effects of metal underlayer grain size on carbon nanotube growth

Effects of Metal Underlayer Grain Size on Carbon Nanotube Growth

David P. Burt,† W. Murray Whyte,‡ John M. R. Weaver,† Andrew Glidle,†

Jonathan P. Edgeworth,§ Julie V. Macpherson,*,§ and Phillip S. Dobson*,†,‡

Departments of Electronics and Electrical Engineering, and Mechanical Engineering, UniVersity of Glasgow,Glasgow, U.K., and Department of Chemistry, UniVersity of Warwick, CoVentry, U.K.

ReceiVed: March 9, 2009; ReVised Manuscript ReceiVed: July 12, 2009

In this paper we demonstrate that the nucleation density of single-walled carbon nanotubes (SWNTs), formedby thermal catalytic chemical vapor deposition, strongly depends on the grain size of Al underlayers coveredwith a native oxide (Al/Al2O3). By varying the substrate temperature during Al sputter deposition it waspossible to investigate the effect of Al grain size on growth without inducing changes in the underlayerthickness, surface chemistry, or any other growth parameter. The resulting SWNT growth structures rangedfrom low-density 2D nanotube networks that lay across the surface of the substrate to high density 3D nucleationwhich gave rise to vertical “forest” growth. The height of the SWNT “forest” was observed to increase withincreasing Al deposition temperature as follows, 200 > 100 > 60 > 20 °C on Si/Al but in the order 100 >200 > 60 > 20 °C on SiO2/Al substrates for fixed growth conditions. The differences in the SWNT growthtrends on Si and SiO2 substrates are believed to be due to the existence of an optimal Al/Al2O3 underlayergrain size for the formation of active catalytic nanoparticles, with larger Al/Al2O3 grains forming on SiO2

than Si at a fixed substrate temperature. Numerous surface analysis techniques including AFM, XPS, FESEM,TEM, and Raman spectroscopy have been employed to ascertain that the observed changes in nanotube growthfor this system are related primarily to changes in underlayer morphology.

Introduction

Carbon nanotubes exhibit remarkable properties making themsuitable for a wide range of applications such as chemicalsensors,1 microelectromechanical devices,2 composite materials,3

tips in scanned probe microscopy,4 and as promising materialsfor future electronic devices.5 However, the full potential ofcarbon nanotubes will only be realized if the fundamentalprocesses responsible for their growth are understood andcontrolled.

Carbon nanotube synthesis by catalyzed chemical vapordeposition (cCVD) can produce high-quality carbon nanotubeswith few contaminants or defects and also allows for spatialpositioning, which is important for device fabrication.6 Recently,considerable efforts to understand the mechanisms that underpincarbon nanotube synthesis via cCVD have been made, resultingin the identification of many influential factors. These includethe growth times,7 furnace temperature,7,8 reaction chamberpressure,9 feedstock composition,9,10 gas flow,11 catalyst (ma-terial,12 thickness,9,13 and chemical state14,15), and substrateporosity.16 Numerous different process conditions have beenreported that give rise to carbon nanotube growth,17 but the largenumber of process variables has made it difficult to isolatespecific contributions to the growth mechanism.

It is widely recognized that the formation of discrete catalyticmetal nanoparticles (typically Fe, Ni, or Co) is critical forcontrolled carbon nanotube cCVD growth18 and that in manycases the carbon nanotube diameter is largely governed by thecatalyst particle size.19,20 Catalytic nanoparticles can be created

by annealing thin (∼nm) metal films formed by plasmadeposition techniques such as dc or rf sputtering, thermal orelectron beam evaporation;12 or derived from solutions contain-ing a dispersion of catalytic particles in the form of ionic salts(e.g., ferric nitrate),21 dendrimer macromolecules,22 or artificialferritin.23 However, the catalyst particle size and dispersion canalso be influenced by the material on which it is deposited.

Recently, there have been numerous studies that reportchanges in the carbon nanotube growth (e.g., nanotube structureand nucleation density) brought about by the inclusion of catalystsupport materials or metallic underlayers (e.g., Cr, Ir, W, Ta,and Ti), positioned between the bulk substrate (usually Si) andcatalytic nanoparticles.24-27 Interestingly, aluminum (with anative surface oxide) (Al/Al2O3) or alumina (Al2O3) filmsdeposited by high-vacuum processes or from solution have beenparticularly effective at promoting the growth of carbon nano-tubes on substrates such as Si, SiO2 or metals (e.g., Au, Ag,W, NiCr, steel) for a range of catalytic materials (e.g., Co, Ni,Fe).25,28-33 Various explanations have been offered to accountfor the remarkable efficacy of Al/Al2O3 in promoting carbonnanotube growth. Notable suggestions are that the Al/Al2O3 actsas a diffusion barrier preventing catalyst from being eliminatedto the substrate28,32,34-39 or that the key property of the oxidizedaluminum is its thickness,25,28,34,40,41 morphology,14,25,28,40,42-44

reactivity,14,35,39,40,45 or surface energy/wettability.44,46-49 Thesestudies indicate that the action of the buffer layer extends beyondtopographic effects alone. However, since the effect of surfacemorphology has not been investigated in isolation, the influenceof underlayers of both different chemical composition andstructure on carbon nanotube growth does not allow topographiceffects to be readily ascertained.

To address this, we investigate the fundamental effect of theunderlayer surface morphology while maintaining the samesurface chemistry. By employing a variety of surface analysis

* To whom corespondence should be addressed. E-mail: [email protected] (P.S.D.); [email protected] (J.V.M.).

† Department of Electronics and Electrical Engineering, University ofGlasgow.

‡ Department of Mechanical Engineering, University of Glasgow.§ Department of Chemistry, University of Warwick.

J. Phys. Chem. C 2009, 113, 15133–15139 15133

10.1021/jp902117g CCC: $40.75 2009 American Chemical SocietyPublished on Web 08/05/2009

techniques, we demonstrate that the nucleation yield of carbonnanotubes produced by thermal cCVD can be strongly influ-enced by the morphology of pregrowth Al/Al2O3 grains on bothSi and SiO2 substrates.

Experimental Section

Substrate Preparation. Three inch diameter n-type (phos-phorus) silicon (100) wafers (Compart Technology Ltd., Peter-borough, U.K., resistivity 1-10 Ω.cm) with a 380 ( 50 µmthickness were used throughout. SiO2 barrier layers were formedby wet thermal oxidation (Scottish Microelectronic Centre,Edinburgh) to give an oxide thickness of 300 nm. Al films weredeposited on Si and SiO2 substrates using a Plassys (MP 900S)sputtering tool with a base pressure of ∼5 × 10-8 mbar, Arpressure of 20 µbar, and a current of 1 A for 1 min. This gaverise to a consistent Al thickness of 10 nm at 20 °C.50 The Algrain size was varied by controlling the substrate temperature(20, 60, 100, 200 °C) while all other parameters remainedconstant. From this section onward, the substrate temperatureduring Al sputter deposition is indicated in parentheses, e.g.,Al(20 °C). The exposure of the Al film to atmosphere after Aldeposition ensured that the surface of the Al was oxidized;crucial for successful carbon nanotube growth.36 A 0.5 nm thickcatalytic film of Ni was deposited on samples using a Plassys(MEB 400S) electron beam evaporator at a deposition rate of0.01 nm s-1, monitored by a quartz crystal sensor (chamberpressure ∼1 × 10-7 mbar).

Instrumentation. Field emission scanning electron micros-copy (FESEM) was conducted under high-vacuum conditionsusing either a Hitachi S4700 microscope with 1 kV acceleratingvoltage and secondary electron detection or a FEI NovaNanoSEM 630 with an accelerating voltage of 10 kV withsecondary (through lens detector) and backscattered electrondetection. Transmission electron microscopy (TEM) was carriedout using a Energy Filtered FEI T20 Technai electron micro-scope operating at 200 keV with an emission current of ∼6 µAfrom a LaB6 filament. Raman spectroscopy was carried out usinga LabRam INV Raman spectrometer (Jobin Yvon Ltd.) with a633 nm HeNe laser at a magnification of 50× over a samplingspot diameter ∼20 µm. For each site measured, spectra wereacquired for 30 s and averaged over two cycles. X-rayphotoelectron spectroscopy (XPS) was carried out at theNational Centre for Electron Spectroscopy and Surface Analysis(NCESS), Daresbury Laboratory, UK. A Scienta ESCA 300photoelectron spectrometer with an analyzer pass energy set to150 eV and slit width of 0.8 mm was employed. A flood gungenerating low-energy electrons (4 eV) was used to reducesubstrate charging when required. A rotating anode X-ray sourceprovided monochromated Al KR radiation (1486.6 eV) andelectrons were detected at a takeoff angle of 90° (normal to thesubstrate). To investigate substrate surface structure, atomicforce microscopy (AFM) was carried out using a DigitalInstruments (Veeco) MultiMode AFM with a Nanoscope IIIacontroller, Extender Module and J scanner. AFM measurementswere performed in air at room temperature (20 °C) using thetapping mode operation with Veeco FESP (resonance frequency≈75 kHz, force constant ≈3 N/m) silicon probes (radius ofcurvature <10 nm). Probes were changed regularly to minimizeeffects of tip convolution.51,52

Carbon Nanotube Synthesis and Thermal Annealing.Carbon nanotubes were synthesized by alcohol assisted thermalchemical vapor deposition (AACVD)53,54 in a tube furnace(Model STF55666C, Lindberg/Blue-M) with a four inch diam-eter quartz tube at atmospheric pressure. All substrates were

∼1 × 1 cm squares and positioned on a 3 in. oxidized siliconwafer supported on a quartz boat. The AACVD growth schemeparameters are outlined in Table 1. Gas flow rates of Ar(99.999%, BOC Gases) and H2 (99.995%, BOC Gases) wereregulated by mass flow controllers (Type 1179, MKS Instru-ments Inc.). The concentration of the carbon source (absoluteethanol 99.8+%, Fisher Scientific) was controlled by passingAr through ethanol contained in a bubbler positioned in athermostatic bath (20 °C). The ethanol temperature inside thebubbler was monitored independently using a liquid thermometer.

Thermal annealing of substrates employed in the XPSmeasurements was carried out according to the scheme shownin Table 1 but without the introduction of ethanol vapor duringstage 4. Samples were stored under ambient conditions prior toand post nanotube synthesis and annealing.

Results and Discussion

Adatoms and nanoparticles become mobile at temperaturessignificantly lower than their bulk material melting points.9,55

The surface particle diffusion rate and hence particle sizedistribution is also highly dependent on the thickness of thedeposited metal.56 Therefore, to investigate the effects ofunderlayer morphology; control of the substrate temperatureduring Al deposition allowed samples with different Al grainsizes to be prepared. Other parameters associated with carbonnanotube growth such as the amount of Al underlayer deposited,catalyst material, and growth parameters were kept constantthroughout this study. Figure 1 shows AFM tapping mode heightimages (1 × 1 µm) of Al grains deposited on (i-iv) Si and(v-viii) SiO2 substrates held at different temperatures duringAl deposition. Figure 1(i-iv) and (v-viii) show that the grainsize of the Al increased when higher substrate temperatures wereemployed. Figure 2 shows FESEM images of carbon nanotubesgrown on (i-iv) Si and (v-viii) SiO2 substrates with an Alunderlayer deposited at substrate temperatures of 20, 60, 100,and 200 °C coated with 0.5 nm Ni deposited by electron beamevaporation. The center of each sample was scratched to exposethe underlying substrate and the carbon nanotube growththickness was measured using a viewing angle of 45°. Figure2(i-iv) shows that on Si substrates the carbon nanotube growthregime switches from a “spaghetti-like” surface network formedon Al(20 °C) to three-dimensional “forest” growth on Al(g60°C) as a result of the “crowding effect” previously describedby Dai et al.6 The “crowding effect” refers to the ordering ofcarbon nanotubes in the direction perpendicular to the substratesurface brought about by a high density of nanotube nucleationand strong van der Waals forces of attraction between thenanotubes. In Figure 2(i-iv), it is clear that the height of thecarbon nanotube forest continues to increase on Si substrateswith Al deposited at higher substrate temperatures. The forestheight does not necessarily give a direct measurement of the

TABLE 1: Carbon Nanotube cCVD Growth SystemParameters

stagetime(min)

furnacetemperature (°C) process gases (flow rate)

1. warming up 10 20-700 Ar (1920 sccm)2. warming up 15 700-800 Ar (1920 sccm)3. anneal 5 800 Ar (1920 sccm,

H2 (80 sccm)4. growth 20 800 Ar (1200 sccm),

H2 (80 sccm),Ar/EtOH(vap) (720 sccm)

5. cooling down 180+ 800-20 Ar (1920 sccm)

15134 J. Phys. Chem. C, Vol. 113, No. 34, 2009 Burt et al.

carbon nanotube length, but instead we believe it to be a productof the nanotube alignment, which in turn relates to the nanotubenucleation density, i.e., carbon nanotube alignment in highnucleation density nanotube forests has more order thanlow nucleation density forests that comprise meandering nano-tubes. This is also considered in recent studies carried out bySugime et al.57

Figure 2(v-viii) shows typical nanotube growth results onSiO2 with Al deposited on substrates at temperatures of 20, 60,100, and 200 °C. Again, the influence of substrate temperatureduring Al deposition on carbon nanotube growth is clear.Throughout this study, an increase in the growth height for SiO2

samples with underlayers of Al(20 °C) to Al(100 °C), wasobserved. However, the height of the nanotube growth on Al(200°C) was consistently (10 repetitions) lower than substrates withAl(100 °C) underlayers. The same trend in nanotube forestheight could also be visually observed by the blackening of thesubstrate after growth, indicating the presence of carbonnanotubes. We believe further increases in the Al grain size onSiO2 would lead to subsequent reductions in nanotube growthyield. However, this study has been restricted to a maximumsubstrate heater temperature of 200 °C due to limitations of thesputter coater.

To determine the type of nanotubes present on the samples,high-resolution TEM was employed as shown in Figure 3.Samples were prepared by contacting copper TEM grids withthe nanotube covered substrates. More than 20 regions contain-ing carbon nanotubes were studied, and the nanotubes on bothSi and SiO2 substrates were identified as bundles of SWNTswith perhaps some double-walled carbon nanotubes (DWNTs)present. Some amorphous carbon was present in regions, andthe nanotubes appeared to have defects and/or kinks, althoughthis may have been caused by sample preparation. The diametersof 38 nanotubes of typical appearance were measured usingTEM and ranged from 1.6 to 4.0 nm. No consistent change in

the nanotube diameter was observed with changes in the Algrain size. The mean nanotube diameters on SiO2 substrateswith Al deposited at 20, 60, 100, and 200 °C were 2.5 (standarddeviation (SD) 0.5 nm), 2.7 (SD 0.2 nm), 3.0 (SD 0.5 nm), and2.9 nm (SD 0.2 nm), respectively. There was also no appreciabledifference between nanotubes grown on different substratematerials, as the mean nanotube diameter was 2.9 (SD 0.6 nm,13 samples) and 2.8 nm (SD 0.4 nm, 25 samples) on Si andSiO2 substrates, respectively.

Figure 4 shows Raman spectra of carbon nanotubes grownon Si with Al(20 °C) and 0.5 nm Ni deposited. The peaks inthe main spectrum due to the substrate are identified with ’S’.The inset shows the radial breathing mode (RBM) region moreclearly; identifying peaks at 104, 139, and 187 cm-1. The modefrequency, ωRBM, of an isolated SWNT on a Si/SiO2 substraterelates to the nanotube diameter,58 dt, by ωRBM ) R/dt, where R) 248 cm-1. Applying this relationship gives nanotube diam-eters in the range of 1.3-2.4 nm. The SWNT diameter range,determined from the Raman signal, was close to the valuesmeasured by TEM. The main spectrum in Figure 4 shows aclear G peak at 1589 cm-1 and a relatively strong ‘disorder’,D, band at 1325 cm-1. A significant D-band region on samples

Figure 1. AFM tapping mode height images (same scale) of 10 nm Al sputter deposited on (i-iv) Si (native oxide) and (v-viii) SiO2 (300 nmthermal oxide) at different substrate temperatures (as shown).

Figure 2. FESEM images of SWNT growth on (i-iv) Si (same scale) and (v-viii) SiO2 (same scale) with 10 nm Al deposited at substratetemperatures of (i,v) 20, (ii,vi) 60, (iii,vii) 100, and (iv,viii) 200 °C. Catalyst (Ni) thickness is 0.5 nm evaporated at a 0.01 nm/s deposition rate.The viewing angle is 45°.

Figure 3. TEM image of SWNTs grown on SiO2 (300 nm) with 10nm Al(100 °C) and 0.5 nm Ni.

Effects of Metal Underlayers on Carbon Nanotube Growth J. Phys. Chem. C, Vol. 113, No. 34, 2009 15135

containing isolated nanotubes is normally attributed to thepresence of sp3-hybridized carbon bonds due to either defectsalong the length of the tube, exposed ends of the nanotube oramorphous carbon within the sampling region.12 This agreeswith our TEM observations that showed kinks in the nanotubesthat are indicative of defects (e.g., 5-7 pairs59) in the nanotubestructure. The overall intensity of the Raman signal decreasedas the SWNT forest height increased, possibly due to reabsorp-tion of the Raman scattered light within the nanotube layer.

Table 2 summarizes the mean Al grain size, nanotube forestheight and Raman D/G ratio for the samples reported herein.The mean Al grain size was consistently smaller when depositedon Si than on SiO2 for a given substrate temperature. The meangrain size, for each of the AFM images in Figure 1, wascalculated by manual identification, and then labeling andcounting the grains using ImageJ software.60 Despite somevariation in the height of the nanotube forests that occurredbetween growth cycles for the Si and SiO2 substrates, the growthyield (height) consistently followed the order: 200 > 100 > 60> 20 °C for Si substrates and 100 > 200 > 60 > 20 °C for SiO2

substrates. Importantly, the nanotube forest height trend on SiO2

substrates shows that an Al grain size of 2564 nm2 (Al(100 °C)produced the highest forest growth but also identified that arelatively small increase in the Al grain size (2681 nm2 forAl(200 °C) can have a significant detrimental effect on the forestheight. It has been reported that deceleration of the nanotubeforest growth rate (assuming root growth) can result due toresistance to feedstock diffusion through the forest structure tothe catalyst at the substrate.61,62 However, this mechanism appliesto tall (∼mm) SWNT forests and should not significantlyinfluence the nanotube growth mechanism for the forests withthe heights reported herein.62 Instead, we believe the nucleationdensity of the nanotubes influences the forest height by structural

(crowding) effects rather than preventing the diffusion offeedstock (see below).

The D/G ratio is indicative of the quality and purity of thedeposited carbon nanotubes. We have found that for sparse 2Dnetworks grown using our system D/G ratios ∼0.16 were obtained(not shown), which identified that SWNT with low defects couldbe produced using the conditions described herein. The TEMsample preparation method is not capable of transferring nanotubesfrom samples of such low density, thus preventing structuralanalysis of these nanotubes. However, for the samples shown inFigure 2 with high density 2D or 3D growth, high D/G ratios ∼0.5were typical (Table 2). TEM revealed no obvious difference inthe small amount of amorphous carbon present in the samples(Figure 2). Since TEM identified that the amount of amorphouscarbon was low for all samples with a high D/G ratio, this suggeststhat defects such as 5-7 pairs59 or vacancies in the nanotubestructure are the likely source of the large D band.63

To investigate chemical changes in the substrates that mightinfluence the nanotube growth characteristics, we examined thesurface composition with XPS before and after subjecting thesamples to the heating conditions associated with nanotube growth(Table 1) in the absence of the carbon feedstock to preventthe formation of nanotubes. To account for substrate charging, thebinding energy for each region was shifted (∼3 eV) to align thecarbon (1s) peak to a reference of 285.0 eV.64 The XPS signal forthe Al region recorded on a Si/SiO2/Al(100 °C) sample is shownin Figure 5(i). It can be seen that Al exists in both the oxide andmetallic form before thermal annealing. This is due to the surfaceoxide that forms in ambient conditions and the metallic Al in thecore of the Al grains. The loss of the Al 2p3/2 metal peak andenhancement of the Al 2p3/2 oxide peak after annealing (red line)occurred for this and all other cases. Hence, Al was completelyoxidized to Al2O3 by thermal annealing, probably due to residualoxygen present in the process tube.34,65 This is supported by theunderlayer film becoming transparent after thermal annealing,suggesting the absence of aluminum in the metallic form. Thedegree of Al oxidation before annealing showed slight variationbetween samples but did not reflect the growth trends observedherein. We believe this variation in Al oxidation to be due to thereduction in vacuum (∼1 × 10-7 mbar) in the sputtering toolchamber caused by outgassing as a result of substrate holderheating.

Figure 4. Raman microscopy spectra of SWNTs grown on Si/Al(20°C) with 0.5 nm Ni. The inset shows the RBM region in more detail.

TABLE 2: Summary of Mean Al Grain Size, NanotubeForest Height, and Raman D/G Ratio for SWNTs Grown onSiO2 and Si Substrates Sputter-Coated with 10 nm AlDeposited at 20, 60, 100, and 200 °C

substratedeposition

temperature (°C)mean grainsize (nm2)

forestheight (µm) D/G

SiO2/Al(20 °C) 20 1704 0.0 0.40SiO2/Al(60 °C) 60 1845 6.0 0.46SiO2/Al(100 °C) 100 2564 18.0 0.53SiO2/Al(200 °C) 200 2681 9.0 0.54Si/Al(20 °C) 20 733 0.0 0.43Si/Al(60 °C) 60 1087 2.6 0.45Si/Al(100 °C) 100 1185 5.1 0.42Si/Al(200 °C) 200 1613 9.6 0.45

Figure 5. XPS signal for the (i) Al 2p region of Si/SiO2/Al(100)substrates and the (ii) Ni 2p region of Si/SiO2/Al(100)/Ni substratesthat were annealed (red) and not annealed (black). (iii) XPS ratio ofthe integrated Ni (2p) region annealed/not annealed for Si/Ni, Si/SiO2/Ni, Si/Al(100)/Ni, and Si/SiO2/Al(100)/Ni substrates. Substrates wereannealed at 800 °C in Ar/H2(4%) for 25 min.


An increase in the Si 2p peak amplitude (not shown) withincreasing temperature of Al deposition on both Si and SiO2

(oxidized silicon, binding energy +3.9 eV relative to Si peak)is ascribed to the presence of gaps or thin regions in theunderlayer film as a result of changes in the Al/Al2O3 grainsize (see Figure 1). Figure 5(ii) shows a plot of XPS data forthe Ni 2p region measured on a Si/SiO2/Al(100 °C)/Ni substratebefore (black line) and after (red line) thermal annealing. Thereis a clear reduction in the amplitude of the Ni signal as a resultof the annealing process. The peak separation of 17.4 eVbetween the Ni (2p3/2) and Ni(2p1/2) peaks agree closely toliterature values (17.3 eV).66 Shake-up peaks found at bindingenergy values ∼6-7 eV higher than both the Ni (2p3/2) and Ni(2p1/2) peaks are especially prominent, probably due to the smallsize of the Ni catalyst particles generated in this study.67

To investigate the behavior of Ni on Al/Al2O3 and the underlyingsubstrate material (Si and SiO2), the ratio of the areas of the Ni 2ppeaks after/before thermal annealing was determined for Si/Ni,SiO2/Ni, Si/Al/Al2O3/Ni, and SiO2/Al/Al2O3/Ni substrates, and issummarized in Figure 5(iii). For all substrates, the annealing processdecreased the total amount of Ni present. This is believed to bedue to two contributing factors, diffusion of the catalyst into theunderlying surface and/or Ni evaporation due to the high temper-atures employed during the anneal step.18,68 The annealing processresulted in a large decrease in the Ni (2p) signal on Si substrates(reduced to 3.5% of the preanneal value), which is probably dueto evaporation and significant losses by Ni diffusion into the Si.Native SiO2 will not act as a diffusion barrier at elevatedtemperatures (>600 °C).38 In contrast, a smaller decrease in the Ni(2p) signal was observed on SiO2 (reduced to 44%) probably dueonly to evaporation, consistent with Ni not reacting with SiO2 andthicker (>4 nm) SiO2 films being a good diffusion barrier.38,69 Thisobservation is not surprising given the diffusion coefficients of Niin Si (∼10-7 cm2 s-1)70 and SiO2 (∼10-17 cm2 s-1)71 at 800 °Cand that 300 nm thick SiO2 films were employed herein. Si/Al/Al2O3 and SiO2/Al/Al2O3 substrates also show Ni losses afterannealing but significantly less than Si substrates with a nativeoxide. The small discrepancy between Ni losses on Si/Al/Al2O3

(reduced to 23%) and SiO2/Al/Al2O3 (reduced to 29%) substratesis believed to be due to Ni removal by diffusion into the regionsof Si between the Al grains. Since SiO2/Al/Al2O3 and SiO2

substrates are both good diffusion barriers, the difference in theNi loss is likely to be due to different Ni evaporation rates fromeach substrate. This could originate from the size (specifically thesurface area-to-volume ratio) of the Ni nanoparticles on eachsubstrate. For example, due to weaker attractive forces actingbetween metallic nanoparticles and SiO2 than Al2O3,46 larger Niparticles form on SiO2 than SiO2/Al/Al2O3.

With knowledge of the substrate preparation and XPS surfaceanalysis, AFM was employed to investigate surface morphologyand identify the Ni nanoparticle distribution on substrates after thethermal anneal process. Figure 6 shows AFM height images ofannealed Ni (0.5 nm) deposited on (i) Si and (ii) SiO2. Annealingproduced discrete nanoparticles on the (i) Si substrate, whereasclusters of nanoparticles formed on the (ii) SiO2 surface obscuringthe underlying substrate. The particles present on these surfacesare believed to be Ni, as they were absent on annealed Si and SiO2

substrates without a layer of deposited Ni (not shown). After theannealing process there were significantly more Ni particles onthe SiO2 substrate than the Si substrate, consistent with the XPSresults presented herein and previous studies reporting considerablecatalyst losses on silicon substrates.68

AFM measurements of Al grains on Si and SiO2 substratesshowed no appreciable change in the grain structure after the

thermal annealing process, indicative of a thermally stable substratesuch as Al2O3 (mp ≈ 2000 °C) rather than metallic Al (mp ≈ 660°C). We investigated the Ni nanoparticle formation on Al-coatedsubstrates using AFM. Due to the inherent roughness of the Alfilm, amplitude (error) images are presented to accentuate the grainboundaries and distinguish the features of the surface morphologydue to the presence of Ni. For clarity, we present AFM images forannealed Ni on the substrate with the largest Al grains. Figure 7(i)shows an AFM amplitude image of an annealed SiO2/Al(200 °C)substrate, without a deposited Ni catalyst layer. An AFM amplitudeimage of an annealed SiO2/Al(200 °C)/Ni (0.5 nm) substrate isshown in Figure 7(ii). The formation of smaller grains on top ofthe Al2O3 grains is clear and was only observed on annealedsubstrates with a deposited Ni layer. Hence, these particles are mostlikely Ni agglomerates that have formed during the annealing stage.Cross-sectional analysis (in height image, not shown) of thenanoparticles on top of the Al2O3 grains showed them to be ∼2nm in height, which is on the order of the nanotube diametersreported herein.

Panels (iii) and (iv) of Figure 7 show FESEM images of thesample used in Figure 7(ii) and identify the Al2O3 as isolated islands(on the SiO2 surface) that were not completely resolved by theAFM tip due to convolution effects.51,52 The backscat-tered signal is sensitive to the atomic mass of the element andtherefore provides information about the elemental compositionof the sample.14,72 Particles giving high contrast are observable ontop of the Al2O3 grains, as identified in Figure 7(iv). Smaller

Figure 6. AFM tapping mode height images of 0.5 nm Ni on (i) Siand (ii) SiO2 substrates after annealing at 800 °C in Ar/H2 (4%) for 25min.

Figure 7. AFM tapping mode amplitude images of Al(200 °C) onSiO2 (i) without and (ii) with 0.5 nm Ni deposited. FESEM images ofthe same region of a SiO2 substrate with an Al(200 °C) film and 0.5nm Ni using a (iii) secondary electron through lens detector and (iv)backscattered electron detector. Ni nanoparticles on Al grains areidentified by their high contrast and circled in (iv). All samples wereannealed at 800 °C in Ar/H2 (4%) for 25 min.


particles coat the surface of the Al2O3 grains, but due to thedetection limit of the SEM backscattered electron signal we wereunable to resolve these particles, thus preventing a statisticalanalysis. However, we believe them to be Ni and the samenanoparticles identified in the AFM image in Figure 7(ii). Thesedeductions are supported by the detection of backscattered electronsfrom annealed Al/Al2O3 samples without Ni deposited (samesample as in Figure 7(i)) that gave poor image contrast (not shown).There was no evidence of Ni between the Al/Al2O3 grains in shownin Figure 7(iv) (i.e., on the SiO2 surface), which is indicative of aminimum energy configuration existing for the Ni particles on theAl/Al2O3 grains. This is similar to the recent findings of Matteviet al. who described a stronger interaction occurring betweenFe-Al2O3 than Fe-SiO2.46

We have demonstrated that the as-deposited Al/Al2O3 grain sizestrongly influences the height of the nanotube forest, which relatesto the nucleation density of the SWNTs.57 To account for thechange in nanotube nucleation density, we believe the underlayergrain size influences the number of active catalyst nanoparticleswhich reside on its surface. Figure 8 shows schematics of aproposed mechanism, based on changes in the Ni nanoparticle sizedistribution brought about by variations in the Al/Al2O3 grain size,which can account for the observed nanotube growth trends. Wehave identified that the Ni nanoparticles are sited on the surface ofthe oxidized Al grains (vide infra). Hence, we believe that the Al/Al2O3 grain size and the interaction between the Al2O3 and Nistrongly influence the Ni nanoparticle size and hence their catalyticactivity.18 Since it is energetically favorable for the Ni to resideon the Al2O3 surface,46 the amount of Ni is largely confined to theAl/Al2O3 grain on which it was deposited and is unlikely to crossgrain boundaries. Therefore, the amount of Ni on each Al/Al2O3

grain that is available for catalyst particle agglomeration duringthe anneal stage is highly likely to relate to the Al/Al2O3 grainsize. Figure 8(i-iii) shows that as the Al/Al2O3 grain size increases(from left to middle), more catalyst material is available foragglomeration, which increases the average size of the Ni catalyst

particles to a size where they become active for SWNT growth.18

This increase in the number of active particles results in an increasein the nanotube nucleation density, as observed experimentally forgrowth on Al/Al2O3 on Si for deposition temperatures of 20-200°C and on SiO2 for deposition temperatures of 20-100 °C. Forlarger Al/Al2O3 grains (e.g., Al/Al2O3 (200 °C) on SiO2), there ismore Ni available which gives rise to Ni nanoparticles (right) thatare larger and hence inactive. Therefore, the presence of feweractive Ni nanoparticles results in a lower nanotube nucleationdensity and due to the “crowding effect”,6 a reduction in the heightof the nanotube forest is observed.57

Conclusions

We have demonstrated that the grain size of thin Al/Al2O3 filmsis influenced by the substrate material (Si or SiO2) but can also becontrolled by the surface temperature during Al sputter deposition,without changing the mean film thickness. The Al/Al2O3 grain sizewas consistently larger on SiO2 than Si surfaces for a givendeposition temperature, probably due to wetting characteristics.XPS surface analysis proved the Al/Al2O3 films employed hereinto be effective barriers to Ni diffusion during subsequent thermalCVD. Analysis of the catalytic surfaces by the combination of XPS,AFM, and FESEM techniques confirmed the presence of Ninanoparticles on the surface of the Al/Al2O3 grains after thermalannealing. Although chemically the same, the Al/Al2O3 grain sizewas shown to have a marked effect on the SWNT nucleationdensity formed by thermal AACVD, without effecting nanotubediameter. The variation in the nanotube nucleation density wasreflected in the SWNT growth morphology; evidenced by a switchfrom ∼2D surface networks to 3D vertically aligned nanotube“forests” as the Al/Al2O3 grain size was increased. The trend inSWNT “forest” height varied with Al deposition temperature asfollows, 200 > 100 > 60 > 20 °C on Si/Al but in the order 100 >200 > 60 > 20 °C on SiO2/Al substrates. It was found that thetrend in SWNT “forest” height correlated with the grain sizeirrespective of the substrate material and that Al deposited on SiO2

at 100 °C, produced the tallest nanotube growth (Table 2). Wepresent a mechanism to explain the observed variation of nanotubenucleation density, which is attributed to a change in the catalystparticle size distribution for the different underlayer morphologiesthat, in turn, governs catalyst particle activity. This study illustratesthat in order to understand the effect of different substrate materialson nanotube growth the morphology of the catalyst support shouldalso be considered.

Acknowledgment. The authors thank the Engineering andPhysical Sciences Research Council (EPSRC number EP/C518276/1) for funding. We acknowledge Dr. Graham Beamson (NCESS)for assistance with XPS measurements, the Daresbury NCESSfacility, and the EPSRC (EP/E025722/1) for associated funding.We thank Mr. Colin How (Department of Physics, University ofGlasgow) for help with TEM and Mrs. Lesley Donaldson (Depart-ment of Electronics and Electrical Engineering, University ofGlasgow) and Dr. Douglas MacIntyre (Department of Electronicsand Electrical Engineering, University of Glasgow) for experi-mental assistance.

References and Notes

(1) (a) See for example; Kong, J.; Franklin, N. R.; Zhou, C.; Chapline,M. G.; Peng, S.; Cho, K.; Dai, H. Science 2000, 287, 622. (b) Dumitrescu, I.;Unwin, P. R.; Wilson, N. R.; Macpherson, J. V. Anal. Chem. 2008, 80,3598.

(2) Hayamizu, Y.; Yamada, T.; Mizuno, K.; Davis, R. C.; Futaba, D. N.;Yumura, M.; Hata, K. Nat. Nanotechnol. 2008, 3, 289.

Figure 8. Schematic (not to scale) showing the effect of Al/Al2O3

underlayer grain size on the formation of active Ni nanoparticles andconsequently SWNTs. (i) A layer of Ni (shown in dark gray) evaporatedon to different sized Al/Al2O3 grains. (ii) After the thermal CVD processthe evaporated Ni agglomerates to form nanoparticles. SWNTs growfrom active catalyst particles that are within a narrow size range. (iii)Nanoparticle activity is highly size dependent.


(3) Thostenson, E. T.; Ren, Z.; Chou, T.-W. Composite Sci. Technol.2001, 61, 1899.

(4) (a) See for example; Dai, H.; Hafner, J. H.; Rinzler, A. G.; Colbert,D. T.; Smalley, R. E. Nature 1996, 384, 147. (b) Wong, S. S.; Harper,J. D.; Lansbury, P. T.; Lieber, C. M. J. Am. Chem. Soc. 1998, 120, 603.

(5) (a) See for example; Yamada, T. Carbon Nanotubes: Science andApplications; Meyyappan, M., Ed.; CRC Press: Boca Raton, FL, 2005; pp163-187. (b) Baughman, R. H.; Zakhidov, A. A.; de Heer, W. A. Science2002, 297, 787.

(6) Dai, H. Acc. Chem. Res. 2002, 35, 1035.(7) Cui, H.; Eres, G.; Howe, J. Y.; Puretkzy, A.; Varela, M.; Geohegan,

D. B.; Lowndes, D. H. Chem. Phys. Lett. 2003, 374, 222.(8) Ducati, C.; Alexandrou, I.; Chhowalla, M.; Amaratunga, G. A. J.;

Robertson, J. J. Appl. Phys. 2002, 92, 3299.(9) Pisana, S.; Cantoro, M.; Parvez, A.; Hofmann, S.; Ferrari, A. C.;

Robertson, J. Physica E 2007, 37, 1.(10) Veronese, G. P.; Rizzoli, R.; Angelucci, R.; Cuffiani, M.; Malferrari,

L.; Montanari, A.; Odorici, F. Physica E 2007, 37, 21.(11) Huang, S.; Cai, X.; Liu, J. J. Am. Chem. Soc. 2003, 125, 5636.(12) Mizuno, K.; Hata, K.; Saito, T.; Ohshima, S.; Yumura, M.; Iijima,

S. J. Phys. Chem. B 2005, 109, 2632.(13) Kakehi, K.; Noda, S.; Maruyama, S.; Yamaguchi, Jpn. J. Appl.

Phys. 2008, 47, 1961.(14) de los Arcos, T.; Garnier, M. G.; Seo, J. W.; Oelhafen, P.;

Thommen, V.; Mathys, D. J. Phys. Chem. B 2004, 108, 7728.(15) de los Arcos, T.; Oelhafen, P.; Thommen, V.; Mathys, D. J. Phys.

Chem. C 2007, 111, 16392.(16) Fan, S.; Chapline, M. G.; Franklin, N. R.; Tombler, T. W.; Cassell,

A. M.; Dai, H. Science 1999, 283, 512.(17) Meyyappan, M. Carbon Nanotubes: Science and Applications;

Meyyappan, M., Ed.; CRC Press: Boca Raton, FL, 2005; pp 99-112.(18) Takagi, D.; Homma, Y.; Hibino, H.; Suzuki, S.; Kobayashi, Y. Nano

Lett. 2006, 6, 2642.(19) Cheung, C. L.; Kurtz, A.; Park, H.; Lieber, C. M. J. Phys. Chem.

B 2002, 106, 2429.(20) Choi, H. C.; Kim, W.; Wang, D.; Dai, H. J. Phys. Chem. B 2002,

106, 12361.(21) Cheung, C. L.; Hafner, J. H.; Odom, T. W.; Kim, K.; Lieber, C. M.

Appl. Phys. Lett. 2000, 76, 3136.(22) Choi, H. C.; Kim, W.; Dunwei, W.; Dai, H. J. Phys. Chem. B 2002,

106, 12361.(23) Li, Y.; Kim, W.; Zhang, Y.; Rolandi, M.; Wang, D.; Dai, H. J.

Phys. Chem. B 2001, 105, 11424.(24) Han, J.-H.; Kim, H. J.; Yang, M.-H.; Yang, C. W.; Yoo, J.-B.;

Park, C.-Y.; Song, Y.-H.; Nam, K.-S. Mater. Sci. Eng. C. 2001, 16, 65.(25) Delzeit, L.; Chen, B.; Cassell, A.; Stevens, R.; Nguyen, C.;

Meyyappan, M. Chem. Phys. Lett. 2001, 348, 368.(26) Ng, H. T.; Chen, B.; Koehne, J. E.; Cassell, A. M.; Li, J.; Han, J.;

Meyyappan, M. J. Phys. Chem. B 2003, 107, 8484.(27) Kabir, M. S.; Morjan, R. E.; Nerushev, O. A.; Lundgren, P.;

Bengtsson, S.; Enokson, P.; Campbell, E. E. B. Nanotechnology 2005, 16,458.

(28) Delzeit, L.; Nguyen, C. V.; Chen, B.; Stevens, R.; Cassell, A.; Han,J.; Meyyappan, M. J. Phys. Chem. B 2002, 106, 5629.

(29) Delzeit, L.; Nguyen, C. V.; Stevens, R. M.; Han, J.; Meyyappan,M. Nanotechnology 2002, 13, 280.

(30) Cassell, A. M.; Raymakers, J. A.; Kong, J.; Dai, H. J. Phys. Chem.B 1999, 103, 6484.

(31) Kong, J.; Soh, H. T.; Cassell, A. M.; Quate, C. F.; Dai, H. Nature1998, 395, 878.

(32) Parthangal, P. M.; Cavicchi, R. E.; Zachariah, M. R. Nanotech-nology 2007, 18, 185605.

(33) Matthews, K. D.; Lemaitre, M. G.; Kim, T.; Chen, H.; Shim, M.;Zuo, J.-M. J. Appl. Phys. 2006, 100, 044309.

(34) Zhang, R. Y.; Amlani, I.; Baker, J.; Tresek, J.; Tsui, R. K. NanoLett. 2003, 3, 731.

(35) de los Arcos, T.; Garnier, M. G.; Oelhafen, P.; Mathys, D.; Seo,J. W.; Domingo, C.; Garcıa-Ramos, J. V.; Sanchez-Cortes, S. Carbon 2004,42, 187.

(36) de los Arcos, T.; Wu, Z. M.; Oelhafen, P. Chem. Phys. Lett. 2003,380, 419.

(37) de los Arcos, T.; Vonau, F.; Garnier, M. G.; Thommen, V.; Boyen,H.-G.; Oelhafen, P.; Duggelin, M.; Mathis, D.; Guggenheim, R. Appl. Phys.Lett. 2002, 80, 2383.

(38) Simmons, J. M.; Nichols, B. M.; Marcus, M. S.; Castellini, O. M.;Hamers, R. J.; Eriksson, M. A. Small 2006, 2, 902.

(39) Lee, H. C.; Alegaonkar, P. S.; Kim, D. Y.; Lee, J. H.; Yoo, J. B.Philos. Mag. Lett. 2007, 87, 767.

(40) Komukai, T.; Aoki, K.; Furuta, H.; Furuta, M.; Oura, K.; Hirao, T.Jap. J. Appl. Phys. 2006, 45, 6043.

(41) In light of these studies, we have found that changing film thicknessalso results in changes in the Al grain size. Hence, the effects of increasingthe Al barrier thickness to prevent catalyst-substrate interactions and theincreased roughness associated with thicker metallic films cannot beindependently addressed.

(42) Hongo, H.; Nihey, F.; Ichihashi, T.; Ochiai, Y.; Yudasaka, M.;Iijima, S. Chem. Phys. Lett. 2003, 380, 158.

(43) Chakrabarti, S.; Kume, H.; Pan, L.; Nagasaka, T.; Nakayama, Y.J. Phys. Chem. C 2007, 111, 1929.

(44) Ward, J. W.; Wei, B. Q.; Ajayan, P. M. Chem. Phys. Lett. 2003,376, 717.

(45) Noda, S.; Hasegawa, K.; Sugime, H.; Kakehi, K.; Zhang, Z.;Maruyama, S.; Yamagichi, Y. Jpn. J. Appl. Phys. 2007, 46, L399.

(46) Mattevi, C.; Wirth, C. T.; Hofmann, S.; Blume, R.; Cantoro, M.;Ducati, C.; Cepek, C.; Knop-Gericke, A.; Milne, S.; Castellarin-Cudia, C.;Dolafi, S.; Goldoni, A.; Schloegl, R.; Robertson, J. J. Phys. Chem. C 2008,112, 12207.

(47) Ohno, H.; Takagi, D.; Yamada, K.; Chiashi, S.; Tokura, A.; Homma,Y. Jpn. J. Appl. Phys. 2008, 47, 1956.

(48) Hongo, H.; Yudasaka, M.; Ichihasi, T.; Nihey, F.; Iijima, S. Chem.Phys. Lett. 2002, 361, 349.

(49) Wang, Y.; Luo, Z.; Li, B.; Ho, P. S.; Yao, Z.; Shi, L.; Bryan, E. N.;Nemanich, R. J. J. Appl. Phys. 2007, 101, 124310.

(50) Difficulties arise in identifying the film thickness because the grainsize is comparable to the film thickness. Hence, the sputtered Al filmthickness was determined by extrapolation of measurements from thickerfilms formed over longer deposition times measured using AFM. It isimportant to note that the deposition time, hence the amount of Al depositedremained constant for all Al films.

(51) Goss, C. A.; Brumfield, J. C.; Irene, E. A.; Murray, R. W. Langmuir1993, 9, 2986.

(52) Zenhausern, F.; Adrian, M.; Emch, R.; Tamborelli, M.; Jobin, M.;Descouts, P. Ultramicroscopy 1992, 42-44, 1168.

(53) Maruyama, S.; Kojima, R.; Miyauchi, Y.; Chiashi, S.; Kohno, M.Chem. Phys. Lett. 2002, 360, 229.

(54) Huang, L.; Cui, X.; White, B.; O’Brien, S. P. O. J. Phys. Chem. B2004, 108, 16451.

(55) Moulijn, J. A.; van Diepen, A. E.; Kapteijn, F. Appl. Catal. A 2001,212, 3.

(56) Carey, J. D.; Ong, L. L.; Silva, S. R. P. Nanotechnology 2003, 14,1223.

(57) Sugime, H.; Noda, S.; Maruyama, S.; Yamaguchi, Y. Carbon 2009,47, 234.

(58) Jorio, A.; Saito, R.; Hafner, J. H.; Lieber, C. M.; Hunter, M.;McClure, T.; Dresselhaus, G.; Dresselhaus, M. S. Phys. ReV. Lett. 2001,86, 1118.

(59) Suenaga, K.; Wakabayashi, H.; Koshino, M.; Sato, Y.; Urita, K.;Iijima, S. Nat. Nanotechnol. 2007, 2, 358.

(60) http://rsbweb.nih.gov/ij/index.html.(61) Zhu, L.; Hess, D. W.; Wong, C.-P. J. Phys. Chem. B 2006, 110,

5445.(62) Xiang, R.; Yang, Z.; Zhang, Q.; Luo, G.; Qian, W.; Wei, F.;

Kadowaki, M.; Einarsson, E.; Maruyama, S. J. Phys. Chem. C 2008, 112,4892.

(63) Jorio, A.; Dresselhaus, M. S.; Dresselhaus, G. Carbon Nano-tubes - AdVanced Topics in the Synthesis, Structure, Properties andApplications: Topics in Applied Physics; Springer: Berlin; Heidelberg, 2008;Vol. 111, p 89.

(64) Crist, B. V. Fundamental XPS Data from pure element, pure oxidesand chemical compounds; XPS International, LLC: Mountain View, CA,1999.

(65) Lacerda, R. G.; Teo, K. B. K.; The, A. S.; Yang, M. H.; Dalal,S. H.; Jefferson, D. A.; Durrell, J. H.; Rupesinghe, N. L.; Roy, D.;Amaratunga, G. A. J.; Milne, W. I.; Wyczisk, F.; Legagneux, P.; Chhowalla,M. J. Appl. Phys. 2004, 96, 4456.

(66) Handbook of The Elements and NatiVe Oxides; XPS InternationalLLC: Mountain View, CA, 1999.

(67) Biju, V.; Abdul Khadar, M. J. Nanopart. Res. 2002, 4, 247.(68) Chakraborty, A. K.; Jacobs, J.; Anderson, C.; Roberts, C. J.; Hunt,

M. R. C. J. Appl. Phys. 2006, 100, 084321.(69) Pretorius, R.; Harris, J. M.; Nicolet, M.-A. Solid-State Electron.

1978, 21, 667.(70) Ridgeway, J. W. T.; Haneman, D. Phys. Status Solidi 1970, 38,

K-31.(71) Ghoshtagore, R. N. J. Appl. Phys. 1969, 40, 4374.(72) Jasek, O.; Elias, M.; Zajıekova, L.; Kudrle, V.; Bublan, M.;

Matıjkova, J.; Rek, A.; Bursık, J.; Kadleeıkova, M. Mater. Sci. Eng. C 2006,26, 1189.

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effects of metal underlayer grain size on carbon nanotube growth

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