Layer-by-layer quasi-epitaxial growth of a crystalline organic thin film

j. . . . . . . .

ELSEVIER

CRYSTAL GROWTH

Journal of Crystal Growth 152 (1995) 65-72

Layer-by-layer quasi-epitaxial growth of a crystalline organic thin film P. Fenter

a,* P.E. Burrows b,c p. Eisenberger a S.R.

Forrest

b,c

a Princeton Materials Institute, Department of Physics, Princeton University, Princeton, New Jersey 08544, USA b Princeton Materials Institute, Department of Electrical Engineering, Princeton University, Princeton, New Jersey 08544, USA c Princeton Materials Institute, Advanced Technology Center for Photonic and Optoelectronic Materials, Princeton University, Princeton, New Jersey 08544, USA

Received 5 December 1994; manuscript received in final form 19 January 1995

Abstract

We use reflection high energy electron diffraction and grazing incidence X-ray diffraction to study films of the archetype crystalline organic material: 3,4,9,10-perylenetetracarboxylic dianhydride (PTCDA) grown on Au(lll) by organic molecular beam deposition. Although the PTCDA lattice is significantly mismatched to the Au(lll) substrate, the films are orientationally aligned and have a well defined film thickness, providing evidence for layer-by-layer "quasi-epitaxial" growth. Furthermore, we have performed the first characterization of the molecular level disorder in these films, and have determined the presence of a significant (3.8%) strain in the 2D lattice parameters.

I. Introduction

A wealth of new electronic and optoelectronic devices have been realized by the atomic-level control of the composition and structure of elemental and compound semiconductors [1]. Further control of material properties is possible using supramolecular structures such as multiple quantum wells [2]. Unfortunately, requirements for close lattice-matching at semiconductor heterointerfaces restrict these structures to only a few materials combinations since a strained, epitaxial layer cannot be grown beyond a certain critical thickness [3]. Recently, however, a new class of engineered materials has been demon* Corresponding author.

strated - ordered thin films consisting of molecular compounds bonded primarily by the relatively flexible van der Waals (vdW) interaction [4-6]. Even in the absence of a precise lattice match, such films can be ordered over extended distances without a high density of strain-induced defects. This "quasi-epitaxial" (QE) growth has been demonstrated by the ultrahigh vacuum ( U H V ) process of organic molecular b e a m deposition (OMBD) [6-8], and more recently by Langmuir-Blodgett deposition [9]. Furthermore, theoretical work [10,11] has shown that the primary condition enabling ordered growth of lattice mismatched vdW-bonded materials is an asymmetry between the intra- and inter-layer elastic con,, (,~,, stants, where the elastic c o n s t a n t ~intra "e'inter-"is the second spatial derivative (evaluated parallel

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66

P. Fenter et al. /Journal of Crystal Growth 152 (1995) 65-72

to the substrate plane) of the intralayer potential within (between) the contacting layers. That is, Q E growth is favored for systems where ~b'i'ntra> ~)'itnter (independent of the magnitudes of ~)intra and 4~imer)" For systems where 4]i'ntra< 4~'i'nter, incorporation of strain into the organic layer results in a loss of order after only a few monolayers. This ability to grow crystalline multilayers of molecularly designed, organic materials, without regard to lattice matching, offers the potential to engineer a wide range of new materials with properties which are inaccessible through conventional routes. In this paper, we describe a study employing both reflection high energy electron diffraction ( R H E E D ) and high resolution grazing incidence X-ray diffraction (GIXD) [12] to reveal several new and interesting features of the film structure and orientation of the archetype planar molecule: 3,4,9,10-perylenetetracarboxylic dianhydride (PTCDA), grown on A u ( l l l ) by OMBD. Here, Au substrates are used due to their welldefined surface structures and their practical applications as contacts in organic-based optoelectronic devices. Thus, Au is a model substrate for understanding the kinetics of growth leading to Q E of planar-stacking organic molecules typified by PTCDA. We find that P T C D A grows in a (102) surface orientation (that is with the P T C D A lamella oriented parallel to the gold surface), with a single-crystal domain size ( ~ 2000 .~) comparable to that of the Au substrate, and a uniform crystalline thickness of 60 ,~ ( ~ 17 monolayers) over extended distances. We also quantify both disorder within the film, and perturbations of the P T C D A lattice structure due to interactions between the film and substrate. This is, apparently, the first high resolution GIXD study of Q E growth, and provides direct evidence for layer-by-layer crystalline growth, 2D lattice strain, and molecular level disorder in an O M B D thin film.

2. Experimental method

Prior to growth, the Au substrates were cleaned by flame annealing in air with a propane/oxygen torch, and immediately transferred to the U H V

system with a base pressure of < 10-10 Torr, via a load-lock. Each substrate was baked under U H V for 20 min at 600 K to remove any physisorbed material, and cooled to 300 K before growth. The PTCDA, prepurified by repeated temperature gradient sublimation [13], was loaded into a Knudsen cell and thoroughly degassed at 530 K prior to film deposition. The P T C D A film was gorown by OMBD at a deposition rate of 0.1-0.2 A/s, as monitored by a quartz crystal thickness monitor. During growth, the temperature of the Knudsen cell was 600 K and the chamber pressure was approximately 2 x 10 - 9 Torr, and the sample was continuously rotated to avoid any anisotropies related to the deposition geometry. We have characterized the structure of the A u ( l l l ) substrate after film deposition using GIXD, and find that the surface retains a (23 × f3-) A u ( l l l ) surface reconstruction [14], indicating that the bonding between the Au substrate and the P T C D A involves a negligible amount of charge transfer [15]. After growth, the samples were either studied in situ with R H E E D , or were transferred to the X-ray diffractometer under dry N 2. The GIXD studies were performed at the Exxon X10A beam line at the National Synchrotron Light Source (NSLS) at Brookhaven National Laboratories under a vacuum of ~ 10 .8 Tort using the "z-axis" surface spectrometer [12] and an X-ray wavelength of A = 1.285 A. The experimental resolution was determined by slits placed before the detector. The data reported here, taken over a period of 48 h, show no measurable X-ray beam damage.

3. Results

T h e two-dimensional (2D) structure of P T C D A / A u ( l l l ) can be determined from the Bragg peak positions as a function of momentum transfer parallel to the surface, QII, and the azimuthal orientation, 4~, of the diffraction peaks with respect to the Au surface Bragg rods (determined by both R H E E D and GIXD). The spacing and stacking of the 2D lamella are determined by the magnitude of the momentum trans-

P. Fenter et al. / Journal of Crystal Growth 152 (1995) 65-72

Y Fig. 1. R H E E D image (left) from 60 A of PTCDA on A u ( l l l ) . Beam conditions are ~ 1/~A at 10 kV incident at ~ 2° to the substrate plane. The diagram at the right identifies the principal reciprocal lattice rods.

fer perpendicular to the surface plane, Qz, of these peaks as determined by GIXD. In Fig. 1, we show a typical in situ R H E E D pattern from a 60 A thick film of PTCDA on Au. The long, unbroken streaks suggest a flat PTCDA surface with no island growth. We note that due to the ease with which the films are damaged by the high energy R H E E D beam, we were not able to quantitatively obtain data as to the growth mode (i.e. layer-by-layer or islanding) using R H E E D oscillations. However, structural details of the films are obtained from the streak spacings. The principal streaks correspond otO Bragog spacingos (momentum transfer) of 11.0 A (0.57 A - 1 ), 8.0 A (0.78 A-l), 6.1 A (1.0 A - ' ) , 4.7 A (1.3 ,~-1) and 3.9 A (1.6 .~-1) with an estimated uncertainty of _+5%, and are consistent with previous R H E E D

67

data [4]. We index these as the (0,1,1), (0,1,2), (0,2,0), (0,2,2) and (0,2,4) reciprocal lattice rods, based upon the previously determined bulk structure of PTCDA [16,17]. Table 1 summarizes the GIXD data (Figs. 2a-2d and 4) for a similar PTCDA film which are consistent with the R H E E D data, and in which there are no discernible differences in the structure of these films due to atmospheric exposure. The finite Qz of the 'in-plane' X-ray diffraction peaks are a result of the monoclinic stacking of the (102) planes in the bulk structure [16]. Taken together, these measurements show that the PTCDA film grows with a (102) surface orientation. In Figs. 2a and 2b, we show the azimuthal dependence of the (0,1,1) and (0,1,2) diffraction peak intensities. The film/substrate azimuthal orientation is defined (in this sample) by the alignment of the (0,1,2) diffraction peak along the azimuth containing the Au surface diffraction at ~b = 0 (see Fig. 2b). The multiplicity of these peaks is due to the quasi-hexagonal 2D structure within the PTCDA lamella; for instance, while the (0,1,2) peak is aligned along the Au symmetry axis, the (0,1, - 2) peak is not. Consequently, while the data appear complex, the film has a unique orientational relationship with respect to the Au substrate (Fig. 3). Despite the alignment, the Bragg spacing of the (0,1,2) plane is incommensurate by 2% with respect to the nearest multiple of the substrate lattice spacing. From both the radial and azimuthal peak positions in Figs. 2a-2d, we determine that the dimensions of the surface unit mesh are (12.21 _+ 0.03 A)(19.58_+ 0.04 A) (see Fig. 3), and differ from the known bulk structure [16] having dimen-

Table 1 G I X D diffraction peak positions and uncorrected widths

(h,k,I)

QII ( ~ - 1 )

AQII (.~- l)

Qz ( ~ - i )

~b (deg)

(0,1,2) (0,1,- 2) a (0,2,4) (0,1,- 1) a (0,1,1) (1,0,2)

0.8230 + 0.0005 0.8213 _+ 0.0004 1.641 _+ 0.001 0.6064 + 0.0005 0.6065 + 0.0005 0.0 + 0.003

0.015 _+ 0.0006 0.0115 + 0.0006 0.0186 + 0.0017 0.012 + 0.001 0.0139 + 0.001 --

0.367 + 0.003 0.372 _+ 0.004 0.74 b 0.2087 _+ 0.005 0.209 b 1.949 _+ 0.0006

0.05 + 0.06 17.19 _+ 0.05 0.3 + 0.4 23.39 + 0.07 19.64 _+ 0.06 --

a These diffraction peaks come from different sublattices, whose orientation differs only by multiples of 60 ° with respect to the substrate. b The exact Qz position of these peaks was not measured.

P. Fenter et aL /Journal of Crystal Growth 152 (1995) 65-72

68

sions of 11.96 A x 19.91 .A. While these differences preserve (within 0.4%) the area of the unit cell, they result in a change of the aspect ratio (i.e., a/b) of the unit cell by 3.8%. That is, the PTCDA film exhibits an average lattice strain in the form of a rectangular distortion, presumably as a result of the film-substrate interaction• In Fig. 4, we show the X-ray diffraction intensity near the (1,0,2) PTCDA Bragg peak as a

I

-30

I

I

I

function of the momentum transfer perpendicular to the surface, Qoz- From the peak position at Qz = 1.949 + 0.006 A - t , we find a lamella spacing of dl02 = 3.22 + 0.01 A, which agrees with previous results [16]. A rocking scan through this peak (Fig. 4, inset) indicates that the (1,0,2) direction is aligned with the A u ( l l l ) surface normal to within < 0.1°, implying that the P T C D A molecules lie flat on the A u ( l l l ) surface. The o

I

I

I

I

I

I

-20

-10

0

10

20

30

-30

I

I

I

I

I

-20

-10

0

10

20

phi (o)

i

I

I

(0,1,1)

(C)

,

I

0.56

,

0.58

'

'

'

I

(e)

1

,

I

0.6

Q//(fi/1) '

'

'

i

'

.~

,

0.62

'

'

I

0.64

'

'

0.78

0.8

0.8

0.82

Q//(~I)

0.84

0.82

0.84

Q//(/~I)

0.86

'

(0,1,2)

(f)

,

0.78

30

Phi (o)

0.86

1.6

,

(0,2,4)

,

i

1.62

,

i

i

I

1.64

i

I

Q//(~I)

i

I

1.66

i

,

l

1.68

Fig. 2. G I X D data versus m o m e n t u m transfer parallel to the substrate, QII, and the azimuthal orientation, ~b. Azimuthal scan of the (a) (0,1,1) diffraction peak, and (b) (0,1,2) peak. (c) Radial scan of the (0,1,1) diffraction peak for ~ = 23.4 (closed circles) and ~b = 19.6 (open circles). (d) Radial scan of the (0,1,2) diffraction peak for ~b = 0 ° (closed circles) and ~b = 17.2 ° (open circles). (e) and (f): High resolution radial scans t h r o u g h the (0,1,2) and (0,2,4) peaks, respectively ( ¢ = 17.2). The solid lines in (c)-(f) are fits to the data (using a Gaussian and a linear background), while the lines in (a) and (b) simply connect the data points. The resolution was AQI I = 0.009 A i and AQz = 0.18 ,~ 1. In (a), (c), and (f), the intensities are multiplied by a factor of 2.

P. Fenter et al. /Journal of Crystal Growth 152 (1995) 65-72 (a)

oo

•

° o° o 4 o

o °

o #,

I°

" o

o °° o o o~,O~ooO o

o

. .o-o.. .

"-o-

1 o°°o o o

•

(0,1,2)

/, ° oS oo

o

o

- o . . o°%-° o ° o

- ~o-

;o- e(o,1,-1~

o °

(b)

I Fig. 3. A schematic of the (a) reciprocal, and (b) real space structure of the P T C D A film, with the unit cell noted in each. In (a) the open (closed) spots indicate the positions of the P T C D A (Au surface) diffraction conditions as a function of QII" The dashed and solid lines in (a) denote the azimuths in which the (0,1,+1) and (0,1,+2) peaks are found, respectively. Note that the azimuth containing the (012) diffraction peak is aligned with the Au surface diffraction peak (solid spot). The vertical solid lines in (b) denote the (0,1,2) Bragg planes.

existence of fringes on either side of the P T C D A Bragg peak implies that the film has a well defined crystalline thickness over the ~ 1 mm 2 area of the X-ray beam spot. In contrast, films grown under different growth conditions do not show these oscillations, due to 3D islanding. Consequently, these data, together with the R H E E D data which also rule out 3D islanding of the P T C D A film, provide strong evidence that QE film growth can proceed via a layer-by-layer mode. A quantitative analysis of the shape of the (102) Bragg peak provides a more detailed picture of the film structure. While the (102) peak width, Aq, directly determines the average thickness, L, of the films (Aq = 27r/L), the shape of this diffraction peak (in particular, near the shoulders) provides a direct measure of the uni-

69

formity a n d / o r disorder in the film thickness. Surprisingly, these data are well described by an incoherent superposition of the scattering from a uniform, 17 layer P T C D A film and the Au surface (which results in the sloping "background" due the Au crystal truncation rod), as shown in Fig. 4. Since inclusion of interference between the PTCDA and Au results in destructive interference conditions which are not observed in the data, it is apparent that there exists some structural variation in the film. We have attempted to fit the data to a variety of alternative structures which include disorder in the P T C D A film (such as incoherent Gaussian roughness, or roughness of the P T C D A planes due to steps on the Au surface). While these models do not agree as well with the data, they are nevertheless physically reasonable. It is natural then, to estimate the roughness of the film by allowing a range of thicknesses following a Gaussian distribution. Within this formalism, we find an average thickness of 17 layers and a roughness (FWHM) of 1.5 layers. This possible small variation in thickness does not change our conclusion regarding the layer-by-layer growth mode. To investigate the broadening mechanism for the data in Fig. 2, we compare the radial width of the (0,1,2) and (0,2,4) peaks, as shown in Figs. 2e and 2f. The (0,2,4) peak is significantly broader than the (0,1,2) peak, indicative of a non-uniformi

v

'~

f "A -1

0.01

,

1.5

,

,

,

I

2.0

,

0o

,

,

i

,

2.5

Qz(i-b Fig. 4. Radial scan data as a function of Qz of the (1,0,2) Bragg peak corresponding to the lamella spacing. The slowly increasing "background" is due to the Au surface truncation rod. The inset shows the rocking curve of the (1,0,2) peak, showing alignment with the surface normal of the A u substrate. Resolution was AQI I = 0.09 A 1, and AQz = 0.009

70

P. Fenter et al./ Journal of Crystal Growth 152 (1995) 65-72

ity in the lattice spacing within the P T C D A film. Taking into account the finite experimental resolution (0.009 ,~-1), we determine that the intrinsic domain size of the film is ~ 2000 A (comparable to the substrate domain size), and that the Q-dependent broadening is due to a mean variation of the lattice parameter of ~ 2%. In addition we find that the azimuthal width of these peaks (A& ~ 2 ° ) is not determined by the domain size, but instead is due to angular broadening. That is, we find that there is limited in-plane orientational disorder of the P T C D A lattice with respect to the substrate (although the domains remain aligned on average). Through a comparison of the peak widths and positions of the data in Fig. 2, we find that the (0,1,2) peak position is 0.12% larger than the ( 0 , 1 , - 2) peak, and the width of the (0,1,2) peak is also 30% larger than that of the ( 0 , 1 , - 2) peak (Figs. 2e and 2f). It is unlikely that the broadening is due to a macroscopic anisotropy in the sample (such as a sample miscut) since these diffraction peaks are found within a small ( ~ 25° ) angular range on the sample. Given the limited data set, we are unable to unambiguously determine the nature of this broadening. Yet, since the (0,1, + 2) Bragg planes differ only in their orientation with respect to the A u ( l l l ) substrate, these small but significant structural differences should reflect the perturbation of the P T C D A lattice due to its interaction with the Au substrate. We suggest that the broadening is consistent with a mass density wave (MDW) [18]. That is, it is strongest near the (0,1,2) direction, and is weakest at large angles with respect to this alignment direction, where the ( 0 , 1 , - 2) and ( 0 , 1 , - 1) peaks are found at ~ 77° and ~ 83°, respectively, relative to the (0,1,2) peak orientation. MDWs have been predicted to exist in van der Waals bonded adlayers, although they are not expected to play a dominant role in determining the Q E film orientation [10,11]. Since we have found that the P T C D A film is strained, it is anticipated that this strain should be relieved for sufficiently thick films. To study mechanisms of strain relief in these organic thin films, we have also investigated the structure and disorder in thicker films. In Fig. 5, we show

0.822

,

~

,

L

,

0.1 0.01

CY 0.8211

0.001 ~

0.8203

"~ > 10-4

~

0.8194

10_5

0.8185

10-6 0.2

0.4 0.6 0.8 1 e~i, incident angle (°)

1.2

Fig. 5. Plot of the peak radial position, Q, of the (0,1,2) Bragg peak as a function of the incident angle, a. Also shown is the inverse penetration depth, 1 / A ( I / A ) , as a function of a. The shift in the Bragg peak position appears w h e n the incident angle is decreased below the critical angle, a c (i.e., w h e n the penetration depth b e c o m e s smaller than the film thickness).

GIXD data for a thick ( ~ 1100 .&) film grown under conditions similar those in Figs. 1-4. While this film has the same 2D structure as the thin film we find that the (102) peak (not shown) does not exhibit pronounced oscillations in Qz, indicating that this thicker P T C D A film is not as flat as the 60 ,~ film. This is consistent with previous observations of discontinuous R H E E D streaks for thick films [4]. In Fig. 5, we show the (0,1,2) peak as a function of the incident angle of the X-ray beam. Since X-rays exhibit total external reflection, the penetration depth, A, of X-rays is greatly reduced at incident angles below the critical angle, a c (for PTCDA, a~ = 0.16 °, and the penetration depth is A ~ 40 A at angles < ac). These data clearly show a monotonic shift as a function of incident angle, which appears to be correlated with the changing penetration depth. This implies that there is some strain relaxation in the top ~ 40 ,~ of the film. At present, it is not clear if the film roughness and the strain relief are coupled phenomena, or if they occur independently.

4. Discussion

In interpreting these results, it is important to note that we have not yet determined the degree

P. Fenter et a l . / Journal of Crystal Growth 152 (1995) 65-72

to which these findings reflect either an equilibrium or a kinetically limited state. In addition to the data shown here, we have also grown P T C D A films under different conditions which were characterized by 3D islanding. This shows that there may exist more than one film configuration (e.g., meta-stable and equilibrium), and that the film configuration is sensitive to the growth conditions. Yet, from the alignment of the P T C D A film and the observed structural perturbations (e.g., the uniform 2D strain, and the apparent distortions along the alignment direction), we conclude that the perturbations are a result of the interaction between the film and the substrate. In addition, while we have studied many samples (with varying degrees of order and uniformity), we have always found some degree of 2D angular disorder. This is fully consistent with the previously determined favored condition for quasi-epitaxy; i.e., a weak elastic constant between the organic layer and substrate (~b'i'ntra> These observations lead to a detailed picture of quasi-epitaxy: on one hand Q E films may exhibit a significant lattice strain (3.8% change in the 2D aspect ratio). On the other hand, we find that there is only a weak film/substrate interaction as evidenced by the 2D angular disorder. These two observations can be, at least qualitatively, reconciled by noting that organic materials are generally much "softer" than inorganic materials. Consequently, for a given distortion magnitude, the strain energy per layer will be comparatively smaller for organic films than for mismatched inorganic films. Both observations may then be consistent with a relatively weak film/substrate interaction. We anticipate that future studies will reveal the relationship between the observed strain, the critical thickness, and strain relief mechanisms in these films. Finally, it is useful to place these results into the context of the epitaxy of organics in general. It has been found that epitaxy can be achieved in one of two ways. One is by growing a highly strained film. Here, the organic molecules can be forced into an epitaxial structure where latticematching to the substrate results in a film structure far from that obtained in the bulk. A particu-

71

lar example of this is the case of the extremely slow growth of metal-phthalocyanines on semiconductors [6]. Due to the large degree of strain associated with the mismatch of the film with the substrate lattice, the growth becomes disordered after only a few monolayers, which is significantly different than our observations of Q E film growth. On the other hand, it has also been demonstrated that under some conditions, unstrained film structures appear to be achieved. An example of this is P T C D A on alkali-halide substrates held at high temperatures during growth [19]. This resuits in irregular polycrystalline, columnar growth, where the organic is in its relaxed, bulk structure. Our results suggest that Q E films are obtained using conditions intermediate between these two extremes. That is, by growing under moderately non-equilibrium conditions (e.g. low substrate temperature a n d / o r high growth rate), highly organized, crystalline thin films can be obtained on many different substrate structures. Due to the weak, but nevertheless significant interactions between the film and substrate lattices, the overlayer grows in a preferred orientation with respect to the substrate. The strain energy associated with the mismatch is, however, not sufficient to induce large-scale crystalline defects, although lattice distortions are observed which apparently reduce the strain over distances of several tens of nanometers. We are currently examining the factors which appear to connect epitaxial and Q E growth by studying structures which result from varying the growth conditions over a broad range. As has been inferred in previous work [20], recent results in our own laboratory strongly suggest that layer-by-layer QE growth and the achievement of unstrained (although polycrystalline) 3D films are indeed closely related modes of film growth. Which "type" of growth is achieved, however, goes beyond the fundamental condition for Q E (i.e. ~titntra> t~titnter) since our studies indicate that film structure also clearly depends on the growth kinetics. 5. Conclusions

In summary, in situ R H E E D and high resolution GIXD studies demonstrate that quasi-epi-

72

P. Fenter et al. /Journal of Crystal Growth 152 (1995) 65-72

taxial films of P T C D A grown on A u ( l l l ) by U H V OMBD, exhibit a remarkably high degree of crystalline order. A molecularly flat film can be grown in a layer-by-layer fashion over a macroscopic area without lattice-matching constraints. Furthermore, we have performed the first characterization of disorder and strain in a quasi-epitaxial system. These results are consistent with the conditions necessary for Q E of vdW bonded systems [10,11] whose intralayer elastic constants are significantly stronger than those between the film and the substrate. From these and previous observations, we conclude that quasi-epitaxial growth allows for the engineering of a new class of materials with optical and electronic properties tailored to meet the needs of many advanced optoelectronic device applications.

Acknowledgements The authors acknowledge the A F O S R (G. Pomrenke and C. Lee) and the N S F / M R S E C program for partial support of this work. Part of this work was performed at the NSLS, which is supported by D O E Contract No. DE-AC027600016. We thank Andrew Jackson for his assistance during the initial phases of this experiment.

References [1] S.M. Sze, Physics of Semiconductor Devices, 2nd ed. (Wiley, New York, 19xx). [2] R. Dingle, Ed., Semiconductors and Semimetals, Vol. 24, Applications of Multiquantum Wells, Selective Doping, and Superlattices (Academic Press, New York, 1987).

[3] J.W. Mathews and A.E. Blakeslee, J. Crystal Growth 27 (1974) 118. [4] S.R. Forrest, P.E. Burrows, E.I. Haskal and F.F. So, Phys. Rev. B 49 (1994) 11309. [5] F.F. So, S.R. Forrest, Y.Q. Shi and W.H. Steier, Appl. Phys. Letters 56 (1990) 674. [6] H. Yamamoto, H. Tada, T. Kawaguchi and A. Koma, Appl. Phys. Letters 64 (1994) 2099. [7] P.E. Burrows, Y. Zhang, E.I. Haskal and S.R. Forrest, Appl. Phys. Letters 61 (1992) 2417. [8] C. Ludwig, B. Gomph, W. Glatz, J. Petersen, W. Eisenmenger, M. Mobus, U. Zimmerman and N. Karl, Z. Phys. B 86 (1992) 397. [9] R. Viswanathan, J.A. Zasadzinski and D.K. Schwartz, Science 261 (1993) 449. [10] S.R. Forrest and Y. Zhang, Phys. Rev. B 49 (1994) 11297. [11] Y. Zhang and S.R. Forrest, Phys. Rev. Letters 71 (1993) 2765. [12] P. Fuoss, K.S. Liang and P. Eisenberger, Synchrotron Radiation Research Advances in Surface Science, Ed. R.Z. Bachrach (Plenum, New York, 1991). [13] S.R. Forrest, M.L. Kaplan and P.M. Schmidt, Ann. Rev. Mater. Sci. 17 (1987) 189. [14] K.G. Huang, D. Gibbs, D.M. Zehner, A.R. Sandy and S.G.J. Mochrie, Phys. Rev. Letters 65 (1990) 3313. [15] J.P. Bellier and A. Hamelin, Compt. Rend. Paris 280 (1975) 1489. [16] A.J. Lovinger, S.R. Forrest, M.L. Kaplan, P.H. Schmidt and T. Venkatesan, J. Appl. Phys. 55 (1984) 476. [17] The indexing of the RHEED streaks is different from that in Ref. [4]. While those data were interpreted in terms of an expanded surface unit cell for a monolayer of PTCDA on graphite as observed by STM (Ref. [7]), the present data are consistent with a (102) surface orientation of a bulk PTCDA structure. At present we cannot preclude the possibility that the structure for the first monolayer differs from that of the multilayer film. [18] A.D. Novaco and J.P. McTague, Phys. Rev. Letters 38 (1977) 1286. [19] M. Mobus, N. Karl and T. Kobayashi, J. Crystal Growth 116 (1992) 495. [20] M. Mobus and N. Karl, Thin Solid Films 215 (1992) 213.