Controlled Synthesis: Nucleation and Growth in Solution

Abstract

The controlled synthesis of metallic nanomaterials in solution is central to realize many applications that arise from their fascinating properties. As properties in metal nanomaterials are strongly dependent upon size, shape, composition, structure (solid versus hollow interiors), and surface functionality, controlled synthesis is a powerful approach to tailor and optimize properties as well as to establish how they are dependent on the several physical and chemical parameters that define a nanostructure. In this context, this chapter focuses on the fundamentals of the controlled synthesis of metal nanomaterials in solution phase in terms of the available theoretical framework. Specifically, it starts by introducing the mechanisms employed for the stabilization of nanomaterials during solution-phase synthesis (Sect. 2.2). The basics of nucleation and growth in solution will be discussed in Sect. 2.3. After that, the shape-controlled synthesis of Ag nanomaterials will be employed as proof-of-concept example of how thermodynamic versus kinetic considerations, oxidative etching, and surface capping can be employed to effectively maneuver the shape of a metal nanocrystal in solution (Sect. 2.4). Finally, some of the current challenges and outlook regarding the controlled synthesis of metal-based nanomaterials will be presented (Sect. 2.5).

2.1 Motivation for Controlled Synthesis of Metallic Nanomaterials

The motivation for the controlled synthesis of metallic nanomaterials comes from the fact that their properties are strongly dependent upon size, shape, composition, structure (solid versus hollow interiors), and surface functionality [1]. Thus, the precise control over these parameters opens up the possibility of designing nanomaterials with desired or optimized performances for a given application. Taking catalysis as an example, in addition to small sizes (that are imperative to achieve high-surface areas), the control over the shape enables the exposure of specific surface facets that can be more catalytically active or selective for a reaction of interest [2]. The shape-dependent catalytic activity of Pd nanocrystals towards the electroxidation of formic acid is shown in Fig. 2.1 and represents a notable example on how the control over shape can be employed for optimizing catalytic performance. As {100} surface facets are more active relative to {111} for the electroxidation of formic acid, a gradual increase in the current densities was observed as the shape was changed from cubes to truncated cubes, cuboctrahedrons, truncated octahedrons, and finally octahedrons, as these shapes gradually enable an increased exposure of {111} relative to {100} surface facets [3]. In addition to catalytic activity, optical, magnetic, and electronic properties have also shown shape/size-dependency [4].

Fig. 2.1

Maximum current densities for formic acid oxidation employing Pd nanocrystals displaying controlled shapes as electrocatalysts. The current densities (electrocatalytic activities) were strongly dependent on the relative exposure of {100} and {111} surface facets, decreasing as the shapes transitioned from cubes to octahedrons [3]. Copyright 2012 American Chemical Society

It is noteworthy that, in order to reach its full potential, the controlled synthesis of metallic nanomaterials must meet several requirements that include monodispersity in terms of size, shape, structure, and composition. Among different methods for the controlled synthesis of metallic nanomaterials, solution-phase approaches offer many advantages that include the facile stabilization by the addition of proper capping agents, easy extraction/separation from the reaction mixture (by centrifugation, for example), straightforward surface modification/functionalization, potential for large-scale production, and versatility regarding the several experimental parameters that can be controlled during the synthesis (temperature, nature of precursors, stabilizers, solvent, reducing agents, their molar ratios, concentrations, etc.) [5]. Surprisingly, substantial developments on the solution-phase-controlled synthesis of metallic nanomaterials with a variety of shapes and sizes were made only recently. Figure 2.2 illustrates a variety of shapes that have been realized by the controlled synthesis of metal nanostructures [1]. Despite this progress, a detailed atomic understanding of the mechanisms of nucleation and growth stages during nanoparticle formation is still not completely understood, and most fundamentals are currently borrowed from classical colloidal theories [6].

Fig. 2.2

Different shapes that have been enabled for a range of metals nanocrystals by solution-phase synthesis

In this chapter, rather than describing a variety of protocols for the controlled synthesis of different classes of metal nanomaterials, we will focus on the main fundamentals and concepts behind the controlled synthesis of metal nanomaterials in solution phase. Specifically, we will start by introducing the mechanisms employed for the stabilization of nanomaterials in the context of solution-phase synthesis (Sect. 2.2). The fundamentals of nucleation and growth in solution will be discussed in Sect. 2.3. After that, the shape-controlled synthesis of Ag nanomaterials will be employed as proof-of-concept example of how thermodynamic versus kinetic considerations, oxidative etching, and surface capping can be employed to effectively maneuver the shape of a metal nanocrystal in solution (Sect. 2.4). Finally, we will present some of the current challenges and outlook regarding the controlled synthesis of metal-based nanomaterials (Sect. 2.5).

2.2 Stabilization in Solution-Phase Synthesis

The main feature that sets nanomaterials apart from their micro- and macro-counterparts is their large surface to volume ratios, which is responsible for many unique or improved properties that nanomaterials possess [7]. A large surface to volume ratio implies that nanomaterials display high-surface energies (as surface energy increases with surface area), which makes them thermodynamically unstable or metastable [8]. Consequently, the synthesis of nanomaterials in solution requires one to overcome their large surface energies and prevent them from agglomerating. Agglomeration is thermodynamically driven by the reduction in surface area and thus surface energy. In fact, the reduction of surface energy is also the driving force for other processes in solution-phase syntheses such as surface restructuring, formation of facets, and Ostwald ripening [5]. In the following section, we will discuss two main approaches that are commonly employed to stabilize metallic nanomaterials against agglomeration: the electrostatic and steric (or polymeric) stabilization. Interestingly, as will be discussed in Sect. 3 and 4, the steric stabilization can also be put to work for favoring the formation of nanomaterials with monodisperse sizes and controlled shapes.

2.2.1 Electrostatic Stabilization

The concept behind the electrostatic stabilization is illustrated in Fig. 2.3 and it is based on the repulsion of electrical charges having the same sign present in the electrical double layers between neighboring particles [9]. The electrostatic stabilization can also be understood in terms of osmotic flow, in which solvent from the suspension flows into the region where the double layers overlap until the distance between the nanoparticles equals to or become larger than the sum of their individual double layers. As the overlap of double layers lead to a substantial increase in the concentration of ions in the inter-particle region, the osmotic flow works to restore the equilibrium concentration.

Fig. 2.3

Electrostatic stabilization between two approaching particles: the electrostatic repulsion of like-charges drive particles apart as their double layers overlap

The formation of a double layer arises from the adsorption of charged species at the surface, dissociation of surface-charged species, and/or accumulation or depletion of electrons at the surface following the formation of a solid surface in a polar solvent or electrolyte. An equal number of counter-ions with the opposite charge will then surround the system to produce a charge-neutral double layer, as depicted in Fig. 2.3. In addition to the contribution from columbic forces, Brownian motion and entropy also contribute to the location of species in solution, resulting in an inhomogeneous distribution of surface ions and counter ions [10]. The DLVO theory, which was named after Derjaguin, Landau, Verwey and Overbeek, successfully describes how the electrostatic stabilization between neighboring nanoparticles in suspension works [11]. The DLVO theory assumes that the interaction between two electrically charged approaching particles in suspension consists of a combination of Van der Waals attraction and electrostatic repulsion arising from the neighboring double layers [12, 13]. Figure 2.4 shows a plot of the overall potential (V, solid line) illustrated as a sum of Van der Waals attraction (V_A, dashed line) and electrostatic repulsion (V_R, dashed line) as a function of the distance between the approaching particles. When the two particles are far away from each other (distance is large), both V_A and V_R tend to zero. As the particles start to approach each other, the electrostatic repulsion is stronger than the Van der Waals attraction (V_R > V_A), leading to an increase in the overall potential (V) with the decrease in distance until a maximum is reached at V_MAX, which corresponds to the repulsive energy barrier. As the distance is further decreased, the Van der Waals attraction surpasses the electrostatic repulsion (V_A > V_R), and the overall potential becomes strongly dominated by V_A, which leads to the agglomeration of the particles. If V_MAX> ~ 10 kT (k is the Boltzmann constant and T the temperature), collision from Brownian motion will not overcome V_MAX and agglomeration will not take place. On the other hand, if V_MAX< ~ 10 kT, agglomeration takes place and the particle suspension would not be stable.

Fig. 2.4

Overall potential energy (V, solid line) expressed as a combination of the attractive Van der Waals and repulsive electrostatic potentials (V_A and V_R, respectively, dashed lines) [9]. Copyright © 2002, John Wiley and Sons

It is important to note that these considerations indicate that electrostatic stabilization mechanism only applies to dilute systems, is dependent on the concentration of ions in suspension, agglomerated particles cannot be re-dispersed, and comprises a kinetic stabilization method. As an example, the widely employed synthesis of Au nanoparticles by the reduction of AuCl⁻ _4(aq) by trisodium citrate pioneered by Turchevich [14] and refined by Frens [15] is based on electrostatic stabilization, in which citrate anions (negative charge) bind to the surface of the Au NPs during the synthesis.

2.2.2 Steric (or Polymeric) Stabilization

The steric (also called polymeric) stabilization consists in preventing neighboring particles of getting close to each other in the range of attractive forces where Van der Waals attraction would lead to agglomeration by the presence of polymers at the particles surface [16]. In general, a polymer can interact with a solid surface by one-end binding (anchored polymer) or via weak interactions from random points along its backbone (adsorbing polymer) [9]. When two particles covered with polymer approach each other, the polymer layers from each particle will interact only when the distance between particles become lower than the sum of the thicknesses of the neighboring polymer layers. Figure 2.5 depicts two approaching nanoparticles presenting an anchored polymer at their surface. As the distance between the particles becomes lower than the sum of the thicknesses of the neighboring polymer layers, they can overlap or become compressed resulting in strong repulsion between the neighboring nanoparticles. When the polymer is at good solvent conditions and neighboring adsorbed layers can overlap (under low-surface coverage), stabilization occurs as a result of the unfavorable mixing of the adsorbed layers, as an increase in entropy accompanied by an increase in the Gibbs free energy results from the interpenetration of polymers (assuming the enthalpy variations are negligible due to the interpenetration). At higher coverage (and/or under poor solvent conditions), the approaching adsorbed layers become compressed at the surface as the nanoparticles approach each other, and therefore, cannot interpenetrate. This process leads to an increase in the Gibbs free energy, driving the nanoparticles apart. Hence, steric stabilization emerges from the volume restriction effect (exclusion by space) due to the decrease in the possible configurations in the region between the particles and from the osmotic effect as the concentration of polymeric molecules in the region between the two particles becomes high [11].

Fig. 2.5

Steric stabilization between two approaching particles containing anchored polymers at their surface. Under both low- and high-surface coverage conditions, the polymer layers repel one another as they approach and/or interact

It is important to mention that the steric stabilization is not so well understood in comparison to the electrostatic stabilization. Nevertheless, it is widely employed in the solution-phase synthesis of nanomaterials, offering many advantages when compared with electrostatic stabilization. For instance, steric stabilization is thermodynamic (agglomerated particles can be redispersed), it is not restricted to dilute systems, and is not sensitive to the presence of electrolytes, which is highly desirable for several applications. Moreover, it contributes to the production of monodisperse nanoparticles as the polymer layer adsorbed at the nanoparticle surface can also work as a diffusion barrier to the addition of growth species, favoring the diffusion-limited growth (Sect. 3.2) [5]. Polymers employed for steric stabilization can also play other roles in solution-phase syntheses such as serving as reducing agents, helping in manipulating precursor reduction rates, and interacting with distinct nanocrystals surface facets with different binding affinities, which is central in the context of shape-controlled synthesis [1, 17].

2.3 Fundamentals of Nucleation and Growth

The controlled synthesis of metal nanomaterials requires (and refers to) nanomaterials with monodisperse sizes and shapes. Therefore, all nanostructures produced from a reaction mixture must have sizes and shapes as similar as possible (narrow size and shape distribution). The controlled synthesis of metal nanomaterials from a homogeneous solution comprises two main stages: nucleation and growth. In this section, we will discuss the theory behind nucleation and growth stages in order to understand how these steps can be employed to obtain nanomaterials with monodisperse sizes.

2.3.1 Homogeneous Nucleation

Nucleation refers to the extremely localized building of a distinct thermodynamic phase and represents the first stages during a crystallization process. More specifically, it can be defined as the process by which building blocks (metal atoms in the synthesis of metal nanomaterials) arrange themselves according to their crystalline structure to form a site upon which additional building blocks can deposit over and undergo subsequent growth [18]. In this context, nuclei correspond to infinitesimal clusters consisting of very few atoms of the growth species.

Lamer et al. pioneered the study on the synthesis of uniform colloids and the related nucleation and growth mechanisms [6]. These studies date back from the 1940s and are still the basis of our understanding on the synthesis of uniform nanomaterials (including metallic, polymeric, semiconducting, and oxide systems) [19]. From these studies, the concept of “burst nucleation” has been established to be crucial for the synthesis of monodispersed particles. Burst nucleation refers to the formation of a large number of nuclei in a short period of time, followed by growth without additional nucleation. This allows nuclei to have similar growth histories and, therefore, yield nanoparticles with same sizes. This concept is normally referred to as the separation of nucleation and growth stages during the synthesis [19]. For instance, if nucleation and growth could occur simultaneously, the nuclei growth histories would be significantly different and heterogeneous size distributions would be obtained.

The supersaturation of growth species comprises the first requirement for the occurrence of homogeneous nucleation. In the synthesis of metal nanoparticles, supersaturation may be achieved via the in situ generation of the solute growth species (metal atoms) through the conversion of soluble precursors into metal atoms by reduction or decomposition reactions. Alternatively, supersaturation could be induced by the decrease in the temperature of an equilibrium mixture containing the growth species. Although nucleation can take place in liquid, gas, and solid phases, we will center our discussion on the liquid phase, as the fundamentals are essentially the same in all cases.

The reduction of the overall Gibbs free energy to generate a new solid phase from a supersaturated solution containing the growth species as the solute represents the driving force for homogenous nucleation. The change in Gibbs free energy per unit volume of the solid phase (∆G_v) is dependent on the concentration of the solute and can be expressed as follows:

$$ \Delta {{\text{G}}_{\text{v}}}=- {\text{kT}}/{\ln{}}({{\text{C}}/{{\text{C}}_0}}) $$

(2.1)

where T is the temperature, Ω the atomic volume, C the concentration of the solute, and C₀ the solubility or equilibrium concentration [20]. When C > C₀, ∆G_v becomes negative and nucleation is spontaneously favored. However, this decrease in the Gibbs free energy is accompanied by an increase in the surface energy per unit area (γ) of the newly formed solid phase. Thus, the overall Gibbs free energy change (∆G) for the formation of spherical nuclei with radius r from a supersaturated solution corresponds to:

$$\Delta \text{G}=\left( 4\text{/}3 \right)\text{ }\!\!\pi\!\!\text{ }{{\text{r}}^{3}}\Delta {{\text{G}}_{\text{v}}}+4\text{ }\!\!\pi\!\!\text{ }{{\text{r}}^{2}}\text{ }\!\!\gamma\!\!\text{ }$$

(2.2)

where the first term corresponds to the change of volume free energy (volume term, ∆μ_v) and the second to the change in surface free energy (surface term, ∆μ_s). Figure 2.6 displays a plot of ∆G as a function of r (solid line) [20]. The surface and volume terms are also shown as dashed lines. As γ is always positive and ∆G_v negative as long as the solution is supersaturated, the plot of ∆G as a function of r has a maximum that corresponds to the critical radius r*. The critical radius (r*) can be defined as the minimum radius that a nucleus must possess in order to spontaneously grow from a supersaturated solution, or the minimum size required for the formation of a stable nucleus by precipitation from a supersaturated solution. When r < r*, the surface term is higher than the volume term and any nucleus smaller than r* will dissolve into the solution. In other words, the increase in the surface energy due to the nucleus formation is higher than the stabilization due to the formation of a new solid phase from the supersaturated solution, and nucleation is not favored. On the other hand, when r > r*, the volume term exceeds the contribution from the surface term so that the created nucleus is stable and undergoes subsequent growth (the stabilization due to the formation of the new solid phase from the supersaturated solution becomes higher than the destabilization due to the increase in surface energy). Setting d∆G/dr allows us to find r* and ∆G* according to the following equations:

Fig. 2.6

Change in the Gibbs free energy (G, solid line) as a function of the nucleus radius (r) as a sum of the volume and surface free energies [20]. Copyright © 2001, Elsevier

$$ \text{r}*=-2\text{ }\!\!\gamma\!\!\text{ }/\Delta {{\text{G}}_{\text{v}}}$$

(2.3)

$$\Delta \text{G}*=16\text{ }\!\!\pi\!\!\text{ }\!\!\gamma\!\!\text{ }/3{{\left( \Delta {{\text{G}}_{\text{v}}} \right)}^{2}}$$

(2.4)

where ∆G* corresponds to the critical Gibbs free energy, i.e., the minimum energy barrier that must be reached for nucleation to take place. These equations illustrate that r* can be decreased by lowering the γ of the new solid phase. Also, increasing ∆G_v will lead to a decrease in both r* and ∆G* (∆G_v increases with the supersaturation, Eq. 1). Surface energy, on the other hand, can be manipulated by changes in the solvent, temperature, and the use of “impurities” in the reaction mixture such as capping agents [19]. As an example, Fig. 2.7 shows how r* and ∆G* for a spherical nucleus can vary as a function of temperature. It can be observed as an increase in both r* and ∆G* with temperature. This is because supersaturation (and thus ∆G_v) decreases as the temperature is increased (Eq. 1) [5].

Fig. 2.7

Change in the Gibbs free energy (G) as a function of the nucleus radius (r) for two different temperatures, where T₁ < T₂ [20]. Copyright © 2001, Elsevier

The nucleation rate (N_R, per unit volume and unit time) is defined as the increase in the number of particles as a function of time and can be written as a function of ∆G* in the Arrhenius form:

$$ {\text{dN/dt}}= {{\text{N}}_{\text{R}}}= {\text{ exp}}({- \Delta G*/k{\text{T}}}) $$

((2.5))

N_R is also dependent on the number of growth species per unit volume (number of nucleation centers, which is determined by the initial concentration C₀) and the jump frequency of growth species from one site to another. Thus, N_R can be written as:

$$ {{\text{N}}_{\text{R}}}=\{({{\text{C}}_{0}}k\text{T}/(3\text{ }\!\!\pi\!\!\text{ }\!\!\eta\!\!\text{ }{{\text{d}}^{3}})\}\text{exp}\left( -\Delta \text{G}*/k\text{T} \right) $$

(2.6)

where η corresponds to the viscosity of the solution and d to the diameter of growth species. This equation shows that high nucleation rates are promoted by high initial concentration of growth species (high supersaturation), low viscosity, and low ∆G*. Moreover, in order to start the accumulation of nuclei for subsequent growth, N_R must be high enough to surpass the re-dissolution of nuclei. It is important to mention, however, that the thermodynamic model discussed above has some limitations when the synthesis of nanocrystals is regarded, as both γ and ∆G_v should not be constant and vary with size [19].

Owing to their sub-nanometer sizes, there is actually little information and knowledge regarding the true identity of nuclei during the synthesis of metal nanomaterials due to the lack of experimental tools capable of precisely capturing and characterizing them. According to theoretical investigations and evidence from mass spectrometry as well as adsorption and emission spectroscopies, it has been proposed that the crystal embryos seem to be closely related to their corresponding metal clusters [1, 21].

Figure 2.8 depicts the Lamer plot, which shows the change in the atomic concentration of the solute (growth species) during the nucleation and growth processes as a function of time. This plot is also very useful for illustrating the concept of burst nucleation. In the synthesis of metal nanocrystals, the concentration of solute (metal atoms) is increased by the decomposition or reduction of the corresponding soluble precursors. In the first stage, although the concentration of metal atoms increases, no nucleation/precipitation occurs even after the concentration of solute is above the equilibrium concentration (C_s). At this stage, although the solution is supersaturated, the energy barrier for nucleation (∆G*) is high due to the contribution from surface energy. Nucleation takes place only when the supersaturation reaches a specific value above the solubility (C_min = minimum solute concentration required for precipitation), which corresponds to the point where ∆G* is reached for the formation of nuclei. At this point, atoms start to aggregate into nuclei and homogeneous nucleation begins. As nucleation takes place, the concentration or supersaturation of metal atoms in solution starts to decrease, leading to a decrease in ∆G_v. When the concentration is below C_min (where ∆G* is reached), nucleation stops and growth starts. Growth then occurs under a continuous supply of metal atoms (solute) as long as their concentration is above the equilibrium solubility (C_s). In addition to growth via atomic addition and Ostwald ripening, the nuclei can also merge into larger objects via coalescence and oriented attachment [22]. The stabilization mechanisms discussed in Sect. 2 are essential to avoid extensive agglomeration that would hamper the production of nanoscaled materials. According to the Lamer plot, the concentration of the growth species should be increased abruptly to a very high supersaturation level followed by quickly bringing it below C_min in order to achieve burst nucleation [6].

Fig. 2.8

Variations in the atomic concentration of growth species in solution as a function of time during a the generation of atoms, b nucleation, and c growth stages [1]. Copyright © 2009 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

2.3.2 Growth

In addition to the “burst nucleation” concept and the subsequent separation of nucleation and growth during the synthesis, the size distribution of nanoparticles in solution-phase synthesis is dependent on the growth stage. More specifically, the size distribution of the initial nuclei may increase or decrease depending on the kinetics of growth [23]. The growth of nuclei can be subdivided in four major steps as illustrated in Fig. 2.9: generation of growth species (Fig. 2.9a), their diffusion (Fig. 2.9b), adsorption (Fig. 2.9c), and irreversible incorporation to the growth surface (Fig. 2.9d). Steps (a–c) can be further grouped and regarded as diffusion (supplying the growth species to the growth surface), while (d) corresponds to growth itself (by surface processes). Therefore, growth can be controlled either by diffusion or surface processes, and each of these regimes lead to differences in the final size distribution.

Fig. 2.9

Schematics of the main steps that occur during growth: generation of growth species (a), followed by their diffusion (b), adsorption (c), and irreversible incorporation (d) to the growth surface. Steps (a–c) comprise diffusion while (d) is regarded as growth itself

First, let us consider the case in which the growth is controlled by diffusion, i.e., controlled by the diffusion of growth species from the bulk of the solution to the growth surface. This situation is favored when diffusion is the slowest step during growth. In this condition, the growth rate (dr/dt) of a spherical nucleus of radius r depends on the flux of growth species to its surface (J) according to the equation:

$$ \text{dr/dt}=\text{J}{{\text{V}}_{\text{m}}}\text{/}4\text{ }\!\!\pi\!\!\text{ }{{\text{r}}^{2}} $$

(2.7)

where V_m is the molar volume of the nuclei. If the distance among nuclei is large enough, their growth can be considered independent from one another, and thus the diffusion layer formed at the surface of each growing particle will be undisturbed. The Fick’s law of diffusion expresses the flux of growing species (J) through the diffusion layer of each growing particle as:

$$ \text{J}=4\text{ }\!\!\pi\!\!\text{ }{{\text{x}}^{2}}\text{D}\left( \text{dC/dx} \right) $$

(2.8)

where D is the diffusion coefficient, C the concentration, and x the distance from the center of the growing particle. Assuming that J is constant with respect to x and integrating C(x) from x to x + δ, we get:

$$ \text{J}=4\text{ }\!\!\pi\!\!\text{ Dr}(\text{r}+\text{ }\!\!\delta\!\!\text{ })/\delta [\text{C}(\text{r}+\text{ }\!\!\delta\!\!\text{ }){{\text{C}}_{\text{s}}}] $$

(2.9)

where C_s is the concentration at the surface of the growing particle. At large δ values (δ > r), Eq. 9 becomes:

$$ \text{J}=4\text{ }\!\!\pi\!\!\text{ Dr}\left( {{\text{C}}_{\text{bulk}}}{{\text{C}}_{\text{s}}} \right) $$

(2.10)

where C_bulk is the concentration of the bulk solution [19]. Hence, from Eq. 7 and 10 we get the growth rate as:

$$ \text{dr/dt}={{\text{V}}_{\text{m}}}\text{D}\left( {{\text{C}}_{\text{bulk}}}{{\text{C}}_{\text{s}}} \right)/\text{r}$$

(2.11)

indicating that the growth rate of a particle is inversely proportional to its radius. Assuming the initial radius of the nuclei to be r₀ and the change in C_bulk to be negligible, this equation can be solved as:

$$ {{\text{r}}^{2}}={{\text{k}}_{\text{D}}}\text{t}+{{\text{r}}_{0}}^{2} $$

(2.12)

where k_D = 2DV_m(C−C_s). Thus, for two particles having an initial radius difference δr₀, the radius difference δr as the particles grow decreases with time according to Eq. 13 and 14, demonstrating that the growth limited by diffusion induces the formation of nanoparticles with monodisperse sizes [5].

$$ \text{ }\!\!\delta\!\!\text{ r}={{\text{r}}_{0}}\delta {{\text{r}}_{0}}/\text{r}$$

(2.13)

$$ \delta \text{r}={{\text{r}}_{0}}\delta {{\text{r}}_{0}}/{{\left( {{\text{k}}_{\text{D}}}\text{t}+{{\text{r}}_{0}}^{2} \right)}^{1/2}}$$

(2.14)

From the experimental point of view, growth limited by diffusion can be favored by: (i) keeping the concentration of growth species very low, so that the distance for diffusion from the solution to the growth surface would be large; (ii) increasing the viscosity of the solvent; (iii) controlling the supply of growth species to the growth surface; and (iv) introducing a diffusion barrier at the particle’s surface. Interestingly, in the controlled synthesis of metallic nanomaterials, many polymers employed to stabilize the nanoparticles during the synthesis (by steric stabilization) can also serve as a diffusion barrier that aids to promote diffusion-limited growth.

Now let us turn our attention to the opposite scenario in which the diffusion of growth species from the solution to the surface is fast, and the growth rate becomes controlled by surface processes [24]. The surface processes in which the adsorbed growth species becomes irreversibly incorporated at the surface (growth itself, Fig. 2.9d) can follow a layer-by-layer or polynuclear pathway. In the layer-by-layer pathway, the growth species have enough time to diffuse on the surface so that a second layer if formed only after a complete layer is produced. In this condition, the growth rate is proportional to the surface area and can be written as:

$$ \text{dr/dt}={{\text{k}}_{\text{m}}}{{\text{r}}^{2}} $$

(2.15)

where k_m is a proportionality constant. Solving this equation we have:

$$ 1/r=1/{{\text{r}}_{0}}-{{\text{k}}_{\text{m}}}\text{t}$$

(2.16)

Thus, the radius difference between two growing particles increases with size (r) and time (t) according to the following expressions, showing that this mechanism does not promote the formation of monodisperse sizes:

$$ \text{ }\!\!\delta\!\!\text{ r}={{\text{r}}^{2}}\text{ }\!\!\delta\!\!\text{ }{{\text{r}}_{0}}/{{\text{r}}_{0}}^{2} $$

(2.17)

$$ \text{ }\!\!\delta\!\!\text{ r}=\text{ }\!\!\delta\!\!\text{ }{{\text{r}}_{0}}/{{\left( 1-{{\text{k}}_{\text{m}}}\text{t}{{\text{r}}_{0}} \right)}^{2}} $$

(2.18)

In the polynuclear mechanism, the concentration of growth species at the surface becomes so high that the formation of a second growing layer is favored before first growth layer is complete. In this case, growth species do not have enough time to diffuse on the surface and the growth rate becomes independent of both particle size and time. Particles then grow linearly with time and the radius difference between distinct growing particles remains constant. Although the absolute radius difference is constant, relative radius difference becomes smaller with the increase in size and time, also favoring the formation of monodisperse nanoparticles (to a lesser extent relative to growth controlled by diffusion). Therefore, the above discussion indicates that growth limited by diffusion represents the best growth condition for achieving monodispersed nanoparticles.

In addition to the classical crystal growth, which can be roughly described by diffusion and incorporation of metal atoms to the growth surface as well as the diffusion of metal atoms from more soluble to the less soluble crystals (Ostwald ripening), the so-called non-classical mechanisms also play an important role during the growth of metal nanostructures [25, 26]. In fact, recent advances regarding in situ transmission electron microscopy techniques and time-resolved small-angle X-ray scattering measurements have allowed the direct observation of solution-phase growth in metal nanoparticles and demonstrated that non-classical mechanisms such as aggregation, coalescence, and oriented attachment affect the growth route and final morphology of metal nanoparticles [22, 27–29]. Specifically, these non-classical mechanisms refer to growth in which primary crystallites assemble into secondary crystals, and can be understood as follows: first, the reversible formation of loose assemblies of primary particles is observed, which are free to rotate and rearrange with respect to one another through Brownian motion. These primary particles can either dissociate or irreversibly attach to each other along some preferential crystallographic orientations to produce a new secondary crystal. In this process, the formation of unique architectures, the incorporation of stacking faults, twin defects, and edge dislocations can be observed [26, 30, 31]. It is important to note that an aggregate corresponds to an irreversible assembly, and thus oriented attachment refers to an irreversible assembly of oriented primary particles. Although the detailed assembly process and the driving forces for the oriented attachment of primary particles along certain directions are still unclear, reduction in the overall surface energy by eliminating some high-energy facets represents one plausible explanation [26].

2.3.3 Growth by Heterogeneous Nucleation (or Seeded Growth)

Heterogeneous nucleation refers to the formation of the new solid phase (nucleation) onto the surface of another material [32]. Heterogeneous nucleation possess lower ∆G* (critical Gibbs free energy barrier for nucleation), and thus are more thermodynamically favored relative to homogeneous nucleation [18]. While in homogeneous nucleation the nucleus is approximated by a sphere whose surface area is 4πr², the nuclei in heterogeneous nucleation have a surface area lower than 4πr² due to the introduction of the interface between the nucleus and the surface, which leads to a decrease in the surface free energy (recall that the increase in ∆G is due to the surface free energy, Eq. 2). The synthesis of metal nanomaterials by heterogeneous nucleation (seeded growth) in solution employing nanoparticles obtained in a separate/previous synthetic step as seeds is very helpful, as it enables one to effectively separate nucleation and growth processes as well as to apply a much wider range of growth conditions, milder reducing agents, lower temperatures, and different solvents. This approach has been extensively employed for the shape-controlled synthesis of a variety of metal nanocrystals, including bimetallic (core-shell) systems where the composition of seeds differs from the growth species [32–34]. For example, if the seed particles possess well-defined shapes and are monodisperse, nanocrystal growth via heterogeneous nucleation can not only produce monodisperse nanocrystal sizes but can also influence (and in some cases template) the shape of the resulting nanocrystal after the final growth step [35].

2.4 Manipulating Nucleation and Growth for Shape-Control

2.4.1 Growth Mechanisms

We have mentioned in the previous sections that the solution-phase synthesis of metal nanomaterials is comprised of two main stages: nucleation and growth. In the shape-controlled synthesis of metal nanomaterials, these stages can be further subdivided into (i) nucleation; (ii) growth of nuclei into seeds; (iii) and evolution of the seeds into the final nanocrystal. These steps are schematically illustrated in Fig. 2.10 [1]. As nuclei correspond to infinitesimal clusters consisting of very few atoms formed as the new solid phase in solution, structure fluctuations may occur in the nuclei with the incorporation of defects in order to minimize surface energy until they evolve into seeds. The seeds can then be defined as species that are larger than the nuclei that are locked into a specific structure, in which structure fluctuation does not occur as they become energetically more costly as the size increases. Interestingly, for several systems (including Ag, Au, and Pd), it has been established that the final shape assumed by the nanocrystal is determined by both the structure of its corresponding seed and the relative binding affinities of capping agents to its distinct surface facets [36–38].

Fig. 2.10

Main stages of controlled synthesis of metal nanomaterials (with face-centered cubic (fcc) structure): (i) nucleation; (ii) growth of nuclei into seeds; (iii) and evolution of the seeds into the final nanocrystal. The shape assumed by the nanocrystal depends on the structure of its corresponding seed and the relative binding affinities of capping agents to its distinct surface facets. R, for example, refers to the ratio between the growth rates along the {100} and {111} directions [38]. Copyright © 2007 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Figure 2.10 illustrates how the structure of the seeds play an important role in determining the final shape of a metal nanocrystal in solution-phase synthesis [38]. For example, a single-crystalline seed (truncated octahedron) can evolve into an octahedron, cuboctahedron, or cube depending on the relative growth rate of {111} and {100} surface facets (expressed as R in Fig. 2.10). Considering their structure, seeds often assume single-crystalline, single-twinned, and/or multiple-twinned structures. Thus, the key for the synthesis of metal nanocrystals with controlled shapes relies on maneuvering nucleation and growth processes so that only one kind of seed, with the proper structure, is present in the reaction mixture. The structure of seeds during solution-phase synthesis (single crystalline, single-twinned, or multiple-twinned) are dependent upon many parameters that include thermodynamic and kinetic considerations, while the relative population distribution of seeds can be further influenced by oxidative etching processes. Moreover, surface capping plays an important role in controlling the shape as the final nanocrystal evolves from its corresponding seed during nanocrystal growth.

When the synthesis in under thermodynamic control, the population of seeds displaying different structures will be determined by the statistical thermodynamics of their free energies. The most stable (lowest energy) seeds will be favored, and surface energy considerations are crucial in understanding and predicting the structure of the produced seeds. Surface energy (γ) can be defined as the energy required for creating a unit area of “new” surface, or as the excess free energy per unit area for a particular crystallographic face according to Eq. 19 and 20:

$$ \text{ }\!\!\gamma\!\!\text{ }={{(\text{ }\!\!\delta\!\!\text{ G}/\text{ }\!\!\delta\!\!\text{ A})}_{\text{n,T,P}}}$$

(2.19)

$$ \text{ }\!\!\gamma\!\!\text{ }=\left( 1\text{/}2 \right){{\text{N}}_{\text{b}}}\text{ }\!\!\varepsilon\!\!\text{ }\!\!\rho\!\!\text{ }$$

(2.20)

where G is the free energy, A the surface area, N_b the number of bonds that need to the broken to produce the new surface, ε the bond strength, and ρ the density of surface atoms. [39] When we consider an amorphous solid or liquid droplet, these equations show that simply decreasing the surface area can lower the total surface energy, and the perfectly symmetric sphere is the resulting shape. However, noble metals assume a face-centered cubic (fcc) crystal structure, in which the distinct surface facets possess different atomic arrangements and surface energy values as illustrated in Fig. 2.11. Stable morphologies that enable the minimization of free energy can be achieved by the exposure of low-index crystal planes that exhibit closest atomic packing (for fcc metals, γ{111}< γ{100} < γ{110}). Although these surface energy values suggest that single-crystal seeds should assume octahedral or tetrahedral shapes (which enables the exposure of {111} facets), they have higher surface areas relative to a cube of the same size. The formation of single-crystal seeds can then be considered in the context of Wulff’s theorem, which represents a method to determine the equilibrium shape of a droplet or crystal having a fixed volume [40]. Energy minimization arguments are used to show that certain crystal planes are preferred over others, giving the crystal its final shape. According to the Wulff construction, single-crystal seeds are expected to exist as truncated octahedrons enclosed by a mix of {111} and {100} facets. This shape, which presents a nearly spherical profile, displays the smallest surface area and leads to the minimization of the total surface energy. Single- and multiple-twinned seeds have also been observed in the synthesis of metal nanocrystals under thermodynamic control. Here, a twin defect refers to single atomic layer in the form of a (111) mirror plane. The surface of a single-twinned seed tends to be enclosed by a mix of {111} and {100} facets to lower the surface free energy, while a multiple-twinned seed is enclosed by {111} facets [41]. As the strain energy caused by the twin defects greatly increases as the seeds grow in size, multiple-twinned seeds are stable only at smaller sizes [42].

Fig. 2.11

Atomic arrangements and surface energy values for {100}, {110}, and {111} surface facets (from left to right, respectively)

Kinetic control relies on the manipulation over the structure/population of seeds containing different numbers of twin defects by decreasing the reduction or decomposition rate of a precursor that is responsible for providing metal atoms (growth species). In fact, when the decomposition or reduction becomes considerably low, atoms tend to form nuclei and seeds through random hexagonal close packing together with the inclusion of stacking faults, and the resulting seeds typically takes a shape that deviates (having higher surface energies) from those favored by thermodynamics [43]. In practice, structures that deviate from thermodynamic considerations can be achieved by a set of special conditions that enable an extremely low concentration of metal atoms in solution so the nuclei will not be able to grow into polyhedral structures. These include substantially slowing down precursor decomposition or reduction rates, coupling the reduction to an oxidation process, and/or by taking advantage of Ostwald ripening (process by which relatively large structures grow at the expense of smaller ones) [1].

In addition to thermodynamic and kinetic control, which dictate the identity and structure of the produced seeds, the relative population of single-crystal and twinned seeds can be influenced by oxidative etching processes, in which metal atoms can be oxidized back to the ionic form by O₂ from air (present in the synthesis) and a ligand such as Cl⁻ and Br⁻ [44]. The defect zones in twinned seeds have higher energy relative to single-crystal regions, making them more susceptible to undergo oxidation. Thus, multiply-twinned seeds are more susceptible to oxidation than single-twinned and single-crystal seeds, respectively [45]. Consequently, if properly controlled, oxidative etching can enable the selective removal of a specific kind of seed from the reaction mixture as well as to control the number of seeds [46, 47].

When we consider the evolution of the seeds into the nanocrystal, capping agents that facilitate the interaction of distinct surface facets with different binding affinities can be put to work to control the shape of the produced nanocrystal. Upon interaction with the surface, capping agents can change the order of free energies from distinct facets, changing their relative growth rates. Therefore, the use of a capping agent can be considered as a thermodynamic approach. For example, a capping agent that binds more strongly to a specific facet can slow down its growth relative to the other facets, allowing for the control over its relative surface area in the final nanocrystal. This is illustrated by the possibility of producing octahedron, cuboctahedron, or cube nanocrystals from a single-crystal seed as shown in Fig. 2.10. If a capping agent that binds more strongly to the {111} relative to the {100} facet from the single-crystal seed is employed, the growth along {111} is significantly slowed down relative to the {100} direction producing octahedrons (contain only {111} surface facets). In the opposite scenario, cubes would be obtained. Intermediary conditions would lead to the formation of cubeoctahedrons (enclosed by a mix of {111} and {100} surface facets). It is important to emphasize that although many successful examples have been reported on the use of capping agents to control the shape of metal nanocrystals, the exact mechanisms are still not completely understood due to the lack of experimental techniques to precisely probe the molecular structure and resulting surface energy of different facets upon the binding of a capping agent.

Now that we have discussed the fundamentals of nucleation and growth, as well as the general concepts involved in synthesis of metal nanomaterials with controlled shapes, let us discuss the controlled synthesis of Ag nanostructures as an example on how thermodynamic and kinetic control, oxidative etching, and surface capping can be employed to give birth to a myriad of nanocrystals with monodispersed sizes and well-defined shapes.

2.4.2 Controlled Synthesis of Silver Nanomaterials

In the last decade, many methods have been developed for the synthesis of Ag nanocrystals displaying uniform size distribution and variety of shapes that include cubes, bypiramides, wires, octahedrons, triangular plates, bars, among others [1, 17, 37, 48]. Representative scanning electron microscopy images of some of these shapes are as illustrated in Fig. 2.12. Among the various solution-phase methods for obtaining Ag nanocrystals, we will focus our discussion on the polyol reduction. Polyol reduction is probably the best-established protocol for generating Ag nanomaterials with controllable shapes [37]. This approach is illustrated in Fig. 2.13 and it is based on the utilization of ethylene glycol as both a solvent and source of reducing agent (glycoaldehyde), high temperatures (~ 140–160 °C), Ag⁺ as the Ag precursor, and polyvinylpyrrolidone (PVP) as the capping agent. As the reduction of Ag⁺ to Ag starts, the concentration Ag⁰ in solution starts to increase. As discussed in Sect. 3.1, nucleation takes place only when the supersaturation with Ag reaches a value above the solubility where ∆G* is satisfied. As the nuclei are formed, they can undergo structure fluctuations at small sizes due to the input of thermal energy so that defects can be incorporated to reduce their surface energy [49]. For example, the formation of a twin defect can be thermodynamically favored at small sizes as they result in the exposure of {111} (lowest energy) facets. As nuclei grow, these structural fluctuations start to become energetically more costly. The birth of the seeds is reached when these structure fluctuations stop (thermal energy is not enough to enable structure fluctuation). In the polyol synthesis of Ag nanomaterials, a mixture of multiple-twinned, single-twinned, and single-crystal seeds is produced in the reaction mixture as shown in Fig. 2.13, which can evolve into wires, bypiramids, and cubes as shown in Fig. 2.13a–c, respectively [50].

Fig. 2.12

Scanning electron micrographs of Ag wires (a), cubes (b), truncated octahedrons (c), octahedrons (d), bypiramids (e), and plates (f) [36, 47, 48, 55]. a Reproduced by permission of the PCCP Owner Societies, (c–d ) Copyright © 2006 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim, (e) Copyright 2006 American Chemical Society, (f) Copyright © 2006 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim

Fig. 2.13

Polyol approach for the synthesis of Ag nanocrystals displaying controlled shapes. A mixture of multiple-twinned, single-twinned, and single-crystal seeds is produced in the reaction mixture, which can grow into wires (a), bypiramids (b), and cubes (c) in the presence of polyvinylpyrrolidone (PVP) as the capping agent [50]. Copyright 2006 American Chemical Society

The multiple-twinned (decahedral) seeds are the most abundant ones (relative to single crystal and single-twinned seeds) in the reaction mixture as introduction of twin planes is thermodynamically favored at small sizes [51]. Therefore, the removal of the other seeds is not required to achieve the synthesis of Ag nanowires. As their twin defects are more reactive relative to single-crystal regions, Ag atoms (growth species) preferentially add to the twin defects of the multiple-twinned seeds during growth, leading to the elongation of the seeds producing small pentagonal rods. As PVP is employed as a surface capping agent and has a higher binding affinity with the {100} side facets relative to the {111} end facets, growth along the side facets is hindered, leading anisotropic growth into nanowires tens of micrometers long (Fig. 2.12a and 2.13a) [47, 52]. Although the other seeds (single crystal and single-twinned) may also grow in the same reaction mixture, their smaller sizes relative to nanowires facilitate their separation by centrifugation.

The removal of twinned seeds is required to achieve the synthesis of Ag nanocubes (Fig. 2.12b and 2.13c). This selective removal of twinned seeds can be accomplished by oxidative etching upon the addition of Cl⁻ ions in the reaction mixture [53]. In this case, the O₂/Cl⁻ pair becomes an efficient etchant of twinned seeds owing to their higher susceptibility to oxidation relative to the single-crystal seeds [54]. As PVP binds more strongly {100} relative to the {111} facets in the single-crystal seeds that are left behind, cubes are produced. Interestingly, it has been observed that under a continued supply of Ag atoms by extending the polyol reaction, the Ag nanocrystal shape can evolve from a cube to a cuboctahedron and then to an octahedron, with an increasing ratio of {111} to {100} facets (Fig. 2.12c and d, respectively) [36]. In this case, rather than a layer by layer growth, Ag selectively deposits over the {100} cube facets, slowing down the growth rate along the {111} direction and enabling its increase exposure at the surface of the final nanocrystal [36].

In order to selectively produce single-twinned seeds, the degree of oxidative etching must be moderated so that only the multiple-twinned seeds (that are the most reactive) are removed from the reaction mixture by oxidation. This can be performed by replacing Cl⁻ with the milder etchant Br⁻ and has enabled the synthesis of right bypiramids (Fig. 2.12e and 2.13b) [48].

All these shapes discussed so far (wires, cubes, octahedrons, and bypiramids) are thermodynamically favored. In order to obtain thermodynamically unfavorable shapes, such as triangular nanoplates (Fig. 2.12f), different strategies must be employed, such as the substantial reduction on the supply of Ag atoms by greatly reducing the precursor reduction rate [1]. In this case, Ag atoms may form plate-like seeds through random hexagonal close packing with the inclusion of stacking faults, that can subsequently grow to yield Ag triangular plates [43]. The use of the hydroxyl end groups of PVP as an extremely mild reducing agent has been employed as an effective approach to slow down the Ag⁺ reduction rate and yield the formation of Ag triangular plates [55].

It is important to mention that these Ag nanocrystals could be further employed as seeds for overgrowth by heterogeneous nucleation (seeded growth) in order to obtain nanomaterials with monodisperse sizes. Moreover, the shape can be further manipulated by the proper choice of capping agents. In a typical example, it has been demonstrated that Ag cuboctahedrons (single-crystal structures enclosed by a mix of {111} and {100} surface facets) could be employed as seeds for Ag heterogeneous nucleation and growth, as shown in Fig. 2.14a [35]. In this seeded-growth process, the utilization of sodium citrate (Na₃CA) as capping agent led to the production of Ag octahedrons enclosed by {111} facets (Fig. 2.14b), while the utilization of PVP led to the formation of cubes/bars enclosed by {100} facets (Fig. 2.14c). It has been established that, in the case of Ag, citrate binds more strongly to {111} relative to {100} facets.[56] Therefore, the utilization of citrate as capping agent allows for the reduction on the growth rate along the [111] direction and thus leading to the exposure of {111} facets in the final nanocrystal (formation of octahedrons). Conversely, PVP binds more strongly to {100} than {111} facets of Ag, reducing the growth rate along the [100] direction and resulting in the exposure of {100} facets in the final nanocrystal (formation of cubes and bars) [52]. Other Ag nanocrystals, such as Ag triangular plates, can also be employed as seeds for Ag overgrowth using this principle as depicted in Fig. 2.15a [57]. Ag triangular plates are covered by {111} faces in both top and bottom surfaces, while their sides contain twin planes and stacking faults. When Na₃CA is employed as capping agent, growth along the side facets (lateral growth) occurs preferentially relative to growth along the top and bottom surfaces (Fig. 2.15b and c). Even though the side surface of the plates also contains {111} planes, the stacking faults and twin defects make them more reactive towards Ag nucleation relative to the top and bottom surfaces in the presence of Na₃CA. On the other hand, vertical growth is induced in the presence of PVP (Fig. 2.15d and e). As PVP can bind more strongly to {100} facets, the growth along the top and bottom {111} surfaces becomes faster relative to lateral growth (that also occurs).

Fig. 2.14

a Scheme for the formation of Ag octahedrons or cubes by overgrowth using Ag cuboctahedrons as seeds and Na₃CA or PVP as capping agents. b and c Show transmission electron micrographs for the produced octahedrons and cubes, respectively [35]. Copyright 2010 American Chemical Society

Fig. 2.15

a Scheme for the lateral or vertical overgrowth of Ag employing Ag triangular plates as seeds and Na₃CA or PVP as capping agents. b–c and d–e Show scanning electron micrographs for the produced plates after lateral and vertical growth, respectively [57]. Copyright © 2011 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim

Seeded growth has been widely applied for producing a variety of controlled bimetallic nanomaterials with a myriad of architectures by the utilization of Ag, Au, Pd, and Pt nanocrystals of different shapes as seeds for the overgrowth of Ag, Au, Pd, and Pt, for example [32, 34]. In addition to the utilization of pre-formed nanocrystals as seeds for heterogeneous nucleation and growth, they can also be employed as chemical templates for the synthesis of controlled nanomaterials. One example is the utilization of Ag nanocrystals as sacrificial templates in a galvanic replacement reaction with ions AuCl₄ ⁻, PdCl₄ ²⁻, and PtCl₆ ²⁻ to produce Ag-Au, Ag-Pd, and Ag-Pt nanocrystals, respectively, having bimetallic compositions, ultrathin walls, hollow interiors, and different morphologies (boxes, cages, tubes, frames, shells, etc.) [58].

2.5 Final Remarks

This chapter described the fundamentals of the controlled synthesis of metal nanomaterials in solution phase in terms of the available theoretical framework. Our discussion encompassed the mechanisms for the stabilization of nanostructures in solution, the theory behind nucleation and growth stages, and the role that thermodynamic and kinetic considerations, oxidative etching, and surface capping can play in the synthesis of metal nanomaterials with controlled shapes using Ag as a representative example. It is obvious that solution-phase synthesis represents a powerful approach to produce metal nanomaterials with monodisperse and controlled sizes, shapes, architectures, and composition. Solution-phase methods enable the control over several experimental parameters (temperature, choice of precursors, stabilizers, solvent, reducing agents, their molar ratios, concentrations, among others) and have, at least in theory, the potential for large-scale fabrication. However, progress on the truly controlled synthesis of metal nanomaterials has only been achieved in the last decade, and the synthesis of many metal nanocrystals in a variety of shapes has been realized. Despite these attractive features, several challenges (both fundamental and practical) still remain and need to be addressed so that solution-phase synthesis can reach its full potential. These include: (i) a better understanding on the atomistic details and mechanisms behind the nucleation and growth processes is required; (ii) the precise roles played by solvent molecules, capping agents, and trace impurities must be unraveled; (iii) the scope of controlled synthesis needs to be increased, i.e., the development of protocols for the controlled synthesis of several metals still need to be developed; (iv) protocols must be reproducible; (v) it is desirable that solution-phase protocols are facile, environmentally friendly, and carried out in water as the solvent; and (vi) the controlled synthesis of metal nanomaterials in large scale must be realized in order to allow their transition from the lab to industry. It is safe to say that addressing these shortcomings will allow not only the synthesis variety of metal nanocrystals with controlled shapes/sizes by design, but also enable them to have a more profound impact in various fields of knowledge.