We describe an atomic layer etching (ALE) method for copper that involves cyclic exposure to an oxidant and hexafluoroacetylacetone (Hhfac) at 275 • C. The process does not attack dielectrics such as SiO 2 or SiN x , and the surface reactions are kinetically self-limiting to afford a precise etch depth that is spatially uniform. Exposure of a copper surface to molecular oxygen, O 2 , a weak oxidant, forms a ∼0.3 nm thick layer of Cu 2 O, which is removed in a subsequent step by exposure to Hhfac. The etch reaction involves disproportionation of Cu(hfac) intermediates, such that ∼0.09 nm copper is removed per cycle. Exposure of copper to ozone, a stronger oxidant, affords ∼15 nm of CuO; when this oxidized surface is exposed to Hhfac, 8.4 nm of copper is removed per cycle. The etch products, Cu(hfac) 2 and H 2 O, are efficiently pumped away; H 2 O, a poor oxidant, does not attack the bare Cu surface. The roughness of the copper surface increases slowly over successive etch cycles. Thermochemical and bulk etching data indicate that this approach should work for a variety of other metals.

This paper provides a detailed analysis of the deposition of iron by chemical vapor deposition from the well-known precursor iron pentacarbonyl, Fe(CO)5. The authors show that at a constant temperature (e.g., 300 °C) the growth rate decreases monotonically with time. Growth eventually ceases altogether at a certain film thickness and cannot restart, even under conditions that are favorable for nucleation. The authors propose that the reduction in Fe deposition rate observed here and in previous studies results from surface poisoning: the dissociative chemisorption of CO molecules on the Fe surface at elevated temperature forms inactive surface species, especially graphitic carbon, which accumulate on the surface and eventually stop Fe growth. Remarkably, the surface poisoning effect can be inhibited, so that Fe deposition occurs at a constant rate with no self-limiting growth behavior, by coflowing NH3 along with the Fe(CO)5 precursor during growth. The adsorbed NH3 inhibits CO chemisorption by displacing CO from the growth surface and inhibiting CO chemisorption. The resulting Fe films are of high purity, i.e., carbon and nitrogen contents each below 1 at. %.

The authors demonstrate that the addition of an ammonia coflow during the chemical vapor deposition of MoCxNy, Fe, or Ru thin films at ≤200 °C from the metal carbonyl precursors Mo(CO)6, Fe(CO)5, or Ru3(CO)12 affords area-selective growth: film grows readily on titanium metal or vanadium nitride substrate surfaces, but no nucleation occurs on air-exposed SiO2, TiO2, Al2O3, or MgO within the investigated times of 1–2 h. By contrast, in the absence of ammonia, nucleation and deposition on these oxide surfaces can either be slow or rapid, depending strongly on the oxide surface preparation. NH3 is also the source of N in MoCxNy, which has a resistivity of 200 μΩ cm and becomes superconducting at a critical temperature of 4 K. The authors hypothesize that the passivating effect of NH3 on oxide surfaces involves site blocking to prevent precursor adsorption, or an acid–base interaction to stabilize surface-bound metal subcarbonyl intermediates, or a combination of these mechanisms. A key finding is that surface selective growth is often crucially dependent on the sample history of the substrate, which must be specified in detail if reproducible results are to be obtained.

We report a method to control the surface morphology of thin copper films during growth by chemical vapor deposition from the precursor Cu(hfac)VTMS. A molecular inhibitor -an additive that modifies the surface attachment kinetics but does not decompose and contribute impurity atoms to the film -is added during the nucleation and/or growth stages of the film. Here we show that the reaction by-product VTMS can serve as such an inhibitor. If the inhibitor is added during the nucleation stage, when bare substrate surface is still exposed, the inhibitor greatly reduces the rate of coalescence and promotes the formation of a large density of uniformly-sized copper islands. Alternatively, if the film is allowed to nucleate in the absence of the inhibitor, subsequent addition of the inhibitor leads to a continuous copper film that is remarkably smooth on the nm scale. © 2014 The Electrochemical Society. [DOI: 10.1149/2.009405jss] All rights reserved.Manuscript submitted January 14, 2014; revised manuscript received February 27, 2014. Published March 11, 2014 Copper is used in many advanced nanoscale technologies due to its high electrical and thermal conductivity, and its strong surface plasmon resonance when in the form of nanoparticles.1-5 For continuous films, such as those used as the seed layer for electrodeposition in integrated circuits, the film must be less than 10 nm thick, pinholefree, and extremely smooth, with an rms roughness of less than 1 nm. For optical devices based on copper nanoparticles, it is important to control the nanoparticle size and morphology. 4,6 Rigorous control of copper growth can be difficult: the surface energy of copper is high and the atomic diffusion rate is significant, so that dewetting often occurs during growth or subsequent annealing. 7-12Thin films of copper can be deposited by a wide variety of techniques including wet chemical growth, physical vapor deposition, chemical vapor deposition (CVD) and atomic layer deposition (ALD). To deposit copper conformally in substrate architectures such as trenches and vias that have re-entrant or high aspect ratio features, ALD and CVD are preferred techniques because of the ability of the precursor molecules to diffuse throughout the structure.13-17 A general difficulty arises when the substrate is relatively unreactive, such as an oxide surface: the resulting films tend to be rough owing to a combination of sparse nucleation and the tendency of the deposited material to agglomerate. 18 Once surface roughness on the length scale of the island separation is formed, it cannot be eliminated by the overgrowth of more material. 18The use of additives to enhance film smoothness is well established in the electrochemical deposition of copper 19,20 but is not common in CVD. For CVD, the morphology of copper films can sometimes be improved by adding a second component to the growth gas. For example, addition of H 2 O to a flux of Cu(hfac)VTMS (hfac = hexafluoroacetylacetonate and VTMS = vinyltrimethylsilane) enhances the wettability of the surfac...

We report a simple process for the selective deposition of copper films on RuO 2 , while no Cu nucleation occurs on thermal SiO 2 or porous carbon doped oxide (CDO). Using the precursor Cu(hfac)VTMS, selectivity is attained by adding a co-flow of excess VTMS to act as a growth inhibitor. With precursor alone, 52 nm of Cu grows on RuO 2 ; on CDO or on thermal SiO 2 , nucleation is delayed such that 41 or 1.3 nm are deposited, respectively. Repeating the experiment with the co-flow of VTMS affords a 12 nm thick Cu film on RuO 2 with roughness of 1.8 nm. But on CDO or thermal SiO 2 , the Cu deposition is only 0.10 or ∼0.04 nm, respectively. AFM scans of the CDO and SiO 2 surfaces are identical to the bare substrates. The small quantity of Cu that is deposited must be finely distributed, presumably on defect sites; it can be etched to below the RBS detection limit using a co-flow of Hhfac and VTMS for few minutes at the end of the growth. The process window is wide: selective growth occurs for a range of VTMS pressures (0.5-2.0 mTorr), growth times (up to 90 min), and growth temperatures (up to 180 • C).