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In the extreme case whereno oriented alignment exists, the attachment is trans-formed to a random aggregation.In real growth processes of inorganic nanocrys-tals in solution, ion-by-ion addition always takesplace at the initial stage. If the solubility of thegrown nanocrystals is not too low, usually, the growthis associated with OR mechanism, for which a ki-netic model was rigorously developed by Lifshitz, Sly-ozov, and Wagner, and is also known as the LSWmodel[100;101]. The LSW model predicts that themean particle radius should evolve as a function oftime according to Eq.(9)D ¡ D0 = k(t ¡ t0)1n (9)Sometimes it is also read asDn ¡ Dn0 = kt (90)where D is the average particle diameter at any time,D0 is the initial average particle diameter, k is therate constant for OR process, t is time, and n is a co-efficient dependent of the limiting step to the growth.This model is a direct result of the combination ofGibbs-Thompson equation and Fick0s first law. Thepower law coefficient n=3 is obtained by consideringdilute conditions, where diffusion of ions in solutionis the limiting step. In concentrated conditions, whenn=2, the growth is controlled by the volume diffusionof ions in the matrix; and when n=4, it is deducedthat growth is controlled by dissolution kinetics atthe particle matrix interface[102].However, recent work has demonstrated that thismodel cannot be utilized universally to account forthe growth process in some systems. Instead, OAmechanism was proposed as another significant pro-cess that may occur during nanocrystal growth whenthe viscosity of the solution is not too high[103–105].Recent studies indicate that OA mechanism is verysignificant, even in the early stages of nanocrystalgrowth. Huang et al:[106]in their work showed thatthe coarsening kinetics of mercaptoethanol-stabilizedZnS nanoparticles can be pided into two steps. Atthe beginning of the hydrothermal treatment, themercaptoethanol capping limits diffusion of ions, andthus prevents particle coarsening via OR mechanism.Thus, in the first step, the coarsening is controlled byOA and the growth can be described by an OA kineticequationD = D0(3p2k1t + 1)k1t + 1(10)The second step of crystal growth can be described us-ing LSWmodel (Eq.(9)). However, their work showedthat both OA and OR processes were important in thesecond stage, and the particle size at any time was theresult of the two processes occurring simultaneouslyD = D0(3p2k1t + 1)k1t + 1+ k2t1n (11)Note that the component of growth described by thefirst term (describing pure OA growth) diminishesrapidly so that the second term predominates overlonger times.ASG mechanism, as we have mentioned, is readilyaffected by experimental parameters. For example,we can stabilize nanoparticles sterically by cappingthem with surfactants or electrically by surroundingthem with charges. Inversely, we can also destabilizenanoparticles for growth purpose by removing the sur-face capping layers. Further more, we can selectivelyblock the growth of a certain surface plane by sta-bilizing it, while allow the growth of a naturally lessactive surface plane. By employing this technique, wecan realize controlled synthesis by ASG mechanism.According to Vayssieres[107], colloid particles can bestabilized after surface capping only if the variationof interfacial tension due to adsorption in negative.This parameter is directly related to pH value andionic strength of the solution:∆° = ° ¡ °0= 25:7¾maxlnh1 ¡ I³0:117sh(1:15∆pH)¾max´2i¡6pI[ch(1:15∆pH) ¡ 1] (12)PZIT = PZC + 2:46 + 2log¾max ¡ logI (13)∆pH = PZIT ¡ PZC (14)where ° and °0 are interfacial tension due to adsorp-tion and the initial interfacial tension, respectively,¾max is the maximum surface charge, I is the ionicstrength of the solution, PZIT is the point of zero in-terfacial tension, and PZC is the point of zero charge.For oxide colloidal particles in solution, they are sta-ble against precipitation when pH>PZIT, and unsta-ble when PZC·pH<PZIT. Therefore, we can adjustpH value and ionic strength to tune the colloidal parti-cles in stable or unstable regime. Generally speaking,when the limiting step is surface reaction, random ag-gregation dominates; when the process is the diffusioncontrolled, oriented attachment is prevalent.Interface tension and surface charges are trulykey factors that determine the growth behavior ofnanoscale entities. According to the knowledge weintroduced previously, we can design synthetic ap-proaches to grow a series of novel nanostructures aswill be shown in the following section.3. Synthesis and Characterization of ZnO 1DNanomaterialsSynthesis of ZnO 1D nanomaterials has attracteda lot of research attention in recent years. We list therepresentative synthetic methods developed in thisfield in literature [108–130]. In this section, we mainlyfocus on the work carried out in our group. 3.1 Vapor phase growth of ZnO 1D nanomaterialsFor vapor phase growth of ZnO 1D nanomaterials,we will give two examples to show how to control themorphology of the materials by varying experimentalparameters, and we will elucidate the roles of the pa-rameters that affect the growth behavior according tothe theories we introduced in the previous section.In the first example, commercial ZnO and graphitepowders mixed in a molar ratio of 4:1 were used asthe source material. Several pieces of Si wafers wereplaced downstream in a conventional horizontal tubefurnace side by side, and another Si wafer was putabove the source materials horizontally, as the de-position substrates[115]. The morphologies of ZnOnanostructures vary with the distance from the sourcematerial to the substrate from mirorods, throughnanoplatelets, nanoplatelet flowers, nanobelt flowers,to nanowire flowers with each morphology appearingin a wide region. The morphologies of the structureson the substrates above the source material are dis-played in scanning electron microscopy (SEM) imagesin Fig.1(a). It is noteworthy that the aspect ratio (de-fined by length pided by width for nanoplatelets andnanobelts, and length pided by diameter for rodsand nanowires) of the structures increases with in-creasing the distance from the source material to thesubstrate (Fig.1). Transmission electron microscopy(TEM) images in Fig.2 reveal that ZnO nanoplateletis irregular in shape, and the top and bottom surfacesare fixed as §(2¯ 1¯ 10) planes. The nanobelts also pos-sess §(2¯ 1¯ 10) top and bottom surfaces, and have fixedgrowth directions along [0001] axis, the same as ZnOnanowires.In the second example, zinc powders used as thesource materials and a tin-coated silicon wafer wasplaced downstream as the substrate[116]. During theheating process water was introduced into the reac-tion. Figure 3(a) shows the morphology of Sn-dopedZnO, and Fig.3(b) exhibits a TEM image, which re-veal that the grows along the a axis, with its top andbottom surfaces as §(0001). Naturally in growth ofZnO nanomaterials, the highest growth rate is alongthe c axis, however in this example we found that itwas possible to change the growth behavior of ZnOnanobelts by doping.Now, we discuss the dependence of morphologieson experimental parameters.3.1.1 Supersaturation and surface energy Inthe effort to grow nanoscale structures, the supersat-uration is generally much larger than unity, and thecrystal grows under conditions far away from ther-mal equilibrium. Under such a circumstance, growthkinetics takes a part, or even determines the growthbehavior of the surface planes and the final morphol-ogy of the structure.Similar to what Brenner and Sears[81]and Namet al:[131]proposed, the forced diffusion of ZnO vaporby Ar is more important than the molecular diffu-sion driven by a concentration gradient for the lowReynolds number of our system. Under this advec-tion condition, there generally exists a maximum inthe ZnO vapor concentration and the supersaturationdecreases downstream from the source material, andas the flow rate of the Ar carrier gas increases, themaximum shifts further downstream, as illustrated inFig.4. In region I in the first example, the supersat-uration is low, and rods with well-defined facets off10¯ 10g major plane, f01¯ 11g capping planes, (000¯ 1)base plane, and (0001) minor tip plane are formedsimilar to Laudise and Ballman[132]reported. Theplanes of the nanobeltsFig.4 Supersaturation profile of ZnO vapor in the re-actor under different flow rate of carrier gasalong the distance downstream from the sourcematerial[115]crystal growth in this region best reflects the crys-tallographic symmetry of the hexagonal ZnO. In re-gion II, the supersaturation is the highest, and thegrowth takes place under circumstances far from ther-mal equilibrium, where the kinetics plays a more sig-nificant role in determining the morphology of thestructure. When a molecule is adsorbed on a sur-face plane, for example, the top surface of a plateshape structure, it will either incorporate into latticesor desorb from the surfaces. When the top surface isthe lowest in surface energy, the diffusion distance issmall for the molecules adsorbed on this surface, andthus, the molecules have little chance to incorporateinto the lattice lattices on this surface. Meanwhile,the molecules have good chance to diffuse to the frontand the side surfaces, where the surface energies arehigher and diffusion distances are longer, and thus themolecules are much easier to incorporate into the lat-tice planes on these surfaces. Furthermore, it shouldbe kept in mind that the higher the surface energy,the easier for the nucleation on that surface. There isalso a positive feedback effect in this competition pro-cess, namely that, the more aggressive in the compe-tition for the capturing of the adsorbed molecules, thelarger the growth rate for the front and side surfaces,and then the larger surface area for the top surface;the larger area for the top surface, the larger num-ber collection of impinged molecules on this surfaceto support the growth of the front and side surfaces.The failure in looting the adsorbed molecules leads tothe much slower growth of the top surface. The ag-gression characteristics of the front and the side sur-faces similarly determine their proportion in the finalmorphology. When these two surfaces are similar insurface energy or belong to the same family of latticeplanes, the competition ability of these surfaces is alsosimilar. Then, the morphology may take the form ofa nanosheet or a nanoplatelet structure. In contrary,when these two surfaces differ greatly in energy, the fi-nal morphology may be more like a nanobelt as shownin region III, where the supersaturation is lower thanthat in region II, and the competition for capturingthe adsorbed molecules is more severe, namely that,the front surface loots the molecules not only from thetop surface, but also from the side surface. When thecrystal has several fast growth directions, for example,ZnO shows the fast growth directions along [0001],[2¯ 1¯ 10], and [01¯ 10] axes, and has polar surfaces of(0001) and (01¯ 10). Although the nanoplatelets are ir-regular in shape, they show preferred growth directionalong the packing of the polar (01¯ 10) planes. Whenthe substrate is far away from the source material, thesupersaturation is very low. More severe the competi-tion for the adsorbed molecules than in region III, themorphology develops its quasi-one-dimensional char-acter, and nanowires are formed.3.1.2 Temperature at the source material and the sub-strate, temperature gradient in the tube furnace, dis-tance from the source material and the substrate, andthe heating rate of the reactor The temperatureat the source material determines the vapor pressureof it. The pressure increases exponentially with heat-ing temperature. When the vapor advects within thecarrier gas downstream to the substrates, there willbuild a local supersaturation higher than in the casewith lower heating temperature. Therefore, the mor-phology will change accordingly. The temperature atthe substrate and the temperature gradient similarlyinfluence the supersaturation profile, and hence themorphology. Distance from the source material andthe substate implies the selection of a local super-saturation and temperature. The heating rate mayhave influence on the initial nucleation process: higherheating rate makes it possible the homogeneous nu-cleation and low dispersity of the morphology, whilelower heating rate generally causes high dispersity.Temperature at the substrate has also effect on thenucleation density.3.1.3 The gas flow rate and the inner diameter of theceramic tube The gas flow rate and the innerdiameter of the ceramic tube have fundamental influence on the supersaturation profile through theinteraction with the Reynolds number. In Fig.4, su-persaturation profile under different gas flow rate hasbeen displayed. With higher gas flow rate and smallinner diameter of the ceramic tube, the maximum su-persaturation will move downstream from the sourcematerial, and the morphology of the structure at afixed position on the substrate will change accord-ingly.3.1.4 The starting material Materials withnanometer scaled sizes have been well-known to showa suppressed melting temperature. Therefore, usingnanomaterial as the source corresponds to elevatingthe heating temperature of the source material, andsimilarly changes the supersaturation profile and thefinal morphology of the structure. As we have shownpreviously, ZnO is a hexagonal phase with inherentpolarity. Naturally, ZnO grows along c axis muchfaster than along a and b axes. The final morphologytends to minimize the surface area of (0002) plane.Therefore, the nanobelts and nanoplatelets in the firstexample possess top and bottom surfaces of §(2¯ 1¯ 10).Nanobelts and nanoplatelets with §(0001) top andbottom surfaces are believed unstable. However, sim-ply by adding Sn in the source materials we can obtain§(0001) surface bounded ZnO nanobelts. Recently,Ding et al:[133]have proposed a mechanism for thestabilization of this structure by ion dopants.3.2 Solution phase growth of ZnO 1D nanomaterialsIn this section we will show how to control mor-phologies of ZnO nanostructures by ASG mechanism.By combining the crystallography of ZnO and crys-tal growth knowledge we introduced in the previoussection, we synthesized a series of ZnO nanomaterialssimply by adjusting experimental parameters. Therepresentative ZnO nanostructures we have grown in-clude ZnO nanorods, nanoplatelets, nanostacks, nan-otubes, nanorings, nanotrees, and nanoforests. Gen-erally, growth of ZnO nanorods was carried out aspreviously reported[122–124]. In brief, single crystallinesilicon wafers were cleaned and coated with zinc ac-etate dihydrate solution in ethanol with different con-centrations and coating rounds, which yielded ZnOseed crystals with different density and aggregationstatus after annealing in air at 350±C. The substrateswere placed upside down in a solution containing0.02 mol/L Zn(NO3)2¢6H2O and 0.02 mol/L hexam-ethyltetramine (HMTA) and reacted at 90±C for acertain time. ZnO nanotubes were converted fromZnO nanorods by selective removal of nanorod corein an HMTA solution. We grew ZnO nanoplateletsby using seed-coated polycrystalline aluminum sub-strates. The pH values of the solution were adjustedwith the addition of NH3¢H2O in this experiment.ZnO nanostacks were synthesized using bare poly-crystalline aluminum substrates by increasing ionicstrength of the solution. By adding a surfactant ofcetyltrimethylammonium bromide (CTAB), we ob-tained ZnO nanorings. To synthesize ZnO nanotrees,we adjusted the ionic strength of the solution by usingZnCl2, NaCl, and urea as source materials. Finally,ZnO nanoforests were synthesized by using a modifiedseed coating techniques. In the following paragraphs,we will detail the experimental results and discuss themechanism for morphology evolution.3.2.1 ZnO nanorods It is well-known that cylindershaped rods are the most general morphology of ZnOwith the wurtzite crystal lattice. It is both thermody-namically and kinetically favorable for developing thec-axis oriented rods. To facilitate the nucleation, wehave coated Si substrate with ZnO nanocrystal seedlayers. The diameter and length of ZnO nanorods canbe varied by changing the reaction time, temperature,concentration of source materials, and the structureof the seed layer. Due to the concentration depletionduring the growth (both homogeneous and heteroge-neous nucleation), ZnO nanorods will cease growthafter a certain time (within 10 h), and the length willreach a limiting maximum value even with extendingthe growth time. The growth temperature does playa direct role in growth kinetics. With increasing thetemperature, ZnO nanorods grow faster and the di-ameter gets larger while the length is similar or evenshorter since no enough source materials are availableduring the later part of the growth. It is noteworthythat when the growth temperature is higher, the ho-mogeneous nucleation rate of ZnO in solution is alsolarger and their further growth can even compete withthe growth on substrates. Diameter of ZnO nanorodscan be decreased by lowering the concentration of thesource materials. Although under such conditions, thegrowth rate decreases, the length of the nanorods canactually be larger since the homonucleation is sloweddown and a higher percentage of source materials cangrow on the substrates (see Fig.5). The effect of theseed layer is more complex, and we will return to thisissue later.3.2.2 ZnO nanotubes Thermodynamics and kinet-ics favor the growth of cylinder shaped ZnO nanorodsalong c-axis. Inversely, ZnO nanorods also dissolvepreferentially along this axis. Under hydrothermal treatment in a pure HMTA solution with other exper-imental parameters the same to that used for growingZnO nanorods, the core of ZnO nanorods can be selec-tively removed to form hollow nanotubes as shown inFig.6. From SEM images, it is apparent that the nan-otubes are shorter than the original nanorods whilethe diameter is almost the same. The results indi-cate that the dissolution along c-axis is much fasterthan along other axes. However, the different dissolu-tion rates cannot fully explain the selective etching offof the core of ZnO nanorods. Therefore, there mustbe other essential factors that affect the etching rate.Under hydrothermal treatment, the dissolution pro-ceeds rapidly in the core part, if it is more defectivewith NH4+ molecule trapping. The selective removalof core part of ZnO nanorods has to be carried outin experimental conditions similar to that used forgrowth. Under strong acidic or alkaline conditions,the etching takes place everywhere rapidly, and thenanorods become porous before complete dissolution.Actually, we have found that the dissolution behaviorsbetween Zn polar and O polar surfaces differ dramat-ically (data not shown here).3.2.3 ZnO nanoplatelets As we just men-tioned, the free growth mode favors the formation ofc-axis oriented nanorods. We can use external pa-rameters to change the growth kinetics. In this case,we use Al(OH)¡4 (generated by reaction between Aland OH¡) to passivate the positively charged Zn ter-minated (0001) surface. After the passivation, thesurface energy of (0001) plane decreases dramaticallyand the growth of this surface is suppressed. Then,f10¯ 10g surfaces become the favorable growth planes.The consequence is the developing of platelet mor-phology. The thickness of the nanoplatelets can betuned by adjusting the passivating strength, whichcan be achieved by changing the pH values in solution.As shown in Fig.7, ZnO nanoplatelets with thicknesssmaller than about 10 nm can be achieved by thismethod. Though the thickness of ZnO nanoplateletscan be very small, they are still highly crystallized.The reason that the (0001) surface can be effectivelypassivated by Al(OH)¡4 are two fold: firstly, thissurface is inherently Zn terminated and positivelycharged, which can attract negative ions, such asAl(OH)¡4 ; secondly, thermodynamic chemical poten- tial decreases by eliminating the electrostatic energyat the surface owing to the neutralization of posi-tive and negative charges, and the more complete thepassivation, the larger the potential decrease. Withstrong passivation, the growth of (0001) surface canbe suppressed dramatically. Therefore, it is a natu-ral outcome that we can tune the thickness of ZnOnanoplatelets by adjusting the passivation, which canbe simply realized by changing the pH values of thesolution.3.2.4 ZnO nanostacks Similar to growth ofZnO nanoplatelets, ZnO nanostacks can be synthe-sized on a polycrystalline Al substrate with or with-out seed layers. It is essential to increase the ionicstrength of the solution a little bit by adding 1 mmolof NaCl. As we have discussed in previous sections,ionic strength is one of the key factors that affectthe stability of colloidal particles in solution. Theconsequence is the growth proceeds more like in adiffusion-limited OA process. As exhibited in Fig.8,ZnO nanoplatelets formed by oriented attachment ofnanoparticles. Since in this case, the side surfaces ofthe nanoplatelets are also passivated to a certain ex-tent, the growth along c-axis can take place, however,in this case, by a layer-by-layer assembly fashion, thenanoparticles have a certain size. The thickness ofthe layer, equally, the size of the nanoparticles, canbe controlled by adjusting the ionic strength of thesolution. It is interesting that the layer-by-layer as-sembly is orientationally registered. The significantimplication of this registry is the concept of three-dimensional oriented attachment, that is to say, thewhole structure is a single-crystal-like entity. The ori-entational registry becomes apparent from the SAEDpattern taken from the nanostacks. The growth phe-nomena of ZnO nanostacks is similar to the growth ofcalcite and other minerals by biomimetic approach insoft chemistry, and has potential applications as ad-vanced ceramics and functional nanocomposites aftersuitable modification of the synthetic routes.3.2.5 ZnO nanorings ZnO nanoparticles can as-semble by OA mechanism to form nanoplatelets. Inthat case, we adjusted the ionic strength of the solu-tion to slightly destabilize the colloidal particles. Inthis example, we use a surfactant-CTAB to passivatethe surface of ZnO nanoparticles. Owning to the dif-ference in surface energy for (0001) and f10¯ 10g sur-faces, the passivation strength is also different. TheCTAB passivated ZnO nanoparticles can graduallyassemble via OA mechanism, however the assemblyproceeds in a different manner than that for ZnOnanoplatelets. As shown in Fig.9, ZnO nanorings withhexagonal shape are exhibited. The mechanism of theassembly is similar to what Liu and Zeng[134]reportedfor the growth of CdS nanorings[134]. The key fac-tor that affects the OA process is the passivation anddestabilization of ZnO nanoparticles. Too weak or toostrong of the passivation is unfavorable for the for-mation of nanorings. Naturally, the hexagonal sym-metry of crystal lattice of ZnO is the reason for theformation of hexagonal shaped nanorings. Althoughthe inherent symmetry determines external shape ofthe nanostructures we prepared, the OA is not neces-sarily perfect. As we have stated that imperfect at-tachment sometimes did take place when the attachedcrystals were not properly registered. As shown inFig.9(b), the periphery of a nanoring is not closedand the sides bend inward similar what suggested byLiu and Zeng[134]in their work. We have also foundsome intermediate products that consists only two orthree sides, whereas the angle between neighboringsides keeps a fixed angle of 120±. It is also interestingthat the nanorings can be enclosed either with hollowtube sides never shown before or with solid sides as reported by other researchers, which may be relatedto the particular growth conditions in this example.Furthermore, the nanorings can be completely filledby over-growth (Fig.9(d)).3.2.6 ZnO nanotrees The nanotree structurescan be synthesized at an intermediate ionic strength.In this example, we will show how the passivationstrength can be used to tune the morphology of ZnOnanostructures. We used ZnCl2 and urea as sourcematerials and added NaCl to adjust the ionic strength.In this case, a-axis oriented ZnO nanorods grew asexpected since the passivation of (0001) surface wasrather strong. However, the uniform nucleation andgrowth from the side surfaces were not possible be-cause diffusion is the rate limiting step, even thoughthe surface energies are comparable for the side sur-faces of a cylinder. The side surfaces of ZnO nanorodsbecame rough and nucleation of branches took placerandomly and the top and bottom surfaces kept nearlyfree from growth since they are strongly passivatedsimilar to the case of ZnO nanoplatelets. As shownin Fig.10, branches and sub-branches grew randomlyon the stems of ZnO nanotrees. The branches andsub-branches all inclined with a certain angle to thestems. After high-resolution TEM imaging, we foundthat the branches and sub-branches grew with an im-perfect OA, and the inclination to the stems wereapproximately 60± within 10± of misfit angle. Theinclination at 60± is understandable from the hexag-onal lattice symmetry of ZnO. The misfit angle couldnot be due to the growth kinetics since it is not achemical reaction limited process. We believe thatthe electrostatic interaction among the branches andsub-branches caused the misfit in the nanotree struc-ture.3.2.7 ZnO nanoforests Similar to the synthe-sis of ZnO nanorods, growth of ZnO nanoforests wasalso carried out in the same precursor solution by us-ing a seed-coated Si substrate. The only difference isthat we used a modified method to synthesize ZnOseed layers. As shown previously, ZnO seed crys-tals were synthesized by decomposing zinc acetateprecursors in ethanol solution on Si substrates. Bythis method, ZnO seed layers can be synthesized assmooth films. However, owing to the high evapo-rating rate of ethanol, high-density continuous filmis not possible to be obtained by this method. Al-ternatively, we used a mixed solution of ethanol andpropanol to prepare the seed layer. By this approach,we could readily synthesize ZnO nanocrystal seed lay-ers with high density over large areas owing to thesuppressed evaporating rate of the mixed solution.Different from the synthesis of ZnO nanorods, in thisexample, there are two steps in the growth of ZnOnanoforests, where the first step is the growth of ZnOnanorods and the second step is the further growthof thin nanowires on the top of the nanorods. Fig-ure 11 displays the so-called ZnO nanoforests. Thenanowires are much thinner than the nanorods. Thereis only one nanowire standing on the top of somenanorods, whereas several nanowires can also grow onthe top of one nanorod. The two-step growth processcan be understood by considering the consumptionof the source materials. As has been discussed forthe growth of ZnO nanorods, the growth will cease inseveral hours because of the complete consumption ofsource materials. However, in this example, the seedcrystals are deposited with high density and have finerparticle sizes, which can supply source materials byOR mechanism during the second growth step. Sincethe concentration of source materials in the secondgrowth step is much lower than in the first one, con-tinuous growth of ZnO nanorods with the same diam-eter is not possible. Therefore, nucleation of nanorodswith smaller diameter and final gradual growth intonanowires caused the formation of ZnO nanoforests.It is now clear that by controlling the experimen-tal parameters the growth thermodynamics and kinet-ics can be varied, and therefore, morphology selectivesynthesis of ZnO 1D nanomaterials is possible. How-ever, more work is needed to figure out quantitativeinformation about the interfacial energy and morphol-ogy evolution.4. ApplicationsThere have been many reports on the applicationof ZnO 1D nanomaterials. Similar to the previoussection, we again list the representative applicationsdemonstrated in literature [9–30, 135–138]. In thissection, we focus on the work carried out in our grouponly.4.1 Photocatalytic decomposition of organic dyes inwater[15]ZnO nanoplatelets have been investigated for theapplicability in photodegrading organic dye of eosinB. Figure 12(a) shows the absorption spectra of aque-ous solutions of eosin B (initial concentration of1.0£10¡5 mol/L, 30 mL) in the presence of ZnOnanoplatelets III (2 mg) under exposure to UV lightfor a different period of time, where the absorption peak corresponding to the eosin B molecules at517 nm decreases in intensity rapidly with the ex-tension of the exposure time, and disappears almostcompletely after about 30 min. No new absorptionpeaks appear in the whole spectrum.In a further experiment, the solution of eosin B ex-perienced a series of experimental conditions: (a) withZnO nanoplatelets III (2 mg), in the dark; (b) withoutcatalyst, with UV light; (c) with ZnO nanorods (2 mg)and UV light; (d) with ZnO nanoplatelets I (2 mg)and V light; (e) with ZnO nanoplatelets II (2 mg)and UV light; (f) with ZnO nanoplatelets III (2 mg)and UV light. The results are illustrated in Fig.12(b).Under the experimental conditions of (a) and (b),the photocatalytic effect on eosin degradation withoutcatalysts or without exposure to UV light is similarlylow, where only a slight decrease in the concentrationof eosin B was detected. However, from the data incurves c (ZnO nanorods), d (ZnO nanoplatelets I),e (ZnO nanoplatelets II), and f (ZnO nanoplateletsIII), it is apparent that under identical conditionswith exposure to UV light, the ZnO nanoplateletsdemonstrate much greater photocatalytic activity indegradating eosin B than ZnO nanorods do. It isworth noting that thinner nanoplatelets become, thehigher their performance. Such trend in the pho-tocatalytic activity of the ZnO nanomaterials is inline with the higher surface to volume ratio of thenanoplatelets than the nanorods, and the thinnernanoplatelets compared to the thicker ones. Anotherpossibility to account for the better performance inthe photocatalytic activity of the ZnO nanoplateletsthan the ZnO nanorods is surface dependent photo-catalytic efficiency, as has been demonstrated that Ptnanocrystals[139]. From the structural point of view,the §(1000) planes of ZnO are polar ones. Theypossibly adsorb the organic dyes more strongly thanother nonpolar planes such as the f10¯ 10g planes inthe nanorods and nanoparticles. Since organic dyescould be degraded only if they adsorb on the surfaceof the catalyst, the §(1000) planes of ZnO may dom-inate, to some extent, the photochemical processes indegrading the organic dyes. We notice that the perfor-mance of the ZnO nanoplatelets is much better thancommercial Degussa P25 titania, and the thinnestnanoplatelets perform better than to the porous ZnSnanoparticles.4.2 Enhanced field emission property[20]ZnO nanorods have been investigated as electronfield emitters (FE) due to their high-temperaturechemical stability. The material morphology effecton FE was extensively explored[140–142]. However, ef-fective pathways for FE improvement have remainedundiscovered, and there have been many obstacles toachieve ideal results. On one hand, producing the tipradii smaller than 5 nm is still a challenge. On theother hand, even though narrow diameter tips couldbe produced, their FE performance still has a roomfor improvement. Therefore, one has to find alterna-tive pathways to reach good FE parameters.We adopted two alternative approaches to dra-matically improve FE performance of ZnO nanorodarrays. The first one is decreasing the tip radiiof ZnO nanorods by controlling their growth, andthe second one is lowering ZnO work function bydecorating nanorods with metal nanoparticles hav-ing a lower work function than that of ZnO. Fig-ure 13(a) shows the comparative J-E curves of ZnOnanorods, nanocones, and nanorods decorated withPt, and Ag nanoparticles, at a separation between the emitting tips and the anode of 100 ¹m. Theturn-on fields are measured as about 9.2, 2.6, 1.9, and3.7 V/¹m, respectively, for these four samples. Theachieved current densities of ZnO nanocones and ZnOnanorods decorated with Pt nanoparticles are largerthan 1 mA/cm2. By plotting ln(J=E2) vs 1/E andfitting the data to Fowler-Nordheim relation[143], weget linear curves, as shown in Fig.13(b). We mayconclude that a decrease in ZnO nanorod tip radiusand decoration with Pt and Ag nanoparticles drasti-cally enhance the FE performances. We also extractthe field-enhancement factors as 120 and 2600, forZnO nanorods and nanocones, respectively, which aresmaller than the calculated FE factors, as expected.5. ConclusionsIn this article we have introduced growth mech-anisms, synthesis methods, and applications of ZnO1D nanomaterials, mainly focusing on the work byour group. We briefly summarize the main contentsof this work as follows.(1) There are three significant growth mechanismsfor synthesizing ZnO 1D nanomaterials, VLS, VS, andASG mechanism. The VLS and VL mechanisms arevapor phase growth processes, where at least a va-por phase is involved. The difference between thesetwo mechanisms is that the former needs a catalyst todirect the growth, whereas the latter does not. Sizecontrol is easier for the former, whereas morphologycontrol is better by the latter. The ASG mechanism isa solution growth process, which can also be called asolution-solid mechanism, or if a catalyst is involved, asolution-solid (liquid)-solid mechanism is possible. InASG mechanism, morphology control is more complexsince many experimental parameters are involved.(2) A series of ZnO 1D nanomaterials have beensynthesized using these three growth mechanisms andunderlying crystal growth theories. In this arti-cle, we have demonstrated how to controllably syn-thesize ZnO nanomaterials with desired morpholo-gies by tuning experimental parameters, such as thesupersaturation ratio in vapor phase growth andionic strength in solution phase growth. By ourmethods, nanorods, nanowires, nanotubes,
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