapphire substrate (the Nbi-Al203 interface), one between the MgO deposited on the Nb film (the MgO-Nb I interface), and one between the Nb deposited on the MgO film (the Nby-MgO interface). Electron diffraction studies of the multilayer showed that each layer was deposited in an epitactic manner with respect to its corresponding substrate. The ¢pitactic orientation relationships between Nbl-Al203, MgO-Nbl, and N b 2 - M g O were determined from SADPs. These are shown in Fig. 2(a) which shows a composite S A D P obtained with the tMeasurement Group, Inc., Raleigh, North Carolina.
, ~ _ ~ ) M g O (~~-
• Nbl
01"lN.b21~'ONb I
O Nb2
2 0ONbl
• MgO [] AI203
Fig. 2. SADPs showing the orientation relationships between N'b2, MgO, Nb= and A1203: (a) SADP along the [lll]~r~,[011]Mt~,[001]Nb~and [~20r]A~o3zone axes;(b) a schematicdiagram of the SADP indexedwith four crystals.
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results in the following epitactic orientation relationships in the multilayer
(110)Nb2 II(100)MgO II(100)Nbl [I(0112)AI~O, [11 l]Nb2II[01 l]~go II[001]Nb, l[[520I]~a2o,
(1)
In the above, we have used the convention that the parallel planes in the two crystals are taken to be those which are parallel to the interface, and the parallel directions lie in the two parallel planes. In fact, in Fig. 2, the [001] Nb~ zone axis is misoriented by about 2.4 ° with respect to the [~20I] zone axis of sapphire, in agreement with earlier studies [3, 4]. In the following each of the three systems NblA1203, MgO-Nb~, and Nb2-MgO will be discussed separately. 3.2. Interfaces in the N b - M g O - N b - A l e O ~ multilayer 3.2.1. The Nb I-AI eO ~ interface. Bright-field micrographs (see Fig. 1) and HREM observations indicated that the Nb~ film deposited on the (0112) plane of sapphire had the following epitactic orientation relationship
(100)~, 11(0112)~a,o,
[01 l]Nb~[I[2110]~a2o3
(2a)
(roughly) equivalent orientation relationship would be A
(100)Nb, II(0112)~a~o,
[00l]Nb, II[220I]~a2o,
plane of sapphire substrate using a magnetron sputter deposition technique. Consequently, all the thin films prepared by vapor deposition of Nb on sapphire (MBE or sputter deposition) are characterized by a unique orientation relationship between Nb and A1203 [equations (2) or (3)]. Note, however, that the Nb-A1203 interfaces prepared by internal oxidation gave a different orientation relationship [6, 7]
(110)N b II(0001)~O3 [110]Nb II[2H0ba~o~
However, in this case, apparently the precipitates were not c~-Al203 but rather a closely related material denote as ~'-A1203 by Mader [6]. The X-ray studies on the Nb--A1203 system showed that Nb could be deposited epitaxially on (R) plane of sapphire with either a (100) or (110) texture depending on the substrate temperature [4]. Of these two possible orientations, the (100) films fabricated at the higher deposition temperature of 850°C was a more perfect single crystal [4]. It is worth noting that silicon is also usually grown by CVD or MBE on the (01T2)ju2o3 (R) plane of sapphire to produce silicon-on-sapphire devices in the electronics industry (see, e.g. [17]). In contrast to b.c.c, niobium, silicon has a diamond cubic structure with an f.c.c, lattice. This system, which has been extensively studied, shows the following epitactic relationship [18]
(2b)
(100)s i II(01 À2),a=o,
Orientation relationships (2a) and (2b) are only approximately equivalent because the angle between [001]m and [011]Nb is 45 ° while the angle between [220I~2o~ and [~110]~a2o3directions is 47.1 °. It should also be mentioned that, despite the predominantly single crystalline nature of the Nb~ film, some small misoriented Nb grains were observed within the film. The observed orientation relationship for MBEgrown Nb film on the (0112) plane of sapphire is equivalent to the results previously obtained for the MBE growth of Nb on sapphire substrates [3,4, 10, 11]. This orientation relationship has been confirmed by HREM for Nb films deposited on differently oriented sapphire substrates: (0001)Aa2o3, (lI00)~a2o~ and (1~10)~a2o3[10, 11]. We note that in the work of Mayer et ai. [10, 11], the orientation relationship is written as
[001]si 11[~110ba2o3
(111)N b II(0001)m~o, [1 ]'0]r~b H[2TI0]~a2o3
(3)
which is nearly equivalent to equation (2); the angle between (0112)~a~o, and (0001)~a~o, is 57.6 ° while that between (100)r~b, and (11 l)Nb is 54.7 °. Knowles et al. [12] also reported the same orientation relationship, equation (3), for Nb films deposited on (0112) (R)
(4)
(5a)
or, sometimes [19] (100)si II(0112)~o3 [001]si II[0111]A12O3.
(5b)
3.2.2. Misfit dislocations at the NbI-AleO 3 interface. For interpretable HREM images of interfaces,
the electron beam should be parallel to low index zone axes of the two crystals and, also, parallel to the interface plane (see e.g. [20]). However, as mentioned above, the [001]Sb, zone axis of Nb~ was not precisely parallel to the [~20I]m2oJ zone axis of sapphire but rather was misoriented from it by ~ 2.4 °. Hence the low index planes in the two crystals, e.g. (100)Nbt and (0112)~2o 3, could not be simultaneously oriented parallel to the electron beam and a compromise had to be made in orienting the interface for high resolution imaging. It was found best to orient the [220I],j~o, zone axis of sapphire as accurately as possible parallel to the electron beam and allow for the slight misorientation of the corresponding Nb I zone axis (i.e. the [001]Nb~). Such a HREM mierograph is shown in Fig. 3 at a relatively low magnification. The interface plane is parallel to (0112)~2o, and approximately parallel to (100)Nbt. Well-defined atomic rows and
LI et aL: CHARACTERIZATION OF Nb-AI203 AND Nb-MgO INTERFACES
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Fig. 3. HREM image of a Nbt-AhO 3 -mtetfa~, The interface plane is parallel to (I00)~, and (01'D,)~o, !tlad the incident electron beam is parallel to [001]w~,a n d ~ 0 ~ o 3. The interface is marked by double arrows and a step is indicated by the symbol S. lattice planes can be identified. The NbI-AI203 interface lying parallel to (01T2)~2o3 planes, is atomically flat, sharp, and free of any interfacial phase. Interfacial steps of height equal to the (01"[2) interplanar spacing are clearly visible; one of these is indicated by S. The average separation between steps is about 30nm. To study the film/substrate coherency, the continuity of close-packed planes, d~ and d2, in the two crystals was examined [21]. In the [001]Nb1111220I]~O, projection of Fig. 3, the traces of closest-packed planes in Nbl and Al203 are (110)Nbl and (1120)~2o 3 respectively, which are inclined to the interface at angles of 45 ° and 43 °, respectively. Hence, in order to evaluate the misfit parameter 2(d2 - d, ) f = - -
(6)
the projections of(1 lO)Nb, and (1120)~a~o3planes on the interface, i.e. d l - - d ~ ) c s c 4 5 ° and d2~-d~?~csc43 ° were compared. Using this criterion, a lattice mismatch of 5.5% is obtained. It would be expected that this misfit strain be accommodated by an array of misfit dislocations of Burgers vectors b spaced a distance s apart at the interface. However, contrary to expectations, the experimental micrograph in Fig. 3 does not show any interfacial misfit dislocations. A clearer view is shown in Fig. 4, which is a higher magnification HREM image of the Nbi-Al203 interface. The clearly resolved white dots in the image of both crystals around the interface shows the smooth continuity of (ll0)sbl planes with the (ll~0)Aa2o3 planes. The absence of misfit dislocations, and the smooth continuity of the close-packed planes at the interface implies that atomic matching between Nb, and A1203 has been achieved locally by a distortion of the Nb lattice in the vicinity of the interface (i.e. this region of Nb is pseudomorphic with the sapphire substrate). This has been achieved, presumably, by the elastic distortion of the region of the Nb lattice next to the interface. It should be mentioned that the AI20 s lattice, in contrast to that of Nb, is undisturbed up to the interface. This result is not surprising, as
Fig. 4. A high magnification HREM image of the Nb,AIzO3 interface projected along [001]~, and [220T]~a2o~.The insets are computer simulated images of perfect crystals of Nb and A1203(thickness-- 18 rim, Af= - 6 0 nm). The unit cell of Nb is outlined.
A1203 possesses much higher elastic moduli than Nb. Although no misfit dislocations are visible at the interface, there are a series of dislocations at a small distance from the interface in the Nb lattice. This is in agreement with the work of Rfihle, Mader and colleagues [6-9], who first reported the existence of "stand-off" dislocations in the Nb-A1203 system obtained by internal oxidation. In Fig. 5, the core regions of some of these dislocations close to the Nb1-A1203 interface are indicated by inverted "T" symbols. The extra half plane of {110}sb forming the dislocations does not reach the interface but ends at a distance of about eight {ll0}Nb lattice planes. In fact, the stand-off distance from the interface varies from one place to another, but, on the average, it is about seven {200}Nb interplanar spacings. These dislocations are found to be spaced approximately
Fig. 5. HREM image of Nb,-Al203 interface with the incident beam parallel to [001]wo, and [220I]~2o3. The core regions of dislocations are indicated by inverted "T" symbols. The insets are computer simulated images of perfect crystals of Nb and Al203 (thickness = 21 nm, Af = - 60 nm).
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periodically, about one every twelve (020)~, planes (i.e. 2 rim) which is sufficient to accommodate the lattice mismatch between Nb and A1203. Therefore, while the re#on of niobium next to the interface is highly strained, the rest of the Nb lattice at a distance of 1.2 um from the interface is strain-free and the elastic strains are accommodated by "stand-off" misfit dislocations. Given the findings of stand-off dislocations at Nb--AI203 interface, Hirth and collaborators suggested that, because of mismatch in elastic moduli in a bimaterial, there is a repulsive image force on the interracial dislocations, tending to push them away from the interface into the material with the lower elastic modulus [22]. They calculated the equilibrium "stand-off" distance of misfit dislocations at an interface and obtained remarkably good agreement with the experimental results of Mader et al. [6-9] on the Nb-AI2 O3 system. It is interesting that Mayer et al. [10, 11] report that the misfit dislocations observed in their MBE-grown Nb films do not show "stand-off" positions. They attribute this to the kinetics of growth, i.e. to the fact that MBE deposition of Nb on sapphire was performed at lower temperatures compared to the manufacture of Nb-AI203 interfaces by either diffusion bonding or internal oxidation. Thus, they suggest that the geometrical arrangement of misfit dislocations at the interface in the MBE grown system might not be thermally stable, and that the dislocations could climb from the interface to stable stand-off distances if the samples are heat-treated or deposited at higher temperatures. This suggestion, however, implies that the misfit dislocations are generated at the interface' and, because of the low growth temperature, cannot climb to their equilibrium "stand-off" positions away from the interface. The generation of misfit dislocations at the interface may happen, for example, at the edges of the island deposits in a three-dimensional growth situation [23]. More commonly, however, misfit dislocation generation takes place by the nucleation of dislocation semi-loops at the film surface and the glide of these t o t h e interface (see, e.g. [24]). On reaching the interface, the interfacial segments of these ghssile lattice dislocations become "misfit" dislocations. In the case of Nb-AI203, the mechanism by which misfit dislocations are generated is not clear. However, if they are generated as lattice dislocations at the film surface, then it is possible that they never reach the interface but rather stop their glide when they reach the equilibrium distance at which there is no further driving force. This is an alternative to the one discussed by previous investigators, who suggested that dislocations are generated at the interface and then climb to their equilibrium position [10, 11].
It is also noteworthy that, in the present investigation, the misfit dislocation do show "stand-off"
r
Fig. 6. The core region of a dislocation in Nb I close to the Nb l-Al20 ~interface. A Burgers circuit is performed around a dislocation. positions, similar to the case of (110)Nb II(0001)m2o3 interface manufactured by diffusion bonding. This is despite the fact that, in the present MBE growth, the substrate temperature wag 920°C, similar to the temperature used in the experiments of Mayer et al. [10, 11]. It is clear that issues associated with "standoff" dislocations are unresolved. An attempt was made to determine the Burgers vector of the "stand-off" dislocations near the N b A12Oa interface. In Fig. 6 a Burgers circuit is drawn around the image of one of the "stand-off" dislocations. Since the projected Burgers vector of the "standoff" dislocation determined from the HREM image (Fig. 6) is ½[110]Nb,it is likely that the actual Burgers vector is ½(liD, as the latter is the usual lattice dislocation in a b.c.c, crystal. The insets in Fig. 4 are computer simulated images for both the Nb and the AlcOa at a foil thickness of 18 nm and a refocus Af = - 60 nm. The simulated images with these values of thickness and defocus agree well with the observed HREM images. The bright spots in the HREM images on the two sides of the interface correspond to AI atom positions in sapphire and Nb tunnels in the niobium, respectively. In sapphire, the aluminum and oxygen atoms are located on two different, but adjacent (01T2) atomic planes, which are shown in Fig. 7(a). In Fig. 7(b), the (a) •
•
•
j
•
•
•
•
•
•
•
•
•
•
;...;.;.; •
•
.
;.;.;.;.o •
o
•
AI
•
o
.
; . o . ; . ; . ; •
(b) •
•
Oaypn
Fig. 7. (a) The atomic configuration in a (01]'2) plane of sapphire, (b) the atomic configuration in a (100) plane of Nb.
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a/.: CHARACTERIZATION OF Nb-A1203 AND Nb-MgO INTERFACES
Nb atom positions on the (100) plane of Nb are shown at the same scale as in Fig. 7(a). A comparison of these two figures shows the similarity of a (!O0)~r~ plane to a (0112)~azo~ plane: with soqae small ahifls, the AI atoms on the (0!T2)Ahosplane correspond well to the Nb atoms on the ( 1 ~ ) ~ plane, Attempts were made to determine the atomic structure of Nbz-Al203 interface from its H ~ images of the interface. Although totally satisfactory image simulations of the Nb-A1203 interface have not yet been obtained, Fig. 7 suggests that it is likely that an oxygen layer is adjacent to the (100)Nb plane. Because an explanation of the three-dimensional relationship between A1203 and epitaxially grown Nb must presumably lie in the relationship of the Nb atom to the O sublattice of sapphire, the eventual elucidation of the detailed interface structure is a matter of some considerable interest. 3.2.3. The MgO-Nb~ interface. The MgO film grown at 650°C on the Nb~ layer was single crystal, with an epitactic orientation relationship given by equation (1)
(lO0)MgO [[(100)Nb, [011]MsO II[001]Nb,
•
(5)
Figure 8 is a H R E M image of thisinterfaceobserved along the [011]MgO and [001]Nb,zone axes, and shows a "saw-tooth" morphology with the two lattices jutting into each other. The average M g O - N b ~ interface plane is tiltedby 3.5° with respect to (100)Nb,. A close inspection of the structural details of the interfaceshows that this "saw-tooth" interfacearises because of faceting,with a number of latticeplanes on the two sidesof the interfacejoining each other in a smooth fashion (see inset of Fig. 9). The interface at one side of the corner is formed by (IIT)MsO, while the interfaceat the other side is formed by (IT1)MsO; these are continuous with (ll0)Nbl and ('[10)Nbl planes, respectively.The development of {I I I}UtO facets at the M g O - N b ~ interfaceundoubtedly took place after growth, as evidence for the faceted structure was absent in R H E E D studies of the M g O grown at 650°C. It is likelythat thisinterfacefaceted during the latergrowth of Nb2 at 920°C. This struc-
Fig. 8. HREM image ofa MgO-Nb I interface viewed along [011]MgO and [001]Nb,. The misfit dislocations are indicated by small dots.
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Fig. 9' A hillh ~ HREM ~ o f ~ MgO-Nb, interface o~-raed along [011]~o and [O01]m,. The misfit dislocatiom are indicated by an inverted "T" symbol. The HREM image of the faeeted MgO-Nb I interface is shown in the top inset. The insets are computer simulated images of perfect crystals of MgO and Nb (thickness = 10 nm, Af-- -60nm). The unit cell of each crystal is outlined. ture may minimize the interfacial energy, as { I I I } are the close-packed oxygen planes in MgO, even though the plane of easy cleavage is (001). Note that the faceted interface, with facets parallel to (l IT)MsOand (ITl)MsO , accounts for a 141% increase in the total interface area compared to a fiat (100)MiO/(100)Nb, interface. Therefore, the interracial energy of the {111}MsO faceted MgO-Nbl interface is significantly lower than that of the (100)~O/(100)Nbl interface. Note also that strong strain contrast is absent at the (11])MsO facets, while strong strain contrast appears at the (ITl)MsO facets. In the [011]MsOI[[001]Sb~ projection of Fig. 9, the traces of closest-packed planes in MgO are (1T 1)MsO, while in Nb, they are (T10)Nb. The (IT1)MsO and (TI0)Nbl planes are inclined to the interface at angles of 54.7 ° and 45 °, respectively. Therefore, the projection of (I[I)Mso and (]10)Nbl planes on the interface, i.e. d I = d ~ c s c 5 4 . 7 ° and d2 = d~)csc45 ° were compared in evaluating the misfit parameter of the (100)MsO/ (100)Sb, interface, which gives a lattice mismatch of 10.4%. The misfit dislocations are shown by small dots in Fig. 8 and inverted "T"s in Fig. 9. Unlike the Nb-AIzO3 interface, these misfit dislocations lie at the MgO-Nb~ interface, spaced periodically about one every nine (1T1)MsOplanes (i.e. 2.7 nm). In order to determine the Burgers vector of the dislocations, a Burgers circuit was constructed as shown in Fig. 10. From this Burgers circuit, the projected Burgers vector was ~[~11]MsO, which makes this a 60 ° dislocation. The actual Burgers vector is therefore probably ½[101] or I[1]0]. On geometrical grounds, using such a Burgers vector, a regular array of misfit dislocations separated by five (lIl)MsO planes (i.e. 1.5 nm) would he required to completely accommodate the lattice mismatch of 10.4%. As the observed average distance between misfit dislocations is 2.7 nm [i.e. one every nine (l'[1)M~t~ planes], the large mismatch of 10.4% between (IlI)MsO and (II0)sb, at the interface does not appear to be completely accommodated by misfit
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Fig. 11. HREM image of the Nb2-MgO interface viewed along [lll]N~ and [011]MsO. The misfit dislocations are marked by small dots and the interface is marked by double Fig. 10. A Burgers circuit around a misfit dislocation at the MgO-Nb I interface.
arrows.
very nearly parallelto the (ITI)MgO facets.However, the (01"[)Nb2planes are misoriented by ~10 ° with dislocations.A n explanation of such a large discrep- respect to (Ill)Mso. ancy between the expected and observed average In general, the facet lengths arc not equal and are separation of misfit dislocations could be due to often associated with interracialstcps. The sense of the faceting of the interface although the R H E E D this facet asymmetry, and hence the sense of the studies referredto above suggest that facetingoccurs steps, is the same for nearly all facets. These steps afterdeposition.Taking the continuityof (IT l))~o and accommodate the 3.3° average deviation between the (~10)blb,planes at the (lII)MgO facet,the mismatch is (100)MgO plane and the N b E - M g O interface plane. only 4%, which may be assumed to be the "effective" A large local deviation results cither in large steps associatedwith one or two facetedcomers, or in small mismatch at the (100)MgO/(100)Nb,interface. The insetsin Fig. 9 are computer simulated images steps associated with a few short faceted segments. of M g O and N b I for a thickness of 10 n m and a In the [II l]Nb~II[01 I]MsO projection of Fig. I I, sets defocus Af= -60 nm. The simulated images agree of close-packed planes which arc compared in N b well with the observed H R E M images. The bright and M g O are (IIT)MgO (the oxygen close-packed spots in the H R E M images of the M g O and Nb~ planes) and (0]'l)Nb2.These planes arc inclinedto the represent the positionsof M g and N b atom columns, interface at 54.7° and 60 ° respectively.Therefore, respectively. the projection of (I l]')MgOand (0TI)Nb2planes on the 3.2.4.The Nb2-MgO interface.The Nb2 filmgrown interface, i.e. d~ = d ~ c s c 5 4 . 7 ° and a2 = d~0b~)csc60° on the M g O layer was not a single crystal, but are compared in evaluating the misfit parameter, contained a characteristicsubgrain cellstructure,with which gives a lattice mismatch of 9.8%. The core an average celldiameter of about 70 n m and a strong regions of a few dislocations are indicated by small texture.The N b 2 - M g O orientationrelationshipgiven dots in Fig. 11 and inverted "T"s in Fig. 13. Unlike in equation (1), i.e. the MgO-Nb I interface, these misfit dislocations are not located at the interface itself and again show (110)Nb2 II(100)Mgo
[111]Nbz[I[011]MaO
(6)
refers to this preferential orientation of the textured Nb 2 film with respect to the MgO layer. Figure 11 is a HREM image of the Nb2-MgO interface observed along the [111]m2 and [011]Mgo zone axes. The average interface plane is tilted by 3.3 ° with respect to the (100)UgOplane. Steps with heights equal to one to four (200)MsOinterplanar spacing were frequently observed. The HREM images in Figs 11 and 12 show some of these steps and also the faceting of the interface along two sets of MgO lattice planes. The flat region of the interface is parallel to (lI0)Sb2ll(100)UgO planes, while the faceted regions Fig. 12. HREM image of the Nb2-MgO interface viewed of the interface are parallel to (1 lI)M~o or (lll)MaO along [11 l]Nb2and [011]Mso.The faceted interface is marked planes. The (10I)Nb, planes of the Nb film are in turn by small dots.
LI et al.: CHARACTERIZATION OF Nb-A1203 AND Nb-MgO INTERFACES
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Fig. 13. HREM image of the Nb2-M ~ ~ a c e viewed along [111]~ and [011]MsO.The core regionsof dislocations are marked by the inverted "T" symbol. The insets are computer simulated images of perfect crystals of Nb and MgO (thickness = I0 nm, Af = - 6 0 nm).
(b)
•
•
•
a "stand-off" position within the Nb lattice, similar to that of the Nb~-A1203 interface. They are spaced periodically about one every fourteen (0TT)Nb2planes (i.e. 3.8 rim). From Burgers circuits drawn around the core region of the misfit dislocations, the projected Burgers vector of the latter was ~[12T]Nb(see Fig. 13). On geometrical grounds, using this Burgers vector, a regular array of misfit dislocations separated by ten (0]'[)Sb2 planes (i.e. 2.8 nm) would completely accommodate the lattice mismatch. As the observed average distance between misfit dislocations is 3.8 nm [i.e. one at every fourteenth (0TT)N~ plane], the large m i s m a t c h of 9.8% between (111)M~o and (0TI)N~planes at the interface does not appear to be completely accommodated by misfit dislocations. It is worth noting that the MgO-Nbl interface formed by deposition of MgO film on the Nbl layer, and the N b , - M g O interface formed by deposition of Nb film on the MgO layer, have very different characteristics. The misfit dislocations were located at the interface in the first case while they were "standoff" dislocations in the second case. In addition, the MgO-Nbl interface was completely faceted parallel to the close packed (oxygen) planes of MgO, i.e. {111}Mo . On the other hand, the Nb2-MgO interface consisted of fiat regions parallel to (lI0)wo2-(100)Mso planes and faceted regions parallel to {lll}Mso. These differences might be related to the conditions of growth: in one case Nb is the substrate and MgO is the deposit, while in the other case, MgO is the substrate and Nb is the deposit. In terms of the misfit dislocation arrangement in the two cases, dislocations can presumably be nucleated more easily in Nb than in MgO. More importantly, the MgO substrate upon which the Nb 2 film grows may not be so perfect as the Nbn film upon which the MgO grows. Finally, the difference in growth temperature (920°C for Nb, 650°C for MgO) may be important. However, the origin of these differences needs further study. The insets in Fig. 13 are computer simulated images of MgO and Nb 2 for a thickness of 10 nm and
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Fig. 14. (a) HREM image of the Nb2-MgO interface region, showing a rotated microdomain of MgO. (b) Schematic diagram of the microdomain structure at the Nb2-MgO interface region. a defocus value of Af = - 60 nm. There is again good agreement between the simulated images and the observed HREM images of the MgO and the Nb2. As in the case of MgO-Nbl, the bright spots in the HREM images of MgO and Nb2 represent the positions of Mg and Nb atom columns, respectively. Very narrow regions with a domain structure were also observed at the Nb2-MgO interface, as shown in Fig. 14(a). One possibility is that this narrow region is NbO. However, since the lattice parameters of NbO, which is also f.c.c., and MgO are quite similar, and the region is too narrow to obtain clear compositional information, it has not been possible to identify the nature of the narrow interracial region. One possibility is that the narrow region contains a rotation microdomain of MgO, designated as (MgO)2, in Fig. 14: this microdomain is rotated by 54.7 ° around the [011]ugo axis with respect to the other domain, (MgO)l, as shown in Fig. 14(b). 3.3. Interfacial reaction during H R E M irradiation
During the course of HREM investigations, time-dependent changes occurred in the Nb~-A1203 interface region. In order to monitor this electron
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LI et al.: CHARACTERIZATION OF Nb-AI203 AND Nb-MgO INTERFACES
Fig. 15. HREM images of the Nbi-Al203 interface showing the evolution of the interface with electron beam irradiation: (a) initial NbI-AI203 interface; Co) a modified layer formed after 210 minutes of electron beam irradiation. The incident electron beam is parallel to the [O01]NblII[220I]m2o,direction.
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beam-induced solid-state reaction, in situ H R E M observations were performed on cross-sectional N b r A l 2 0 3 specimens as a function of the duration of electron beam irradiation. During the H R E M observations, an electron beam with a current density of 80 pA/cm 2 (read from the fluorescent screen of the JEOL 4000EX electron microscope) was focused on the same area of specimen. It can be seen from Figs 3, 4 and 5 that the original NbI-AI203 interface is perfectly smooth, sharp, and free of any interfacial phase in the initial stages of H R E M observation. After electron beam irradiation, however, a narrow region with some modification of the interface was easily recognizable. Figure 15 is a H R E M image of the interface after electron beam irradiation for 210 rain. A 2.1 rum layer with slightly different image characteristics can be distinguished between the N b 1 film and the sapphire substrate.
Fig. 17. HREM image of the Nbl-Al203 interface showing grain growth during electron beam irradiation: (a) initial small grain embedded in the Nb I film at the Nbl-A1203 interface region, (b) the same grain in the Nb I film after 210 min of electron irradiation. As seen in Fig. 15, the thickness of the region which has been modified is not uniform all along the interface, and, in fact, its HREM image is quite similar to that of the AI2Oa substrate. Consequently, we have not been able to positively identify the nature of change which has taken place. The kinetics of such a modification of the interface can be determined by varying the duration of electron beam irradiation. Figure 16 shows the change in the
LI et al.: CHARACTERIZATION OF Nb-Al203 AND Nb-MgO INTERFACES thickness, x, of the modified layer, as determined from H R E M images, with time, t. Although relatively large error bars are associated with the experimental points, it is reasonable to fit a parabolic curve through the points, i.e. the thickening of the modified layer follows a parabolic law x 2 = K ( t - to) where K is the rate constant for the reaction and to is the incubation period. The parabolic fit to the experimental results (Fig. 16) gives K ~ 4 A 2 rain- l and to = 50 rain. Spinel formation by the solid state reaction of MgO and AI203 at a MBE grown MgO-AI203 interface under electron beam irradiation has recently been investigated, and the orientation relationships, and kinetics and mechanisms of spinel formation under electron irradiation has been discussed in detail in a previous paper [14]. Growth kinetics of the modified layer at the Nbl-Al203 interface under electron irradiation is very similar to that of MgAI204 formation at the MgO-AI203 interface. However, as' mentioned previously, the nature of the modified layer at the interface is not known. Electron beam-induced grain growth in the Nb 1 film was also observed. Figure 17(a) is a HREM image of a small grain embedded in the Nbl film at the Nbt-AI203 interface region. Figure 17(b) shows the migration of the grain boundary which occurred during electron irradiation. In this figure, the original interface between Nb~ and the embedded grain is indicated by dashed lines and the new interface after electron beam irradiation for 210 rain is marked by small dots. It is clear that the Nb~ grain consumes the embedded grain under electron irradiation.
4. CONCLUSION 1. The orientation relationships between Nb2, MgO, Nbt and AI203 layers are:
( 110)Nb2 II( 100)MsO II (100)Nb, II(01T2)~a2o, [111]N~ II[011]MsO II[001]Nb, II[2201],a2o3
•
2. The initial Nb1-A1203 interface produced by MBE growth is atomically flat and sharp. The misfit dislocations are not located at the interface; there are "stand-off" dislocations in the N b lattice at a distance of about seven {200}Nb interplanar spacings from the interface. These are spaced periodically about one every twelve (020)Sb planes. 3. The MgO-Nb~ interface was faceted with a "saw-tooth" morphology; the faceting resulted in near-parallelism of close-packed planes of Nb] and MgO (the oxygen close-packed planes), i.e. { 111 }uao [[{ 110}Nb,. The faceting most probably took place after the growth of MgO, presumably because of the lower N b - M g O interfacial energy along { 110}Sb and {111}uso. Misfit dislocations were located at the interface with an average separation of ~2.7 nm. 4. The Nb2-MgO interface had many steps and was faceted with two sets of {I 11}Mgo lattice planes. The close-packed {110}Nb2 lattice planes of Nb, are
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either very nearly parallel, or have a small misorientation, to {11 l}MgO. Misfit dislocations are again not located at the interface but a set of "standoff" dislocations in the N b lattice at a short distance away from the interface are spaced periodically about one every 3.8 nm. 5. A modified layer formed at the Nb~-Al203 interface by a solid state reaction occurring under electron irradiation. The modified layer thickened parabolically. In addition to the formation of a modified layer at the interface, electron beam-induced grain growth within N b was also. observed. ,4cknowledgements--The work at Case Western Reserve
University was supported by the DARPA/ONR University Research Initiative at CWRU under contract N00014-86-K0773. The work at the University of Illinois was supported by the Department of Energy under contract DE-AC0276ERO1198; use of the Materials Research Laboratory facilities at the University of Illinois is acknowledged.
REFERENCES 1. M. Rfihle, A. G. Evans, M. F. Ashby and J. P. Hirth (editors), Metal-Ceramic Interfaces. Pergamon Press, Oxford (1990). 2. These proceedings. ,4cta metall, mater., Vol. 40, Suppl. (1992). 3. S. M. Durbin, J. E. Cunningham and C. P. Flynn, J. Phys. F, Met. Phys. 12, L75 (1985). 4. D. B. McWhan, Mater. Res. Syrup. Proc. 35, 493 (1985). 5. M. Florjancic, W. Mader, M. Rfikle and M. Turwitt, J. Physique 46, C4-129 (1985). 6. W. Mader, Mater. Res. Soc. Syrup. Proc. 82, 403 (1987). 7. W. Mader, Z. Metallk. 80, 139 (1989). 8. M. Kuwabara, J. C. H. Spence and M. R/ihle, J. Mater. Res. 4, 972 (1989). 9. W. Mader and M. Rill'de, ,4cta metall. 37, 853 (1989). 10. J. Mayer, C. P. Flynn and M. R/ihle, Ultramicrosc. 33, 51 0990). 11. J. Mayer, W. Mader, D. Knauss, F. Ernst and M. Rfihle, Mater. Res. Soc. Syrup. Proc. 183, 55 (1990). 12. K. M. Knowles, K. B. Alexander, R. E. Somekh and W. M. Stobbs, Inst. Phys. Conf. Set. 90, 245 (1987). 13. C. P. Flynn, in Metal-Ceramic Interfaces (edited by M. PriMe, A. G. Evans, M. F. Ashby and J. P. Hirth), pp. 168-175. Pergamon Press, Oxford (1990). 14. D. X. Li, P. Pirouz, A. H. Heuer, S. Yadavalli and C. P. Flynn, Phil. Mag..4 65, 403 (1992). 15. M. A. O'Keefe and P. R. Buseck, Trans. Am. Crystallogr. Assoc. 15, 24 (1979). 16. D. X. Li, P. pirouz, A. H. Heuer, S. Yadavalli and C. P. Flynn, Mater Res. Soc. Syrup. Proc. 221, 93 (1991). 17. M. S. Abrahams, C. J. Bniocchi, J. F. Corboy Jr and G. W. Cullen, ,4ppl. Phys. Lett. 28, 275 (1976). 18. K. C. Pans, J. C. Barry, G. R. Booker, T. B. Peters and M. G. Pitt, Inst. Phys. Conf. Ser. No. 76, 35 (1985). 19. J. L. Batstone, Phil. Mag. B 63, 1037 (1991). 20. P. Pirouz and F. Ernst, in Metal-Ceramic Interfaces (edited by M. Rfikle, A. G. Evans, M. F. Ashby and J. P. Hirth), pp. 199-221. Pergamon Press, Oxford (1990). 21. G. B. Olson and M. Cohen, Acta metall. 27, 1907 (1979). 22. S. V. Kamat, J. P. Hirth and B. Carnahan, Mater. Res. Soc. Syrup. Proc. 103, 55 (1988). 23. R. Vincent, Phil. Mag. 19, 1127 (1969). 24. J. W. Matthews, in Epitaxial Growth (edited by J. W. Matthews), Part B, pp. 559-609. Academic Press, New York (1975).
0956-7151/92 $5.00+ 0.00 Copyright © 1992Pergamon Press Ltd
THE CHARACTERIZATION OF Nb-A1203 AND Nb-MgO INTERFACES IN MBE GROWN Nb-MgO-Nb--A1203 MULTILAYERS D. X. LI 1, P. PIROUZ 1 and A. H. HEUER l, S. YADAVALLI2 and C. P. FLYNN 2 IDepartment of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106 and 2Dfpartlnent of Physics and Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana, IL 61801, U.S.A. /dmtrlctwNb-MgO--Nb-A1203 multilayer stacks were prepared by sequential deposition of Nb and MgO films on the (0]'12) (R) plane of sapphire, using Molecular Beam Epitaxy (MBE). Cross-sectional specimens of this multilayer were prepared and various interfaces investigated by high resolution transmission electron microscopy (HREM). Orientation relationships between Nb and AI203 and between Nb and MgO were determined and misfit dislocations, steps and facets at the Nb-AI203 and Nb-MgO interfaces were analyzed. The HREM observations show that, under electron beam irradiation, a modified layer forms at the Nb-AJ203 interface by a solid state reaction. Electron beam-induced grain growth in the Nb film also takes place. g/mmnt---Des empilements de multicouches Nb/MgO/Nb/A1203 sont pr6par~s par d6p6t S&luentiel de films de Nb et de MgO sur le plan (01].2) (R) du saphir, par 6pitaxie par jet mol~daire (EJM). Des 6chantillons de sections droites de ees milticouches sont pr6par~s et diff~rentes interfaces sont 6tudi/~s en micrmcopie 61ectronique par transmission en haute r~tsohition. Les relations d'orientation entre Nb et AI203 et entre Nb et MgO sont d6termin6es et les dislocations de d~,saccord, les marches et les facettes montrent que, sous irradiation 61ectronique, une couche modifi6e se forme ~i l'interface Nb/A1203 par r6action ~ l'6tat solide. I1 se produit aussi une croissasnce des grains indulte par le faisceau d'61ectrons daus le film de niobium. Zammmenfamung---Durch aufeinander folgende Abscheidung yon Nb/MgO-Filmen mittels Molekularstrahlepitaxie werden Nb/MgO/Nb/AI203-Schichtefolgen auf der (0112) (R)-Ebene yon Saphir hergestellt. An Querschnittspriparaten dieser Vielfachschicht werden die verschiedenen Grenz~chen mit hochaufl6sender Elektronenmikroskopie untersucht. Die Orientierungsbeziehungen zwischen Nb und A1203 und zwischen Nb und MgO werden bestimmt; die Feldpassungsvereetzungen, Stufen und Facetten an den beiden Grenzfl~chen werden analysiert. Aulkrdem ergibt sich, dab sich w~thrend der starken Elektronenbestrahlung bei der hochaufl6senden Abbildung eine modifizierte Schicht an der Nb/A]203-C_vrenztiiche dutch eine Festk6rperreaktion bildet. Elektronenstrahi-induziertes Kornwachstum lauft zusatzlich in den Nb-Filmen ab.
1. INTRODUCTION The combination of two or more materials as a way of making a composite has resulted in many new materials with novel properties. One class of such composites is the structural type whereby the purpose of making such a material is to achieve better mechanical properties than the constituent materials from which the composite is made. Such composities are generally made from a combination of metals, ceramics or polymers and the important issues are mechanical properties such as strength, fracture toughness, creep, etc. A second class is the electronic composites in which the constituent materials are either different semiconductors, or combinations of a semiconductor and an insulator (or a metal). The properties of interest in the latter type of composites are the electronic properties where they may have unique properties such as two-dimensional conductivity.
The methods employed in the production of different types of composites are often very different. In the case of electronic composites, techniques such as Chemical Vapor Deposition (CVD), Liquid Phase Epitaxy (LPE), or Molecular Beam Epitazy (MBE) are used where a film of one material is deposited from the vapor or liquid phase onto a different substrate of another material. F o r structural composites, the method of production is often a technique such as diffusion bonding, eutectic bonding, liquid infiltration, internal oxidation/reduction, etc. Because of the cleanliness of the systems and the control of growth parameters such as temperature and pressure, production of electronic composites is usually better controlled than is production of structural composites. Such strict control is essential in electronic composites because small changes in the purity of the crystals, or small variations in the growth parameters, can have drastic effects on the electronic properties of the composite, while in structural composites, small changes
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in the crystal purity and manufacturing parameters can generally be tolerated. In addition, because of the dimensional scale of the different types of composites, economical factors can play a determining role in the choice of the technique of composite production. An important part of any composite structure is the interface between its different constituents. This problem is well known in electronic composites where considerable effort has been invested in the investigation of interface structure. X-ray diffraction, electron diffraction, and transmission electron microscopy (TEM) are the usual methods employed for investigating the interface structure. An interface may have its own unique properties, which often will have a controlling influence on the properties of the composite as a whole. Thus, in a ceramic-matrix (structural) composite, the interface between two strong ceramics may be weak because of weak chemical bonding between the two constituent ceramics. In this case, the fracture toughness of the composite is controlled primarily by the weak interface rather than the toughness of the constituents. A different case is the presence of a conducting interface in a semiconductor-on-insulator electronic composite (such as silicon-on-sapphire, SOS) which dramatically changes the electronic properties of the material. The interface properties often depend on the atomic structure of the interface, For example, it is widely believed that the Scbottky or ohmic behavior of a semiconductor junction is determined by surface states arising from strained or broken bonds at the interface. Another example is the mechanical strength of an interface in a structural composite, which may be determined by the arrangement of the different atomic species on the two sides of the interface and the way that they bond to each other. The investigation of a composite interface is much easier when it is made in a controlled fashion, so that effects of temperature and purity on important properties can be studied. Partly because of this, investigation of interfaces in electronic composites has progressed more than for structural composites, resulting in a richer literature. Recently, however, interfaces in structural composites have received increasing attention. One of the more useful techniques for studying bimaterial interfaces is high resolution transmission electron microscopy (HREM) which, in favorable cases, can determine the detailed atomic arrangements at interfaces. This technique has been applied to various metal-ceramic and ceramic-ceramic systems (for a review see, e.g. [1, 2]). Nb-Al203 is an extensively investigated metalceramic "model" system. Nb and Al203 possess nearly the same thermal expansion coefficients, so that interfacial bonding can be studied without the added complication of gross thermal stress. The pioneering work was carried out by Durbin et aL who grew artificial superlattices of b.c.c, metals on different sapphire substrates by MBE [3]. McWhan first stud-
ied the growth of Nb, and other metals, on the (01"[2)Aa2o3 (R) plane of sapphire [4]. Both Durbin et al. and McWhan used RHEED and X-ray diffraction for characterizing the interfaces of these systems [3,4]. HREM study of the Nb-AI203 system was pioneered by Rfihle, Mader, and colleagues [5-11]. The first Nb-AI203 interface studied by this group was made by diffusion bonding [5]. Subsequently Nb-AI203 interfaces were obtained by precipitation of AI203 particles during internal oxidation of a ]~a--Al alloy [6-9]. Later, HREM studies were made on other Nb-AI203 interfaces obtained by MBE growth of single crystal Nb on different sapphire substrates including (0001)~a2o3, (1T00)~a2o3 and (1120)~a2o~[10, 11]. Other HREM studies include that of Knowles et al. who investigated the atomic structure of the (001)Nb/(01T2)~a203 interface [12]. The aim of the present work was to characterize the orientation relationships, atomic structures, and interface reactions of several interfaces in a metalceramic multilayer grown by MBE. This is important because, as McWhan has suggested, MBE growth appears to reflect the three-dimensional symmetry of the substrate rather than the symmetry of the surface [4]. The (001)Nb/(01T2)~a203 interface studied in the present work is similar to the interface studied by Knowles et al. [12]. This orientation is of special interest in connection with the three-dimensional relationship which links the Nb and A1203 lattices when Nb is grown on principal sapphire planes at an optimal temperature. In particular, (001)Nb grOWStilted by 2.5-3 ° from (01"[2)m2o3 in order to conform with this relationship [3,4]. In addition to the (001)Nb/ (01T2)m2o3 interface, a previously unstudied metalceramic interface, Nb-MgO, was also investigated here. For this purpose, a multilayer of Nb-MgONb-A1203 was grown by sequential deposition of Nb and MgO films on a sapphire substrate. Subsequently, the multilayer composite was studied by selected area electron diffraction (SADP), conventional transmission electron microscopy (CTEM) and HREM. 2. EXPERIMENTAL Nb and MgO were sequentially deposited on the (01"[2) (R) plane of an epitaxial grade sapphire substrate by MBE to produce Nb-MgO-Nb-A1203 multilayer stacks. The thicknesses of Nb and MgO films were, respectively, ~100 and ~60nm. The substrate was preheated and cleaned by annealing at 1150°C for one hour in a vacuum of about 10- ~0torr. The Nb and MgO films were deposited at about 0.05 nm/s in the same vacuum onto the sapphire substrate which was kept at a temperature of 920°C for Nb and 650°C for MgO [13]. The cross-sectional HREM specimens were prepared by the following procedure. First the multilayer stack was oriented by the X-ray Laue back-reflection technique from the sapphire substrate side, and
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LI et aL: CHARACTERIZATION OF Nb-Al203 AND Nb-MgO INTERFACES parallel slices were cut by a diamond saw in the proper orientation.Then, two pieces of the sapphire substrate carrying the N b and M g O films were glued together, with the film surfaces facing each other, using M-bond 610 adhesive.t These "sandwiches" were cut in strips 1.5 mm long and embedded in conducting epoxy contained in a copper ring. The cross-sectional specimens were prepared from these couples using standard techniques, involving mechanical grinding to a thickness of 0.13 mm, dimpling to a thickness of 20 #m and ion beam milling at 6 kV [14]. The cross-sectional TEM specimens were prepared with the foil normal parallel to the [220]]~,o3 direction of the sapphire substrate. The samples were subsequently examined in a top-entry JEOL 4000EX high resolution electron microscope with a point-to-point resolution of ~0.18nm. Image simulations were performed using the SHRLI programs developed by O'Keefe [15] for the operating parameters of the JEOL 4000EX electron microscope, viz. coefficient of spherical aberration Cs = 1.0mm; defocus spread 6f= 8 rim; divergence angle half-width =--0.75mrad. The Scherzer defocus at 400 kV for this microscope is ~ - 50 nm. In order to overcome specimen charging which causes image drift and, sometimes, microscope instabilities resulting in discharges, the top surface of the specimen was covered with a 400 mesh copper grid after it was placed in the microscope sample holder.
Fig. 1. Bright-field image of Nb-MgO-Nb-AI203 multilayer.
electron beam simultaneously parallel (or nearly parallel) to the zone axes of Nbt, Nby, M g O and Al203. A schematic diagram of this SADP, indexed with respect to the four crystals,is shown in Fig. 2(b). According to this composite diffraction pattern, deposition on the (0112) (R) plane of sapphire
3. RESULTS AND DISCUSSION 3. I. Growth morphology and orientation relationship
Figure 1 is a low magnification cross-sectional TEM micrograph of the multilayer showing the growth morphology of the Nb and MgO films. The thicknesses of the Nb and MgO layers determined from bright-field TEM micrographs were 106 and 63 nm, respectively. The niobium film deposited on the sapphire substrate is denoted as Nb I while the niobium film deposited on MgO is designated as Nb2. As Fig. 1 shows, three interfaces are present in the multilayer: one between the Nb deposited on the sapphire substrate (the Nbi-Al203 interface), one between the MgO deposited on the Nb film (the MgO-Nb I interface), and one between the Nb deposited on the MgO film (the Nby-MgO interface). Electron diffraction studies of the multilayer showed that each layer was deposited in an epitactic manner with respect to its corresponding substrate. The ¢pitactic orientation relationships between Nbl-Al203, MgO-Nbl, and N b 2 - M g O were determined from SADPs. These are shown in Fig. 2(a) which shows a composite S A D P obtained with the tMeasurement Group, Inc., Raleigh, North Carolina.
, ~ _ ~ ) M g O (~~-
• Nbl
01"lN.b21~'ONb I
O Nb2
2 0ONbl
• MgO [] AI203
Fig. 2. SADPs showing the orientation relationships between N'b2, MgO, Nb= and A1203: (a) SADP along the [lll]~r~,[011]Mt~,[001]Nb~and [~20r]A~o3zone axes;(b) a schematicdiagram of the SADP indexedwith four crystals.
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results in the following epitactic orientation relationships in the multilayer
(110)Nb2 II(100)MgO II(100)Nbl [I(0112)AI~O, [11 l]Nb2II[01 l]~go II[001]Nb, l[[520I]~a2o,
(1)
In the above, we have used the convention that the parallel planes in the two crystals are taken to be those which are parallel to the interface, and the parallel directions lie in the two parallel planes. In fact, in Fig. 2, the [001] Nb~ zone axis is misoriented by about 2.4 ° with respect to the [~20I] zone axis of sapphire, in agreement with earlier studies [3, 4]. In the following each of the three systems NblA1203, MgO-Nb~, and Nb2-MgO will be discussed separately. 3.2. Interfaces in the N b - M g O - N b - A l e O ~ multilayer 3.2.1. The Nb I-AI eO ~ interface. Bright-field micrographs (see Fig. 1) and HREM observations indicated that the Nb~ film deposited on the (0112) plane of sapphire had the following epitactic orientation relationship
(100)~, 11(0112)~a,o,
[01 l]Nb~[I[2110]~a2o3
(2a)
(roughly) equivalent orientation relationship would be A
(100)Nb, II(0112)~a~o,
[00l]Nb, II[220I]~a2o,
plane of sapphire substrate using a magnetron sputter deposition technique. Consequently, all the thin films prepared by vapor deposition of Nb on sapphire (MBE or sputter deposition) are characterized by a unique orientation relationship between Nb and A1203 [equations (2) or (3)]. Note, however, that the Nb-A1203 interfaces prepared by internal oxidation gave a different orientation relationship [6, 7]
(110)N b II(0001)~O3 [110]Nb II[2H0ba~o~
However, in this case, apparently the precipitates were not c~-Al203 but rather a closely related material denote as ~'-A1203 by Mader [6]. The X-ray studies on the Nb--A1203 system showed that Nb could be deposited epitaxially on (R) plane of sapphire with either a (100) or (110) texture depending on the substrate temperature [4]. Of these two possible orientations, the (100) films fabricated at the higher deposition temperature of 850°C was a more perfect single crystal [4]. It is worth noting that silicon is also usually grown by CVD or MBE on the (01T2)ju2o3 (R) plane of sapphire to produce silicon-on-sapphire devices in the electronics industry (see, e.g. [17]). In contrast to b.c.c, niobium, silicon has a diamond cubic structure with an f.c.c, lattice. This system, which has been extensively studied, shows the following epitactic relationship [18]
(2b)
(100)s i II(01 À2),a=o,
Orientation relationships (2a) and (2b) are only approximately equivalent because the angle between [001]m and [011]Nb is 45 ° while the angle between [220I~2o~ and [~110]~a2o3directions is 47.1 °. It should also be mentioned that, despite the predominantly single crystalline nature of the Nb~ film, some small misoriented Nb grains were observed within the film. The observed orientation relationship for MBEgrown Nb film on the (0112) plane of sapphire is equivalent to the results previously obtained for the MBE growth of Nb on sapphire substrates [3,4, 10, 11]. This orientation relationship has been confirmed by HREM for Nb films deposited on differently oriented sapphire substrates: (0001)Aa2o3, (lI00)~a2o~ and (1~10)~a2o3[10, 11]. We note that in the work of Mayer et ai. [10, 11], the orientation relationship is written as
[001]si 11[~110ba2o3
(111)N b II(0001)m~o, [1 ]'0]r~b H[2TI0]~a2o3
(3)
which is nearly equivalent to equation (2); the angle between (0112)~a~o, and (0001)~a~o, is 57.6 ° while that between (100)r~b, and (11 l)Nb is 54.7 °. Knowles et al. [12] also reported the same orientation relationship, equation (3), for Nb films deposited on (0112) (R)
(4)
(5a)
or, sometimes [19] (100)si II(0112)~o3 [001]si II[0111]A12O3.
(5b)
3.2.2. Misfit dislocations at the NbI-AleO 3 interface. For interpretable HREM images of interfaces,
the electron beam should be parallel to low index zone axes of the two crystals and, also, parallel to the interface plane (see e.g. [20]). However, as mentioned above, the [001]Sb, zone axis of Nb~ was not precisely parallel to the [~20I]m2oJ zone axis of sapphire but rather was misoriented from it by ~ 2.4 °. Hence the low index planes in the two crystals, e.g. (100)Nbt and (0112)~2o 3, could not be simultaneously oriented parallel to the electron beam and a compromise had to be made in orienting the interface for high resolution imaging. It was found best to orient the [220I],j~o, zone axis of sapphire as accurately as possible parallel to the electron beam and allow for the slight misorientation of the corresponding Nb I zone axis (i.e. the [001]Nb~). Such a HREM mierograph is shown in Fig. 3 at a relatively low magnification. The interface plane is parallel to (0112)~2o, and approximately parallel to (100)Nbt. Well-defined atomic rows and
LI et aL: CHARACTERIZATION OF Nb-AI203 AND Nb-MgO INTERFACES
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Fig. 3. HREM image of a Nbt-AhO 3 -mtetfa~, The interface plane is parallel to (I00)~, and (01'D,)~o, !tlad the incident electron beam is parallel to [001]w~,a n d ~ 0 ~ o 3. The interface is marked by double arrows and a step is indicated by the symbol S. lattice planes can be identified. The NbI-AI203 interface lying parallel to (01T2)~2o3 planes, is atomically flat, sharp, and free of any interfacial phase. Interfacial steps of height equal to the (01"[2) interplanar spacing are clearly visible; one of these is indicated by S. The average separation between steps is about 30nm. To study the film/substrate coherency, the continuity of close-packed planes, d~ and d2, in the two crystals was examined [21]. In the [001]Nb1111220I]~O, projection of Fig. 3, the traces of closest-packed planes in Nbl and Al203 are (110)Nbl and (1120)~2o 3 respectively, which are inclined to the interface at angles of 45 ° and 43 °, respectively. Hence, in order to evaluate the misfit parameter 2(d2 - d, ) f = - -
(6)
the projections of(1 lO)Nb, and (1120)~a~o3planes on the interface, i.e. d l - - d ~ ) c s c 4 5 ° and d2~-d~?~csc43 ° were compared. Using this criterion, a lattice mismatch of 5.5% is obtained. It would be expected that this misfit strain be accommodated by an array of misfit dislocations of Burgers vectors b spaced a distance s apart at the interface. However, contrary to expectations, the experimental micrograph in Fig. 3 does not show any interfacial misfit dislocations. A clearer view is shown in Fig. 4, which is a higher magnification HREM image of the Nbi-Al203 interface. The clearly resolved white dots in the image of both crystals around the interface shows the smooth continuity of (ll0)sbl planes with the (ll~0)Aa2o3 planes. The absence of misfit dislocations, and the smooth continuity of the close-packed planes at the interface implies that atomic matching between Nb, and A1203 has been achieved locally by a distortion of the Nb lattice in the vicinity of the interface (i.e. this region of Nb is pseudomorphic with the sapphire substrate). This has been achieved, presumably, by the elastic distortion of the region of the Nb lattice next to the interface. It should be mentioned that the AI20 s lattice, in contrast to that of Nb, is undisturbed up to the interface. This result is not surprising, as
Fig. 4. A high magnification HREM image of the Nb,AIzO3 interface projected along [001]~, and [220T]~a2o~.The insets are computer simulated images of perfect crystals of Nb and A1203(thickness-- 18 rim, Af= - 6 0 nm). The unit cell of Nb is outlined.
A1203 possesses much higher elastic moduli than Nb. Although no misfit dislocations are visible at the interface, there are a series of dislocations at a small distance from the interface in the Nb lattice. This is in agreement with the work of Rfihle, Mader and colleagues [6-9], who first reported the existence of "stand-off" dislocations in the Nb-A1203 system obtained by internal oxidation. In Fig. 5, the core regions of some of these dislocations close to the Nb1-A1203 interface are indicated by inverted "T" symbols. The extra half plane of {110}sb forming the dislocations does not reach the interface but ends at a distance of about eight {ll0}Nb lattice planes. In fact, the stand-off distance from the interface varies from one place to another, but, on the average, it is about seven {200}Nb interplanar spacings. These dislocations are found to be spaced approximately
Fig. 5. HREM image of Nb,-Al203 interface with the incident beam parallel to [001]wo, and [220I]~2o3. The core regions of dislocations are indicated by inverted "T" symbols. The insets are computer simulated images of perfect crystals of Nb and Al203 (thickness = 21 nm, Af = - 60 nm).
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periodically, about one every twelve (020)~, planes (i.e. 2 rim) which is sufficient to accommodate the lattice mismatch between Nb and A1203. Therefore, while the re#on of niobium next to the interface is highly strained, the rest of the Nb lattice at a distance of 1.2 um from the interface is strain-free and the elastic strains are accommodated by "stand-off" misfit dislocations. Given the findings of stand-off dislocations at Nb--AI203 interface, Hirth and collaborators suggested that, because of mismatch in elastic moduli in a bimaterial, there is a repulsive image force on the interracial dislocations, tending to push them away from the interface into the material with the lower elastic modulus [22]. They calculated the equilibrium "stand-off" distance of misfit dislocations at an interface and obtained remarkably good agreement with the experimental results of Mader et al. [6-9] on the Nb-AI2 O3 system. It is interesting that Mayer et al. [10, 11] report that the misfit dislocations observed in their MBE-grown Nb films do not show "stand-off" positions. They attribute this to the kinetics of growth, i.e. to the fact that MBE deposition of Nb on sapphire was performed at lower temperatures compared to the manufacture of Nb-AI203 interfaces by either diffusion bonding or internal oxidation. Thus, they suggest that the geometrical arrangement of misfit dislocations at the interface in the MBE grown system might not be thermally stable, and that the dislocations could climb from the interface to stable stand-off distances if the samples are heat-treated or deposited at higher temperatures. This suggestion, however, implies that the misfit dislocations are generated at the interface' and, because of the low growth temperature, cannot climb to their equilibrium "stand-off" positions away from the interface. The generation of misfit dislocations at the interface may happen, for example, at the edges of the island deposits in a three-dimensional growth situation [23]. More commonly, however, misfit dislocation generation takes place by the nucleation of dislocation semi-loops at the film surface and the glide of these t o t h e interface (see, e.g. [24]). On reaching the interface, the interfacial segments of these ghssile lattice dislocations become "misfit" dislocations. In the case of Nb-AI203, the mechanism by which misfit dislocations are generated is not clear. However, if they are generated as lattice dislocations at the film surface, then it is possible that they never reach the interface but rather stop their glide when they reach the equilibrium distance at which there is no further driving force. This is an alternative to the one discussed by previous investigators, who suggested that dislocations are generated at the interface and then climb to their equilibrium position [10, 11].
It is also noteworthy that, in the present investigation, the misfit dislocation do show "stand-off"
r
Fig. 6. The core region of a dislocation in Nb I close to the Nb l-Al20 ~interface. A Burgers circuit is performed around a dislocation. positions, similar to the case of (110)Nb II(0001)m2o3 interface manufactured by diffusion bonding. This is despite the fact that, in the present MBE growth, the substrate temperature wag 920°C, similar to the temperature used in the experiments of Mayer et al. [10, 11]. It is clear that issues associated with "standoff" dislocations are unresolved. An attempt was made to determine the Burgers vector of the "stand-off" dislocations near the N b A12Oa interface. In Fig. 6 a Burgers circuit is drawn around the image of one of the "stand-off" dislocations. Since the projected Burgers vector of the "standoff" dislocation determined from the HREM image (Fig. 6) is ½[110]Nb,it is likely that the actual Burgers vector is ½(liD, as the latter is the usual lattice dislocation in a b.c.c, crystal. The insets in Fig. 4 are computer simulated images for both the Nb and the AlcOa at a foil thickness of 18 nm and a refocus Af = - 60 nm. The simulated images with these values of thickness and defocus agree well with the observed HREM images. The bright spots in the HREM images on the two sides of the interface correspond to AI atom positions in sapphire and Nb tunnels in the niobium, respectively. In sapphire, the aluminum and oxygen atoms are located on two different, but adjacent (01T2) atomic planes, which are shown in Fig. 7(a). In Fig. 7(b), the (a) •
•
•
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•
•
•
•
•
•
•
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•
•
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•
.
;.;.;.;.o •
o
•
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•
o
.
; . o . ; . ; . ; •
(b) •
•
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Fig. 7. (a) The atomic configuration in a (01]'2) plane of sapphire, (b) the atomic configuration in a (100) plane of Nb.
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Nb atom positions on the (100) plane of Nb are shown at the same scale as in Fig. 7(a). A comparison of these two figures shows the similarity of a (!O0)~r~ plane to a (0112)~azo~ plane: with soqae small ahifls, the AI atoms on the (0!T2)Ahosplane correspond well to the Nb atoms on the ( 1 ~ ) ~ plane, Attempts were made to determine the atomic structure of Nbz-Al203 interface from its H ~ images of the interface. Although totally satisfactory image simulations of the Nb-A1203 interface have not yet been obtained, Fig. 7 suggests that it is likely that an oxygen layer is adjacent to the (100)Nb plane. Because an explanation of the three-dimensional relationship between A1203 and epitaxially grown Nb must presumably lie in the relationship of the Nb atom to the O sublattice of sapphire, the eventual elucidation of the detailed interface structure is a matter of some considerable interest. 3.2.3. The MgO-Nb~ interface. The MgO film grown at 650°C on the Nb~ layer was single crystal, with an epitactic orientation relationship given by equation (1)
(lO0)MgO [[(100)Nb, [011]MsO II[001]Nb,
•
(5)
Figure 8 is a H R E M image of thisinterfaceobserved along the [011]MgO and [001]Nb,zone axes, and shows a "saw-tooth" morphology with the two lattices jutting into each other. The average M g O - N b ~ interface plane is tiltedby 3.5° with respect to (100)Nb,. A close inspection of the structural details of the interfaceshows that this "saw-tooth" interfacearises because of faceting,with a number of latticeplanes on the two sidesof the interfacejoining each other in a smooth fashion (see inset of Fig. 9). The interface at one side of the corner is formed by (IIT)MsO, while the interfaceat the other side is formed by (IT1)MsO; these are continuous with (ll0)Nbl and ('[10)Nbl planes, respectively.The development of {I I I}UtO facets at the M g O - N b ~ interfaceundoubtedly took place after growth, as evidence for the faceted structure was absent in R H E E D studies of the M g O grown at 650°C. It is likelythat thisinterfacefaceted during the latergrowth of Nb2 at 920°C. This struc-
Fig. 8. HREM image ofa MgO-Nb I interface viewed along [011]MgO and [001]Nb,. The misfit dislocations are indicated by small dots.
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Fig. 9' A hillh ~ HREM ~ o f ~ MgO-Nb, interface o~-raed along [011]~o and [O01]m,. The misfit dislocatiom are indicated by an inverted "T" symbol. The HREM image of the faeeted MgO-Nb I interface is shown in the top inset. The insets are computer simulated images of perfect crystals of MgO and Nb (thickness = 10 nm, Af-- -60nm). The unit cell of each crystal is outlined. ture may minimize the interfacial energy, as { I I I } are the close-packed oxygen planes in MgO, even though the plane of easy cleavage is (001). Note that the faceted interface, with facets parallel to (l IT)MsOand (ITl)MsO , accounts for a 141% increase in the total interface area compared to a fiat (100)MiO/(100)Nb, interface. Therefore, the interracial energy of the {111}MsO faceted MgO-Nbl interface is significantly lower than that of the (100)~O/(100)Nbl interface. Note also that strong strain contrast is absent at the (11])MsO facets, while strong strain contrast appears at the (ITl)MsO facets. In the [011]MsOI[[001]Sb~ projection of Fig. 9, the traces of closest-packed planes in MgO are (1T 1)MsO, while in Nb, they are (T10)Nb. The (IT1)MsO and (TI0)Nbl planes are inclined to the interface at angles of 54.7 ° and 45 °, respectively. Therefore, the projection of (I[I)Mso and (]10)Nbl planes on the interface, i.e. d I = d ~ c s c 5 4 . 7 ° and d2 = d~)csc45 ° were compared in evaluating the misfit parameter of the (100)MsO/ (100)Sb, interface, which gives a lattice mismatch of 10.4%. The misfit dislocations are shown by small dots in Fig. 8 and inverted "T"s in Fig. 9. Unlike the Nb-AIzO3 interface, these misfit dislocations lie at the MgO-Nb~ interface, spaced periodically about one every nine (1T1)MsOplanes (i.e. 2.7 nm). In order to determine the Burgers vector of the dislocations, a Burgers circuit was constructed as shown in Fig. 10. From this Burgers circuit, the projected Burgers vector was ~[~11]MsO, which makes this a 60 ° dislocation. The actual Burgers vector is therefore probably ½[101] or I[1]0]. On geometrical grounds, using such a Burgers vector, a regular array of misfit dislocations separated by five (lIl)MsO planes (i.e. 1.5 nm) would he required to completely accommodate the lattice mismatch of 10.4%. As the observed average distance between misfit dislocations is 2.7 nm [i.e. one every nine (l'[1)M~t~ planes], the large mismatch of 10.4% between (IlI)MsO and (II0)sb, at the interface does not appear to be completely accommodated by misfit
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Fig. 11. HREM image of the Nb2-MgO interface viewed along [lll]N~ and [011]MsO. The misfit dislocations are marked by small dots and the interface is marked by double Fig. 10. A Burgers circuit around a misfit dislocation at the MgO-Nb I interface.
arrows.
very nearly parallelto the (ITI)MgO facets.However, the (01"[)Nb2planes are misoriented by ~10 ° with dislocations.A n explanation of such a large discrep- respect to (Ill)Mso. ancy between the expected and observed average In general, the facet lengths arc not equal and are separation of misfit dislocations could be due to often associated with interracialstcps. The sense of the faceting of the interface although the R H E E D this facet asymmetry, and hence the sense of the studies referredto above suggest that facetingoccurs steps, is the same for nearly all facets. These steps afterdeposition.Taking the continuityof (IT l))~o and accommodate the 3.3° average deviation between the (~10)blb,planes at the (lII)MgO facet,the mismatch is (100)MgO plane and the N b E - M g O interface plane. only 4%, which may be assumed to be the "effective" A large local deviation results cither in large steps associatedwith one or two facetedcomers, or in small mismatch at the (100)MgO/(100)Nb,interface. The insetsin Fig. 9 are computer simulated images steps associated with a few short faceted segments. of M g O and N b I for a thickness of 10 n m and a In the [II l]Nb~II[01 I]MsO projection of Fig. I I, sets defocus Af= -60 nm. The simulated images agree of close-packed planes which arc compared in N b well with the observed H R E M images. The bright and M g O are (IIT)MgO (the oxygen close-packed spots in the H R E M images of the M g O and Nb~ planes) and (0]'l)Nb2.These planes arc inclinedto the represent the positionsof M g and N b atom columns, interface at 54.7° and 60 ° respectively.Therefore, respectively. the projection of (I l]')MgOand (0TI)Nb2planes on the 3.2.4.The Nb2-MgO interface.The Nb2 filmgrown interface, i.e. d~ = d ~ c s c 5 4 . 7 ° and a2 = d~0b~)csc60° on the M g O layer was not a single crystal, but are compared in evaluating the misfit parameter, contained a characteristicsubgrain cellstructure,with which gives a lattice mismatch of 9.8%. The core an average celldiameter of about 70 n m and a strong regions of a few dislocations are indicated by small texture.The N b 2 - M g O orientationrelationshipgiven dots in Fig. 11 and inverted "T"s in Fig. 13. Unlike in equation (1), i.e. the MgO-Nb I interface, these misfit dislocations are not located at the interface itself and again show (110)Nb2 II(100)Mgo
[111]Nbz[I[011]MaO
(6)
refers to this preferential orientation of the textured Nb 2 film with respect to the MgO layer. Figure 11 is a HREM image of the Nb2-MgO interface observed along the [111]m2 and [011]Mgo zone axes. The average interface plane is tilted by 3.3 ° with respect to the (100)UgOplane. Steps with heights equal to one to four (200)MsOinterplanar spacing were frequently observed. The HREM images in Figs 11 and 12 show some of these steps and also the faceting of the interface along two sets of MgO lattice planes. The flat region of the interface is parallel to (lI0)Sb2ll(100)UgO planes, while the faceted regions Fig. 12. HREM image of the Nb2-MgO interface viewed of the interface are parallel to (1 lI)M~o or (lll)MaO along [11 l]Nb2and [011]Mso.The faceted interface is marked planes. The (10I)Nb, planes of the Nb film are in turn by small dots.
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Fig. 13. HREM image of the Nb2-M ~ ~ a c e viewed along [111]~ and [011]MsO.The core regionsof dislocations are marked by the inverted "T" symbol. The insets are computer simulated images of perfect crystals of Nb and MgO (thickness = I0 nm, Af = - 6 0 nm).
(b)
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a "stand-off" position within the Nb lattice, similar to that of the Nb~-A1203 interface. They are spaced periodically about one every fourteen (0TT)Nb2planes (i.e. 3.8 rim). From Burgers circuits drawn around the core region of the misfit dislocations, the projected Burgers vector of the latter was ~[12T]Nb(see Fig. 13). On geometrical grounds, using this Burgers vector, a regular array of misfit dislocations separated by ten (0]'[)Sb2 planes (i.e. 2.8 nm) would completely accommodate the lattice mismatch. As the observed average distance between misfit dislocations is 3.8 nm [i.e. one at every fourteenth (0TT)N~ plane], the large m i s m a t c h of 9.8% between (111)M~o and (0TI)N~planes at the interface does not appear to be completely accommodated by misfit dislocations. It is worth noting that the MgO-Nbl interface formed by deposition of MgO film on the Nbl layer, and the N b , - M g O interface formed by deposition of Nb film on the MgO layer, have very different characteristics. The misfit dislocations were located at the interface in the first case while they were "standoff" dislocations in the second case. In addition, the MgO-Nbl interface was completely faceted parallel to the close packed (oxygen) planes of MgO, i.e. {111}Mo . On the other hand, the Nb2-MgO interface consisted of fiat regions parallel to (lI0)wo2-(100)Mso planes and faceted regions parallel to {lll}Mso. These differences might be related to the conditions of growth: in one case Nb is the substrate and MgO is the deposit, while in the other case, MgO is the substrate and Nb is the deposit. In terms of the misfit dislocation arrangement in the two cases, dislocations can presumably be nucleated more easily in Nb than in MgO. More importantly, the MgO substrate upon which the Nb 2 film grows may not be so perfect as the Nbn film upon which the MgO grows. Finally, the difference in growth temperature (920°C for Nb, 650°C for MgO) may be important. However, the origin of these differences needs further study. The insets in Fig. 13 are computer simulated images of MgO and Nb 2 for a thickness of 10 nm and
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Fig. 14. (a) HREM image of the Nb2-MgO interface region, showing a rotated microdomain of MgO. (b) Schematic diagram of the microdomain structure at the Nb2-MgO interface region. a defocus value of Af = - 60 nm. There is again good agreement between the simulated images and the observed HREM images of the MgO and the Nb2. As in the case of MgO-Nbl, the bright spots in the HREM images of MgO and Nb2 represent the positions of Mg and Nb atom columns, respectively. Very narrow regions with a domain structure were also observed at the Nb2-MgO interface, as shown in Fig. 14(a). One possibility is that this narrow region is NbO. However, since the lattice parameters of NbO, which is also f.c.c., and MgO are quite similar, and the region is too narrow to obtain clear compositional information, it has not been possible to identify the nature of the narrow interracial region. One possibility is that the narrow region contains a rotation microdomain of MgO, designated as (MgO)2, in Fig. 14: this microdomain is rotated by 54.7 ° around the [011]ugo axis with respect to the other domain, (MgO)l, as shown in Fig. 14(b). 3.3. Interfacial reaction during H R E M irradiation
During the course of HREM investigations, time-dependent changes occurred in the Nb~-A1203 interface region. In order to monitor this electron
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LI et al.: CHARACTERIZATION OF Nb-AI203 AND Nb-MgO INTERFACES
Fig. 15. HREM images of the Nbi-Al203 interface showing the evolution of the interface with electron beam irradiation: (a) initial NbI-AI203 interface; Co) a modified layer formed after 210 minutes of electron beam irradiation. The incident electron beam is parallel to the [O01]NblII[220I]m2o,direction.
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beam-induced solid-state reaction, in situ H R E M observations were performed on cross-sectional N b r A l 2 0 3 specimens as a function of the duration of electron beam irradiation. During the H R E M observations, an electron beam with a current density of 80 pA/cm 2 (read from the fluorescent screen of the JEOL 4000EX electron microscope) was focused on the same area of specimen. It can be seen from Figs 3, 4 and 5 that the original NbI-AI203 interface is perfectly smooth, sharp, and free of any interfacial phase in the initial stages of H R E M observation. After electron beam irradiation, however, a narrow region with some modification of the interface was easily recognizable. Figure 15 is a H R E M image of the interface after electron beam irradiation for 210 rain. A 2.1 rum layer with slightly different image characteristics can be distinguished between the N b 1 film and the sapphire substrate.
Fig. 17. HREM image of the Nbl-Al203 interface showing grain growth during electron beam irradiation: (a) initial small grain embedded in the Nb I film at the Nbl-A1203 interface region, (b) the same grain in the Nb I film after 210 min of electron irradiation. As seen in Fig. 15, the thickness of the region which has been modified is not uniform all along the interface, and, in fact, its HREM image is quite similar to that of the AI2Oa substrate. Consequently, we have not been able to positively identify the nature of change which has taken place. The kinetics of such a modification of the interface can be determined by varying the duration of electron beam irradiation. Figure 16 shows the change in the
LI et al.: CHARACTERIZATION OF Nb-Al203 AND Nb-MgO INTERFACES thickness, x, of the modified layer, as determined from H R E M images, with time, t. Although relatively large error bars are associated with the experimental points, it is reasonable to fit a parabolic curve through the points, i.e. the thickening of the modified layer follows a parabolic law x 2 = K ( t - to) where K is the rate constant for the reaction and to is the incubation period. The parabolic fit to the experimental results (Fig. 16) gives K ~ 4 A 2 rain- l and to = 50 rain. Spinel formation by the solid state reaction of MgO and AI203 at a MBE grown MgO-AI203 interface under electron beam irradiation has recently been investigated, and the orientation relationships, and kinetics and mechanisms of spinel formation under electron irradiation has been discussed in detail in a previous paper [14]. Growth kinetics of the modified layer at the Nbl-Al203 interface under electron irradiation is very similar to that of MgAI204 formation at the MgO-AI203 interface. However, as' mentioned previously, the nature of the modified layer at the interface is not known. Electron beam-induced grain growth in the Nb 1 film was also observed. Figure 17(a) is a HREM image of a small grain embedded in the Nbl film at the Nbt-AI203 interface region. Figure 17(b) shows the migration of the grain boundary which occurred during electron irradiation. In this figure, the original interface between Nb~ and the embedded grain is indicated by dashed lines and the new interface after electron beam irradiation for 210 rain is marked by small dots. It is clear that the Nb~ grain consumes the embedded grain under electron irradiation.
4. CONCLUSION 1. The orientation relationships between Nb2, MgO, Nbt and AI203 layers are:
( 110)Nb2 II( 100)MsO II (100)Nb, II(01T2)~a2o, [111]N~ II[011]MsO II[001]Nb, II[2201],a2o3
•
2. The initial Nb1-A1203 interface produced by MBE growth is atomically flat and sharp. The misfit dislocations are not located at the interface; there are "stand-off" dislocations in the N b lattice at a distance of about seven {200}Nb interplanar spacings from the interface. These are spaced periodically about one every twelve (020)Sb planes. 3. The MgO-Nb~ interface was faceted with a "saw-tooth" morphology; the faceting resulted in near-parallelism of close-packed planes of Nb] and MgO (the oxygen close-packed planes), i.e. { 111 }uao [[{ 110}Nb,. The faceting most probably took place after the growth of MgO, presumably because of the lower N b - M g O interfacial energy along { 110}Sb and {111}uso. Misfit dislocations were located at the interface with an average separation of ~2.7 nm. 4. The Nb2-MgO interface had many steps and was faceted with two sets of {I 11}Mgo lattice planes. The close-packed {110}Nb2 lattice planes of Nb, are
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either very nearly parallel, or have a small misorientation, to {11 l}MgO. Misfit dislocations are again not located at the interface but a set of "standoff" dislocations in the N b lattice at a short distance away from the interface are spaced periodically about one every 3.8 nm. 5. A modified layer formed at the Nb~-Al203 interface by a solid state reaction occurring under electron irradiation. The modified layer thickened parabolically. In addition to the formation of a modified layer at the interface, electron beam-induced grain growth within N b was also. observed. ,4cknowledgements--The work at Case Western Reserve
University was supported by the DARPA/ONR University Research Initiative at CWRU under contract N00014-86-K0773. The work at the University of Illinois was supported by the Department of Energy under contract DE-AC0276ERO1198; use of the Materials Research Laboratory facilities at the University of Illinois is acknowledged.
REFERENCES 1. M. Rfihle, A. G. Evans, M. F. Ashby and J. P. Hirth (editors), Metal-Ceramic Interfaces. Pergamon Press, Oxford (1990). 2. These proceedings. ,4cta metall, mater., Vol. 40, Suppl. (1992). 3. S. M. Durbin, J. E. Cunningham and C. P. Flynn, J. Phys. F, Met. Phys. 12, L75 (1985). 4. D. B. McWhan, Mater. Res. Syrup. Proc. 35, 493 (1985). 5. M. Florjancic, W. Mader, M. Rfikle and M. Turwitt, J. Physique 46, C4-129 (1985). 6. W. Mader, Mater. Res. Soc. Syrup. Proc. 82, 403 (1987). 7. W. Mader, Z. Metallk. 80, 139 (1989). 8. M. Kuwabara, J. C. H. Spence and M. R/ihle, J. Mater. Res. 4, 972 (1989). 9. W. Mader and M. Rill'de, ,4cta metall. 37, 853 (1989). 10. J. Mayer, C. P. Flynn and M. R/ihle, Ultramicrosc. 33, 51 0990). 11. J. Mayer, W. Mader, D. Knauss, F. Ernst and M. Rfihle, Mater. Res. Soc. Syrup. Proc. 183, 55 (1990). 12. K. M. Knowles, K. B. Alexander, R. E. Somekh and W. M. Stobbs, Inst. Phys. Conf. Set. 90, 245 (1987). 13. C. P. Flynn, in Metal-Ceramic Interfaces (edited by M. PriMe, A. G. Evans, M. F. Ashby and J. P. Hirth), pp. 168-175. Pergamon Press, Oxford (1990). 14. D. X. Li, P. Pirouz, A. H. Heuer, S. Yadavalli and C. P. Flynn, Phil. Mag..4 65, 403 (1992). 15. M. A. O'Keefe and P. R. Buseck, Trans. Am. Crystallogr. Assoc. 15, 24 (1979). 16. D. X. Li, P. pirouz, A. H. Heuer, S. Yadavalli and C. P. Flynn, Mater Res. Soc. Syrup. Proc. 221, 93 (1991). 17. M. S. Abrahams, C. J. Bniocchi, J. F. Corboy Jr and G. W. Cullen, ,4ppl. Phys. Lett. 28, 275 (1976). 18. K. C. Pans, J. C. Barry, G. R. Booker, T. B. Peters and M. G. Pitt, Inst. Phys. Conf. Ser. No. 76, 35 (1985). 19. J. L. Batstone, Phil. Mag. B 63, 1037 (1991). 20. P. Pirouz and F. Ernst, in Metal-Ceramic Interfaces (edited by M. Rfikle, A. G. Evans, M. F. Ashby and J. P. Hirth), pp. 199-221. Pergamon Press, Oxford (1990). 21. G. B. Olson and M. Cohen, Acta metall. 27, 1907 (1979). 22. S. V. Kamat, J. P. Hirth and B. Carnahan, Mater. Res. Soc. Syrup. Proc. 103, 55 (1988). 23. R. Vincent, Phil. Mag. 19, 1127 (1969). 24. J. W. Matthews, in Epitaxial Growth (edited by J. W. Matthews), Part B, pp. 559-609. Academic Press, New York (1975).