The early stages of MgO epitaxy on lattice-matched Cr0.7Mo0.3(001)

surface science :.:: ........

ELSEVIER

. .........

Surface Science 339 (1995) 297-309

The early stages of MgO epitaxy on lattice-matched

Cr0.7Mo0.3(001) S.A. Chambers *, Y. Gao, Y. Liang Environmental Molecular Sciences Laboratory, Pacific Northwest Laboratory 1, P.O. Box 999, MS K2-12, Richland, WA 99352, USA Received 14 March 1995; accepted for publication 3 May 1995

Abstract We have investigated the structure and morphology of thin epitaxial overlayers of MgO on lattice-matched interlayers of Cr0.7Moo. 3 grown on MgO(001) substrates. Lattice matching results in unstrained MgO epitaxial films with short- and

long-range structural order of comparable quality to those of buffer layers of MgO grown on MgO(001) under optimized step-flow growth conditions. In addition, the structural coherence of the oxide/metal interface is excellent due to lattice matching, in contrast to lattice-mismatched systems. However, MgO does not grow in a perfectly laminar fashion on Cro.7Mo0. 3 due to surface and interface free-energy imbalances; thin MgO films exhibit mean roughnesses of ~ 4 A with terrace widths of several hundred ,~. Nevertheless, the use of lattice-matched conducting interlayers between a bulk oxide substrate and an epitaxial thin film of the same material allows high-quality insulting oxides to be prepared on conducting substrates. This approach in turn permits the full arsenal of charged-particle spectroscopic, diffraction and imaging probes to be applied to investigations of the surface science of insulators. Keywords: Atomic force microscopy; Electron-solid diffraction; Epitaxy; Magnesium oxides; Metal-insulator interfaces; Molecular beam epitaxy; Photoelectron diffraction; Single-crystal epitaxy

1. Introduction The epitaxial growth of ceramics on metals is of interest for a number of scientific and technological reasons. At a fundamental level, interfacial adhesion depends on the nature of the localized bonds that

* Corresponding author. Fax: +1 509 375 6916; E-mail: [email protected]. 1 Pacific Northwest Laboratory is a multiprogram national laboratory operated for the US Department of Energy by Battelle Memorial Institute under contract number DE-AC06-76RLO 1830.

form [1-13]. The detailed nature of the bonding in turn depends on the extent of interface strain, the extent of mutual solubility and interfacial reactivity, and the presence of impurities at the interface. Determining the relationship between adhesion and bonding is critically dependent upon being able to control and characterize the interfacial structure. Moreover, the planarity of the epitaxial film and, thus, the step density on the top surface depends on the surface and interface free energies, the presence of impurities, and lattice match. These scientific issues have direct bearing on technologies ranging from protec-

0039-6028/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved

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S.A. Chambers et al. / Surface Science 339 (1995) 297-309

tive ceramic coatings on metals [14] to single-crystal tunnel barriers [15] and magnetic memory structures [16,17] to model catalyst, separations material, and mineral surface design [18-21]. From the point of view of surface science, the ability to prepare well-defined ceramic surfaces by epitaxial growth of thin-film ceramics on conducting substrates makes possible the use of charged-particle spectroscopic, diffraction, and imaging probes that would otherwise be accompanied by deleterious charging effects [22]. In addition, fabrication of ceramic surfaces by this means significantly reduces vacancy and dislocation defects frequently found as a result of bulk-crystal surface preparation. However, it is essential that the criteria for successful epitaxy, as determined by extensive experience with semiconductors, be applied to the case of ceramics on metals in order to produce excellent surfaces, films, and interfaces [23]. These criteria include lattice and crystal symmetry matching, the use of buffer layers grown on the substrate immediately prior to heteroepitaxy in order to create the best possible template, and growth under conditions that maximize step flow. In this paper, we present results for the epitaxial growth of a cubic rocksalt oxide, MgO, on a bodycentered-cubic metal alloy to which MgO is exactly lattice-matched, Cr0.7Mo0.3. This alloy was in turn grown on a MgO(001) substrate with a buffer layer, in order to generate an atomically flat crystalline template with exceedingly good short- and long-range order and very low defect density [24]. The remainder of the paper is organized as follows. Experimental details are provided in Section 2, molecular beam epitaxial (MBE) growth of Cr0.7Mo0.3 interlayers is discussed in Section 3.1, the structure and morphology of thin MgO regrown films are presented in Section 3.2, the interface structure is discussed in Section 3.3, discussion of the results is given in Section 4, and conclusions are drawn in Section 5.

2. Experimental details The MBE system used in this study consists of two electron beam evaporation sources for Mo and Cr, an effusion cell for Mg, and an electron cy-

clotron resonance (ECR) source for oxygen. Metal fluxes were monitored by means of quartz crystal oscillators (QCO) adjacent to the sources. Oxygen partial pressure was measured with a residual gas analyzer (RGA). Film composition was determined with a small, high-speed Auger electron spectrometer (AES) that collects Auger electrons excited by the same 15 keV electron beam used to generate reflection high-energy electron diffraction (RHEED) patterns [25]. More precise post-growth compositional and structural measurements were carried out in an appended analytical chamber that houses an X-ray photoelectron spectrometer (XPS). The XPS system was also used to perform scanned-angle X-ray photoelectron and Auger electron diffraction ( X P D / AED) measurements with an angular resolution of _ 1°. The analysis chamber also contains reverseview low-energy electron diffraction (LEED) optics. Ex-situ atomic force microscopy (AFM) and crosssectional transmission electron microscopy (TEM) measurements were done on a Digital Nanoscope III AFM and a JEOL 2010 TEM, respectively. Polished MgO(001) substrates obtained from Princeton Scientific Corporation were cleaned in acetone and methanol prior to insertion into the MBE chamber. Once under UHV, the substrates were further cleaned by heating to 750°C while being exposed to an oxygen plasma at a pressure of ~ 2 X 10 -5 Torr for 15 min. The resulting surfaces possessed at most ~ 0.1 monolayer of carbon and exhibited a 1 X 1 RHEED pattern. A 300 .~ MgO buffer layer was then grown at a substrate temperature of 750°C using the oxygen plasma source. The RHEED pattern improved considerably upon buffer layer growth due to the elimination of surface defects and roughness on the polished substrate. Once the oxygen partial pressure had dropped into the 10 -11 Torr range, a layer of CrxMOl_ x (x = 0.7) several hundred A thick was grown at room temperature. In-situ AES measurements were carried out during growth to ensure that x remained between 0.56 and 0.80. After growth, the film was annealed for a few minutes at 550°C in order to improve the structural quality and surface morphology. MgO regrowth was then carried out at substrate temperatures of 25°C and 550°C. Growth at the higher temperature was done by admitting molecular oxygen into the MBE chamber at a partial pressure of ~ 5 × 10 -7

S.A. Chamberset aL /Surface Science 339 (1995) 297-309

Torr while exposing the surface to the Mg atomic beam at a flux of ~ 0.20 monolayer equivalents of MgO per second. Incident Mg does not adsorb on the hot surface prior to oxygen admission. Admitting molecular oxygen slowly into the chamber while under Mg rich conditions results in the nucleation of epitaxial MgO and negligible oxidation of either Mo or Cr. The ECR plasma source was not used when growing very thin films in order to avoid oxidation of the C r / M o alloy by the activated oxygen. We have found that leaving the ECR plasma source off during growth results in a lower growth rate compared to when the source is on, but that the film quality is the same. Growth at room temperature was carried out by alternating ~ 1 monolayer-equivalent depositions of Mg with exposure to molecular oxygen with the plasma off. The exact procedure used

CrxMol.x(O01)

299

for the room-temperature growths is described in more detail in Section 4.

3. Results 3.1. M B E growth o f Cro.7Moo. 3

There are no pure metals to which MgO(001) is lattice-matched. The two metals which are most closely lattice-matched to MgO(001) are Ag(001) and Cr(001), with mismatches of 3.3% and 3.5%, respectively. Thus, strained-layer epitaxy is in principle possible up to some critical thickness, at which point the overlayer relaxes and either agglomerates, forms misfit dislocations, or both. The critical thickness for any given materials system can be estimated

MgO/CrxMol.x(O01 )

Cr Mo

Q

Mg O

Fig. 1. Interface crystallographyfor the MgO/CrxMOl_x(001) system. Cr and Mo constitute a random bcc alloy for which the lattice parameter scales approximatelylinearly with x. The alloy is lattice-matchedto MgO to within + 1% for x = 0.56 to 0.80.

300

S.A. Chambers et al. /Surface Science 339 (1995) 297-309

by means of elastic theory, and the characteristic value depends on the elastic constants for the two materials. However, the critical thickness is typically ~ 8 to 15 .~ for a ~ 3% mismatch. Trampert et al. have obtained high-resolution TEM lattice images of the Ag/MgO(001) system and dislocations are evident along the interface plane, as are bent lattice planes several tens of .~ away from the dislocation core [26]. We have grown MgO on Cr(001) and the overlayer is only weakly ordered, as judged by RHEED, LEED and XPD [22]. In order to achieve a lattice match, one must employ a compound metal with the appropriate lattice parameter. The lattice parameters for Cr and Mo are such that act / v~- = 0.965(amgo/2) and arao/v/2 = 1.052(aragO/2). Furthermore, both Cr and Mo are Group-Via transition metals and form a solid solution at elevated temperatures for which the lattice parameter varies approximately linearly with composition. Trzebiatowski et al. did an X-ray analysis of alloys prepared over the full range of composition and found a nearly linear dependence of the lattice parameter on composition for temperatures from 600°C to 1700°C [27]. A miscibility gap has been reported for temperatures below 880°C, but the position of the gap could not be determined experimentally because of the low diffusivity of Cr and Mo in the solid solution [28]. Thus, it is in principle possible to obtain a >__99% lattice match to MgO by using a substrate of CrxMo 1_x(001), where x = 0.56 to 0.80, provided the kinetics of phase segregation of the alloy at lower temperatures can be prevented. This task can in principle be accomplished by coevaporating Cr and Mo in the proper proportions on a MgO(001) substrate and not subjecting the specimen to extensive thermal treatments at temperatures below ~ 900°C [28]. Monitoring the fluxes of the two metals in the gas phase and the film composition simultaneously affords the best opportunity to obtain the desired composition. We have done so by simultaneously employing real-time AES and separate QCOs for the two atomic beams. We illustrate the expected crystallography for the MgO/CrxMol_x(O01) interface in Fig. 1. The bcc surface unit mesh for the C r / M o alloy is shown in the upper left. A partial monolayer of MgO is superimposed on the alloy substrate in the upper right, and the 45 ° rotation about [001] required to bring the

/

0

lattice matched / I to MgO to / L4- withinI_.°/,._/=

1.6

+ <

"

m

~ 1.2

o ~ 0.8

0.4

o

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

Mole Fraction of Cr (x)

Fig. 2. CrL3M2,3M4,5 to MoM4,sN2,3M4,5 peak-to-peakAuger ratio as a functionof Cr mole fraction, x, in CrxMOl_x(001).

rocksalt surface unit mesh into registry with that of the alloy surface is evident. Also shown is a crosssectional view of the interface (lower right). The assumed registry for the interface has oxygen atoms in the interfacial MgO layer in four-fold hollow sites of the bcc surface mesh of the C r / M o alloy. We will comment further on this assumed registry and the associated implications for interfacial bonding in Section 4. We show in Fig. 2 an AES calibration curve which relates Cr to Mo peak-to-peak ratios obtained from derivative Auger spectra to alloy composition [25]. This curve was obtained by first measuring spectra for pure films of nominally single-crystal Cr and Mo films on MgO(001), and then taking linear combinations of these spectra to simulate spectra for the full range of alloy composition. We also show the range of composition over which the alloy is lattice-matched to MgO(001) to within 1%. During alloy growth, we obtain repetitive, back-to-back spectra for the evolving epitaxial film in the N(E) mode, differentiate the spectra, compute the Cr to Mo peak-to-peak ratio, and determine the film composition in real time during growth from the calibration curve. The time required to obtain an energy

S.A. Chambers et aL / Surface Science 339 (1995) 297-309

301

Cro.TMOo.3/MgO(001) Growth Rate = 0.35 ML/sec

~,,~Mo MNN •~

. "-~..~"--J-~.

Cr LMM t = 84 sec ............

~..~.e~.. ._, . . . . . . .

f

\\./-~

.---,. j.~\ "~

~--

0 sec

"

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I 150

I 200

I 250

-42 sec

I 300

I 350

u

I 400

'., ~..~--.L

, .~\

11)

I 450

\

I 500

\ .~_........ %

I 550

I 600

Kinetic Energy (eV)

Fig. 3. Auger spectra as a function of time obtained during MBE growth of a Cr0.7Mo0.3 alloy film on MgO(001). Each spectrum required 42 seconds of acquisition time, and spectral acquisition started at the higher kinetic energy. The shutters were opened at t = 0.

scan that encompasses both the M o M N N and Cr L M M manifold with good statistics is 42 s. Four such spectra, obtained at the beginning of a C r / M o growth, are shown in Fig. 3. The shutters were opened at t = 0, at which point a scan was initiated starting at a kinetic energy of 600 eV. The scan marked t = - 42 s was obtained just before opening the shutters. The O KLL manifold associated with the MgO(001) buffer layer is seen in this spectrum. Within ~ 10 seconds of opening the shutter, both the Cr LMM and O KLL manifolds are visible in the region from 450 to 550 eV. During this short-lived period of the growth, interference of the Cr L M M manifold with the O KLL peak makes quantitative determination of the alloy more difficult. However, within a minute of opening the shutter, and growing at a rate of ~ 0.35 monolayers per second, the O KLL peak is completely attenuated due to the fact that the sampling depth is limited to a few monolayers. This result stems from a combination of the glancing primary electron beam condition (angle of incidence ~ 3 °) and the detection angle of the spectrometer (35 ° off the surface plane) [25]. Several spectra are obtained throughout each growth by this means, yielding a film composition profile, and insuring that the desired stoichiometry is achieved. In Fig. 4 we show XPD rsults for a 200 ,~

epitaxial film of Cr0.66Mo0.34 grown on MgO(001). Here we plot the Cr2p3/2 and M o 3 d core-level intensities, along with Fe2p3/2 intensities from i

,

i

,

i

,

i

'

i

,

200 A Cr o 6sMoo34/Mg0(001)

r

,

i

'

,

[0011~

o

_¢ >

¢r

10

20

30

40 50 60 70 80 Polar Angle, e (degrees)

90

100

Fig. 4. Cr2p3/2 and Mo3d photoelectron angular distributions in the (100) azimuthal plane for a 200 ~ epitaxial film of the bee alloy Cro.66Moo.34 on MgO(001). Also shown is an Fe2p3/2 angular distribution from Fe(001), which is known to have a bcc structure.

302

&A. Chamberset al. / Surface Science339 (1995) 297-309

Fe(001) [29], as a function of polar angle in the (100) azimuth. The Fe(001) scan is included as a structural "fingerprint" for a well-defined bcc crystal. This scan encompasses the [011] and [001] forwardscattering directions in the bcc alloy structure. The anisotropy over the [001] peak is an excellent probe of short-range structural order in the alloy film. Well-ordered, (001)-oriented, cubic single crystals, such as the Fe(001) crystal used as a bcc reference material, typically show anisotropies over the surface normal (defined as 100(/0=90- Io=8o)/Io=9o) of 50 to 60%. In this case, the Fe(001) reference surface exhibits an anisotropy of ~ 60%. In comparison, the C r / M o alloy film shows anisotropies of ~ 55% for both Cr2p3/2 and M o 3 d core-level intensities, indicating that the short-range structural order in the epitaxial film is comparable to that of a bulk single crystal of the same structure. In addition, there is no evidence of strain in the film, as judged by the fact

that the [011] forward-scattering peak falls precisely at 45.0 ° . Any deviation from 45.0 ° would be indicative of tetragonal distortion, as might be expected for a lattice-mismatched overlayer. In Fig. 5 we show RHEED patterns with the primary beam aligned along the [110] direction in MgO(001) for a MgO buffer layer (a), and a 500 ,~ Cr0.7Mo0. 3 layer (b). The streaks seen in the zerothorder Laue zone for the C r / M o alloy are at the same positions as those in the MgO buffer layer pattern, indicating no detectable change in lattice parameter. Thus, the alloy films exhibit excellent long-range structural order in addition to the excellent shortrange structural order revealed by the XPD scans. 3.2. Structure and morphology of thin MgO regrown films Thin epitaxial films of MgO were regrown on Cr0.7Mo0. 3 interlayers at 550°C as described in Sec-

(a)

(b)

(c)

(d)

(e)

(f)

Fig. 5. RHEED patterns at 15 keY with the primary beam aligned along [110] in the MgO(001) surface for (a) a 300 .~ MgO buffer layer on MgO(001), (b) a 500 ,~ Cro.7Moo.3 interlayer on MgO(001), and regrown MgO overlayers of thickness equal to (c) 5 A, (d) 7 A, (e) 10 A, and (f) 14 ,~.

S.A. Chambers et al. / Surface Science 339 (1995) 297-309 MgO/2OOACr07t~Mo030/MgO(O01) Mg KL2.3L2,3 -~ (110)"azimuth

I[111]

[001]/~. ,/\

0 ls -- (110) azimuth

303

[o01 ._

_

I \

A

o~ o >, CL 2 "E <¢

20

30

40

50

60

70

80 90 20 30 40 Polar Angle, 0 (degrees)

50

60

70

80

9O

Fig. 6. X-ray-excited MgKL2,3L2, 3 Auger and O l s photoelectron angular distributions in the (110) azimuthal plane for various regrown MgO films on Cro.TMoo.a/MgO(001).

tion 2. The ECR plasma source was left off during these thin-film growths. The QCO is not a useful way to measure MgO film thickness from elemental sources for a variety of reasons. First, the QCO is maintained at ~ 25°C by cooling water, whereas the substrate is held at elevated temperature. Thus, both Mg and MgO condense on the QCO, whereas only

MgO/2OOACr0.7(~Moo.3o'MgO(OOt) lMg ~ KL2'3L2'3 ~ l ~ ] "i (m u t h ~

o~ o

[001]

some fraction of the MgO that forms and none of the Mg metal stick to the hot substrate. Furthermore, the mix of materials incident on the QCO creates an ambiguity in programming elastic constants and a density into the QCO controller in order to convert the measured frequency change into a film thickness. Because of these difficulties, we have used the X-

l~Sl~'imuth[0111

tool]

¢

15 ( c o n t . ) ~

2

_S 20

30

40

50

60

70

80 90 20 30 4O Polar Angle, 9 (degrees)

50

60

70

80

90

Fig. 7. X-ray-excited MgKL2,3L2, 3 Auger and O I s photoelectron angular distributions in the (100) azimuthal plane for various regrown M g O films on Cro.7Moo.3//MgO(001).

304

S~A. Chambers et a L / Surface Science 339 (1995) 297-309

ray-excited MgKLL Auger and Mg2s photoemission peaks measured after growth to estimate the MgO film thickness by employing a simple continuum model [30]. Such an estimate provides only an a v e r a g e film thickness if the film grows in an agglomerated fashion. In Fig. 5 we show RHEED patterns with the primary beam aligned along the [110] direction in MgO(001) for regrown MgO films of thickness equal to 5 A (c), 7 .~ (d), 10 .~ (e), and 14 ,~ (f) on a Cr0.7Mo0,3 interlayer. The sharp reflection patterns seen for the MgO buffer and C r / M o alloy layers ((a) and (b)) transform into a mixture of reflection-and transmission-like patterns at the lowest coverage and remain so to a coverage of 14 A. The RHEED patterns for single monolayer-equivalent films (not shown) are completely dominated by transmission and exhibit high background, suggesting very small island formation at low coverage. The transmission component of the pattern decreases with coverage from one to several hundred monolayers, revealing island growth and eventual coalescence [22,23]. The contrast between Bragg peaks and off-peak background does not decrease significantly in going from the buffer-layer pattern to the regrown film patterns, suggesting that the defect densities remain low in the regrown films. This conclusion is corroborated by XPD results to be discussed below. In Fig. 6 and Fig. 7 we show A E D / X P D polarangle intensity distributions in two high-symmetry azimuths for X-ray-excited MgKL2,3L2,3 Auger and O ls photoemission peak intensities films of various thicknesses [31]. The scans marked 15 (cont.) and 300 were measured on films which were grown continuously, rather than incrementally. Looking first at the MgKL2.3L2, 3 scans taken from the 5 A film, forward-focusing peaks along the principal low-index directions [001] (surface normal, 0 = 90°), [011] (0 = 45 ° in the (100) azimuth), and [111] (0 = 35 ° in the (110) azimuth) are seen. This result shows that the rocksalt structure is established within completion of the first few monolayers, as expected based on the crystal structure. Some of the higher-order fine structure is evident after the 5 A-deposit, but these features are not fully developed until the film is > 15 ~, thick. Of more importance in the present context is that there is no evidence of strain at any coverage. The most direct way to follow the tetrago-

MgMgO/Cr°7' M°' 3 3/MgO(001)KL2,3L.2. ~ . [011]peakin ~#1\ ," i ~ S / = (100)azimuth ~ •

[0111

]

o

MgOoverlayer| thickness(A) ]

..~

o

o

....

I ....

I ....

I ....

I ....

I ....

I ....

I ....

I ....

I ....

20 25 30 35 40 45 50 55 60 65 70 PolarAngle, O(degrees) Fig. 8. X-ray-excitedMgKL2,3L2,3 Auger angular distributionsof the [011]forward-scatteringpeak in the (100) azimuthalplane for various regrown MgO films on Cr0.7Mo0.3/MgO(001). The absence of deviation of the peak from 45.0° reveals the lack of detectable strain.

nal distortion accompanying lattice-mismatch-induced strain in cubic materials is by measuring the position of the close-packed [011] forward-focusing peak in the (001) azimuth [32]. This peak will deviate to angles other than 45.0 ° if there is a change in the ratio of perpendicular to in-plane lattice parameter as a result of strain. In the case of epitaxial MgO on lattice-matched Cr0.7Mo0.3, there is no deviation outside experimental error in 0[011] from the value measured for bulk MgO(001). This peak is plotted in Fig. 8 for four thin films and one thick film. In each scan, the [001] peak is used as an internal reference, and is taken to fall precisely at 9 0 . 0 _ 0.2 °. The goniometer has been calibrated so the [001] and [011] peaks measured in the (001) azimuth of MgO(001) are separated by 45.0 ___0.2 °. The measured values of 01011] are 45.2 °, 44.8 °, 44.9 °, 45.0 °, and 45.1 ° for the 5, 7, 10, 14 and 300 ,~ films, respectively. The +0.2 ° uncertainty in 01011] is equivalent to an uncertainty of _ 1% in the ratio of perpendicular to in-plane lattice constants. Thus, there is no measurable strain in these films, at least

S.A. Chambers et al. / Surface Science 339 (1995) 297-309

305

at the bottom of Fig. 9 for the film grown in an incremental fashion. These images, although limited to a one-square-micron area, are representative of the entire film surface. The Z contrast has been greatly expanded to accentuate any irregularities in surface topography. The interlayer alloy film surface is exceedingly fiat with a mean roughness of only ~ 2 .~ and terrace widths of several hundred to a thousand .4,. In contrast, the regrown MgO film is less smooth, with a mean rouaghness of ~ 4 A and terraces of several hundred A. Although the mean roughness is relatively low, there is clear evidence of dispersed islands that are tens of .4, in height and a few hundred ~, across. The shapes of the island images are perturbed by the ~ 500 A radius of curvature of the A F M tip. Therefore, it is not possible to deduce the exact shape of the islands from these images. Nevertheless, the presence of islands explains the transmission component observed in the RHEED patterns in that an appreciable fraction of the islands possess widths which are less than or equal to three times the electron attenuation length foro 15 kV electrons passing through MgO, or ~ 450 A.

!

iJl~

Fig. 9. AFM images of a 50 A. thick interlayer of Cr0.75Mo0.25 grownon MgO(001) (top) and a 14 ,~ thick regrown MgO film on a 500 A thick Cr0.TMO0.3 interlayer grown on MgO(001) (bottom). The z-scale has been expanded by a factor of 1000 relative to the x- and y-scales in order to accentuate any deviation from perfectly laminar growth. The difference in surface morphology is due to differences in surface free energies for the two materials and weak interfacial bonding.

to within the limits of strain detectability available from X P D / A E D . o We show in Fig. 9 AFM images for a 50 A thick interlayer of Cr0.75Mo0.25 (top) and a 14 A thick regrown MgO film (bottom). The C r / M o film was removed from the MBE system for AFM analysis and was not put to any further use. The top MgO film was grown incrementally on a different 500 A thick C r / M o interlayer, as described above. A F M images of MgO films of comparable thickness grown continuously are substantially the same as that shown

3.3. M g O / C r o . 7 M o o . s interface composition and structure

Of obvious concern in the overgrowth of MgO from elemental sources on Cr0.7Mo0. 3 is the possibility of oxidation of the Cr0.7Mo0. 3 surface at the onset of growth. Oxidation of either metal in the alloy will give rise to a disordered oxide phase which is not expected to be a particularly good template for MgO heteroepitaxy. In an earlier publication, we reported preliminary results obtained for thin MgO films grown on Cr0.7Mo0. 3 interlayers at a substrate temperature of 650°C and an oxygen partial pressure in the low 10 -5 Torr range [22]. It was reported that extensive Cr oxidation at the interface occurred under these growth conditions, but that this oxidation did not prevent epitaxial regrowth of MgO. Since that time we have further explored the parameter space of growth conditions and have found a set of conditions that simultaneously enables epitaxial regrowth of high-quality crystalline films, and yet minimizes oxidation at the interface. These conditions have been used in the present study and are

306

S.A. Chambers et aL / Surface Science 339 (1995) 297-309

MgO/200 A Cr o.TMOo.3/MgO(O01)

- 01s

Cr 2p~2 Cr 2p1/2 . ,

k

5 A MgO

/1/ (T,ob=S50"C ~ J l ~ rr)

_c

0

t ,--/

v,

52O

540

~

. . . . . 560 580 Binding Energy (eV)

(T'ub=700°C

p£2=2E.5 ~0. / 600

620

Fig. 10. O ls and Cr2p photoemission spectra obtained at normal emission for thin films of MgO on Cr0.7Mo0.3/MgO(001) grown under more (T~ub = 700 ° and Po 2 = 2 × 10 -5 Torr) or less (Tsub = 550 ° and Po 2 = {2 X 10 -7 Torr) oxidative conditions. The more oxidative conditions result in substantial oxidation of Cr atoms at the interface, as revealed by the Cr2p binding energies.

described in detail in Section 2. To illustrate the difference in extent of interface oxidation, we compare in Fig. 10 XPS spectra obtained at normal emission for thin films of MgO grown under optimized conditions (substrate temperature of 550°C and oxygen partial pressure of 2 X 10 -7 Torr) with a spectrum for a film grown under more oxidative conditions (substrate temperature of 700°C and oxygen partial pressure of 2 X 10 -5 Torr). As reported earlier, higher substrate temperature and oxygen partial pressure result in extensive oxidation of Cr, as judged by Cr2p3/2 and 2Pl/2 binding energies. Cr is more susceptible to oxidation than Mo and is the more sensitive indicator of substrate oxidation. The Cr2p3/2 binding energy under the more oxidative conditions is ~ 577 eV, which is characteristic of oxidized Cr [33]. In contrast, the Cr2p3/2 binding energy measured for films grown at the lower substrate temperature and oxygen partial pressure is ~ 574 eV, which is indicative of Cr metal [33]. It is also interesting to note that although single-crystal epitaxy does occur under more oxidative conditions,

the mean roughness of thin-film MgO surfaces is typically larger by about a factor of two than those of films grown under the less oxidative conditions (compare Fig. 5a in Ref. [22] with Fig. 9b in the present paper). It is also of interest to examine the crystallography of the interface under optimized growth conditions. We have carried out high-resolution TEM lattice imaging and diffraction measurements for these interfaces, and representative results are shown in Fig. 11. This particular specimen consists of a 300 .~ MgO buffer layer, a 200 A interlayer of fro. 7Mo0. 3, and a 1000 ~, top MgO layer. Such a thick regrown MgO layer was used in order to allow specimen preparation by ion milling. The lattice image on the right was obtained with a defocus value of ~ 60 nm and a beam voltage of 200 kV. This defocus value produces excellent images for both the MgO and the C r / M o alloy phases, but does not show a difference in contrast between the two phases due to the sample thickness [23]. The structural coherence across both interfaces is of such quality that the interface cannot be seen in the lattice image. We have used X-ray energy dispersive spectrometry (EDS) in order to determine the approximate locations of the top and bottom interfaces, and these are indicated in the figure. However, in contrast to lattice-mismatched systems such as Ag/MgO(001) in which the inter-

egrown MgO

O0 A of Cro.7Moo.s

IgO buffer layer Fig. 11. TEM selected-area diffraction pattern (left) and high-resolution lattice image (right) for the MgO/Cro.7Moo.3/MgO(001) interface. The image was constructed from the (200) and (110) beams in the MgO and Cr0.7Mo0. 3 phases, respectively. The high degree of structural coherency across the interface precludes locating the interface positions from the lattice image. EDS was used to determine interface positions. The lack of spot splitting or broadening in the diffraction pattern corroborates the conclusion about excellent structural coherence across the interface.

S~A. Chambers et al. / Surface Science 339 (1995) 297-309

face is decorated with dislocations [26], the present system is nearly perfectly lattice-matched and is characterized by a very high degree of structural coherence across the interfaces and no evidence of misfit dislocations. To verify this conclusion, we show at the left of the figure the associated diffraction pattern obtained with a 0.5 /xm selected-area aperture which covers the entire interracial region. The absence of spot splitting or broadening provides further evidence for the high degree of structural coherence across the interface. In addition, there is no evidence whatsoever for phase segregation in the C r / M o alloy. Segregated Cr and Mo grains would be readily apparent in the TEM lattice image because of the nearly 9% lattice mismatch between Cr and Mo. Thus, taken together, Figs. 4, 5 and 11 reveal that that C r / M o interlayer is of exceedingly good crystallographic quality.

4. Discussion We have demonstrated that exceedingly high-quality crystalline films of oxides on metals can be prepared by MBE using elemental sources on lattice-matched metallic substrates. In this case, we have used a metastable alloy film of Cr and Mo. Metastability results because phase segregation, which is suggested by the bulk Cr-Mo phase diagram at temperatures below ~ 880°C, is kinetically hindered by the low diffusivity of the atomic species. We are able to grow the alloy at room temperature, the MgO overlayer at 550°C, and then, in effect, quench the entire system by cooling the specimen in a matter of minutes after MgO regrowth without any phase segregation in the alloy interlayer. By means of lattice and symmetry matching, and with careful manipulation of growth conditions, abrupt, structural coherent interfaces can be achieved. However, obtaining a high degree of surface smoothness is problematic when there is a mismatch in surface free energies for the different materials. In general, for A on B, laminar growth is thermodynamically preferred if "YA+ 'Yi -- 'YB ~ 0, where TA and YB are the surface free energies of A and B, respectively, and Yi is the interfacial free energy [34]. In order for this equation to be valid, local equilibrium must be achieved, which is equivalent to

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having the diffusion rates of deposit material across the growth surface exceed the adatom arrival rates. These conditions are fulfilled in the present experiments by growing at high temperature and low Mg flux in an oxygen overpressure. The surface free energy of the Cr0.7Mo0.3 alloy is determined by a classical thermodynamic model to be 2.3 J / m 2 [35] and that for MgO(001) has been calculated by a total energy formalism to be 2.64 J / m 2 [36]. Unfortunately, the interface free energy is not known, but can be estimated as follows. The fact that MgO does not grow in a completely laminar fashion on Cr0.7Mo0. 3 reveals that YMgO+ ~/i -- YCr/Mo > 0. Similarly, laminar growth of Cro.7Mo0. 3 on MgO(001) implies that YCr/Mo+ Yi -- YMgO < 0. Inserting the above numbers, we find that - 0 . 3 < Yi _< 0.3 J / m 2. A value of Yi in this range suggests weak interfacial bonding. Strong interfacial bonding in the absence of lattice-mismatch-induced strain results in a large, negative value of Yi- Strong bonding generally always accompanies alloy or compound formation, but can also occur at abrupt interfaces for certain interface structures. The fact that Yi is n o t large and negative, as suggested by the lack of laminar growth of MgO on Cr0.TMO0.3, reveals that bonding across this particular abrupt interface is weak. Strong bonding between MgO and Cr0.7Mo0.3(001) is not expected in light of the likely structure of the interface. As seen in Fig. 1, the assumed registry has 0 2- ions in the interfacial MgO layer located in the four-fold hollow sites of the bcc Cr0.7Mo0.3(001) surface. Here, strong overlap is expected between fully-occupied O2pz orbitals, which are normal to the interface, and Cr or Mo d-states. This model is reasonable because previous studies of oxygen on (001)-oriented cubic transition metals, such as Cu, have revealed that chemisorption occurs at the four-fold hollow sites, with O atoms in or slightly above the terminal plane of metal atoms [37-39]. By necessity, Mg 2÷ ions in the interfacial MgO layer are located in the atop positions of the alloy surface. Ion-core repulsion between Mg 2+ ions and Cr or Mo atoms will keep the interplanar spacing sufficiently large that 0 2ions cannot bury as deeply into the hollow site as would be possible for free oxygen. Furthermore, rumpling of the interracial MgO layer, which would

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permit closer approach of O z- ions while maintaining separation between Mg 2÷ ions and Cr or Mo atoms, is not likely due to the rigidity of MgO(001) surfaces. The "rigidity" of MgO(001) is evident by the lack of differential relaxation on the free surface [31]. The net result is weak bonding at the interface. A less plausible geometry has Mg 2÷ ions in the hollow sites and 0 2- ions at the atop sites. This geometry is not likely to lead to strong interracial bonding either because the rigidity of layers of (001)-oriented MgO will likewise prevent strong overlay between Mg 3s orbitals and C r / M o d-states. In order to prevent thermodynamics from driving interfacial MgO layers to a non-laminar morphology, we attempted to kinetically limit surface diffusion by alternating room-temperature deposition of monolayer quantities of Mg with multiple-Langmuir exposures of molecular oxygen with the plasma off. Here, the goal was to form MgO at the interface one layer at a time in a kinetically limited regime where agglomeration would not occur. Subsequent MgO depositions, it was hoped, would then nucleate on a highly planar interracial layer, leading to an atomically flat top surface. Three such layers were grown sequentially, followed by a five-minute anneal at 650°C. Then, several more monolayers of MgO were grown at 650°C in molecular oxygen at 1 X 10 - 6 Torr with the plasma off, followed by the growth of additional layers of MgO with the plasma on and an oxygen pressure of 1 × 10 -5 Torr. The total deposition was ~ 20 monolayer-equivalents of MgO. The RHEED patterns following growth were dominated by transmission and AFM images revealed that the surface was exceedingly rough. The mean roughness averaged over several one-square-micron regions of the surface was 19.5 A, and typiocal terrace widths were at most a few hundred A. Evidently, this approach is hampered by extensive mass transport accompanying oxidation of elemental Mg at room temperature, and is not useful for producing highly laminar films of overgrown MgO on Cr0.7Mo0.3(001).

5. Conclusions We have used molecular beam epitaxy with elemental sources to grow layered structures of MgO and CrxMol_ x (x is ~ 0 . 7 ) on MgO(001) sub-

strates. This C r / M o alloy is lattice-matched to MgO, thereby allowing very high quality crystallographic material to be grown in layered form. A variety of diffraction measurements reveal that the resulting epilayers possess extremely good short-and longrange crystalline order, and that the interfaces are nearly perfectly structurally coherent. Growth conditions have been found which allow high-quality MgO to be overgrown on Cro.TMOo.3 interlayers with negligible oxidation of Cr or Mo atoms at the interface. The surface and interface free energies are such that Cro.7Moo.3 grows in a laminar fashion on MgO, but MgO does not grow in a laminar fashion on Cr0.TMO0. 3. This behavior is a consequence of the slightly higher surface free energy of MgO compared to Cro.7Moo.3 and weak interracial bonding. Nevertheless, unstrained MgO(001) thin films havin~ surfaces with terrace widths of several hundred A and mean roughnesses of ~ 4 A (two interplanar spacings) can be prepared in this fashion on a conducting substrate. This approach will enable definitive surface science measurements to be made on this insulator without undesirable charging by charged-particle probes, and should be readily applicable to other kinds of metal/insulator systems. o

Acknowledgement The authors gratefully acknowledge partial support for this work from the US Department of Energy, Office of Basic Energy Sciences, Materials Science Division.

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