FEATURE ARTICLE
Organization of Matter on Different Size
Scales: Monodisperse Nanocrystals and
Their Superstructures**
By Andrey L. Rogach,* Dmitri V. Talapin,* Elena V. Shevchenko,
Andreas Kornowski, Markus Haase,
and Horst Weller
Advanced colloidal syntheses enable the preparation of monodisperse semiconductors
and magnetic alloy nanocrystals. They can be further used as building blocks for the
fabrication of ordered assemblies: two-dimensional and three-dimensional arrays and
colloidal supercrystals. This article reviews our recent activities in these fields. A theoretical description of the evolution of an ensemble of nanoparticles in a colloidal solution is applied to the problem of
control over the nanocrystal monodispersity.
1. Introduction
Chemistry and physics on the nanometer scale have experienced an enormous development in the last decade leading to
the appearance of the new interdisciplinary field of ªnanoscienceº. The interest in nanoscale materials arises from the
possibility to manipulate them in one (quantum wells), two
(quantum wires), and three (quantum dots) dimensions. In the
last structures, atomic-like electronic energy levels are formed
due to the charge-carrier confinement, so that the properties of
semiconductors and metals become governed by size.[1,2] There
are two distinct routes to produce quantum dots: in the ªphysicalº approach they are grown by lithographic or molecular
beam techniques. In the ªbottom upº, or ªchemicalº approach,
they are synthesized by methods of colloidal chemistry in a solvent medium, and the term nanocrystals is commonly used to
denote them. A famous demonstration of the size-dependent
properties of semiconductor nanocrystals is the continuous
change of their emission color with decreasing particle size
(Fig. 1).
±
Fig. 1. ªTraffic lightsº and ªrainbowsº from luminescent colloidal nanoparticles.
a) Thiol-capped CdTe nanocrystals synthesized in aqueous solution emit green,
yellow, or red light depending on size (2.5, 3.0, and 4.0 nm, correspondingly) with
room temperature quantum yield of up to 40 %. b) Hexadecylamine-trioctylphosphine oxide-trioctylphosphine capped CdSe/ZnS core±shell nanocrystals are
soluble in non-polar organic solvents and emit in the whole visible spectral region
depending on size with a quantum efficiency of 40±70 %. c) and d) CdSe/ZnS
nanocrystals embedded into a polylaurylmethacrylate matrix retain their superior luminescent properties.
[**] The authors gratefully acknowledge all colleagues who have contributed to
this work, especially mentioning Dr. S. Haubold (Nanosolutions GmbH)
for his contribution to the synthesis of the InAs nanocrystals. The financial
support was provided by the SFB 508, research projects BMBF-Philips and
BIOAND, and by the DFG Schwerpunktprogramm ªPhotonic Crystalsº. A
more complete treatment on the topic of semiconductor nanoparticles and
their superstructures can be found in the chapter on ªSemiconductor Nanoparticlesº, which is to appear in the book Colloids and Colloidal Interfaces
(Ed: F. Caruso), Wiley-VCH, Weinheim.
Because the strong bandgap luminescence of colloidally
synthesized semiconductor quantum dots is tunable by size
due to the quantum confinement effect, they are currently
intensively investigated as emitting materials for thin film
light-emitting devices,[3±5] optical amplifier media for telecommunication networks,[6,7] and biological labels.[8±10] The incorporation of luminescent semiconductor nanocrystals into 3D
photonic crystals[11] and microcavities[12] has attracted considerable attention recently as a promising pathway to novel light
sources with controllable emission. The preparation of ordered
[*] Dr. A. L. Rogach, Dr. D. V. Talapin, E. V. Shevchenko, A. Kornowski,
Dr. M. Haase, Prof. H. Weller
Institute of Physical Chemistry, University of Hamburg
D-20146 Hamburg (Germany)
E-mail: [email protected];
[email protected]; [email protected]
Adv. Funct. Mater. 2002, 12, No. 10, October
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FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
arrays of magnetic nanocrystals is attracting growing interest
because of potential magnetic data-storage applications.[13±15]
The synthesis of monodisperse nanocrystals of desired sizes
is the first and very important step being a pre-requisite of any
further investigation and their use in practice. Monodispersity
is strongly required, e.g., for the purity of emission color in the
case of luminescent semiconductor nanocrystals, as well as for
many other potential applications. Colloidal chemistry routes
to several semiconductor[16±20] and metal[13±15,21,22] nanocrystals
with narrow particle size distributions, high crystallinity, and
controllable surface properties have been developed providing
nanoparticles in gram amounts just like ordinary chemical substances. The synthetic efforts are gradually being concentrated
on both the improvement and simplification of existing syntheses[23,24] and the development of reliable approaches to more
and more nanometer-sized compounds.[25±27] Despite of this,
there is still a lack of theoretical understanding of the processes
occurring during the growth of nanoparticles in colloidal solutions and the problem of keeping the particle size distribution
narrow. Only a few reports deal with this important problem.[28±30] Further synthetic progress will definitely depend on
our ability to understand and control the parameters governing
the properties (e.g., size distribution and photoluminescence
quantum efficiency) of colloidally grown nanocrystals.[31,32]
A new field of research has recently emerged, which focuses
on the use of individual monodisperse nanocrystals as building
blocks for the fabrication of superstructures and the investigation of collective properties of these artificial quantum dot solids. Reviews on this topic have appeared[22,33±35] providing both
a summary of the literature at the time of publication and the
developments done by the respective groups. This article covers the very recent activities of our group on the colloidal synthesis of monodisperse semiconductor and magnetic alloy
nanocrystals and their use for the creation of two-dimensional
(2D) and three-dimensional (3D) arrays and colloidal crystals
by means of self-assembly. In colloidal crystals, individual
nanocrystals play the role of building blocksÐartificial atoms
in the next level of hierarchy. A recently developed theoretical
description of the evolution of an ensemble of nanoparticles in
a colloidal solution[30] is applied to the problem of control over
the nanocrystal monodispersity.
2. Synthesis of Monodisperse Nanocrystals
The availability of reliable colloidal syntheses leading to
nanocrystals being uniform in composition, size, shape, and
surface chemistry is crucial for the fabrication of superstructures. The use of thiols as capping agents in the syntheses of
II±VI semiconductor nanocrystals allowed the preparation of
extremely small molecular-like clusters of exact composition,
e.g., [Cd17S4(SC6H5)28]2±,[36] [Cd32S14(SC6H5)36]´(DMF)4,[37]
[Cd17S4(SCH2CH2OH)26],[38] [Cd32S14(SCH2CH(OH)CH3)36]´
(H2O)4,[39] and [Cd54Te32(SCH2CH2OH)52]8±[40] (the last formula was suggested based on extended X-ray absorption fine
structure (EXAFS) data). The stable clusters correspond to
pronounced minima of the free energy vs. particle-size
dependence owing to their closed structural shells (the concept
of so-called clusters of magic size in the earlier literature[41])
and are naturally ª100 % monodisperseº. These nanoparticles
Andrey L. Rogach received his Ph.D. degree in Physical Chemistry from the Belarussian State
University in Minsk in 1995 for his work on the formation and properties of silver nanoparticles in
different media. He was a DAAD Postdoctoral Fellow at the Institute of Physical Chemistry, University of Hamburg (Germany) in the group of Prof. Horst Weller, whereafter he returned to
Belarus to a position of Senior Research Scientist at the Physico-Chemical Research Institute
(Minsk). He was a guest scientist at the British Telecom Laboratories (Ipswich, UK) and the Oklahoma State University (USA) from 1998±1999. He revisited the University of Hamburg as an Alexander von Humboldt Research Fellow in 2000, followed by a position in the group of Prof. Weller
in 2001. In September 2002 he joined the Photonics & Optoelectronics group of Prof. J. Feldmann
at the Ludwig-Maximilians-University of Munich. His research focuses on different aspects of
chemistry, physics, and applications of semiconductor and metal nanocrystals and on colloidal
photonic crystals.
Dmitri V. Talapin studied chemistry at the Belarussian State University in Minsk, Belarus. Since 2000
he has been working at the Institute of Physical Chemistry, University of Hamburg (Germany), in
the group of Prof. Horst Weller. He received his Doctor of Natural Sciences degree in 2002 for his
work on experimental and theoretical studies on the formation of highly luminescent II±VI, III±V,
and core±shell semiconductor nanocrystals. His current area of research focuses on the organometallic colloidal synthesis of semiconductor nanocrystals and the theoretical modeling of the evolution
dynamics of nanometer-sized particles in a colloidal solution.
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can be crystallized in macroscopically large single crystals allowing their investigation by single-crystal X-ray analysis.[36±39]
Moving from molecular-like clusters of exact composition to
larger particles whose crystalline core consists of ~ 102±104
atoms, nearly continuous tunability of the particle size becomes
possible, as an addition or removal of a unit cell requires only a
small variation of the nanocrystal free energy. The colloidal
synthesis of nanocrystalline particles generally involves several
consecutive stages: nucleation from initially homogeneous solution, growth of the pre-formed nuclei, isolation of particles
from the reaction mixture after they reached the desired size,
post-preparative size fractionation, etc. As a rule, temporal
separation of the nucleation event from the growth of the nuclei is required for narrow size distribution.[35,42] The so-called
hot-injection technique, where the precursors are rapidly
injected into a hot solvent with subsequent temperature drop,
satisfies this requirement.[16,35] During further growth of the
nanocrystals, several regimes can be observed depending on
the system and experimental conditions. They are discussed
below.
dr
S e1=r
ˆ
ds
r ‡Kea=r
with r ˆ
2.1. Ostwald Ripening
Ostwald ripening (OR), the growth mechanism where the
smaller particles dissolve and the monomer released thereby is
consumed by the large particles,[43,44] takes place in most colloidal syntheses of both II±VI and III±V semiconductor nanocrystals. As a result, the average nanocrystal size increases with
time, and the particle concentration decreases. The driving
force of OR is a decrease of the particle solubility with increasing size as expressed by the Gibbs±Thompson equation:
2 γVm 
 2 γVm 
0
0 
C ( r ) = Cbulk
exp 
≈ Cbulk
1 +

rRT
rRT 



(1)

where, C(r) and C0bulk are the solubilities of a particle with
radius r and of the bulk material, respectively, c is the surface
tension, and Vm is the molar volume of the solid. Validity of
the Gibbs±Thompson equation was proven for very small
(r ~ 1±2 nm) colloidal particles.[45,46]
In case of ªlargeº (r > 25 nm) particles the kinetics of OR
can be satisfactorily described analytically in the framework of
the Lifshitz±Slyozov±Wagner (LSW) theory.[47,48] However, the
LSW approach takes into account only two terms of the expansion of the Gibbs±Thompson equation and fails in the description of ensembles of particles smaller than ~ 50 nm in radius
due to the large error arising from the truncation of the expansion of Equation 1. The coefficient 2cVm/(RT) called ªcapillary
lengthº is of the order of 1±3 nm for most solid±liquid interfaces[49,50] and in the case of nanoparticles with r = 1±5 nm the
particle solubility is strongly nonlinearly dependent on r±1.
Moreover, the surface tension c of the nanoparticles can be
considerably larger than that of the corresponding bulk material, as was reported for CdS,[51] Pt, and Au.[52] This results in a
value of capillary length of ~ 33 nm for thiophenol-capped CdS
nanocrystals under the particle growth conditions.[51] Additionally, for nanoscale particles the activation energies of the
Adv. Funct. Mater. 2002, 12, No. 10, October
growth and dissolution processes are also size-dependent.[30] In
this case a general analytical solution describing all processes
occurring during the evolution of the particle ensemble could
not be obtained. In our recent study we applied the MonteCarlo simulation technique to describe the evolution of an
ensemble of nanoparticles in a colloidal solution[30] and showed
that OR in this case is characterized by some features that are
not observed for large (sub-micrometer- and micrometer-sized)
particles. For convenience, we will use the term ªnano-ORº in
further discussion to distinct this particular case from OR of
(sub-)micrometer-sized particles adhering to the LSW theory.
OR implies that the largest particles in the ensemble have
positive and the smallest ones have negative growth rates. The
growth/dissolution rate of a single particle of radius r is given
by the following equation, which seems to be valid for both
nano- and (sub-)micrometer-sized particles:[30]
and τ =
(2)
RT
r
2cVm
0
R 2T 2 DCbulk
4γ 2Vm
(3)
t
(4)
where r* and s are the dimensionless particle radius and time,
respectively, and K is a dimensionless parameter describing the
impacts of diffusion and surface reaction as kinetics-limiting
factors (K << 1 corresponds to the diffusion-controlled process
and K >> 1 to the surface-reaction-controlled one).[30] The
dimensionless parameter S = [M]/C0bulk describes the oversaturation of the monomer in a solution, with the monomer concentration [M]. D is the diffusion coefficient for the monomer,
and a is the transfer coefficient of the activated complex
(0 << a << 1).
The size-dependent particle growth/dissolution rates are
shown in Figure 2a for the cases of purely diffusion-controlled
and purely reaction-controlled nano-OR. Figure 2b shows the
temporal evolution of the particle size distribution during diffusion-controlled nano-OR. The LSW theory predicts that OR
finally leads to an asymptotic particle size distribution, which is
independent of the initial conditions.[47] A similar situation is
realized in nano-OR: any initial particle size distribution
evolves in time toward an asymmetric negatively skived one.
The asymptotic particle size distributions inherent to diffusion
and surface reaction-limited nano-OR are shown in Figure 2a
together with the corresponding instantaneous single particle
growth rates. The resulting particle size distribution is narrower
if the particle growth takes place under diffusion control.
Nano-OR proceeds considerably faster than the OR of
ensembles of larger particles as shown in Figure 3a. The evolution of the width of the particle size distribution expressed as
percentage of standard deviation (r %) vs. average particle size
is shown in Figure 3b. Independent of the initial width of parti-
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FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
a
b
Fig. 2. a) Dashed lines: size dependence of a single particle growth rate during
surface reaction- (top) and diffusion- (bottom) limited Ostwald ripening. The
horizontal dotted lines correspond to zero particle growth rate (equilibrium
between nanoparticles and monomer in solution). Solid lines: corresponding
ªstationaryº particle size distributions after prolonged growth of an ensemble of
nanoparticles: the distribution shape becomes independent of time. b) Temporal
evolution of size distribution of the nanoparticle ensemble during the Ostwald
ripening under diffusion control.
cle size distribution, nano-OR results in a steady-state distribution with a standard deviation independent of the average particle size. If the initial size distribution is narrower/broader
than the steady-state one, the nano-OR is accompanied by
broadening/narrowing of the particle size distribution, correspondingly.
Remarkably, the steady-state particle size distribution inherent to the nano-OR is narrower than that of the OR in ensembles of (sub-)micrometer-sized particles. Moreover, for the
nanometer-sized particles the width of the steady state size distribution depends on the surface tension at the particle±solution interface. Figure 3c shows stationary values of the standard deviation corresponding to different surface tensions at
the particle±solution interface and demonstrates the decrease
of the steady state value of r with increasing c in the case of
diffusion-limited nano-OR. In contrast, a constant value of
r = 21.5 % is predicted by LSW for diffusion-limited OR.
The origin of the difference between nano-OR and conventional OR is caused, as may be expected, by the particle radius.
As was mentioned above, particles with r > 25 nm fulfill the
condition 2cVm/(RT) << r and higher terms of the expansion of
the Gibbs±Thompson equation can be omitted, as shown in
Equation 1. However, if 2cVm/(RT) » r or even > r, these terms
play a major role, accelerating the growth and the dissolution
of particles. This also results in an acceleration of the coarsening rate and in a narrowing of the size distribution, as demonstrated by Figure 3. If 2cVm/(RT) is markedly larger than r, a
stationary value of r as small as 15 % can be achieved
(Fig. 3c).
In summary, Ostwald ripening of nanoparticles occurs faster
and can result in narrower size distributions than in ensembles
of (sub-)micrometer-sized particles. The narrowest size distribution can be achieved if particle growth rates are limited by
the diffusion of monomer from the bulk of the solution towards
the particle surface. From an experimental point of view, OR
provides a simple and precise way to achieve a desired particle
size through control of the duration of the nanocrystal growth
and the proper choice of capping agents passivating the surface
of the growing nanocrystals. However, it ultimately results in
rather broad particle size distributions with standard deviations
lying in the range of 15±21 %. To prepare monodisperse colloidal nanocrystals in this way, additional post-preparative size
fractionation is required. These theoretical predictions correlate well with experimental data obtained both on nanometerand micrometer-sized particles of various materials.[18,35,42]
2.2. Transient Regime of Particle Growth: ªFocusingº
and ªDefocusingº of Particle Size Distribution
Fig. 3. Diffusion-controlled Ostwald ripening in the nanoparticle ensemble.
a) Evolution of the mean particle radius and particle concentration with time.
b) Dependence of the standard deviation of the particle size distribution on the
average particle size during the Ostwald ripening for different initial widths of
particle size distributions. c) Dependence of the stationary value of standard
deviation of the size distribution on the surface tension. Dashed lines in all
frames correspond to LSW predicted coarsening.
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As mentioned in the previous section, the OR growth regime
results in the particle size distribution having a width of at least
15 %. Narrower size distributions can be achieved if the growth
of particles occurs in the transient growth regime before OR
takes place, and is terminated before the equilibrium between
particles and monomer is attained. In this transient regime a
considerable amount of precursor is still present in the colloi-
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dal solution in the form of monomer, as only a small amount of
the precursor was converted to nanocrystals at the nucleation
stage. In contrast, during OR the particles are in equilibrium
with the monomer in solution and the chemical potential of the
monomer corresponds to that of particles lying somewhere in
the middle of the size distribution. The corresponding equilibrium oversaturation of monomer (S) can be estimated from
Equation 1 and for realistic experimental conditions falls into
the range of ~ 1.2 to 20. If the value of S is much above this
range, all particles in the colloidal solution will grow simultaneously by consuming the monomer from the solution. In this
case, a fast increase of the particle size accompanied by a
strong narrowing (further referred to as ªfocusingº) of the size
distribution is observed. Figure 4a shows simulations of the
temporal evolution of the particle size distribution during dif-
required for the nucleation event and the equilibrium monomer oversaturation observed in the OR regime. As a result, at
early stages of the synthesis the nucleation rate drops much
faster than the oversaturation decreases and a large amount of
precursor remains in the form of monomer. On the other hand,
conditions favorable for focusing of the size distribution can be
achieved by applying special experimental techniques. Thus,
swift injection of cold reagents into very hot solvent results in
an explosive-like nucleation, which is immediately suppressed
due to the fast temperature drop, thus allowing the achievement of high monomer oversaturation at early stages of nanocrystal growth (Fig. 4b). Similar results can be obtained by
additional injections of precursors to the solution of growing
nanocrystals:[28] an instantaneous increase of the oversaturation of monomer results in focusing of the particle size distribution (Fig. 4b). By applying these approaches and optimizing
the reaction conditions it is possible to synthesize nanocrystals
of desired size with very narrow size distributions as will be
demonstrated below.
2.3. Experimental Remarks
Fig. 4. a) Temporal evolution of the size distribution of the nanoparticle ensemble at high initial oversaturation of monomer (S0 = 900) under diffusion control.
b) Dependence of the standard deviation of particle size distribution on the average particle size for different growth regimes: Ostwald ripening (S0 = 15, upper
curve); focusing and defocusing of the size distribution at high initial oversaturation of monomer (S0 = 500, middle curve); focusing and defocusing of the size distribution induced by additional injection of monomer (S0 = 15, additional injection of 100 % of initial amount of the precursors at the instant of time indicated
by arrow, lower curve). All curves correspond to diffusion-limited particle
growth.
fusion-controlled nanoparticle growth at high initial monomer
oversaturation. Very fast initial growth resulting in focusing of
the size distribution from 20 % standard deviation (initial
value) down to 6 % is followed by slow broadening (further
referred to as ªdefocusingº) of the size distribution. During the
focusing stage the number of particles remains almost constant,
the size distribution is nearly symmetric and is well fitted by a
normal distribution. The oversaturation drops down to some
equilibrium value. Defocusing is accompanied with a transition
from a symmetric towards an asymmetric negatively skived size
distribution whereas the mean particle size remains nearly constant as shown in Figure 4b. The number of nanocrystals starts
to decrease due to the dissolution of the smallest ones. The
Ostwald ripening mechanism governs the further evolution of
the nanoparticle ensemble.
The situation discussed above occurs in the reaction vessel if
the difference is large between the monomer oversaturation
Adv. Funct. Mater. 2002, 12, No. 10, October
The monodisperse nanocrystals of semiconductor and magnetic compounds this paper is dealing with have been prepared
using a general synthetic strategy that is based on the high-temperature thermolysis of organometallic precursors and/or
reduction of metal salts in so-called coordinating solvents. The
hot injection technique was applied to separate the nucleation
regime from particle growth in the regime of strong focusing of
the size distribution. The time of heating at a desired temperature and the proper choice of stabilizing agents provide control
over the particle size. Stabilizing agents bind to the surface of
growing nanocrystals preventing their growth to the bulk
phase. They also do not allow the particles to coagulate and
provide desired surface properties, such as solubility in a
desired solvent. Moreover, stabilizers protect the nanocrystals
from oxidation at ambient conditions and provide electronic
passivation of surface traps in the case of luminescent particles.
Several stabilizing mixtures are known to date fulfilling all the
requirements mentioned above.
CdSe nanocrystals used in this work were synthesized either
in a tri-n-octylphosphine oxide-tri-n-octylphosphine (TOPOTOP)[16] or in a hexadecylamine (HDA)-TOPO-TOP mixture.[23] The temporal evolution of UV-vis spectra displaying the
growth of CdSe nanocrystals in these mixtures at 300 C is shown
in Figure 5a. The introduction of an additional coordinating
component, HDA, into the conventional organometallic synthesis of CdSe nanocrystals in the TOPO-TOP mixture provides
much better control over the growth dynamics resulting in focusing of the size distribution during particle growth.[23] The absorption spectra possess up to five resolved electronic transitions indicating very narrow in-situ size distributions of the CdSe
nanocrystals grown in HDA-TOPO-TOP. Narrow size distributions are also evidenced from a size histogram measured from
high-resolution transmission electron microscopy (HRTEM)
images on a sample taken from the reaction flask without any
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FEATURE ARTICLE
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FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
Fig. 5. a) Room-temperature absorption spectra of CdSe nanocrystals monitored
during their growth in HDA-TOPO-TOP and TOPO-TOP stabilizing mixtures
at 300 C. b) Size histogram of CdSe nanocrystals taken as prepared from HDATOPO-TOP mixture with average size of ~ 4.05 ± 4 % as calculated from
HRTEM images for over 200 particles. c) The relative particle size distribution of
growing CdSe nanocrystals: (*) and (j) during continuous growth at 300 C in
TOPO-TOP and HDA-TOPO-TOP, correspondingly; (~) obtained after stepwise growth in HDA-TOPO-TOP (1 h 250 C + 1 h 280 C) and (1 h 250 C + 1 h
280 C + 1 h 310 C); (!) obtained after slow additional injection of the precursors at 280 C to the stepwise grown nanocrystals in HDA-TOPO-TOP. The relative values of standard deviation (r) normalized to a starting value (r0) observed
immediately after injection of the stock solution were used. In both stabilizing
mixtures the value of r0 was ~ 20 %.
post-preparative narrowing of the size distribution (Fig. 5b).
Moreover, our investigations as well as the recent data of Peng
et al.[31] suggest that the control over the size distribution of
CdSe nanocrystals can be even further improved by introducing
an acidic component, i.e., a long chain n-alkylphosphonic or carboxylic acid, to the HDA-TOPO-TOP mixture. Figure 6a shows
a TEM image of exceptionally monodisperse, as-prepared CdSe
nanocrystals grown in a mixture containing HDA-TOPO-TOP
and n-tetradecyl phosphonic acid (TDPA).
The evolution of the particle size distribution[53] during the
growth of the CdSe nanocrystals in either TOPO-TOP or
HDA-TOPO-TOP stabilizing mixtures at 300 C is shown in
Figure 5c. Growth in TOPO-TOP was firstly accompanied by
focusing of the particle size distribution followed by defocusing
in accord with the theoretical predictions of the previous section. In contrast, in HDA-TOPO-TOP very fast focusing and
no defocusing were observed during the particle growth in the
course of long-term heating at 300 C resulting in a very narrow
particle size distribution. The observed permanent focusing of
the size distribution might be a result of the large difference
between the rate of nanocrystal growth in the focusing regime
and the rate of subsequent defocusing. The growth of nanocrystals almost terminates after the focusing stage and their size
distribution remains narrow for a long time. The absence of
defocusing means that the narrowest size distribution is
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observed at the largest size achievable at a given growth temperature. As a result, stepwise growth (1 h at 250 C + 1 h at
280 C + 1 h at 310 C) reproducibly allowed to reach size distributions with a standard deviation below ~ 5±6 %. Slow (one
drop per 10±30 s) additional injection of precursors at 280 C
into a solution of CdSe nanocrystals primarily grown at 310 C
resulted in further narrowing of the size distribution, being in
accord with the theoretical predictions of Section 2.2.
Synthesis of monodisperse nanocrystals of magnetic alloys
has some specific peculiarities comparing to the synthesis of
semiconductor nanocrystals. Thus, realization of Ostwald
ripening in this case is still under discussion.[15,54] Long-term
heating of the nanocrystals in the reaction mixture did not
result in a detectable change of the average particle size. Probably, no mass transfer between growing particles takes place
under the reaction conditions used, which means, in other
words, that the addition of matter from solution to the particle
surface is irreversible. In this case, particle growth is possible
only at the expense of precursors available in the reaction mixture, and thus proceeds in the focusing mode. On the other
hand, the injection temperature of precursors plays an important role. Figure 6b shows a TEM image of as-prepared magnetic CoPt3 nanocrystals synthesized in the HDA-diphenyl
a
b
Fig. 6. TEM images of as-prepared a) CdSe nanocrystals grown in TDPA-HDATOPO-TOP mixture and b) CoPt3 nanocrystals grown in the HDA-diethyl ether
coordinating mixture in the presence of 1-adamantancarboxylic acid.
ether coordinating mixture in the presence of 1-adamantancarboxylic acid as the stabilizer.[27] The standard deviation of the
size distribution of as-prepared CoPt3 nanocrystals obtained by
injection of precursors at 170 C is ~ 9 % (Fig. 6b), whereas at
100 C it is ~ 12 % (not shown). Both values are still considerably smaller than the values inherent to the nano-OR growth
regime (Sec. 2.1).
2.4. Post-Preparative Narrowing of Particle Size Distribution
Direct synthesis of highly monodisperse nanocrystals is a goal
that cannot always be achieved in the framework of existing
experimental approaches. For instance, to the best of our knowledge, no recipes allowing the preparation of III±V semiconductor nanocrystals with size distributions better than 15±20 % (asprepared) are known to date. In general, the self-assembly of
nanocrystals in 2D and 3D superlattices requires extremely nar-
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row particle size distributions, which often can not be achieved
in situ during the particle growth. Fortunately, the nanocrystal
size distribution can be improved through post-preparative size
fractionation, which is based on the size-dependent variation of
particle properties. Thus, a size-selective precipitation technique
exploits the difference in solubility of smaller and larger particles.[16,55] A typical example of carrying out the size-selective
precipitation on a nanoparticle colloid is as follows. A sample of
as-prepared nanocrystals with a broad size distribution is dispersed in a solvent and a non-solvent is added dropwise under
stirring until the initially optically clear solution becomes
slightly turbid. The largest nanocrystals in the sample exhibit
the greatest attractive van der Waals forces and tend to aggregate before the smaller particles. The aggregates consisting of
the largest nanocrystals can be isolated by centrifugation or
filtration and re-dissolved in any appropriate solvent. The next
portion of non-solvent is added to the supernatant to isolate the
second size-selected fraction, and so on. The procedure can be
repeated several times and allows to obtain up to ~ 20 size-selected fractions from one portion of the crude solution. Moreover, each size-selected fraction can be subjected again to size
selection to further narrow the size distribution.
The technique described above is robust and can be adopted
to different nanocrystals, both water-soluble crystals and those
soluble in organic solvents. The search for appropriate solvent/
non-solvent pairs has to be performed in each particular case.
Figure 7 demonstrates several examples of post-preparative
size fractionation for different II±VI and III±V semiconductor
nanocrystals. All size-selected fractions possess sharp excitonic
3. 2D and 3D Arrays of Monodisperse Nanocrystals
Formed by Self-Assembly
Self-organization is a process leading to spontaneous formation of ordered arrays from monodisperse nanocrystals. Dispersive attractions of nanoparticles caused by van der Waals forces
are responsible for that.[34,35] Both 2D and 3D arrays can be
prepared simply by placing a drop of a colloidal solution of
monodisperse nanocrystals on a suitable support and allowing
the carrier solvent to evaporate slowly. Figures 6 and 8 present
a gallery of TEM and HRTEM images of 2D arrays of CdSe,
a
b
c
d
Fig. 8. TEM images of 2D arrays of InAs (a) and FePt (b) nanocrystals. HRTEM
image of a 2D array of CdSe nanocrystals with random (c) and preferable (d) orientation of the lattice planes. Insets: corresponding FFTs.
Fig. 7. Room-temperature absorption spectra of size-selected fractions (thin
lines) obtained from crude solutions (thick lines) of various II±VI and III±V
nanocrystals by applying size-selective precipitation technique. Details of preparation and solvent/non-solvent pairs used: CdSe, TOPO-TOP synthesis [16],
n-butanol/ethanol; CdTe, aqueous synthesis [18], capped by thioglycolic acid,
water/isopropanol; InP, dehalosilylation reaction [20], capped with TOPO-TOP,
toluene/methanol; InAs, dehalosilylation reaction [19], capped with TOP, toluene/methanol.
transitions in the absorption spectra, which is a direct evidence
of their narrow particle size distributions. TEM and HRTEM
investigations show that carefully performed size-selective
precipitation allows to achieve size distributions as narrow as
~ 4±7 % depending on the material. Details of the nanocrystal
syntheses are given in the Experimental section.
Adv. Funct. Mater. 2002, 12, No. 10, October
InAs, CoPt3, and FePt nanocrystals. The samples were obtained by dropping dilute solutions of monodisperse nanocrystals onto carbon-coated copper TEM grids and evaporating the
solvent in a desiccator. A long-range hexagonal ordering is
observed in all cases, and the regular inter-particle spacing is
due to the capping ligand shells around the nanocrystals. The
spacing between nanocrystals can be adjusted by employing
different ligands at the synthesis stage or by post-preparative
cap exchange.[35]
The orientation of lattice planes of nearly spherical nanocrystals on the TEM grids is random, as a rule (Fig. 8c). However, in some cases a preferred orientation is observed, e.g., for
larger (7 nm) facetted CdSe nanocrystals elongated along the
c-axis of the wurtzite structure (Fig. 8d). These nanocrystals
are aligned on the TEM grid with the c-axes perpendicular to
the substrate.[58] Importantly, there is even a correlation in orientation of the lattice plane directions of single nanocrystals,
which is further confirmed by six distinct reflexes in the corre-
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FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
spondent FFT pattern being characteristic for the (100) zone
(inset in Fig. 8d), in contrast to the isotropic FFT ring of randomly oriented CdSe particles (Fig. 8c).
The above TEM images demonstrate self-organization of
nanocrystals in 2D monolayers. With increasing surface coverage, a transition from 2D array to 3D arrangements can be
observed. Figure 9 shows this tendency for the case of CoPt3
nanocrystals. With increasing surface coverage, nanocrystals of
the second layer occupy positions ªin betweenº the nanocrys-
a
c
b
c
d
b
d
Fig. 9. TEM images of regular arrangements of CoPt3 nanocrystals illustrating
the transition to the 3D structure. a,b) Two layers and c,d) three layers of monodisperse nanoparticles at different magnifications.
tals in the first layer thus being placed over empty inter-particle
spaces maintained by bulky capping ligands (Fig. 9a,b). With
further increase of the surface coverage the third layer of nanocrystals is formed (Fig. 9c,d). The difference in contrast
between two ground layers and a darker third layer allows the
attribution of each nanocrystal to the layer it is placed in and
leads to the conclusion that CoPt3 nanocrystals are packed in a
cubic close-packed (ccp)-like superlattice, where the nanocrystals are separated from each other by relatively thick (2.5 nm)
organic shells.
The ordering of nanocrystals in superstructures is also reflected in the small-angle XRD patterns[35] through the appearance of Bragg diffraction peaks in the region of 2H angles of
~ 1±15. The diffraction signal is averaged over a large number
of nanocrystals in this case, whereas the TEM data provides
information from selected areas only. Figure 10 shows XRD
patterns measured on different superstructures of CdSe and
FePt nanocrystals. In the case of monodisperse CdSe nanocrystals (see Figs. 6 and 8 for TEM images) fast evaporation of a
low-boiling solvent (hexane) results in the formation of a
glassy-like film with a local-range particle order and a liquidlike radial distribution function (Fig. 10a, curve 1). On the
other hand, CdSe nanocrystals precipitated via slow evapora-
660
a
Ó 2002 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
Fig. 10. Small-angle (a,c) and wide-angle (b,d) powder XRD patterns of CdSe
and FePt nanocrystals. a) Small-angle XRD patterns of glassy (curve 1) and
long-range ordered (curve 2) films of CdSe nanocrystals. b) Wide-angle XRD
patterns of a size series of CdSe nanocrystals. Vertical lines indicate bulk CdSe
reflections (top: wurtzite, hexagonal; bottom: zinc blende, cubic). c) Small-angle
XRD patterns of randomly oriented (curve 1) and packed in a layered superstructure (curve 2) FePt nanocrystals. d) Wide-angle XRD pattern of FePt nanocrystals showing enhanced (200) fcc reflex in the long-range ordered sample
(curve 2) in comparison with the glassy film (curve 1). Vertical lines indicate bulk
reflections of the fcc FePt phase.
tion of a higher-boiling solvent (e.g., toluene) form films with a
long-range particle order exhibiting pronounced reflexes in the
small-angle XRD pattern (Fig. 10a, curve 2). Most of the
reflexes can be attributed to the face-centered cubic (fcc) lattice of CdSe nanocrystals.
The wide-angle part of the XRD patterns corresponds to the
diffraction of X-rays on atoms that the nanocrystals consist of,
and allows the estimation of the average size of the crystalline
domains within each nanocrystal. The width of the diffraction
peaks at wide angles is considerably broadened and increases
with decreasing particle size (Fig. 10b). CdSe nanocrystals with
sizes above ~ 4 nm exhibit XRD patterns with diffraction peaks
in accord with those of hexagonal CdSe (wurtzite phase). In
the case of smaller CdSe nanocrystals the XRD patterns do not
permit to distinguish between the cubic and the hexagonal
phases unambiguously.
The preferred orientation of nanocrystals within superlattices is evidenced from the XRD patterns of FePt particles. Figure 10c, curve 2 shows the small-angle XRD pattern of a film
of FePt nanocrystals possessing six equidistant reflexes evidencing its layered structure.[59] A preferential orientation of
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the nanocrystals inside the film is confirmed by a strong enhancement of the (200) fcc reflection in the wide-angle region
of this sample (Fig. 10d, curve 2). This enhancement is not observed in a sample of the glassy FePt particle film (Fig. 10c,
curve 1) whose diffraction pattern is consistent with randomly
oriented nanocrystals of the fcc FePt phase (Fig. 10d, curve 1).
An interesting case is an assembly of nanocrystals with bimodal size distributions. Mixtures of Au[60] or Au and Ag[61] nanoparticles of two different sizes have recently been reported to
form nanoscale colloidal alloy superlattices. Both AB and AB2
alloy phases have been shown to be dependent on the size and
the local particle number ratios. The experimental data obtained by Schiffrin et al.[60,61] for the mixed AB and AB2 phases
of bimodal ensembles of nanoparticles agreed well with the
previously derived geometrical rules for micrometer-size particles[62] governing the formation of either AB (0.27 < RA/
RB < 0.425) or AB2 (0.482 < RA/RB < 0.624) phases from monodisperse spheres of radii RA and RB. When two monodisperse
colloids of CoPt3 nanocrystals (4.5 nm and 2.6 nm diameter)
were mixed together followed by slow evaporation of the
solvent, an AB5-type superlattice analogous to the structure of
intermetallic compound CaCu5[63] was obtained (Fig. 11a). A
similar structure was observed for binary mixtures of latex
spheres of two different sizes.[64] In the first plane of this lattice
(Fig. 11), each 4.5 nm CoPt3 nanocrystal is surrounded with a
hexagon formed by 2.6 nm nanocrystals. The second plane consists only of hexagons of small particles and the third plane
repeats the first one.
Figures 12a and b present TEM images of 3D arrangements
of CoPt3 and CdSe nanocrystals. A close packing of multiple
layers of nanocrystals exhibiting long-range order is observed.
Figures 12c and d provide a closer look at the top layer of the
3D arrays demonstrating two different arrangements of CdSe
nanocrystals whose FFT patterns (given as insets) are consistent with those expected for the (100) and (110) projection of a
fcc superlattice, respectively.
4. Colloidal Crystals of Nanoparticles Grown
by the Three-Layer Oversaturation Technique
in Solution
Another strategy to fabricate ordered superstructures from
monodisperse nanoparticles is a gentle destabilization of the
colloidal dispersion.[35] Again, dispersive attractions of nanoparticles drive their self-organization and the superlattice for-
a
b
a
b
c
d
FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
Fig. 12. TEM images of 3D arrangements of CoPt3 (a) and CdSe (b) nanocrystals, and HRTEM images of (100) (c) and (110) (d) projections along the CdSe
superlattice with corresponding FFTs.
mation. In superstructures formed in this way, individual nanocrystals playing the role of building blocks (artificial atoms in
the next level of hierarchy) are aligned in a regular 3D lattice
(ªartificial solidº). Natural and artificial opals are examples of
such ordered superlattices on the micrometer scale,[65,66] and
colloidal crystals made of SiO2 or latex microspheres are currently attracting a lot of attention as photonic bandgap materials.[67] There are only a few examples of colloidal crystals created from nanoparticles known up to date.[68±74] Very recently,
we have proposed a simple three-layer technique of controlled
oversaturation leading to the crystallization of monodisperse
CdSe[73] and FePt[74] nanocrystals into perfectly faceted colloidal crystals with sizes of 10±200 lm. This method also works
well in the case of CoPt3 nanocrystals[27] indicating its wide or
even general applicability.
Figure 13a presents a schematic outline illustrating the concept of the crystallization procedure. A colloidal solution of
nanocrystals in a solvent like toluene is placed in a vertically
positioned glass tube, and the system is slowly destabilized by
diffusion of a non-solvent (e.g., methanol) into the colloid
resulting in nucleation and growth of colloidal crystals, prefer-
c
Fig. 11. a) Schematic view of an intermetallic compound
CaCu5. b,c) TEM images illustrating AB5-type superstructure formed by CoPt3 nanocrystals with bimodal size distribution.
Adv. Funct. Mater. 2002, 12, No. 10, October
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FEATURE ARTICLE
A. L. Rogach et al./Monodisperse Nanocrystals and Their Superstructures
Fig. 13. a) Schematic outline illustrating the concept of the oversaturation technique used for crystallization of monodisperse nanocrystals. Left tube: the nonsolvent (methanol) diffuses directly into the colloidal solution of CdSe nanocrystals in toluene. Right tube: buffer layer of propanol-2 is used in between to obtain
colloidal crystals of higher quality. b,c) Optical micrographs of colloidal crystals
consisting of CdSe nanocrystals made by a digital camera through an objective of
an optical microscope. Reprinted from [73] with permission.
entially on the walls of the tubes. The spatial distribution of
local oversaturations caused by the non-solvent diffusion determines the quality of the colloidal crystals. To make the oversaturation front not as sharp as in the case of a direct solvent±nonsolvent contact, a third buffer layer of propanol-2 is used between the solution of the nanocrystals and the methanol layer.
The method has been successfully applied for the crystallization of CdSe, FePt, and CoPt3 nanocrystals. For the last two
materials, a three-layer variation of the crystallization technique was used as a more advanced one. Crystalline nuclei
started to form after about one to two weeks, and grew slowly
to colloidal crystals over 1±2 months. Figures 13b and c show
optical micrographs of colloidal crystals of CdSe nanocrystals.
The images were taken with a digital camera through an objective of an optical microscope. Irregular-shaped red-colored colloidal crystals of CdSe nanoparticles formed in the absence of
a buffer layer reach 80±220 lm in size (Fig. 13b). In the presence of a buffer layer, colloidal crystals grew in the form of perfectly faceted hexagonal orange±red colored platelets that
were very similar in size, about 100 lm in lateral dimension
and 20 lm in depth (Fig. 13c).
Thicker black- and thinner brownish-colored colloidal crystals of FePt and CoPt3 nanoparticles grew preferentially in the
form of faceted triangular or hexagonal platelets with 10±
30 lm long edges. Figure 14 presents a gallery of scanning electron microscopy (SEM) images of CoPt3 colloidal crystals pro-
Fig. 14. SEM images of colloidal crystals of CoPt3 nanocrystals illustrating their
different shapes.
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Ó 2002 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
viding a closer look on their morphology. They preferably form
faceted triangular and hexagonal platelets, transition forms
between these two shapes, and tetrahedral crystals. Figure 15a
shows a TEM image of a representative fragment of FePt colloidal crystals obtained by their mechanical grinding and treatment in an ultrasonic bath. Hexagonal arrangements of FePt
nanocrystals, i.e., building blocks of the colloidal crystals, are
clearly seen at the edges of the crystalline pieces. An HRSEM
image taken from the surface of a FePt colloidal crystal further
confirms their high long-range quality showing the perfect hexagonal arrangement of nanocrystal building blocks (Fig. 15b).
a
b
Fig. 15. a) TEM image of the fragment of a FePt colloidal crystal. b) HRSEM
image of the surface of a FePt colloidal crystal. FePt nanocrystalsÐbuilding
blocks of the colloidal crystalsÐare clearly seen in both cases.
5. Conclusions
A large diversity of nano-scale materials can be obtained using
the methods of colloidal chemistry. The advanced syntheses
enable the preparation of semiconductor, metal, and magnetic
alloy nanocrystals on a gram scale, which can be further handled
like ordinary chemical substances. The achievement of high
monodispersity of nanoparticles during the synthetic stage is a
challenge requiring a lot of experimental work and can still be
considered a kind of art. Further progress in this direction is
associated with a better understanding of the processes governing the growth of nanocrystals in colloidal solutions. Based on a
theoretical description of Ostwald ripening in an ensemble of
particles in a colloidal solution, we have shown some peculiarities of this process, which are inherent to nanocrystals, and figured out the synthetic regimes when the narrowing (ªfocusingº)
of the size distribution of growing particles takes place. The proper choice of the reaction conditions and coordinating solvents
allows the preparation of exceptionally monodisperse samples as
demonstrated for CdSe nanocrystals grown in HDA-TOPOTOP. Colloidally synthesized InAs, FePt, and CoPt3 nanocrystals
are introduced as well, and the principles of the post-preparative
size-selective fractionation leading to monodisperse nanocrystals
are discussed. The monodisperse nanocrystals readily organize
themselves in different 2D and 3D superstructures, some of
which are presented in this article. A recently developed threelayer technique of controlled oversaturation reproducibly allows
the crystallization of nanocrystals into 3D colloidal crystals. In
these artificial solids, individual nanoparticles play the role of
building blocksÐartificial atoms in the next level of hierarchy.
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6. Experimental
CdSe Nanocrystals: In the conventional TOPO-TOP synthesis, the stock solution prepared by mixing of 4.1 mL of distilled TOP, 0.1 mL of dimethylcadmium,
and 1.0 mL of 1 M solution of TOPSe [56] in TOP was quickly injected into a
vigorously stirred TOPO (10 g) heated to 360 C. In the HDA-TOPO-TOP synthesis, 1 mmol of TOPSe and 1.35 mmol of dimethylcadmium were dissolved in
5 mL of TOP and rapidly injected into a vigorously stirred mixture of 10 g of
TOPO and 5 g of HDA heated to 300 C. Further growth of CdSe nanocrystals
was carried out at 300 C in both cases and terminated by cooling the reaction
mixture. To precipitate TOPO-TOP capped CdSe nanocrystals size-selectively,
40 mL of 1-butanol were added to 4 mL of the crude CdSe solution. The solution
was filtrated and methanol was slowly added. The precipitated nanocrystals were
isolated by filtration, washed thoroughly with methanol, and re-dissolved in toluene. In the case of HDA-TOPO-TOP capped nanocrystals, no size-selective fractionation was required; the particles were precipitated from the crude solution
with ethanol and redissolved in toluene.
InAs Nanocrystals: These nanocrystals were synthesized by the dehalosilylation reaction between InCl3 and tris(trimethylsilyl)arsine (TMS3As) [19,57]. In
a typical preparation route carried out in inert atmosphere, 0.30 g of InCl3 were
dissolved in 1 mL of TOP, mixed with 0.26 mL of TMS3As, and rapidly injected
into 3.2 g of TOP vigorously stirred at 300 C. Further growth occurred at 260 C
for different periods of time depending on the desired size of the nanocrystals. To
prepare InAs particles larger than ~ 3.5 nm in size, additional injections of precursors were necessary. Aliquots of the crude solution of InAs nanocrystals were
taken from the hot reaction mixture, cooled to room temperature and mixed with
a ~ 10-fold excess of toluene. Nanocrystal fractions were precipitated with ethanol and redissolved in toluene.
FePt Nanocrystals: The synthesis of FePt nanocrystals was done following the
ªpolyolº approach of Sun et al. [14]. Under airless conditions, platinum acetylacetonate (0.25 mmol) and 1,2-hexadecanediol (0.75 mmol) were dissolved in
10 mL of dioctylether and heated to 100 C. Oleic acid (0.25 mmol), oleyl amine
(0.25 mmol), and Fe(CO)5 (0.5 mmol) were added to the mixture, which was subsequently heated for 30 min at 300 C and cooled to room temperature. Postpreparative size-selective precipitation was carried out in air using hexane and
ethanol as a solvent and non-solvent, respectively. The precipitates were redissolved in toluene.
CoPt3 Nanocrystals: A modified ªpolyolº approach with employment of
1-adamantancarboxylic acid as a stabilizer was used [27]. Under airless conditions, 0.0328 g of Pt(acac)2, 0.13 g of 1,2-hexadecanediol, and 0.084 g of 1-adamantancarboxylic acid were dissolved in a mixture of diphenyl ether (2.0 mL)
and HDA (4.0 g) and heated to 65 C until a clear solution was formed. The
reaction mixture was heated to 170 C and cobalt stock solution prepared by
dissolving 0.043 g of Co2(CO)8 in 0.4 mL of 1,2-dichlorobenzene was injected
under vigorous stirring, followed by refluxing for 40 min. No post-preparative
size-selective precipitation was carried out for the samples described in this
paper.
Received: March 7, 2002
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