Vladislav V Krisyuk

Tatyana P Koretskaya

Asiya E Turgambaeva

Sergey V Trubin

Ilya V Korolkov

Olivier Debieu

Thomas Duguet

Igor K Igumenov

Constantin Vahlas

Thermal decomposition of tungsten hexacarbonyl: CVD of W-containing films under Pd codeposition and VUV assistance

Keywords: tungsten hexacarbonyl, chemical vapor deposition, W-containing films

Experiments on chemical vapor deposition of W(CO) 6derived films on silicon substrates were carried out at total pressure of 5-10 Torr within the temperature range of 250-350 o С in Ar or H 2 flow. Metallic, carbide and oxide phases composed the obtained films. Deposition in presence of hydrogen results in the increase of the metal content in the film. Sublimed palladium hexafluoracetylacetonate Pd(hfa) 2 was used for Pd catalytic promotion of the deposition process. Codeposition with Pd(hfa) 2 in hydrogen increases W-metal fraction and oxygen content while the Pd content is up to 10 at.%. Influence of vacuum ultraviolet (VUV) radiation from Xe excimer lamp (λ~172 nm) on the quality of the obtained films was investigated. It was found that VUV irradiation can reduce the oxygen-content in the film while W-metal fraction slightly increases. In all films, oxygen was in the form of WO 3 and carbon was mainly incorporated as a graphite metastable phase. The influence of other chemical additives is discussed.

1 Introduction Research dedicated to the production of tungsten films and particles has been intensively carried out since the middle of the last century. Indeed, in addition to its high melting point, tungsten presents high density and low reactivity/toxicity. For these reasons W-based films find applications as ohmic contacts, diffusion barriers and interconnects in microelectronic devices [START_REF] Kodas | The Chemistry of Metal CVD[END_REF][START_REF] Cao | [END_REF][3][4]. W films on polymers may be useful for electromagnetic or radiation shielding of polymer-packaged devices [START_REF] Tong | Advanced Materials and Design for Electromagnetic Interference Shielding[END_REF]6]. It may also be useful to provide polymers with enhanced mechanical, with diffusion barrier and with optical properties ( [START_REF] Wilson | [END_REF] and references therein). Such films are typically produced by CVD or ALD from tungsten halides, most often tungsten hexafluoride, WF 6 and a suitable reducing agent, e.g. H 2 or silane [START_REF] Kodas | The Chemistry of Metal CVD[END_REF][START_REF] Wilson | [END_REF]. Despite the robustness of such processes, exposure of the substrates and of the deposition equipment to corrosive reaction byproducts, are significant disadvantages. The use of metalorganic precursors, such as cyclopentadienyl-containing molecular [8,9] compounds and more frequently tungsten hexacarbonyl W(CO) 6 [10][11][12][13][14] has been investigated by the past as an alternative route to halides chemistry. The seminal report on the metalor-ganic CVD (MOCVD) of W from W(CO) 6 is the one of Kaplan and d'Heurle in the early seventies [10] while the latest one comes from Lai and Lamb [14], who resumed the state of the art for this process. In essence, the deposition temperature (T d ) range is between 200 and 600 °C. However, films deposited at T d less than 500 o C present significant amounts of carbon and oxygen impurities, even for photochemically and plasma assisted deposition. For this reason, the mechanism of surface decomposition of W(CO) 6 was investigated by several groups [15][16][17].

The summary of the obtained results reveals that T d conditions the composition of the films: 95 % pure W films can be prepared at 1000 o C and above, or under ultra high vacuum, or at high dilution of the precursor gas mixture. When T d is lower than 500 o C there is an accumulation of oxygen and carbon in the films by chemisorption and decomposition of CO. Incomplete decomposition of the precursor leads to entrapment of sufficient amount of oxygen and carbon as well. Oxygen in the film oxidizes W to WO 3 but carbon remains segregated or it forms solid solutions with tungsten. Only post deposition annealing at ca. 900 °C leads to the reduction in carbon and oxygen content.

From this review it appears that it has not been possible to process pure W films at low temperature from single thermal dissociation of W(CO) 6 . Using plasma assisted deposition with H 2 can increase the amount of metallic W at low T d [18]. Photostimulated deposition and chemical methods have been marginally investigated but they have not yet led to positive results. Solving this problem would allow W films processed by cost effective MOCVD processes finding applications in the functionalization of thermally sensitive parts such as polymers containing shielding materials [19].

In this work we apply in situ mass spectrometry (MS) of the gas phase in equivalent to low pressure MOCVD of W conditions, with the aim to gain chemical insight to thermally activated decomposition of the precursor. Then, we investigate different potential solutions for the low temperature MOCVD of W-rich films. First, we apply photostimulation of the process involving vacuum ultraviolet (VUV) radiation. Then, we use chemical additives, namely catalysts to drive the surface reaction towards the clean dissociation of the W-CO bond and the decontamination and removal of CO from the reaction zone. The use of catalysts was tested in a CVD process where W films with a small amount of Pt were deposited from Cp 2 WH 2 and from CpPt(CH 3 ) 3 aiming at reducing the carbon and oxygen content in the film [8]. Following this idea we test the catalytic activity of Pd(hfa) 2 for tungsten deposition from W(CO) 6 in the presence of H 2 since Pd can be a cost effective alternative for Pt precursors. Finally, post deposition remote hydrogen plasma treatment (PT) of selected films is also studied with the aim to decrease contaminants.

2 Experimental Mass spectrometric study of W(CO) 6 thermolysis was performed using a specially designed input system for the time-of-flight (TOF) mass spectrometer [20] consisting of an evaporator and a temperature-controlled cell (thermal reactor). This system ensures that the sampling and analysis are only performed on the vapor phase. The compound was conditioned in an open glass ampoule maintained at 30 o C in the evaporator under dynamic vacuum conditions. The vaporized compound was introduced in the reactor which was heated from 90 to 500 o C at a rate of 5 o /min. Sampling was performed directly into the MS ion source through the 0.2 mm effusive orifice of the reactor. When necessary, hydrogen was added in the reactor. These experimental conditions exclude almost all ion-molecule collisions in the mass spectrometer. Mass spectra (ionization electrons energy ca. 70 eV) showing the composition of the gas phase were recorded every 2 min. Obtained mass spectra of W(CO) 6 recorded in the absence of thermolysis are in a good agreement with reference spectra for this compound presented in NIST database (MS #62145). From the full-range mass spectra the temperature dependences of the intensities of the ion peaks were derived. MOCVD of W-containing films was performed in a cold wall stainless steel CVD reactor described in [21]. Deposition was made on Si(100) 10x10 mm 2 substrates. Despite the fact that the ultimate objective of this study is surface functionalization of polymers, depositions are made on 10x10 mm 2 Si(100) coupons because of their appropriateness for films characterization and of the reluctance to nucleation and to adhesion of a metallic film on a polymer surface. These two latter difficulties should be treated prior transfer of the results of the present investigation to a process dealing with polymer metallization.

For each experiment, ~200 mg of W(CO) 6 was conditioned in a glass ampoule. For Pd assisted deposition ~20 mg of Pd(hfa) 2 was conditioned in a separate evaporator. W(CO) 6 and Pd(hfa) 2 were transported to the deposition chamber by Ar and H 2 , respectively, both flowing at 2 L/h. Pd(hfa) 2 sublimation temperature was 40 °C. W(CO) 6 sublimation temperature was 65 °C for W+Pd codeposition and 60 °C for single W deposition. Operating pressure was 5 Torr and T d was in the range 200 -350 o C. VUV stimulation was ensured by a low pressure Xe excimer lamp (λ~172 nm) (CHROMDET Analytical Instruments). The conversion efficiency of power consumption in the VUV radiation is up to 10%. Deposition time = 30 min.

Films morphology was investigated by scanning electron microscopy (SEM) with a JEOL JSM 6700F instrument equipped with a JEOL EX-23000 BU energydispersive X-ray spectrometer (EDS) analyzer.

Remote plasma treatment (PT) experiments were performed in the separate vacuum chamber made of quartz tube with plasma ignited between outer coaxial electrode and grounded flange. Samples were placed face-up onto quartz stand. Hydrogen pressure was ~1 Torr, output power was ~30W, and exposition time was 20 min.

Surface composition was investigated by X-ray photoelectron spectroscopy (XPS) with a SPECS instrument using AlKα (hν=1486.6 eV) non-monochromatic radiation. Source power was 200W, X-ray beam diameter was 1 mm. The samples were mounted on a sample holder using double side adhesive conductive tape. The resulting spectra were processed with the CasaXPS ® software. Calibration was based on the binding energy of the line 4f 7/2 of tungsten in the metallic state, which is equal to 31.4 eV [START_REF]Handbook of X-ray Photoelectron Spectroscopy[END_REF]. Cleaning the surface was carried by bombardment with argon ions at a chamber pressure of 3.0*10 -6 Torr, at voltage of 2.5 kV, with a sputtering time 3 min at sputtering rate of ca. 1-3 Å/min. XRD analysis of the film was performed on Shimadzu XRD-7000 diffractometer (CuKα radiation, Ni-filter, 5°-60° 2θ range, step 0.03 2θ, 1s per step). Indexing of the diffraction pattern was carried out using data for compounds reported in the PDF database [START_REF]Powder Diffraction File, release 2010[END_REF].

Results and discussion

Mass spectrometric study of W(CO) 6

The temperature dependence of the composition of the gas was evaluated by in situ MS study of vapors of the heated precursor using an input system imitating the LPCVD reactor. Such dependences are similar to classic kinetic concentration ones. Figure 1 resumes the evolution of the intensities of [W(CO) 4 ] + (selected as the most intensive W-containing peak) and [CO] + ion peaks as a function of temperature, with and without H 2 . All W-containing peaks demonstrate similar behavior within the examined temperature range indicating that they result from the fragmentation of the molecular ion [W(CO) 6 ] + in the mass analyzer and not from the thermolysis of W(CO) 6 . Upon reaching decomposition onset temperature, intensities of the peaks of Wcontaining ions decrease while intensity of the peak [CO] + , corresponding to the gaseous product of thermolysis, increases. Neither other organic gaseous products nor reactive intermediates such as W(CO) n (n=1-5) were revealed for the heterogeneous thermolysis of W(CO) 6 . The main difference between the processes occurring with and without H 2 is the decomposition onset temperature. This temperature is ca. 120 °C in the presence of H 2 . It is 210 o C; i.e. substantially higher, in the absence of H 2 . It should be noted that this value is close to 230 o C which is the onset temperature obtained by Patokin and Sagalovich [START_REF] Patokin | [END_REF] during MS study of sampled products of W(CO) 6 thermolysis. These data unambiguously indicate a decrease of the decomposition temperature in the presence of H 2 : increase in the carbon monoxide peaks is observed at 120 o C when the intensity of W-containing peaks began to decrease. It should be noted that no new gaseous products were detected in the presence of H 2 .

CVD from W(CO) 6

Based on the obtained MS data, low temperature CVD experiments from W(CO) 6 were performed within the T d range of 250-350 o C and in the presence of H 2 . VUV irradiation and addition of Pd hexafluoroacetylacetonate Pd(hfa) 2 were investigated for their potential to modify films characteristics.

Experimental conditions and characteristics of the obtained films are presented in Tables 123. Table 1 resumes the consumption of the W and Pd precursors for selected experiments. It appears that, in every experiment performed in similar conditions of carrier gas flow rate (samples #27, 28, 29, 30, 32, 33 for 2 L/min, #34, 35 for 1 L/min), the consumption of each precursor is almost the same and complete, thus allowing for comparison among the processed films in terms of thickness and composition. It is worth noting that the use of Ar instead of H 2 as a carrier gas results in the decrease of the sublimation efficiency (#31 for 2 L/min). All films deposited on Si present grey mirror-like surface. The growth rate at 200 and 250 o C is too low to yield films with satisfactory thickness at reasonable time. Consequently, we focus here on the investigation of films processed at 300, 350 and in part 450 o C.

At 300 o C, when hydrogen is used for the transport of W(CO) 6 and Ar for the transport of Pd(hfa) 2 , the use of VUV-irradiation results in the increase of the thickness, and of the W and Pd contents of the film (Table 2, compare #26-28). The percentage of metallic W is also substantially increased but C and O contamination is still high (Table 3, #27-28). Raising the deposition temperature to 350 o C results in significant increase of the growth rate and of the contamination of the surface with carbon. The percentage of metallic W decreases, probably due to noticeable formation of WC x (Table 3). According to EDS data, the use of VUV-radiation results in the decrease of oxygen content but dose not eliminate carbon from the film, the C/O ratio approaching 1 (Table 2, compare #29-30 and #32-33). From the other hand if we compare surface and bulk composition from XPS and EDS respectively (Table 2, e. g. #27 and 28; #29, 30 and 31) it is noticeable that carbon contamination of the surface is lower in case of UVstimulated deposition Thus, photostimulated deposition is favorable to remove any accidental impurities that can come from e.g. gas line or leaks but it can not remove all chemisorbed CO from the surface. The direct impact of H 2 in the characteristics of W films can be evaluated from the comparison of samples 31 and 32: At 350 o C the combination of W(CO) 6 + H 2 yields higher W content in the film than in the case when Ar replaces H 2 . It was not possible to determine the C and these films by XPS. However, EDS provided similar C and O content for both. The influence of the presence of a small amount of Pd(hfa) 2 in the input gas can be illustrated from the comparison between samples 29 and 32 (without VUV activation) and between samples 30 and 33 (with VUV activation). Two conclusions can be unambiguously drawn: First, Pd(hfa) 2 allows increase of film thickness (487 nm vs. 350 nm, and 431 vs. 290 nm, respectively). Secondly, it results in films with lower W content. It is concluded that Pd(hfa) 2 significantly assists the deposition rate. However, it is detrimental to the processing of W-rich films. This is possibly attributed to the nature of the Pd precursor, whose ligand contains four oxygen atoms. It is worth investigating the impact of oxygen-free Pd precursor instead of Pd(hfa) 2 . To prepare thick films we increased the deposition time up to 1 h. Figure 2 shows typical morphology of the thick film prepared from W(CO) 6 and Pd(hfa) 2 in different media. Films obtained in Ar atmosphere are loose and porous, while those deposited in the presence of hydrogen are more dense and compact. XRD patterns of the same films are presented in Fig. 3. Films deposited with H 2 are better crystallized and are mainly composed of carbon and WO 3 .

Remote plasma treatment (PT) has been done for the selected samples to verify if it possible to increase W-metal content in the films. As it is visible from the Table 2 post plasma treatment can substantially reduce carbon contamination of the film, but W still stay oxidized. Comparison of W4f XPS band structure e.g. shown in Fig. 4 for the sample #26 really confirms that W-metal line (~31.4 eV) remains unchanged, WC/WO x line (~33.2 eV) significantly increases, WO 3 line (~34-36 eV) decreases. It is concluded that post treatment with hydrogen plasma can be favorable to reduce carbon contamination of the film prepared from W(CO) 6 .

Summary

Various directions have been investigated with the aim to obtain W-rich films by low temperature MOCVD from W(CO) 6 . Investigation of the decomposition of the precursor in similar to low pressure MOCVD of W by in situ MS confirms that hydrogen reacts with tungsten carbonyl at T d as low as ca. 120 o C and indicates that H 2 should be systematically used as co-reactant in such processes.

Low temperature (T d ≤350 o C) MOCVD from W(CO) 6 results in the formation of the mixture of metallic, carbide and oxide phases in the film. Deposition in the presence of H 2 results in the increase of the metal content in the film at the same temperature.

VUV photostimulation at 300 °C results in (a) moderate increase of film thickness, (b) increase of the W content and (c) substantial increase of the percentage of metallic W. Nevertheless, it does not allow elimination all chemisorbed CO from the surface.

Addition of a small amount of Pd in the film through codeposition with Pd(hfa) 2 assists the deposition rate but results in films with lower W content. To confirm or rule out efficiency of reactive additives as Pd-codeposits study of oxygen-free Pd compound is needed.

Direct deposition of pure metallic W films from W(CO) 6 at low temperature still remains a challenge. Further investigation of photon and catalytic assistance of the process is necessary to identify possible parametric windows providing such films. Meanwhile, post deposition hydrogen plasma treatment may be used to decrease carbon contamination and thus increase the W content of the films.

Figure 1

 1 Figure 1 Temperature dependences of ion peak intensities for [W(CO) 4 ] + (1) and for [CO] + (2) in the presence or not of H 2 in the input gas

Figure 2

 2 Figure 2 Planar and cross-section SEM images for W-containing films deposited at T d = 350 o C for 1 h in conditions #31 (without H 2 , upper micrographs) and #32 (with H 2 , lower micrographs). Scale bar corresponds to 100 nm for all samples.

Figure 3 X

 3 Figure 3 X-ray diffractograms of the samples presented in Fig. 2, processed in conditions #31 (without H 2 , upper diffractogram) and #32 (with H 2 , lower difractogram).

Figure 4

 4 Figure 4 Comparison of W4f XPS band structure before and after plasma treatment (H + ) for the sample #26. The three lines correspond to energies assigned to metallic W, WC/WO x , and WO 3 from left to right.

Table 1

 1 Experimental conditions for typical CVD runs.

			Precursor con-
	Sample #	T d , o C	W(CO) 6 /Pd(hfa) 2 sumption, mg	Remark
	26	300	161/-
	27		182/18
	28(VUV)	300	185/18
	29		194/21
		350
	30(VUV)		194/18
	31		125/-	Ar W = 2 L/h without H 2
	32	350	186/-	H 2 W =2 L/h
	33(VUV)		191/-	without Ar
	34 35(VUV)	350	117/20 128/18	Ar W/Pd =2L/h (shared for two precursors), w =2L/h (separate Н 2 line)

Table 2

 2 Comparison of films composition before (XPS/EDS) and after hydrogen plasma (H + ) treatment from XPS analysis for the selected samples.

	Sam-ple #	Thick-ness by nm SEM,	W	XPS/EDS, at% Pd O	C
	26	107 3.37/12.83	29.36/37.42 61.27/47.44
	26 (H + )		25.82	74.18
	27	120 5.99/14.01 0.77/11.33 39.51/40.26 53.74/34.30
	28	158 11.38/12.29 9.89/5.88 36.2/35.90 42.53/45.91
	29	487 3.9/28.13	-/6.92 17.12/43.02 78.98/16.98
	30	431 6.84/21.88 0.96/3.65 18.79/29.37 73.41/26.22
	31	328 22.6/21.47	55.44/35.11 21.21/16.11
	31 (H + )		12.74	82.56	4.89
	32	350 6.99/38.09	44.04/37.01 48.96/24.9
	32 (H + )		21.82	66.84	11.31
	33	290 3.64/33.23	60.79/23.26 35.58/29.68
	33 (H + )		23.87	65.74	10.39
	40	150	19.59	57.81	22.59

Table 3

 3 Assignment ofW4f XPS lines for selected samples based on Ref.[13].

	Sample # Е b , eV	W compound	at. %
		32.2	W	5.4
	27	33.4	WO x , WC	29.7
		35.2	WO 3	64.9
		30.7	W	34.3
	28	32.0	N/A	37.0
		34.3	WO 3	28.7
		31.4	W	13.4
	29	33.0	WO x , WC	49.4
		34.9	WO 3	37.2
		31.1	W	29.2
	30	32.3	WO x , WC	36.2
		34.7	WO 3	34.6
		31.4	W	8.2
	31	33.52	WO x , WC	63.1
		35.7	WO 3	28.7
		31.4	W	26
	32	32.86	WO x , WC	35.81
		35.5	WO 3	38.19
		31.3	W	23.9
	34	31.8	N/A	23.3
		33.0	WO x , WC	19.2
		35.4	WO 3	33.5
		31.4	W	27.6
	35	32.6	WO x , WC	32.6
		35.4	WO 3	39.8


Pierre-Luc Etchepare
email: pierreluc.etchepare@ensiacet.fr

Loïc Baggetto

Hugues Vergnes

Diane Samélor

Daniel Sadowski

Brigitte Caussat

Constantin Vahlas
email: constantin.vahlas@ensiacet.fr

Process-structure-properties relationship in direct liquid injection chemical vapor deposition of amorphous alumina from aluminum tri-isopropoxide

Keywords: aluminum isopropoxide, amorphous alumina films, direct liquid injection, hydrothermal ageing, glass surface

We propose a method to apply amorphous alumina films on the inner surface of glass containers aiming to improve their hydrothermal barrier property. We have carried out alumina deposition on Si substrates as a function of deposition temperature to determine the physicochemical properties of the thin film materials, and on glass containers to evaluate the influence of postdeposition hydrothermal ageing on the films properties. Film preparation has been achieved by metal-organic chemical vapor deposition (MOCVD), using triisopropoxide aluminum (ATI) dissolved in anhydrous cyclohexane as precursor, in a temperature range between 360 °C and 600 °C. A direct liquid injection technology is used to feed the reactor in a controlled and reproducible way. The amorphous alumina films have been characterized by several techniques such as XRD, EPMA, XPS, SEM, AFM and scratch-test method. Films are amorphous and hydroxylated at 360 and 420 °C and close to stoichiometric at 490 and 560 °C. Hydrothermal ageing simulated by a standard sterilization cycle modifies the adhesion and surface morphology of the alumina film on glass containers to a rough, porous and nonadhesive layer.

Elemental compositions of alumina films on Si substrates. O/Al atomic ratios (EPMA) for films processed in the horizontal (blue dots) and vertical (red dots) reactors. Carbon concentration for films deposited in the horizontal (black dots) and vertical (white dots) reactors, determined, respectively, by XPS and EPMA. Films processed at high temperatures (490-560 °C) are stoichiometric Al 2 O 3 with a very low amount of hydroxyl groups. Carbon concentrations are lower than 1 at%.

1 Introduction Due to their numerous allotropic phases and their tunable crystallinity, alumina films are materials of major technological and industrial potential. Their various properties match a wide panel of applications such as optical transmittance for optics [1], insulating properties for microelectronic components [2], wear resistance [3], catalyst support [4], protection against corrosion [5] and barrier properties [6]. Specifications often require a dense, robust and adherent coating for protecting metal, glass or polymer surfaces with specific properties. In particular, amorphous alumina films present interesting physico-chemical characteristics whose nature, extent and combination depend on the preparation conditions [7,8].

Different techniques have been applied to process alumina coatings on metal or polymer substrates such as electron beam evaporation [9], rf magnetron sputtering [2], atomic layer deposition [10,11], sol-gel [5] or metalorganic CVD (MOCVD) [7,[12][13][14]. The latter is one of the most attractive techniques for the deposition of such films on complex-in-shape geometries with conformal coverage i.e. uniform thickness along the surface. Composition, allotropic phases, stoichiometry, crystallinity and microstructure of the material can be adjusted by fine tuning of the MOCVD experimental parameters such as reactor design, precursor selection, reactive atmosphere, deposition temperature and pressure. The MOCVD of alumina from aluminum tri-isopropoxide (ATI) is well documented [7,8,[12][13][14][15][16][17][START_REF] Morssinkhof | The deposition of thin alumina films on steels by MOCVD. Synthesis, characterization and protective properties against high temperature corrosion[END_REF][START_REF] Soni | [END_REF]. ATI yields amorphous and stoichiometric alumina coatings with a smooth and dense microstructure at 5 Torr in the temperature range 420 °C to 650 °C [8,15,[START_REF] Soni | [END_REF]. Higher process temperatures lead to the deposition of nano-crystallized γ-Al 2 O 3 [15] or to the homogeneous decomposition of ATI which generates a different microstructure [12]. The ATI molecule is well described in the literature [20]. It is sensitive to water vapor and thus it is subjected to ageing upon exposure to ambient atmosphere, which results in its partial or total hydrolysis in nonvolatile compounds, Al(O i Pr) 3-n (OH) n (n=1, 2) and Al(OH) 3 respectively [13,20]. The precursor is usually melted, maintained in supercooled state and vaporized with a bubbler. Despite the simplicity and the cost effectiveness of this solution, the use of ATI in the supercooled state has two consequences. First, maintaining ATI at the vaporization temperature for a long period impacts the stability of the molecule and subsequently the coating quality [13]. Second, it is difficult to know exactly the generated reactive gas mass flow rate. This poorly controlled situation leads to non-reproducible processes, especially for the present case where low activation energy prevails in the entire temperature range of interest, de facto resulting in mass transport limited process. The direct liquid injection (DLI) technology allows overcoming the previously mentioned drawbacks with the controlled atomization and vaporization of ATI dissolved in a carrier solvent [21,22]. The DLI of ATI is not well documented in the literature except by Song et al. [23] and Krumdieck et al. [14] who implemented it in a pulsed-pressure MOCVD process. The authors used a solution of ATI in n-octane and in dry toluene, respectively. Films obtained between 500 and 550 °C at a nominal base pressure of 0.6 Torr were stoichiometric, amorphous and presented a nodular surface morphology.

We recently investigated a DLI-MOCVD process for the application of amorphous alumina from ATI on the internal surface of glass bottles with the aim to increase their barrier properties. We reported on the process modeling involving the development of a numerical model, using the Computational Fluid Dynamics code FLUENT, to calculate local profiles of gas flow, temperature, concentration and deposition rates inside the bottle. Interplay between simulation and deposition experiments allowed identifying process conditions which led to the deposition of transparent films with uniform thickness profile [24,25].

In the present work we report on the experimental investigation of the process-structure-properties relationship of this film. First, we focus on the dissolution of ATI, with the aim to establish criteria for the selection of the different solvents and define the most appropriate one. Then, we detail results on the deposition of alumina films on silicon wafers and glass containers. We are especially interested on the possible impact of the solvent on the films composition and characteristics. Finally, we report on the hydrothermal ageing of the alumina coated glass substrates using a standard sterilization cycle in terms of surface morphology and adhesion, before providing concluding remarks.

Materials and methods

Thin film preparation

The precursor solution was prepared by mixing ATI and anhydrous cyclohexane inside a Schlenk flask. The viscosity of 0.05 M to 1 M solutions of ATI in cyclohexane was measured with a falling sphere viscometer. The turbidity, i.e. the presence of insoluble particles in the solvent was measured for highly concentrated solutions in cyclohexane and n-octane by the nephelometry method employing a Hach 2100 instrument. ATI powder (98%, Acros Organics) was weighed and sealed in a Schlenk flask inside a glovebox circulated with purified Ar (99.9997%, Air Products). The sealed flask was filled outside of the glovebox with the appropriate quantity of anhydrous cyclohexane (99.5%, H 2 O < 10 ppm, Acros Organics) using an air-tight, Ar-purged, glass syringe to prepare a solution of varying concentration from 0.025 to 1 M. Then, the Schlenk flask was connected with the injection system and purged several times with pressurized N 2 (99.9999%, Praxair). DLI of the solution was carried out with Kemstream Vapbox ® systems. One gas line was connected to the injection system consisting of a N 2pressurized Schlenk flask filled with the precursor solution and connected to a set of two injectors. During the deposition, the ATI solution was injected through an injector to form small droplets with a calculated flow rate from 0.90 to 1.55 mL min -1 . The frequency (2 or 3 Hz) and opening times (from 3 to 30 ms) of the injection system were controlled by a computer interface. The injected solution was blended in the mixing chamber with carrier gas (pure N 2 ) with a pressure differential held at 1 bar. The mixing carrier gas flow of N 2 was fixed to a flow rate of 300 standard cubic centimeters per minute (sccm).

Glass containers (42.5 mm diameter, 73 mm height bottles) and Si substrates (10x10 mm, cut from 4" Si (100) wafers (Sil'tronix)) were cleaned in an ultrasonic bath with acetone and ethanol. DLI-MOCVD was performed in two different reactors.

Deposition onto Si substrates was carried out using a horizontal hot-wall reactor composed of a quartz tube (25 mm diameter, 300 mm length) heated by a resistive fur-nace. Samples were placed on a stainless steel (SS) holder in the center of the quartz tube where the temperature was uniform. The total flow rate of the carrier gas with vaporized solution was maintained unchanged at 814 sccm. The inlet ATI molar fraction was estimated to be 8.2×10 -3 .

Deposition on the inner walls of glass containers was performed in a vertical hot-wall DLI-MOCVD reactor, as described in detail elsewhere [24]. It is recalled that the reactive gas phase was driven with pure N 2 and introduced through a SS nozzle from the upper part of the quartz tube. In this configuration, the total flow rate and the inlet ATI molar fraction were equal to 585 sccm and 1.7×10 -3 , respectively. A SS holder maintained the glass container in the gas feeding tube axis. An inductively heated thermoregulated SS cylinder around the reaction chamber ensured a temperature profile along the container with a set temperature at the bottom external surface on the central axis of the container 480 °C [23].

In both systems, the vaporization chamber and SS transport lines were thermally regulated to allow efficient flash vaporization of the mixture droplets without early ATI decomposition or partial solvent vaporization. Dry pumps (base pressure < 5.10 -2 Torr) and pressure gauges connected to the output of the deposition chambers were used to control the reaction pressure to 5 Torr.

Coated glass containers were submitted to one sterilization cycle in wet atmosphere. Samples were filled with ultra-pure water (18.2 MΩ.cm) and placed during 1 h inside an autoclave (Getinge) operating at 121 °C and 2 bar absolute pressure. Heating and cooling rates were controlled to avoid thermal cracks.

Materials characterization X-ray diffraction

(XRD) was performed on a Bruker D8 Advance using a Cu K α (1.5418 Å) X-ray tube operated at 40 kV and 40 mA, a Ni filter and solid-state Lynxeye detector in θ-3°/θ+3°configuration. Samples were measured on a zero background holder and a θ offset of 3°was applied between the X-ray source and detector arms to suppress the strong diffraction of the Si substrates.

X-ray photoelectron spectroscopy (XPS) was performed on a Thermo Scientific K-Alpha instrument capable of a typical base pressure of 10 -9 Torr using monochromatic Al K α (1486.6 eV). The spectrometer energy calibration was performed using the Au4f7/2 (83.9 ± 0.1 eV) and Cu2p3/2 (932.7 ± 0.1 eV) photoelectron lines. Charging compensation and neutralization were applied by using a dual beam flood gun. The probed areas were about 400 µm diameter. Surface erosion was employed using Ar ions accelerated at 2 kV, resulting in an erosion rate of about 0.08 nm s -1 . Constant pass energy of 30 eV and energy steps of 0.1 eV were used for high resolution scans. The photoelectron peaks were analyzed by Gaussian/Lorentzian (G/L = 70/30) peak fitting and using a Shirley background. The atomic concentrations were determined from photoelectron peak areas using the atomic sensitivity factors reported by Scofield, taking into account the transmission function of the analyzer calculated at different pass energies from Ag3d and AgMNN peaks collected for a reference sample.

The O/Al ratios in the films were determined by electron probe microanalysis (EPMA) with a Cameca SXFive apparatus operated at 10 and 15 keV and calibrated using a high purity alumina standard. Samples were covered by a thin layer of carbon in a vaporization chamber (Leica) to prevent charging effects. Each sample was measured 12 times in different locations to determine the spatial homogeneity of the sample composition.

Scanning electron microscopy (SEM) on coated glass container fragments was performed with a LEO 435VP instrument. Samples were fractured and metalized with silver and graphite before analysis in order to avoid charging effects. Acquisition was performed with an acceleration voltage of 10 kV, a probe current between 40-80 pA and a working distance 11 mm. SEM on alumina films deposited on Si substrates was performed with a JEOL JSM-7800F field emission scanning electron microscope operated at 10 kV in backscattered mode. Samples were fractured and metalized with platinum before imaging.

The roughness of the alumina thin films deposited on the glass surface, before and after sterilization cycle was measured by atomic force microscopy (AFM) in tapping mode using an Agilent 5500 instrument. The probe silicon tip (AppNano ACT-50) with a spring constant of 13-77 N m -1 and resonance frequency between 200 and 40 kHz scanned areas of samples with a scan rate of 2 µm s -1 with 256x256 data points. Pico Image software (Agilent Technology) was used to analyze and process the data.

The adhesion of alumina coatings on glass containers before and after sterilization cycle was investigated by scratch-test method (standard EN 1071) using a CSM Revetest Instruments. Samples were scratched using a diamond indenter with a spherical tip of 200 µm and a cone angle of 120°at a constant loading rate of 15 N min -1 . The load applied on the coating surface increased from 1 to 30 N. The scratch track was fixed at 4 mm.

Results and discussion 3.1 Solubility of ATI

The DLI technology requires dissolving the precursor in an appropriate solvent to inject and transport it to the deposition area. The selection of the solvent is an important step of the process optimization. It must take into account the physicochemical, and health and safety characteristics of the solvents, their cost, their compatibility with the precursor and their stability in the deposition conditions [26].

The solvent must have the same molecular dipole moments as the precursor to ensure maximum dissolution. The dipole moment of the ATI molecule equals 0 debye. For this reason the apolar solvents n-pentane, cyclohexane, n-heptane and n-octane were a priori considered since they also meet health and safety criteria excluding carcinogenic, mutagenic and reprotoxic (CMR) substances. A tradeoff was applied for the final selection of the most appropriate solvent, based on four criteria: (a) The occupational exposure limit value (OELV) at 8 h which is also related to health and security but it depends on the countries legislation. (b) The boiling point which should be high enough to avoid partial vaporization of the solvent at ambient temperature in the Schlenk flask. (c) The saturated vapor pressure at the vaporization temperature (160 °C) which should be close enough to that of the precursor so as to prevent early vaporization of the solvent. (d) The ratio of the cost between each alkane and the less expensive solvent (cyclohexane) for ≥ 99% anhydrous quality. These criteria were rated in terms of importance as secondary, important and preponderant and then were weighed by a factor of 3, 5 and 8, respectively. This classification favors the preponderant criterion. Table 1 summarizes the data for each solvent. As for the four criteria, a coefficient of merit was attributed to each value ranging from 1 to 5 for the weaker to the stronger one. N-pentane presents a low boiling point (36.1 °C) and a high saturated vapor pressure at 160 °C (ratio of 3.35 with ATI); i.e. an important risk of early vaporization. The saturated vapor pressure at 160 °C of n-heptane is high in comparison with that of ATI (ratio of 2.98) and its cost is relatively high. N-octane and cyclohexane present OELV and saturated vapor pressure close to some extent. Preliminary conclusions from this a priori investigation reveal that cyclohexane can be selected to dissolve and transport the ATI. However, additional criteria, such as viscosity and turbidity, dealing with the solubility of ATI per se should be considered to validate the selection of the cyclohexane.

The histogram of Fig. 1 shows the viscosity of the four solvents and of the corresponding solutions of ATI at 0.05 M fixed concentration, which is typical for an injected solution in a DLI-MOCVD process. The viscosity of the four solvents is low, not exceeding that of water (1 mPa s at 20 °C). Viscosities of the solutions are almost equal to those of pure solvents; they are compatible with DLI technology and are too close to each other to allow discrimination. increases from 0.98 mPa s at 0.05 M up to 2.25 mPa s at 1 M, i.e. to a value which is still convenient for DLI. A global consideration of these results reveals that cyclohexane is the most appropriate solvent for ATI. However, the solvent must not be subjected to thermal or reactive decomposition during deposition resulting in the pollution of the film. This criterion requires dedicated experimental investigation not detailed in the present article.

Film characteristics XRD patterns of films deposited on

Si do not show any peaks associated with the formation of a crystalline structure for all deposition temperatures (see Fig. S1 of the Supporting Information, online at: www.pss-c.com).

Figure 3 shows the evolution of the O/Al atomic ratio and carbon concentration in the films deposited on Si, as a function of deposition temperature, determined by EPMA and XPS, respectively. The films grown onto Si substrates inside the horizontal CVD reactor (blue data points) prepared at 360 °C and 420 °C show O/Al atomic ratios equal to 2.12 and 1.80, respectively. These values correspond to compositions close to AlOOH (atomic ratio 2.0) and partially hydroxylated AlO 1+x (OH) 1-2x films. Films prepared at 490 °C and 560 °C show O/Al ratios equal to 1.53 and 1.57, respectively, in good agreement with the formula Al 2 O 3 (atomic ratio 1.5). Films deposited in the vertical CVD reactor (red data points) present O/Al ratios of 1.88 and 1.56 at 420 °C and 600 °C, respectively, in close agreement with those measured for the films grown in the horizontal CVD reactor. Hence, it appears that the inlet ATI molar fraction and the reactor configuration used here have little influence on the film composition. The C content (black datapoints) measured by XPS after surface erosion is below 1 at% for all temperatures, similar to the C content measured by EPMA on films grown in the vertical reactor (black datapoints). Gleizes et al. measured similar O/Al atomic ratio (EPMA) for alumina films deposited by MOCVD with sublimed ATI in the temperature range of 480-600 °C [15].

The evolution of the elemental composition with deposition temperature observed by EPMA was further investigated using XPS. Films deposited on Si substrates were eroded by ca. 100 nm and subsequently measured with XPS (Fig. 4). The amount of aliphatic carbon (C-C and C-H bonds only) is less than 1 at% and likely results from small amounts of solvent residue incorporated in the film during growth. The Al2p signal apparently shows a single symmetric peak corresponding to Al 3+ in an O 2-framework. The signals for Al2p in Al 2 O 3 , AlOOH and Al(OH) 3 are very close in binding energy and are consequently hard to separate [27][28][29]. The O1s signal is a better discriminant for the presence of hydroxyl groups, which have a peak expected at ca. 1.3 eV off the main O 2-peak of Al 2 O 3 [27][28][29]. The measured O1s signals clearly show a main peak around 531.4 eV at the location expected for O 2-in the Al 2 O 3 framework [30,31], which is accompanied with a shoulder around 532.7 eV of increasing intensity with decreasing temperature, attributable to hydroxyl groups. The relative amount of hydroxyl groups increases with decreasing temperature, in agreement with the increasing O/Al ratio measured with EPMA. However the intensity of the response related to hydroxyls groups is not as intense as expected for a composition close to AlOOH obtained at 360 °C. This discrepancy likely results from ion beam damages, which decreases the actual quantity of OH signal. Nonetheless, the resulting spectra show a reasonable trend with the EPMA data. This is confirmed for instance in the case of the films deposited at 490 and 560 °C which show the presence of OH groups. The film morphology is shown in the cross-section SEM micrograph in Fig. 5. The film is smooth but slightly porous with nanometer cavities (see Fig. S2 of the Supporting Information) over the entire height, in contrast with alumina films deposited from sublimed ATI are very smooth and dense (see Fig. S2, [32]). This particular morphology may be attributed to the presence of cyclohexane in the DLI process.

Properties of films grown on the inner walls of glass containers

The evolution of the surface morphology and adhesion of the film on glass substrates in wet atmosphere after a sterilization cycle was investigated. SEM micrographs of alumina films on glass container surfaces, before and after sterilization cycle, are presented in Fig. 6. The as-deposited film looks slightly rougher than the alumina films deposited on Si substrate in the horizontal reactor (Fig. 5). The higher roughness might be attributed to the glass surface topography. After a sterilization cycle, the film surface is significantly rougher and porous with sub-micrometric cavities. Surface 3D topography AFM images of alumina films deposited on glass containers before and after a sterilization cycle are shown in Fig. 7. Results of measured roughness for uncoated and coated containers are summarized in Table 2. The root means square average of the roughness (RMS) of the bare containers is low (0.3 and 0.7 nm); i.e. the glass surface prior to deposition is extremely flat and smooth despite a slight increase of the surface roughness after the sterilization cycle. Before the sterilization cycle, alumina films display slightly rough features with a mean RMS value 17.1 nm. The sterilization cycle yields a strong increase of the roughness value to 61.7 nm which demonstrates the increase of surface porosity and subsequent film deterioration. From the SEM and AFM observations it is concluded that the ageing in such wet environment and relatively high temperature damages the film surface making it more porous and rougher. The effect of a sterilization cycle on the adhesion of the film on glass was investigated by scratch-test. SEM images of the tracks are presented in Fig. 8 before (Figs. 8a, b andc) and after (Figs. 9a, b, c andd) sterilization. These micrographs allow defining the features of four different coating failure modes. The first cracking mode (Figs. 8a and9b) is determined by the presence of the indenter mark; i.e. the first circular crack on the film surface. The feature of the second cracking mode (Figs. 8b and9c) is the propagation of lateral cracks perpendicular to the direction of scratch. The third mode (Figs. 8c and9d) occurs with the coating chipping and the substrate exposure. An X-ray dispersive energy spectrometry analysis confirms the presence of an uncoated surface on the edges of the track; thus the total film delamination. The critical loads were measured at the beginning of each cracking failure mode and are given in Table 3. All failure modes can be defined as Hertzian cracking [33] which corresponds to a through-thickness cracking of a hard coating (compressive stresses generated ahead of the indenter and tensile stresses behind it). Tensile cracking failure mode is also found in the literature for alumina coatings deposited on glass substrate [34]. An additional mode (Fig. 9a) is observed for sterilized coated containers with partial delamination only on the edges of the track at low loads. It can be referenced to recovery spallation and buckling. Comparison of critical loads among samples from different studies is only allowed if the coating-substrate system, the film thickness, the scratch-test apparatus, the scratch parameters and the applied load are the same. In this case, the chipping mode occurs at relatively high load which should lead to a satisfactory adhesion of the thin film. On the other hand, for sterilized coated container, an additional mode occurs (buckling/recovery spallation) for a mean load of 11.9 N. This mode is usually related to interfacial failure. It could be explained by a softer coating on a hard substrate where the adhesion is weak. However, the other cracking failure modes related to hard coating appear at higher loads than those for as processed coated container. These results suggest that the upper part of the alumina film is damaged by the sterilization cycle, making it porous, rough and soft. A lower part could remain non affected forming a hard, dense and adhesive layer on the glass surface.

4 Conclusions Alumina films were deposited by DLI-MOCVD from ATI at 5 Torr. Direct liquid injection technology generally allows a better control of generated reactive gaseous mass flow and vaporization, thereby contributing to a more robust process. Cyclohexane has been selected to dissolve and transport ATI to the deposition area. A very low amount of carbon (below 1 at%) is found in the film composition which proves the solvent does not readily decompose during deposition.

Films are amorphous in the temperature range 360-560 °C. Coatings processed at high temperatures (490-560 °C) correspond to stoichiometric Al 2 O 3 with a very low amount of hydroxyl groups; the composition of those deposited at lower temperatures (360-420 °C) ranges from AlOOH to partially hydroxylated alumina AlO 1+x (OH) 1-2x with increasing temperature. The film morphology is relatively smooth (RMS 17.1 nm) and slightly porous with nanometer cavities.

Amorphous alumina films were deposited on the inner surface of glass containers and heat treated in humid atmosphere. SEM observation and AFM imaging show that the sterilization cycle affects the surface topography and morphology, making the film rougher (RMS 61.7 nm) and more porous. It also affects the film adhesion on the glass surface as revealed by scratch-tests.

Supporting Information

Additional supporting information may be found in the online version of this article at the publisher's website.

Table 1

 1 Weighed criteria and biological and physicochemical data for four apolar solvents. Each value is given a score of merit in parenthesis. The resulting total score for each solvent (last column) is the sum of the products of each weighed criterion times the score of merit. It appears that cyclohexane is the most appropriate solvent to dissolve and transport the ATI.

Figure 1

 1 Figure 1 Viscosity of ATI solutions with n-pentane, n-heptane, n-octane and cyclohexane at 0.05 M fixed concentration (histogram, pure solvents and solutions in light and dark grey, respectively) and evolution of the viscosity of ATI solutions with cyclohexane as a function of concentration.

Figure 2

 2 Figure 2 Turbidity values of 1M ATI solutions with cyclohexane (squares) and n-octane (circles). Limit value between lightly troubled water and troubled water at 30 NTU is indicated.

Figure 2

 2 Figure2shows the evolution of the turbidity of 1 M solutions of ATI in cyclohexane and n-octane for comparison as a function of their ageing up to 7 days. It is worth noting that 1 M concentration and 7 days ageing are relatively high values for a CVD process. The turbidity of both solutions increases with ageing. That of the solutions with cyclohexane is systematically lower revealing a higher stability than the one with n-octane. Finally, the evolution of the viscosity of the ATI/cyclohexane solution as a function of concentration up to 1 M is shown in Fig.1. The viscosity

Figure 3

 3 Figure 3Elemental compositions of alumina films on Si substrates. O/Al atomic ratios were measured using EPMA for films processed in the horizontal (blue dots) and vertical (red dots) reactors. Carbon concentration was determined by XPS for films deposited in the horizontal reactor (black dots) and EPMA for those prepared in the vertical reactor (white dots).

Figure 4

 4 Figure 4 High resolution XPS spectra (Al2p and O1s) of alumina films deposited on Si substrates at 360 °C, 420 °C, 490 °C, and 560 °C. The decomposition of the spectra into the assigned chemical species is shown for each core level.

Figure 5

 5 Figure 5 Cross-section SEM micrograph of amorphous alumina film on Si substrate for a deposition temperature of 480 °C.

Figure 6

 6 Figure 6 Surface (top) and cross-section (bottom) SEM micrographs of amorphous alumina deposited at 480 °C on glass containers before (a and c) and after a sterilization cycle (b and d). The flat part seen in part (d) is a bit of the glass bottle.

Figure 7

 7 Figure 7 3D topography images of the surfaces measured by AFM employing the tapping mode for amorphous alumina films on the inner walls of glass containers before (a) and after a sterilization cycle (b).

Figure 8

 8 Figure 8 SEM micrographs of scratch tracks on the surface of alumina film deposited on glass containers at 480 °C before a sterilization cycle. Cracking failure modes depend on the indenter load and are defined as first cracks mode (a), lateral cracks mode (b), chipping mode (c). Dotted semicircle line and arrows are guide to the eye.

Figure 9

 9 Figure 9 SEM micrographs of scratch tracks on the surface of alumina film deposited on glass containers at 480 °C after a sterilization cycle. Cracking failure modes depend on the indenter load and are defined as buckling/recovery spallation mode (a), first cracks mode (b), lateral cracks mode (c), chipping mode (d). Dotted semicircle line and arrows are guide to the eye.

Table 2

 2 Measured roughness of uncoated and coated containers before and after a sterilization cycle during 1 h at 121 °C.

		Uncoated containers	Coated containers
		As proc-essed	Sterilized	As proc-essed	Sterilized
	Roughness RMS (nm)	0.3	0.7	17.1	61.7

Table 3

 3 Critical load of coating failure modes determined with the scratch-test method for no sterilized and sterilized coated glass containers.

			Critical load (N)
		Buckling-spallation mode	First cracks mode	Lateral cracking mode	Chipping mode
	As processed	-	12.2	16.8	25.5
	Sterilized	11.9	22.2	24.1	27.5

Acknowledgements

We are indebted to Djar Oquab, Yannick Thebault, Cedric Charvillat and Jérôme Esvan, from CIRIMAT for help with SEM, AFM XRD and XPS measurements respectively, to Philippe de Parseval, UMS Castaing, for the EPMA characterization, and to Florent Colomes and Iulian Stan for experimental study of ATI solubility. We also thank Hervé Guillon, Kemstream, Montpellier for advice on the operation of the DLI facility.


Loïc Baggetto
email: loic.baggetto@ensiacet.fr

Jérôme Esvan

Cédric Charvillat

Diane Samélor

Hugues Vergnes

Brigitte Caussat

Alain Gleizes

Constantin Vahlas
email: constantin.vahlas@ensiacet.fr

Alumina thin films prepared by direct liquid injection chemical vapor deposition of dimethylaluminum isopropoxide: a process-structure investigation

Keywords: DLI CVD, DMAI precursor, H 2 O and O 2 atmospheres, amorphous alumina films

Elemental composition of the alumina films deposited in the presence of H 2 O from 150 to 300 °C (circles), in the presence of O 2 from 500 to 700 °C (diamonds) or in the presence of N 2 only at 600 °C (squares). O/Al ratios and C concentrations are shown.

The development of a new process to obtain amorphous alumina thin films is presented. We show for the first time the direct liquid injection chemical vapor deposition (DLI CVD) of alumina thin films using dimethylaluminum isopropoxide (DMAI) precursor in two oxidizing atmospheres. At high process temperature (500-700 °C), the film growth takes place in the presence of O 2 whereas at low temperature (150-300 °C) H 2 O vapor is used. The materials characteristics, such as the surface morphology and roughness (SEM and AFM), crystal structure (XRD), composition (EPMA) and chemistry (XPS) are discussed in detail. Very smooth films, with typical roughness values lower than 2.0 nm are obtained. The thin films are all composed of an amorphous material with varying composition. Supported by both EPMA and XPS results, film composition evolves from a partial oxyhydroxide to a stoichiometric oxide at low deposition temperature (150-300 °C) in the presence of H 2 O. At higher growth temperature (500-700 °C) in the presence of O 2 , the composition changes from that of a stoichiometric oxide to a mixture of an oxide with aluminum carbide. 1 Introduction Alumina in the form of powders or coatings is a material of major technological interest for a wide range of applications in optical and microelectronic components, as well as a wear resistance agent and catalyst support, or for the protection against corrosion and high temperature oxidation [1][2][3]. The protection from corrosion and thermal oxidation of stainless steels, light alloys (Tiand Mg-based) and composite polymer materials for aero-nautic or automotive applications calls for the development of alumina thin films with superior performances to act as an oxidation barrier for moderate temperature applications (200-600 °C) [4,5].

Overall, the number of reports on the processing of alumina films is significant. A practical way to produce amorphous alumina thin films is the MOCVD technique [4][5][6][7][8][9], in which a controlled amount of precursor carried over a surface decomposes and reacts to form a condensate thin layer. Aluminum tri-isopropoxide (ATI) is a precursor with significant advantages since it is cheap, stable in air and only slowly reactive with water. However, the shelflife, polymerization and condensation properties of its tetramer pose problems [10] and can lead to tedious reactor maintenance due to line and injector blockage. Furthermore, the applicable range of temperature for ATI is not only limited by the stability window of the substrate but is also constrained by the nature and the elemental composition of the film. Amorphous alumina films grown from ATI are generally hydroxylated for temperatures below 420 °C and have optimal mechanical, corrosion and oxidation protection properties around 520 °C [11,12]. Trimethyl aluminum (TMA) was also evaluated as a CVD precursor to form alumina films. However, in addition of an oxidizing agent, the process requires the use of a catalytic support or a plasma, and employs deposition temperatures as high if not higher than those employed with ATI [13]. These relatively high temperatures may limit the use of TMA-and ATI-processed alumina films for the protection of materials for which the properties, e.g. mechanical, change importantly above 200 °C. This is the case of Mg alloys and polymer composites.

A few groups proposed the use of dimethylaluminum isopropoxide (DMAI) as an alternative CVD precursor using vaporization in bubbler systems [7,8,[14][15][16][17][18][19]. DMAI, which consists of an Al atom linked to 2 methyl and one isopropyl groups, has an intermediate structure between ATI and TMA. Unlike TMA, it is non pyrophoric, and although it is currently more expensive than ATI and is very reactive with water, it presents a significantly higher vapor pressure, it is liquid at room temperature, has a longer shelf-life and is not prone to polymerization. Amorphous alumina films can be deposited from DMAI in the presence of water vapor between 180 and 270 °C, this temperature range being significantly lower than that of ATI processes [8]. Such low deposition temperature paves the way for the protection of light alloys or polymer composite materials. These advantages, combined with the use of a direct liquid injection (DLI) system instead of a bubbler technology for the transport of the precursor in the deposition area make DLI of DMAI an a priori interesting alternative to vaporize ATI in CVD processes. DMAI was also tested for atomic layer deposition (ALD) of alumina [20][21][22][23][24].

In the present work, we discuss for the first time the development of such CVD processes involving DLI of DMAI. The injection system offers several advantages over classical sublimation or evaporation systems since it provides constant precursor vapor flow and leads to a more stable, robust and reproducible process [25][26][27]. Two oxidizing atmospheres were employed, as selected from former works using evaporation systems [7,8,[14][15][16][17][18][19]. Battiston et al. claimed the fabrication of stoichiometric and carbon-free alumina films using O 2 and H 2 O oxidizing atmospheres [7,8] whereas other authors consistently meas-ured carbon leftovers in their films when using pure N 2 [16][17][18]. Here, at high deposition temperatures, oxygen has been used to compensate for the lower than 1.5 O/Al ratio in DMAI and to possibly prevent the incorporation of carbon in the films. At lower deposition temperatures, water vapor has been used to compensate for the lack of oxygen in DMAI while promoting the reaction kinetics. The physico-chemical properties of the thin films were characterized using a wide set of techniques. Surface morphology and roughness were measured using scanning electron microscopy (SEM) and atomic force microscopy (AFM). The amorphous state of the films was ascertained by X-ray diffraction (XRD). Elemental compositions as well as the chemical environments of the bulk of the film, with emphasis on the formation of partial hydroxyls or carbides, were determined through X-ray photoelectron spectroscopy (XPS) and electron probe microanalysis (EPMA). EPMA is a powerful probe as it is significantly more precise and accurate than XPS and energy dispersive X-ray spectroscopy (EDS) to determine compositions.

Experimental

2.1 Thin film preparation Thin films were prepared using a homemade horizontal hot-wall CVD reactor equipped with three gas lines and mass flow controllers (MFC), a large evaporation stainless steel chamber and a quartz tube (25 mm diameter, 300 mm length) heated by a resistive furnace. A dry pump and pressure gauges connected to the output of the quartz tube were used to control the reaction pressure to 5 Torr. Gas lines and the evaporation steel chamber were heated to ca. 80 °C to prevent precursor condensation. One gas line was connected to the injection system consisting of a N 2 -pressurized Schlenk flask filled with the precursor solution and connected to a DLI equipment set of two injectors (Kemstream), one for liquid injection and one for gas mixing. The precursor solution was prepared in a Schlenk flask inside a glovebox circulated with purified Ar (99.9997%, Air Products) by adding a known amount of DMAI solution (>99%, Air Liquide) into the flask. The sealed flask was taken out of the glovebox and completed with the appropriate quantity of anhydrous cyclohexane (99.5%, H 2 O<10 ppm, Sigma-Aldrich) using an air-tight, Ar-purged, glass syringe to prepare a solution of 0.2 M DMAI. Then, the Schlenk flask was connected to the injection system and purged several times with pressurized N 2 (99.9999%, Praxair). The frequency and opening times of the injection system were controlled to feed the evaporation chamber with small precursor solution droplets mixed with N 2 gas with a flow set from 300 to 450 sccm. For the low temperature process, 120 sccm of N 2 was bubbled through deionized H 2 O (> 20 M:.cm) kept at room temperature. For the high temperature process, 50 sccm of O 2 (99.9995%, Air Products) was added to the dilution gas (300-400 sccm) installed on the third gas line. In both cases, the oxidants were in large excess compared to DMAI. Silicon samples (10x10 mm) cut from 4" Si (001) wafers (Sil'tronix) were placed on a stainless steel holder in the center of the quartz tube where the temperature was uniform. Growth conditions for both the low and high temperature processes are listed in Table 1.

Materials characterization

The film thickness was measured using optical reflectivity (Mikropack, Nanocalc 2000). Thickness was confirmed by measuring the film thickness in SEM cross-sections. Samples surface morphology and roughness were measured using SEM and AFM. SEM micrographs were taken in backscattered mode on a JEOL JSM-7800F field emission scanning electron microscope operated at 10 kV. Samples measured with SEM were fractured then covered by a thin layer of sputtered platinum to prevent charging effects. AFM was used in ambient conditions on an Agilent Technologies 5500. Scanning was performed in contact mode with tips of spring constant of about 0.292 N m -1 (AppNano). Scanning rate was set to 2 µm s -1 . Images were processed with the software Pico Image (Agilent Technologies). The film crystalline structure was measured by XRD on a Bruker D8 Advance using a Cu K D (1.5418 Å) X-ray tube operated at 40 kV and 40 mA, a Ni filter and solid-state Lynxeye detector in T+3°/T-3° configuration. Samples were measured on a zero background holder and a T offset of 3° was applied between the X-ray source and detector arms to suppress the strong (004) diffraction of the samples Si substrates normally measured around 69° (Fig. S1, see Supporting Information, online at: www.pss-c.com). EPMA was used to accurately determine the O/Al ratios in the films. Characterization was conducted with a Cameca SXFive apparatus operated at 10 and 15 keV and calibrated using a high purity alumina standard. Samples were covered by a thin layer of carbon inside an evaporation chamber (Leica) to prevent charging effects. Each sample was measured 4 to 6 times in different locations to determine the spatial homogeneity of the sample composition. XPS measurements were performed on a Thermo Scientific K-Alpha instrument using monochromatic Al K D (1486.6 eV) capable of a base pressure of 10 -9 Torr or less, and equipped with a detector positioned at 90° with respect to the sample surface. The spectrometer energy calibration was performed using the Au4f7/2 (83.9 ± 0.1 eV) and Cu2p3/2 (932.7 ± 0.1 eV) photoelectron lines. Charging compensation and neutralization were applied by using a dual beam flood gun. The probed areas were about 400 Pm diameter. A survey scan was collected before and after surface erosion to measure the composition and check for possible contaminants. Surface erosion was employed using Ar ions accelerated at 2 kV, resulting in an erosion rate of about 0.08 nm s -1 . Constant pass energy of 30 eV and energy steps of 0.1 eV were used for high resolution scans. The energy scale was fixed by setting the adventitious carbon C1s energy to 284.4 eV. The photoelectron peaks were analyzed by Gaussian/Lorentzian (G/L = 70/30) peak fitting and using a Shirley background. The atomic concentrations were determined from photoelectron peak areas using the atomic sensitivity factors reported by Scofield, taking into account the transmission function of the analyzer. This function was calculated at different pass energies from Ag3d and AgMNN peaks collected for a silver reference sample. C concentrations, reported from these quantifications, are estimated to be exact within ± 20%.

3 Results and discussion AFM derived surface morphology and roughness of two films grown at 200 °C and 600 °C are presented in Fig. 1, and the corresponding roughness parameters are listed in Table 2. It is clear from the AFM data that the two films are extremely smooth, with typical roughness values of about 1.0 nm for the film grown at 200 °C with H 2 O, and around 1.5 nm for the film grown at 600 °C with O 2 . The film grown at the higher temperature shows somewhat smaller domains and a slightly higher surface roughness. In both cases, the low roughness and absence of nodules support the absence of homogeneous gas phase reactions. Moreover, the low roughness values suggest the formation of a high density of nucleation centers during the film growth, as was proposed to explain the smoothness of alumina films prepared at low temperatures by evaporation of DMAI [8]. 1). R a and R q stand for the arithmetic and root mean squared roughness of a line profile, respectively. S a and S q stand for the areal arithmetic and root mean squared roughness of an entire area scan, respectively. Values are in nm. The results obtained at high temperatures contrast with those reported for alumina films prepared at high temperatures by evaporation of DMAI [7], where significantly rougher films were obtained. This difference might result from the lower growth rate used in our work, which gives surface diffusion and atomic rearrangement more time to take place. 2.

The samples surfaces imaged by SEM are shown in Fig. 2. As observed with AFM, the films are extremely smooth, without any pinholes and look dense. The measurement of the film thickness in the cross-section views was found to be in very good agreement with the optical reflectivity measurements, with deviations of about 5%.

The amorphous structure of the thin film materials was ascertained by XRD for all deposition temperatures in both oxidizing atmospheres. The diffraction patterns (Fig. 3) do not show any peaks associated with the formation of a crystalline structure. The absence of significant broad humps, except that of the Si substrate near 69° (cf. Fig. S1 of the Supporting Information) suggests the formation of amorphous thin film materials, similar to the results of former works concerning the deposition of alumina films from DMAI [7,8,[14][15][16][17][18][19] or ATI [4][5][6][7][8][9][10]12] precursors for temperatures up to 700 °C. In the absence of XRD phase identification, we investigated the composition of the films using XPS (after systematic erosion of a ca. 100 nm surface layer) and EPMA. EPMA is a powerful probe to measure the O/Al with exactness and precision using internal standards. The O/Al ratios obtained from EPMA measurements are compared to the ratios expected for AlOOH and Al 2 O 3 references. The films prepared at 150 and 200 °C in the presence of H 2 O vapor show O/Al atomic ratios of 1.81 and 1.67, respectively (Fig. 4). These values are comprised between those of AlOOH (2.0) and Al 2 O 3 (1.5), thus revealing partially hydroxylated alumina AlO 1+x (OH) 1-2x . The films prepared at 250 and 300 °C show O/Al ratios of 1.54 and 1.50, respectively, in good agreement with the formula Al 2 O 3 . The C content measured by XPS (Fig. S2, Supporting Information) is below 1 at%. The film prepared at 500 °C in the presence of O 2 has an O/Al ratio of 1.48, which may be ascribed to the formula Al 2 O 3 . The films prepared at 600 and 700 °C have O/Al ratios of 1.25 and 1.10, well below 1.50. The decrease in O/Al ratio is clearly related to an increase in the C concentration, up to about 6.5 at% in the film prepared at 700 °C, as has been also evidenced by Schmidt et al. during the preparation of alumina films from sublimed DMAI in this temperature range [17,18]. This suggests the formation of the aluminum carbide Al The XPS results for the samples prepared at low temperature using H 2 O vapor are presented in Fig. 5. For all samples, the absence of reduced species such as Al 2 O 3-x or metallic Al signal expected around 73.5 and 72.5 eV in Al2p, respectively, suggests that ion sputtering erosion did not significantly alter the bulk structure of the alumina material.The amount of aliphatic carbon (C-C and C-H bonds, Fig. S2, Supporting Information) is less than 1 at% for all samples. The carbon species likely result from partial decomposition of the precursor or from solvent residues incorporated in the film during growth. The Al2p signal around 74.4 eV apparently shows a single symmetric peak. Its main contribution is from Al 3+ in an O 2-framework. A small proportion overlapping with the main peak (not shown) is thought to result from Al oxyhydroxides, for which the response is expected to be very similar to that of the oxide. Indeed, the signals for Al2p in Al 2 O 3 , AlOOH and Al(OH) 3 are very close in binding energy and hard to separate [32][33][34]. The O1s signal is a much better discriminant for the presence of hydroxyl groups, which have a peak expected at about 1.3 eV offset from the main O 2- peak of Al 2 O 3 [33,34]. The measured O1s signals clearly show a main peak around 531.4 eV at the location expected for O 2-in the Al 2 O 3 framework [32,33,35,36] accompanied with a shoulder of increasing intensity with decreasing temperature around 532.7 eV attributable to hydroxyl groups. The relative amount of hydroxyl groups increases as the temperature decreases, in excellent agreement with the increasing O/Al ratio measured with EPMA (Fig. 4). In a nutshell, XPS results for films processed at low temperature in the presence of H 2 O show the quasiabsence of C and the presence of hydroxyl groups for the lowest deposition temperatures whereas no hydroxyls are detected at 300 °C, in good agreement with the EPMA results. The XPS results for samples prepared at high temperature in the presence of O 2 are presented in Fig. 6. The Al2p and O1s signals clearly show the presence of intensity at the locations expected for Al 2 O 3 with Al2p binding energy of about 74.4 eV and O1s binding energy near 531.5 eV, as discussed above. Moreover, the signals observed at about 73.2 eV in Al2p and near 282.0 eV in C1s are normally attributed to aluminum carbide Al 4 C 3 [37,38]. The Al2p binding energy of aluminum carbide is lower than that measured for Al 2 O 3 since C is less electronegative (more electropositive) than O, hence Al in aluminum carbide is more effectively reduced and measured at a lower binding energy. Similarly, the C1s binding energy of the carbide is significantly more negative than that of adventitious C-C bonds as Al is more electropositive than C. The relative atomic contribution of carbon in aluminum carbide over aliphatic carbon is 75% at 500 °C, 60% at 600 °C and 73% at 700 °C, and the absolute amount of carbon related to the carbides clearly increases with increasing deposition temperature from about 1 at% at 500 °C to about 5 at% at 700 °C. This result is in line with a former report, showing that the CVD from trimethylalu-minum and methane leads to pure Al 4 C 3 films only at temperatures as high as 1150 °C [30]. Here, the presence of aluminum carbides can be the result of the reaction of cyclohexane with the growing alumina film. This is suggested by earlier results that showed the absence of C in films grown at 560 °C during the MOCVD of O 2 and DMAI transported by evaporation [7]. However, other reports showed that C was incorporated in alumina films when using DMAI precursor in N 2 atmosphere [17,18]. DMAI is composed of 2 C atoms linked to Al, and can by itself be expected to partly contribute to the formation of carbides. It is interesting to note that a sample deposited at 600 °C in pure N 2 (no O 2 added) had a significantly higher C concentration (9.2 at%) with about 75% in the form of aluminum carbides, concomitant with a much lower O/Al ratio (0.98) than for samples deposited in O 2 atmosphere (Fig. 4). This result further strengthens the need to use an oxygen source when employing DMAI precursor in a DLI system to limit the amount of formed aluminum carbides that originate either from the precursor molecule or from the used solvent.

4 Conclusions Direct liquid injection chemical vapor deposition of alumina thin films from DMAI precursor has been investigated for the first time. O 2 as well as H 2 O oxidizing environments have been used to assist the formation of Al 2 O 3 . The produced films are extremely smooth (R a in the range of 1-2 nm) and have an amorphous structure. The composition of the films evolves from a partial oxyhydroxide AlO 1+x (OH) 1-2x to Al 2 O 3 for films processed at low temperatures (150-300 °C) in the presence of H 2 O. The films prepared at higher temperatures (500-700 °C) in the presence of O 2 , consist of the oxide Al 2 O 3 , either pure (500 °C) or mixed with aluminum carbide (600 and 700 °C) in amounts increasing with the temperature.

Supporting Information Additional supporting information may be found in the online version of this article at the publisher's website.

Figure 1

 1 Figure 1 3D topography images of the surfaces measured by contact AFM for alumina films deposited onto Si substrates at (a) 200 °C in the presence of H 2 O and at (b) 600 °C in the presence of O 2 . Film thicknesses are 440 and 720 nm, respectively. The roughness parameters derived from the measurements are listed in Table2.

Figure 2 Figure 3

 23 Figure 2 SEM cross-sections of the alumina films deposited onto Si substrates at (a) 200 °C in the presence of H 2 O and at (b) 600 °C in the presence of O 2 . Film thicknesses are ca. 440 and 720 nm, respectively.

2 Figure 4

 24 Figure 4Elemental composition of the alumina films, using EPMA (O/Al atomic ratios; blue for H 2 O, red for O 2 , and green for N 2 ) and XPS (C atomic concentration; black). Films were deposited in the presence of H 2 O from 150 to 300 °C (circles), in the presence of O 2 from 500 to 700 °C (diamonds) or in the presence of N 2 only at 600 °C (squares). Each sample was measured with EPMA 4 to 6 times in different locations to check the homogeneity of the composition.

Figure 5

 5 Figure5 High resolution X-ray photoelectron spectra (Al2p and O1s) of alumina films deposited from 150 to 300 °C in the presence of H 2 O. The decomposition of the spectra into the assigned chemical species is shown for each core level.

Figure 6

 6 Figure6 High resolution X-ray photoelectron spectra (Al2p, C1s and O1s) of alumina thin films deposited from 500 to 700 °C in the presence of O 2 . The decomposition of the spectra into the assigned chemical species is shown for each core level.

Table 1

 1 Typical DLI MOCVD process parameters used during the preparation of alumina thin films. DMAI precursor, cyclohexane solvent and a total pressure of 5 Torr were used in all cases.

			Injection	Dilution
	Temperature	Reactive	mixing	flow	Growth rate
	(°C)	gas	flow	(sccm)	(Pm/h)
			(sccm)
	150	H 2 O	300	300	0.4
	200	H 2 O	300	300	0.6
	250	H 2 O	300	300	1.1
	300	H 2 O	300	300	1.8
	500	O 2	400	400	0.3
	600	O 2	400	400	1.2
	600	N 2	450	400	0.9
	700	O 2	400	400	2.5

Table 2

 2 AFM-derived roughness parameters (nm) of alumina films prepared at 200 °C in the presence of H 2 O and at 600 °C in the presence of O 2 (Fig.

Acknowledgements Philippe de Parseval (UMS3623, CNRS) and Yannick Thebault (CIRIMAT) are gratefully acknowledged for conducting EPMA and SEM measurements, respectively. Christophe Capello is gratefully acknowledged for carrying out the sawing of Si substrates. This work was financially supported by the STAE-RTRA foundation (Toulouse, France) under the RTRA-STAE/2014/P/VIMA/12 project grant.