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In-situ observation of stacking fault evolution in vacuum-deposited C60
J. F. M. Hardigree, I. R. Ramirez, G. Mazzotta, C. Nicklin, and M. Riede
Citation: Appl. Phys. Lett. 111, 233305 (2017);
View online: https://doi.org/10.1063/1.4995571
View Table of Contents: http://aip.scitation.org/toc/apl/111/23
Published by the American Institute of Physics
In-situ observation of stacking fault evolution in vacuum-deposited C
60
J. F. M. Hardigree,
1,a)
I. R. Ramirez,
1
G. Mazzotta,
1
C. Nicklin,
2
and M. Riede
1
1
Clarendon Laboratory, Department of Physics, University of Oxford, Oxfordshire OX1 3PU, United Kingdom
2
Diamond Light Source, Didcot, Oxfordshire OX11 0DE, United Kingdom
(Received 12 July 2017; accepted 13 November 2017; published online 8 December 2017)
We report an in-situ study of stacking fault evolution in C
60
thin films using grazing-incidence x-ray
scattering. A Williamson-Hall analysis of the main scattering features during growth of a 15 nm film
on glass indicates lattice strain as high as 6% in the first 5 nm of the film, with a decrease to 2%
beyond 8 nm thickness. Deformation stacking faults along the {220} plane are found to occur with
68% probability and closely linked to the formation of a nanocrystalline powder-like film. Our find-
ings, which capture monolayer-resolution growth, are consistent with previous work on crystalline
and powder C
60
films, and provide a crystallographic context for the real-time study of organic semi-
conductor thin films. V
C2017 Author(s). All article content, except where otherwise noted, is
licensed under a Creative Commons Attribution (CC BY) license (http://creativecommons.org/
licenses/by/4.0/).https://doi.org/10.1063/1.4995571
The structure of small-molecule organic semiconductors
in thin film electronic devices has been the subject of numer-
ous studies, stimulated by the interest in developing low-cost
electronic devices for health monitors,
1,2
displays,
3
solar
cells,
4
and digital logic
5,6
from simple organic building
blocks. Thin films of small-molecule organic semiconductors
can exhibit a wide range of microstructural motifs, which are
intimately linked to properties as diverse as absorption,
7,8
photogenerated exciton diffusion,
9,10
photoluminsescence,
11
charge carrier mobility,
12–14
and analyte gas diffusivity.
15,16
In-situ techniques that enable real-time monitoring of
vacuum-deposited thin films can provide highly granular
spatio-temporal information for probing the underlying phys-
ics of organic film formation
17–20
and help elucidate the gen-
eralisable processing-structure-property relationships sought
for advanced device fabrication.
The molecular and thin film structure of C
60
fullerene, a
ubiquitous organic semiconductor, has been investigated in
detail since its successful synthesis nearly 30 years ago.
21
Among the many mesoscale features of C
60
solids are a high
density of twins and stacking faults along the h111iclose-
packing direction, attributed to the weak van der Waals inter-
actions between fullerenes.
22–24
Stacking faults are a feature
of materials with face-centered cubic (FCC) structures, in
which several sets of tetrahedral positions are available in a
given h111iplane. While FCC stacking consists of planes
following a 3-plane sequence of alternating tetrahedral posi-
tions (A-B-C-A-B-C), a stacking fault occurs when atoms in
the third plane in this sequence occupy tetrahedral positions
corresponding to those of the A planes. A particular type of
stacking fault known as crystal twinning occurs when the
fourth plane takes the position of a B plane instead (e.g.,
A-B-C-B-A). Such faults disrupt translational symmetry
along the close packed h111idirection, impacting macro-
scopic properties such as charge transport through the forma-
tion of gap states and interfacial energy barriers in inorganic
semiconductor thin films.
25
Although studies of bulk crystal-
line
23
and vacuum-deposited thin film
26,27
C
60
have identi-
fied the presence of stacking faults and twins, there has been
much less work aimed at quantifying these and probing how
they evolve during deposition, especially under industrially
relevant processing conditions.
To investigate the evolution of C
60
, we employed a
recently developed multi-source deposition chamber at the
Diamond Light Source that enables grazing-incidence X-ray
scattering (GIXS) measurements of growing molecular thin
films during thermal deposition.
28
Capturing the growth
dynamics of organic molecular films at industrially relevant
deposition rates (0.1–1 A
˚/s)
29
presents various technical chal-
lenges. The low scattering density of organic materials neces-
sitates longer exposure times than that of films of a similar
thickness of high-Z atomic species to achieve similar con-
trast.
30
It is also established that organic thin films—in partic-
ular samples less than 10 nm thick—are susceptible to x-ray
beam damage.
30,31
Consequently, tracking the scattering of
the growing film requires balancing the deposition rate and x-
ray exposure time, such that each image provides a suffi-
ciently discrete snapshot of the film’s current state without
subjecting it to detrimental beam damage that could influence
subsequent layers. To balance these competing needs, we set
a fullerene evaporation rate of 0.1 A
˚/s (1 monolayer/min)
and used a 10 s exposure time on a 2D detector (Pilatus 2M).
To limit the beam exposure, we spaced measurements by 65 s
[Fig. S1(a) in the supplementary material]. Because the angle
of incidence during in-situ measurements was set to 0.072,
the minimum x-ray penetration depth is z
1/e,min-
¼(2冑r
e
pq)
1
¼7.5 nm in C
60
.
32,33
As a result, the images
obtained during the first 7.5 nm capture the accumulating
growing film on the substrate, while images acquired at
greater thickness capture the upper 7.5nm of the film. Images
were calibrated using silver behenate (AgBeh), and data
reduction was performed using DAWN;
34
further details on
sample preparation and alignment procedures can be found in
the supplementary material and in a separate report with the
technical details of the deposition chamber.
28
a)
Author to whom correspondence should be addressed: josue.martinez-
hardigree@physics.ox.ac.uk
0003-6951/2017/111(23)/233305/5 V
CAuthor(s) 2017.111, 233305-1
APPLIED PHYSICS LETTERS 111, 233305 (2017)
During the initial few nm of growth, we observe the nucle-
ation of C
60
crystallites on the glass surface, as revealed by the
rapid changes in the diffuse scatter captured in the images of
Figs. 1(a)–1(d). The convergence of the broad scattering near
q
z
¼0.15 A
˚
1
towards q
z
¼0.138 A
˚
1
in the first 300 s or 3 nm
of growth is consistent with the complete coverage of the sur-
face with 4 monolayers of C
60
. By monitoring the parallel
scattering q
xy
over the first 300 s in Fig. 1(e),anestimateofthe
distance between island centers
35
can be extracted from the
peak position of the diffuse scatter, with the correlation length
D
diffuse
2p/Dq
xy
. Features appear near q
xy
¼0.02 A
˚
1
after
depositing just under 1 nm of C
60
(1 monolayer), consistent
with an island-island spacing of 29.1 nm. The tapering of
D
diffuse
to a constant value of 45 nm after 8 nm thickness
reflects that the film has reached steady-state growth and coin-
cides with the expected penetration depth of x-rays at this
energy and angle of incidence. This value is consistent with
47.69 nm found from the FFT analysis of atomic force micros-
copy (AFM) of a 15 nm C
60
film (Fig. S4 in the supplementary
material).
Simple reflectivity models with constant roughness
r
C60/air
¼0.2 nm for the C
60
layer and glass roughness
r
glass/C60
¼1.0 nm (Fig. S2 in the supplementary material)
are only able to match the oscillation period obtained at
q
z
¼0.138 A
˚
1
, but a full description of the measured off-
specular reflectivity using an adjustment of the model by
Woll et al.
18,19
is required to fully capture the growth behav-
ior of the film. As a consequence, the reported film thickness
is based on the response of a calibrated quartz crystal monitor
(QCM) within the deposition chamber
28
and does not account
for changes in thickness which may arise from differences in
the sticking coefficient of C
60
at room temperature within the
first few layers of the film. For a comprehensive investigation
of the sticking coefficient and thermally assisted dewetting of
C
60
and its implications on observed film growth, the reader
is referred to a recent study that makes use of in-situ x-ray
reflectivity to quantify dewetting and upward mass transport
by monitoring the specular signal from film deposition up to
60 min post-deposition.
36
Complementary in-plane scattering along the substrate
horizon (q
z
¼0.022 A
˚
1
) affords insights into the crystallo-
graphic evolution of the upper 7.5 nm of the growing film.
Reflections from background-subtracted images (details in sup-
plementary material) were fitted to Gaussian peaks using open-
source software tools.
37,38
AsshowninFig.2(b),the{111}
reflection (q
xy
¼0.8 A
˚
1
) appears after the first nanometer of
growth, indicating that the surface is populated by hexagonally
close packed islands of C
60
. Initially, the features near q
xy
¼1.3 A
˚
1
and 1.4 A
˚
1
are indistinguishable from a broad scat-
tering area, but within the first 4 nm, these features quickly
become more readily distinguishable from one another, and
their peak widths begin to narrow consistent with domain
growth. During the first 5 nm of growth, we observe a marked
transition from near the {310}
FCC
plane to a value between the
{311}
FCC
plane and the {021}
HCP
plane. Simultaneously, the
feature near 1.3 A
˚
1
initially appears near the {221}
FCC
plane
but with increasing thickness converges towards a qvalue
between that of the {220}
FCC
and {111}
HCP
planes. These
shifts of the scattering vector for both features to intermediate
FIG. 1. Grazing-incidence x-ray scattering (GIXS) images during nucleation
and growth on glass. (a)–(d) Detector images in the small angle region near
the beamstop. The black band in the images corresponds to intermodular gap
on the Pilatus detector. (e) Diffuse scatter as a function of film thickness mea-
sured in the region denoted by the white box (q
z
¼0.138 A
˚
1
) indicated in (a).
FIG. 2. Comparison of scattering features observed in the GIXS measure-
ments at q
z
¼0.022 A
˚
1
. (a) Williamson-Hall (WH) plot of Dqvs qfor
thickness starting at 5 nm. Note: For each qvalue, large icons indicate val-
ues for the 15 nm film from the post-growth scan. The arrow indicates the
direction of the increasing thickness. (b) Evolution of the peak position for
each scattering feature with the increasing film thickness. Solid lines: FCC-
indexed planes; dashed lines: HCP-indexed planes; and broad dashed lines:
qvalues where both FCC and HCP planes can be indexed. The feature at
0.7 A
˚
1
is the Pilatus intermodular gap, and the step at 1.1 A
˚
1
is a shadow
from the substrate shutter on the Be window. The arrow indicates the direc-
tion of the increasing thickness. (c) Extracted grain size (left) and strain
(right, orange) from individual peak fitting and WH analysis. (d) Stacking
fault probability calculated from Dqfor each plane; green symbols are cal-
culated relative to the feature at 1.3 A
˚
1
and red symbols relative to the
feature at 1.48 A
˚
1
.
233305-2 Hardigree et al. Appl. Phys. Lett. 111, 233305 (2017)
values are consistent with the accumulation of several % stack-
ing faults in FCC systems.
39
According to Warren’s selection rules,
40
stacking faults
in FCC materials modify reflections in planes for which
hþkþl6¼3m. In the frame of an FCC lattice, it has been
shown that shifts in the peak position and increased breadth
(FWHM) of {220} peaks mark the onset of stacking faults
along the close-packing direction.
21,24
In the case of powder-
like samples, a Williamson-Hall (WH) plot can be a useful
tool to decouple the influence of grain size and defects on
shifts in scattering vector and line broadening. In a WH plot,
Dqvs qacross all peaks is fit to a linear or a quadratic func-
tion of the scattering vector using the Williamson-Hall (WH)
relation
41
Dq¼2p=Dþ2eq, where the domain size Dand
lattice strain ecan be extracted from the intercept and slope
of the plot, respectively. Figure 2(a) shows the WH plot for
the film thickness starting at 5 nm, beyond which all three
planes could be accurately fit to Gaussians. Despite only fit-
ting three peaks [20 thickness values, linear least squares fit-
ting with mean (l) and standard deviation (r) for the squared
residuals R
2
:l¼0.84 and r¼0.18], linear fits consistently
yield a non-zero positive intercept, indicating that part of the
line broadening arises from small (10 nm) grain scattering
in the film [Fig. 2(c)]. The larger grain sizes extracted from
the analysis below 8 nm thickness may reflect a greater
extent of in-plane grain connectivity in the film, consistent
with the larger q (and smaller island-island distance D
diffuse
)
at low thickness seen in Fig. 1(e). However, it is important to
note that these grain sizes may also include contributions
from changes in the peak breadth Dqand position q,as
both strain and grain size [Fig. 2(c)] are seen to exhibit simi-
lar changes below 8 nm film thickness where individual
peaks only begin to emerge in the diffraction [Fig. 2(b)].
Additionally, fits of the data for films above 7 nm thickness
converge towards a common slope, indicating a smaller role
of microstrain on the observed q-dependence of the line
broadening away from the substrate.
42
As shown in Fig. 2(c),
this strain decreases from as much as 5% when probing C
60
within 7.5 nm of the substrate interface, to just under 2%
in the upper 7.5 nm of the film away from the substrate.
Previously reported strain values near the substrate interface
using in-situ reflection high energy electron diffraction
(RHEED) measurements of laser-deposited C
60
on mica
43
yielded a value of 3.2%, based on double lattice constant
estimates. Given that amorphous glass is not expected to pro-
vide long-range templating for the C
60
lattice, this extracted
strain value is qualitatively in line with the powder-like
growth observed herein (Fig. S5 in the supplementary
material).
Because the in-situ measurements at this angle of inci-
dence are most sensitive to the upper 7.5 nm of the film, in-
plane scattering effectively captures both 1D line defects
such as dislocations at grain boundaries between nanocrys-
tals and 2D planar defects such as stacking faults within the
7.5 nm (10 monolayer) cross-section of the nanocrystals.
To compare the surface and bulk properties of the thin film,
immediately after deposition, the substrate shutter was
closed and the sample was probed between x¼0and 0.2
(Fig. S4 in the supplementary material). As observed during
the in-situ measurements of the upper 7.5 nm layer, the
{111}
FCC
peak exhibits a larger width and hence a smaller
coherence length Dthan the other planes in the film. When the
full film is measured along the horizon at a sample tilt
x¼h
c,glass
¼0.152,wefindD
111
¼8.6 nm, D
22L/300
¼4.9 nm,
and D
311
¼4.7 nm, all consistent with those derived from the
individual peak fitting of the in-situ data. The good fit (R
2
¼0.999) and negligible intercept suggest that a high degree of
defects is the main cause for the q-dependent peak broadening.
The probability that the 4th plane in an FCC 111 plane
sequence of the form A-B-C-A-B-C incorporates a planar fault
of the form A-B-C-B-[…] is the stacking fault probability a,
and the average number of planes between faults can be esti-
mated as 1/a. The probability of deformation stacking faults a
in a powder sample can be estimated using the Warren and
Warekois formula,
44,45
which is given as
a¼
D2hoff
p2h3
jtan hpeak cos2u270 ffiffiffi
3
p;(1)
where D(2h) is the peak offset in degrees, h
3
¼(h
2
þk
2
þl
2
)/
jhþkþlj,jis the fraction of faulted planes for the family of
planes {hkl}, h
peak
is the scattering angle of the reference
peak, and cosuis the average angle between the [111] plane
and faulting planes within {hkl}. A closer examination of
Fig. 2(b) (and Fig. S3 in the supplementary material) indi-
cates that all three main peaks are shifted relative to their
crystallographic references, and so, it is useful to compare
relative shifts between two separate peaks, using the relation
D2hhkl ¼Gav jtan ha;(2)
where hG
av
i¼690冑3•h
3
/p2l
0
2
, and h
3
is the reciprocal
vector averaged over all faulting hkl planes given by
hcosui¼h
3
/冑3l
0
. To estimate the stacking fault probability,
the two diffraction features at 1.3 A
˚
1
and 1.48 A
˚
1
were ref-
erenced to the {220} and {311} planes, respectively. As
shown in Fig. 2(d), the probability of deformation stacking
faults in this film reaches a value of a¼68% relative to the
{220} planes and a¼4.2% for the {311} planes as measured
in the full 15 nm film. These values correspond to an average
number of {111} planes between stacking faults relative to
these references of 1.5 and 24 planes, respectively. The large
calculated probabilities for the {220} planes below 7 nm are
attributable to the significant uncertainty in the position of
the scattering vector within the broad signal [cf. Fig. 2(b)].
As the film increases beyond z
1/e
¼7.5 nm, the decreasing
values of aare averaged-out with less influence from the
substrate, with the plateau near 12 nm, marking the mea-
surement of just the upper “bulk” of the thin film. Although
a full averaging out would be expected at 2z
1/e
¼15 nm, this
difference may simply reflect the outsize influence of the
2 nm nearest the glass interface, where nucleation is taking
place on the substrate. Although high for FCC metals, these
values for aare consistent with simulations of stacking faults
in vacuum-deposited C
60
by Vaughan et al.
27
in which thin
film samples were best described by models with a 50%
probability of HCP stacking on FCC C
60
underlayers.
Lastly, based on the in-situ and post-growth coherence
lengths for the different reflections, it appears that once
the C
60
film transitions to steady state island growth, the
233305-3 Hardigree et al. Appl. Phys. Lett. 111, 233305 (2017)
crystallites are limited to a size of roughly 10 nm along the
FCC h111iand FCC h22 Li/h300idirections. However, the
WH analysis for the full 15 nm film yields a larger grain size
of 174 nm [square symbol, Fig. 2(c)], a factor 20 greater than
the individual peaks and roughly 8the 19.5 nm size calcu-
lated from AFM images (supplementary material, Table S1).
This larger value arises from the fact that when probing the
full film, the diffracted volume is a superposition of the
microstructure through the vertical composition of the film.
As seen in Fig. 2(a), both the measured peak position and
width for the full film measurement are closer to the average
values than the those of the upper half of the film. Moreover,
comparison of the sizes obtained from integral breadth and
WH analysis for the in-situ measurements and the full film
[Fig. 2(c)] indicate close agreement between the two meth-
ods for all but the full film measurement, suggesting that the
differences arise from the WH fitting of the full film data.
Although the higher deposition rate, and short time interval
between in-situ and post-deposition scans as compared to
other studies suggest that dewetting might not be a main con-
sideration for the observed grain size differences, combining
this analysis with in-situ and post-deposition XRR as done
by Bommel et al.
36
would facilitate a quantitative compari-
son between grain size distribution and film homogeneity
normal to the surface, enabling a further link between the
kinetics of film dewetting and as-grown film microstructure.
Additionally, capturing a greater region of qspace, e.g., by
probing at higher incident x-ray energy, would improve grain
size estimates by using a greater number of diffraction orders
in the analysis. It is well-documented that WH analyses can
overestimate grain sizes by more than an order of magni-
tude,
42
in our specific case we can trace this inconsistency to
the vertical inhomogeneity of the film, which is only
resolved when comparing the analysis of the post-deposition
and in-situ GIXS measurements.
46
In summary, our results comprise in-situ monitoring of
stacking fault evolution in C
60
fullerene thin films while
employing deposition rates consistent with scalable, state-of-
the-art organic electronic device manufacturing. Although ini-
tial room-temperature growth on glass is marked by low
roughness and nearly 6% in-plane microstrain, C
60
quickly
incorporates stacking faults, with a
311
¼4.2% and a
220
¼68%.
The observation that these material parameters vary with the
distance from the substrate underscore the need for in-situ
characterisation to understand the coupling of interfacial and
bulk material properties in organic thin films. Our findings
highlight the potential for in-situ GIXS as a powerful multi-
length-scale probe for elucidating the structural and morpho-
logical evolution of vacuum-deposited molecular thin films for
next-generation organic electronic devices.
See supplementary material for details on GIXS image
reduction, in-situ lattice constant evolution, and post-
deposition AFM analysis of the deposited thin film.
This work was supported by a Science and Technology
Facilities Council (STFC) Challenge-Led Applied Systems
Programme (CLASP, ST/L003309/1) focused on advancing
the commercialization of organic solar cells. J.F.M.H. thanks
Wolfson College, Oxford, for Junior Research Fellowship
support. G.M. was supported by the Energy and Physical
Sciences Research Council (EPSRC) Centre for Doctoral
Training in New and Sustainable Photovoltaics and by the
University College Oxford through the Oxford-Radcliffe
scholarship. The authors thank G. Christoforo, P. R. Warren,
S. V. Kesava, and H. Ye (Univ. of Oxford) and A. Warne
and J. Rawle (Diamond Light Source) for their assistance
with beamline instrumentation.
J.F.M.H. and M.R. developed the thin film research
program. J.F.M.H. wrote the manuscript with contributions
from all authors and directed the thin film characterisation
and data analysis. I.R.R., G.M., and C.N. assisted with GIXS
measurements and data analysis.
The authors declare no competing financial interests.
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