Turing-patterned Ta_2S_3 enables sub-2 nm diffusion barrier for advanced Cu interconnects

Domain: physics
Published: 2026-04-08
Editorial outcome: minor_revision

Abstract
As semiconductor technology advances towards the 15-Å node and beyond, the back-end-of-line (BEOL) metal interconnect technology encounters severe scaling challenge, because the conventional barrier/liner TaN/Ta(Co) bilayer cannot be proportionally thinned down as the pitch size shrinks. This challenge arouses significant resistance-capacitance (RC) delay that is currently the bottleneck for advanced integrated circuits. Herein, we propose a Turing-patterned Ta S (Tp-Ta S ) thin barrier to address the above interconnect scaling challenge. The Tp-Ta S , synthesized by a particularly designed energy-enhanced atomic layer deposition (ALD) process, features a labyrinthine stripe morphology with an intertwined layer structure, which can significantly enhance its Cu diffusion barrier performance by forcing the Cu atoms to migrate a significantly longer path through the 2D interlayer space. The synthesized Tp-Ta S barrier with only 1.2 nm thickness exhibits excellent Cu wettability, remarkably strong adhesion, and much better barrier performance than the state-of-the-art TaN barrier.

Introduction
As semiconductor technology advances towards the 15-Å node and beyond, copper interconnects in the back-end-of-line (BEOL) encounter critical scaling challenges. While the minimum feature size of the interconnect pitch must scale below 20 nm, the thickness of the state-of-the-art barrier/liner (TaN/Ta or Co) bilayer cannot be proportionally reduced and must remain around 2.5 nm to ensure adequate diffusion barrier performance.1, 2 This scaling conflict will drive, for instance, a sixfold increase in the interconnect line resistance for the 15-Å node as compared to the 3-nm node, owing to the reduction of the line effective conductive volume from 75.6% to 61.6% (Figure 1a). The resultant surge in resistance–capacitance (RC) delay severely limits the microprocessor operation frequency and therefore is the critical bottleneck in advanced integrated circuits (ICs).3, 4 As an estimation (Figure 1b), if the barrier layer thickness can be reduced to 1.2 nm, the line effective conductive volume can accordingly be remarkably increased from 61.6% to 80.8%, thereby significantly reducing the interconnect line resistance by 41.3%. This estimation underscores the critical importance of developing thin barriers with superior performance to overcome the RC delay limitation in interconnects for the post-Moore era.

To address the above challenge, atomically thin two-dimensional (2D) layered materials have been considered as good candidates to surpass the thickness limitation of the state-of-the-art TaN-based barrier materials. Both theoretical predictions and experimental studies indicate that certain single-crystalline 2D materials such as graphene5, 6, MoS 7, 8, h-BN7 and TaS 9 outperform TaN in blocking copper diffusion, showing significant potential as sub-nanometer thin barriers. However, most single-crystalline 2D materials have to be synthesized at a very high temperature (>800 °C) by chemical vapor deposition (CVD),5, 6, 7 which significantly exceeds the thermal budget limit of the back-end-of-line (BEOL) processing (≤400°C). Other approaches based on transferring are not compatible to be integrated with modern interconnect technologies.5, 6 Recently, Lo et al.9 proposed a transfer-free conversion approach to prepare the TaS barrier from Ta metal at a BEOL-compatible temperature. However, such approach not only is difficult to implement for the BEOL damascene structure, but also inevitably generates high density of impurities and grain boundaries, which are the fast pathways for copper diffusion.9, 10 Moreover, for the conventional 2D materials with van der Waals (vdW) interlayer interactions, the weak interlayer interactions not only can allow for fast copper ion diffusion through the interlayer space, but also lead to insufficient adhesion strength that subject to mechanical failure in the follow-up processes such as chemical mechanical polishing (CMP).

To address the inherent limitations of the interlayer ion diffusion and weak adhesion for the conventional vdW 2D materials, we herein introduce a Turing-patterned layered structure of Ta S as a thin barrier for copper diffusion. Turing-patterned structures have recently been applied to regulating mass transport across membranes11, 12 and catalyst engineering,13, 14, 15, 16, 17 and being inspired, we herein design a Turing-patterned Ta S (Tp-Ta S ) thin film featuring a labyrinthine stripe morphology with an intertwined layer structure (Figure 1c). This unique microstructure can significantly enhance the copper barrier property, as the intertwined layer architecture can greatly extend the copper ion diffusion path by forcing the ions to transport laterally in between the 2D layers, therefore mitigating the long-lasting issue of fast diffusion/electromigration through the vertical grain boundaries for conventional polycrystalline 2D flakes. Moreover, the Ta S structure can be viewed as a conventional TaS 2D layer structure with extra Ta atoms bridging the neighboring TaS layers, and the bridging Ta atoms can substantially increase the energy barrier for the copper ion lateral diffusion in between the TaS layers and concurrently enhance the interlayer mechanical adhesion strength.

To experimentally realize Tp-Ta S , we invented an energy enhanced atomic layer deposition (ALD) process based on an amino-imino Ta precursor and H S plasma. By utilizing the autocatalytic property of the amino-imino Ta precursor together with the supply of the plasma energy to leverage the rapid lateral diffusion of inhibitors (sulfur clusters), we were able to realize the “local activation and lateral inhibition” Turing mechanism (see next section for details) to afford the Tp-Ta S films with a thickness of only 1.2 nm at temperature sufficiently low (≤350 °C) to be compatible with the back-end-of-line (BEOL) processing. These Tp-Ta S films exhibited much improved copper diffusion barrier performance that outperformed the 2.5 nm TaN standard reference; and at the same time, the interfacial adhesion strength was greatly enhanced to 10.33 J/m2, which well satisfied the stringent requirement (>5 J/m2) for copper CMP.18 The herein Ta S barrier material, featuring the Turing pattern, can enlighten an effective path to tackle the progressively severe RC delay challenge in scaled metal interconnects for advanced semiconductor technology.

Results
Fabrication of Turing-patterned Ta S

In general, the formation of Turing pattern (i.e., Turing mechanism) relies on simultaneous presence of an activator and an inhibitor species in a reaction-diffusion system and the inhibitor diffuses significantly faster than the activator.19, 20 Analogously, in a chemical-reaction-based thin-film growth system (such as CVD or ALD), some reaction intermediates or products can catalyze the reaction (also known as autocatalysis) and therefore can act as the activators;21, 22 whereas some byproducts may inhibit the reaction and therefore can act the inhibitors.23, 24 If the inhibitors diffuse faster than the activators, a Turing pattern may be formed by the characteristic short-range activation and long-range inhibition mechanism. Following this general mechanism, we designed a new ALD approach for growing Turing-patterned TaS films by using a tantalum precursor that can generate autocatalysts together with the energetic H S plasma that can generate sulfur cluster inhibitors and accelerate their surface diffusion (Figure 2a). We chose to use (tert-butylimino)tris(diethylamino)tantalum (TBTDET, Ta(NtBu)(NEt ) ) as the tantalum precursor, as this precursor bears the alkylamino and alkylimino ligands, which can react with surface sulfhydryl groups to generate autocatalytic products for the ALD growth of metal sulfides. The surface sulfhydryl groups are naturally formed on metal sulfides when exposed to H S, and these sulfhydryl groups are Brønsted acidic, so they can protonate the precursor alkylamino and/or alkylimino ligands to afford alkylamines when the metal precursor is dosed and chemisorbed onto the surface. The afforded alkylamines are weak Brønsted/Lewis bases but volatile, so they can possibly stick dynamically on the metal sulfide surface, with some extent of mobility, by forming Brønsted acid-base complexes with sulfhydryls23 and/or Lewis acid-base complexes with the Lewis-acidic metal ions.25, 26 In the follow-up ALD half-cycle of the H S exposure, these alkylamines can dynamically coordinate to the metal ions to labilize the remaining surface metal−ligand (metal−nitrogen) bonds for the H S attack; the attack affords the metal sulfide ALD product along with the byproducts of also the alkylamines. Therefore, overall, the alkylamines can be regarded as the autocatalysts for the film growth process, and thus they can act as the activator for a Turing process. On the other hand, sulfur clusters are naturally formed during the H S plasma step; some sulfur clusters (e.g., S ) are fairly inert and therefore can act as the inhibitor for a Turing process. The energy supplied by the plasma power can further accelerate the surface diffusion of the sulfur clusters by kinetic energy transfer via collisions, and therefore the diffusion of the inhibitors can be much faster than the activators so as to satisfy the Turing conditions.

Figure 2b-f shows the experimental results with the Turing pattern formation. The characteristic Turing pattern structure of the labyrinthine stripe morphology can be clearly seen by transmission electron microscopy (TEM) and atomic force microscopy (AFM). To quantitatively describe the Turing pattern, we employed a method based on height-height correlation function (HHCF) to analyze the TEM and AFM images.27, 28 Details of this method are described in Supplementary Information. The HHCF analysis gives a function of H(r), which represents the correlation of the film height for a given lateral distance of r. The typical labyrinthine-stripe Turing pattern has a lateral periodicity in film height, so the corresponding H(r) function should show damped oscillations with respect to r; whereas for a normal film, the H(r) function is simply a decay function of r, where the correlation length reflects the film grain size.28 The oscillatory H(r) function can be further Fourier transformed to extract the characteristic frequency (f), and the inverse of frequency (1/f) reflects the average periodicity of the labyrinthine stripes. From the TEM images shown in Figure 2b-c, the characteristic periodicities of the Turing patterns are extracted to be 19.6 nm and 25.0 nm for the films deposited by 9 and 13 ALD cycles, respectively. The periodicities were also confirmed by the analysis on the AFM images (Figure 2e and Supplementary Fig. 1). Moreover, since the AFM images can directly give out the pixel-wise film height values, the frequency statistics of the height values are plotted in Figure 2e and analyzed by peak fitting. The frequency statistics can be well fitted by two Gaussian peaks, with a peak separation of 0.70 nm. Notably, this number is approximately the thickness of a single layer TaS ,29 and therefore the growth of TaS can be viewed to follow the layer-by-layer Frank-van der Merwe growth mode. Based on the above analysis and discussion, we propose the growth model of Turing-patterned TaS as illustrated in Figure 2d.

To remove the excess sulfur to form the target Tp-Ta S with a bridged layer structure, the as-deposited Tp-TaS was annealed under H /Ar mixed gas flow at 350°C for 30 min. As the AFM image shown in Figure 2f, the afforded Ta S retained the same Turing morphology after annealing, and the characteristic periodicity from the HHCF analysis was also found to be identical as before (1/f = 25.0 nm) (see Supplementary Fig. 1 for H(r)). The pixel-wise height frequency statistics are also analyzed and plotted in Figure 2f. The frequency statistics can still be well fitted by two Gaussian peaks, with a slightly smaller peak separation of 0.63 nm. This smaller separation suggests that the film became more compact by forming Ta S , which has a bridged layer structure. It is worth noting that both Tp-TaS and Tp-Ta S films were quite smooth, with root-mean-square (rms) roughnesses (R ) of 0.509 and 0.453 nm, respectively. These smooth films are fairly compatible with the advanced interconnect structures (Supplementary Fig. 2).

The elemental compositions of the films were analyzed by X-ray fluorescence (XRF). As shown in Figure 3a, the XRF analysis indicated that the as-deposited TaS and annealed Ta S films (13 ALD cycles) had the S:Ta atomic ratios of 2.08 and 1.49, respectively, which are both in good consistence with their stoichiometric compositions. Assuming the bulk densities of 7.85 g/cm3 (TaS ) and 8.20 g/cm3 (Ta S ),30, 31 the thicknesses of the as-deposited TaS and annealed Ta S films were 1.3 and 1.2 nm, respectively. Therefore, the annealing induced TaS -to-Ta S conversion was experimentally confirmed. In fact, from the atomistic structure point of view, the Ta S structure can be viewed as TaS layers with extra Ta atoms located at the interlayer space occupying the octahedral sites formed by the neighboring-layer sulfur atoms (Figure 3b). The interlayer octahedral sites can be occupied at different percentages, forming in-plane superlattices with various Ta S stoichiometries. We calculated the formation energies by density functional theory (DFT) for a series of Ta S stoichiometries with various octahedral site occupation percentages (i.e., 0%, 25%, 33.3%, 66.7% for TaS , Ta S , Ta S , Ta S , respectively). As shown in Figure 3c, under the S-rich condition, which corresponds to the experimental ALD condition, TaS is the stable stoichiometry; whereas under the Ta-rich condition, which can be achieved by vacuum or H /Ar gas annealing, Ta S becomes more stable. High-resolution TEM (HRTEM) images (Figure 3d,e) further confirmed that both as-deposited TaS and annealed Ta S films had a similar Turing patterned morphology. Moreover, fast Fourier transform (FFT) analysis of the Ta S area of the HRTEM image (boxed area in Figure 3e) shows a pattern consisting of blurry crystalline spots (Figure 3f, see also Supplementary Fig. 3), which are in good consistence with the simulated pattern for the crystalline Ta S (Figure 3g). The layer bridging feature of Ta S was further evidenced by Raman analysis. As shown in Figure 3h, aside from the pronounced Raman bands at approximately 300 and 400 cm-1, which are respectively the characteristic E and A vibrational modes of the TaS layer,32 the annealed Ta S film showed an additional Raman band at the lower wavenumber regime of 100−160 cm-1. According to our DFT calculation of the phonon dispersion diagrams for TaS and Ta S as shown in Figure 3i,j, this low wavenumber band corresponds to the interlayer coupled vibrations induced by the bridging Ta atoms that bond the neighboring TaS layers in the Ta S structure. Cross-sectional scanning transmission electron microscopic (STEM) image was further taken on the annealed Ta S film. As shown in Figure 3k (see also Supplementary Fig. 4), a clear intertwined layer structure can be observed, which confirmed the successful synthesis of the Turing-patterned Ta S film.

Interconnect Performance Evaluation

Wettability, adhesion, and barrier properties are the core metrics for evaluating a potential new barrier material for advanced interconnect technologies. The wettability and adhesion determine the metal filling uniformity and interface mechanical reliability, while the barrier property is the key to block the copper metal from diffusing into the dielectrics.

To evaluate the wettability of the Tp-Ta S film, an approximately 10 nm thick Cu film was thermally evaporated onto the Tp-Ta S surface. As shown in Figure 4a,b, using thermal oxide wafer (denoted as SiO to highlight its surface property) as a benchmark, the Cu film grown on Tp-Ta S was fairly smooth with no visible pinholes, demonstrating an excellent wettability; whereas the Cu film grown on SiO was of rather poor wettability, showing many cracks and pinholes. The much-improved wettability of Tp-Ta S can be ascribed to the formation of the relatively strong Cu−S bond at the Cu/Tp-Ta S interface, and this property is highly beneficial for uniform Cu filling into small-size damascene trench/via structures (e.g., reflow process) in advanced interconnect technologies. Possibly, the Tp-Ta S barrier can be used without the need of a liner, thereby saving more space for metal filling in small-size interconnect structures.

The adhesion property was analyzed quantitatively by performing the four-point bending (4-PB) tests on the Cu/film/SiO stacks.33 To this end, we prepared the Cu/Tp-Ta S /SiO stack with the 1.2 nm Tp-Ta S film, together with a benchmarking Cu/s-TaS /SiO stack with a 1.5 nm s-TaS film prepared by sulfidation of an e-beam evaporated Ta metal film,9 and another benchmarking Cu/SiO stack with no adhesion film for comparison. Note that the Cu film on s-TaS was also of good wettability (Supplementary Fig. 5). Figure 4c shows the load-displacement curves of the 4-PB tests for these samples, where the load initially increased with displacement due to elastic deformation, and then suddenly dropped due to the crack initiated from the notch and quickly propagated, releasing the accumulated bending energy.34, 35 The interfacial adhesion energies were measured to be 1.49, 3.06, and 10.33 J/m² for the Cu/SiO , Cu/s-TaS /SiO , and Cu/Tp-Ta S /SiO samples, respectively. Apparently, the Tp-Ta S film significantly enhanced the adhesion between Cu and SiO , and the enhancement was much larger than the s-TaS film. This phenomenon is because the Ta S structure contains the bridging Ta atoms that bond the S atoms of the neighboring TaS layers so that the adhesive interaction between the neighboring TaS layers was much stronger than the vdW interaction in the TaS case. Moreover, it is particular worth noting that the observed adhesion energy of the Cu/Tp-Ta S /SiO stack was significantly larger than the requirement for the interconnect CMP process (>5 J/m2).18 Therefore, as one of the major concerns of using conventional 2D materials for interconnect barriers, the CMP compatibility issue can be well resolved by using the Tp-Ta S barrier.

The barrier property was evaluated by the standard time-dependent dielectric breakdown (TDDB) measurement.36, 37 As illustrated in the inset of Figure 4d (see also Supplementary Fig. 6), the TDDB measurement was performed on a capacitor structure of Cu/barrier/SiO /n++-Si,38 and a constant positive voltage was applied on the Cu electrode to drive the Cu ions into the SiO dielectric. The Cu ions accumulated in the dielectric and eventually formed a conductive path that caused the breakdown of the dielectric.36 Figure 4d displays the cumulative probability distributions of the time to breakdown (t ) measured at 7 MV/cm for the Cu/barrier/SiO /n++-Si stacks with different barrier layers (i.e., 1.2 nm Tp-Ta S , 2.5 nm TaN, and 1.5 nm s-TaS ) and no barrier. The cumulative probability distributions show approximately parallel shifts with a similar slope for all above stacks, which indicates that the variation of t did not alter with the inserted barrier layer. This behavior can be understood as that in our case the variation of t was dominated by the property of the SiO dielectric and the insertion of the barrier layer uniformly delayed the diffusion of the Cu ions into the dielectric. This behavior also implies that all above barrier layers were quite uniform, because otherwise a gentler slope would have been shown in the cumulative probability distribution. To further quantify the barrier property, we took the time to 63.2% failure (TTF ) as the characteristic metric. The TTF values were found to increase from 10.9 s for the control stack with no barrier to 41.1, 102.0, and 118.5 s for the stacks with 1.5 nm s-TaS , 2.5 nm TaN, and 1.2 nm Tp-Ta S as the barriers, respectively. Remarkably, the Tp-Ta S barrier with only 1.2 nm in thickness even outperformed the state-of-the-art TaN barrier with 2.5 nm in thickness, which demonstrates the significant advantage of the Tp-Ta S barrier in interconnect scaling. Additionally, it is worth noting that the barrier property of 1.2 nm Tp-Ta S was 2.88 times better than that of 1.5 nm s-TaS , although the latter was even thicker. This comparison well demonstrates the crucial importance of the herein synthesized Turing structure that can force the Cu atoms to transport a significantly longer distance with higher energy barrier to across the Tp-Ta S barrier layer. We further took the cross-sectional STEM/EDS characterizations of the Cu/Tp-Ta S /SiO /n++-Si and Cu/s-TaS /SiO /n++-Si stacks after the above TDDB stress. As shown in Supplementary Figs. 7 and 8, the Tp-Ta S barrier layer could well block Cu from diffusing through; whereas the s-TaS layer failed to block Cu, showing pronounced Cu signals at the s-TaS /SiO interface, possibly owing to the facile Cu diffusion in between the vdW 2D TaS layers, as will be discussed later.

It is worth noting that the Tp-Ta S barrier was also fairly stable. We measured the TDDB performance of the Cu/Tp-Ta S /SiO /n++-Si stack after one-month storage, and as shown in Supplementary Fig. 9, no significant change was observed. In addition, we found that this Turing-patterned barrier layer strategy not only worked for Cu but also worked for other two important BEOL metals of Co and Ru. As shown in Supplementary Fig. 10, the characteristic TTF values were also found to significantly increase from 16.1 to 186.2 s for Co and from 78.6 to 504.9 s for Ru, after the 1.2 nm Tp-Ta S barrier layer was inserted. These results well demonstrate that the Turing-patterned barrier strategy is of very good generality.

To further investigate the Cu diffusion energy barriers, we performed the DFT calculations for TaS and Ta S . For the conventional 2D material of TaS , Figure 4e shows the lowest-energy pathway for the Cu atom to migrate from an octahedral site (IS) to a neighboring octahedral site (FS) in the interlayer space of the 2D TaS layers. The migration path includes an intermediate state (MS), where Cu moves to the tetrahedral site with an energy of +0.11 eV. To arrive at this tetrahedral site, the Cu atom needs to pass the center of the shared face of the octahedron and tetrahedron (TS1 or TS2) with an energy barrier of +0.17 eV, and therefore the overall energy barrier for this pathway is +0.17 eV. Notably, this energy barrier is fairly low, and therefore Cu atoms are relatively mobile in between the vdW 2D TaS layers. In contrast, for crystalline Ta S , as shown in Figure 4f, the interlayer space is separated into isolated domains by the bridging Ta atoms, and the alien Cu atom is most stable at the center tetrahedral site of each isolated domain (IS or FS), although the octahedral site (MS1 or MS5) is only +0.02 eV higher in energy. For this Cu atom to migrate from one domain to a neighboring domain, it has to pass across the domain wall, where all the octahedral sites are occupied by the interlayer bridging Ta atoms. The lowest-energy pathway for this migration is to via three neighboring tetrahedral sites around the Ta-occupied octahedral sites (MS2→MS3→MS4). However, the energy barrier for this pathway is extremely high, as the calculated barrier energies are +11.98 eV for the MS3 state and +22.53 eV for the corresponding transition states where Cu is at the shared edge center of two tetrahedrons (Supplementary Fig. 11). Therefore, the long-range migration of Cu atoms is rather blocked in perfect-crystalline Ta S . Even if we consider the possible presence of Ta-vacancy point defect at the domain wall, the calculated energy barrier is still significantly higher as +2.14 eV for the intermediate state where Cu occupies the octahedral vacancy site (Supplementary Fig. 12). Therefore, it is fairly conclusive that the bridging Ta atoms of Ta S can significantly suppress the long-range Cu migration and therefore is pivotal for the remarkable barrier property of Ta S as experimentally observed.

Based on the above experimental results, we further estimated the quantitative benefit of utilizing the Tp-Ta S barrier for next-generation Cu interconnect technology. We calculated the BEOL metal line/via resistances at the M1/V1 level based on the critical dimensions (CD) of the 2 nm to 7 Å technology nodes of the IRDS™ Roadmap.1 As shown in Figure 5a, as the interconnect CD shrinks down for smaller technology nodes (aspect ratio (AR) = 1.5), the metal line resistance (R ) substantially increases. Because the line resistance is dominated by the volume of Cu, R of the 1.2 nm Tp-Ta S case is always smaller than that of the conventional 1.5 nm TaN/1 nm Co case. Apparently, the difference in R becomes larger for smaller CD, because the volume fraction of the barrier progressively increases. Notably, for 1.5 nm TaN/1 nm Co, R exceeds the Roadmap limit for the technology nodes ≤1.5 nm; whereas for 1.2 nm Tp-Ta S , R is always below the Roadmap limit for the technology nodes even down to 7 Å, which suggests the high promise of the Tp-Ta S strategy for the line barrier application. Figure 5b shows the calculated via resistances (R ) (AR = 2). For small CD vias, R is dominated by the highly resistive barrier layer at the via bottom, so a bottom barrierless via structure39 has been proposed to meet the Roadmap requirement. Experimentally, such bottom barrierless structure can be achieved by selective ALD of the barrier layer as shown later. As the calculation results shown in Figure 5b, with the bottom barrierless structure, R for both 1.5 nm TaN/1 nm Co and 1.2 nm Tp-Ta S cases are below the Roadmap limit. Comparatively, R of the Tp-Ta S case is even better, with only half of that of the 1.5 nm TaN/1 nm Co case (21.7 versus 42.0 Ω) for the 7 Å technology node. Therefore, the Tp-Ta S strategy is also highly promising for the via barrier application.

To further demonstrate that the fabrication approach developed herein is also compatible with the BEOL interconnect structures, we fabricated high aspect ratio (AR) nano-trenches with 230 nm in depth and ~50 nm in opening width (AR = 4.6) and assessed the step coverage of the Tp-Ta S barrier layer on these trenches. As shown in Figure 5c-f, the STEM images and the corresponding EDS maps clearly show that the high-AR trench was conformally coated by the Tp-Ta S layer with a uniform thickness of 1.2 nm throughout the entire trench wall. Notably, such high AR of the trench is considerably higher than those of the typical BEOL damascene structures,40 and therefore we speculate that the herein proposed fabrication method can be well applied to the modern BEOL interconnect technology. Moreover, to achieve the bottom barrierless metallization for the BEOL via structures, the Tp-Ta S barrier needs to be selectively coated on the via sidewall but not on the bottom metal. To this end, we further developed a selective ALD TaS method by using aniline as the inhibitor to prevent the TaS growth on the metal surface. We used Co metal surface for this demonstration, as Co is the capping metal for Cu lines. Figure 5g-j shows the results of the selective Tp-Ta S coating on a patterned test structure with exposed Co and SiO surfaces. The EDS results clearly show that the Ta signal was only present on the SiO surface but not on the Co surface. Therefore, selective Tp-Ta S coating could also be achieved. Overall, our results clearly demonstrate that the herein proposed Tp-Ta S barrier with the associated fabrication method is highly promising for the next-generation Cu interconnect technology.

Methods
Fabrication of Tp-Ta S . The ALD of Tp-TaS used the Ta precursor of TBTDET and the H S plasma. During deposition, the TBTDET precursor was kept at 90 °C, and the ALD chamber temperature was set at 150 °C. The H S plasma was generated by a 60 W radio frequency (RF) power. In each ALD cycle, the TBTDET dose time was 5 s, the H₂S plasma pulse time was 20 s, and the purge time was 30 s for each half-cycle. The ALD growth saturation behavior was verified by varying the TBTDET dose time and plasma pulse time (Supplementary Fig. 13). To convert Tp-TaS to Tp-Ta S , the deposited Tp-TaS films were in situ annealed at 350 °C for 30 min under a mixed gas atmosphere of Ar:H = 1:1. It is worth noting that Turing pattern did not show up for regular thermal ALD prepared TaS (Supplementary Fig. 14); also, the Turing pattern did not show in another recent ALD TaS report using Ta[N(CH ) ] [NC(CH ) ] precursor and H S plasma with different deposition conditions,41 which suggests the importance of the ALD conditions to form the Turing pattern. High AR nano-trenches (depth 230 nm, opening width ~50 nm) were fabricated on thermal oxide wafers (Si wafers with 300 nm thermally oxidized SiO on top) by e-beam lithographic patterning and inductively coupled plasma (ICP) etching. On these trench structures were conformally coated the Tp-Ta S layer using the same abovementioned ALD TaS and annealing process. To achieve selective Tp-Ta S coating on a patterned Co/SiO test structure, the sample was immersed in aniline for 30 min prior to ALD. Presumably, aniline could be selectively adsorbed on Co, thereby inhibiting the ALD of TaS on Co.

Material Characterizations. The Turing morphology was examined by TEM (Jeol, JEM-3200FS, 200 kV) and AFM (Bruker, Dimension Icon). The elemental composition was analyzed XRF (Shimadzu, EDX-LE). Electron diffraction pattern was simulated by ReciPro using the kinematical and excitation error model. Raman analysis was performed using the 532 nm laser excitation (Thermo Fischer, DXR). Cross-sectional microstructures were examined by STEM (LFEI, Titan Themis 200, 200 kV) equipped with EDS (Super-X EDS). Wettability and selectivity were examined by SEM ((Zeiss, Supra55) equipped with EDS (Oxford Instruments, X-maxN 20). The adhesion energy was measured by 4-PB test (Lloyd Instruments, LRX Plus).

TDDB measurement. TDDB capacitor structures were fabricated on thermal oxide wafers (50 nm SiO on n++-Si wafers). Top Cu electrodes (thickness 100 nm, diameter 300 μm) were thermally evaporated through a shadow mask, and the Cu electrodes were further capped by 20 nm Al to prevent Cu oxidation. Another 20 nm Al was thermally evaporated on the backside of the n++-Si wafer to ensure good electrical bottom contact. For the benchmarking barrier layers, the s-TaS layer was synthesized by a two-step process, where a Ta pure metal layer was first e-beam evaporated on SiO and then sulfidized under the H S plasma at 400 °C for 20 min (H S flow 10 sccm, RF power 70 W).9 The benchmarking TaN layer was synthesized by thermal ALD from TBTDET and NH (deposition temperature 250 °C).42 The TDDB current traces were monitored by a probe station using an Agilent B1500A semiconductor parameter analyzer.

Interconnect resistance estimation. The BEOL metal line/via resistances were calculated at the M1/V1 level using the CD values from the IRDS™ Roadmap.1 The size dependences of the Cu and Co resistivities were taken into account.43, 44 The resistivity of 0.9 mΩ cm was used for TaN,42 and the resistivity of 5.0 mΩ cm was used for Tp-Ta S , as measured by the four-point probe method (Napson, RT-3000/RG-2000).

DFT calculations. The DFT calculations were performed using the Vienna Ab initio Simulation Package (VASP 6.3.2)45 with the projector-augmented wave (PAW)46 method and Perdew-Burke-Ernzerhof (PBE)47 generalized gradient approximation (GGA) exchange-correlation functional. 5p65d36s2 (Ta), 3s23p4 (S), and 3d104s1 (Cu) were treated as valence electrons. To account for van der Waals interactions, the DFT-D3 dispersion correction was applied.48 A plane-wave energy cutoff of 500 eV and a Monkhorst Pack k-point mesh with a k-spacing of 0.03 Å−1 were adopted, which could ensure the total energy convergence within 10−6 eV/atom and forces below 0.01 eV/Å for structural optimization. Gaussian smearing with 0.02 eV smearing width was used. Spin-polarized calculations were conducted to model the magnetic ordering with the initial atomic coordinates derived from the experimental crystallographic data. A vacuum spacing of 20 Å was added to the slab model to minimize the interaction between adjacent slabs. The formation energies (E) of Ta S (can be rewritten as Ta m+n S 2m ) were calculated by E = E (Ta m+n S 2m ) − E (Ta m S 2m ) − nμ , where E (Ta m+n S 2m ) and E (Ta m S 2m ) are the total energies of Ta m+n S 2m and Ta m S 2m , respectively, n is the number of Ta atoms intercalated into the TaS bilayer structure, and μ is the chemical potential of Ta atom, which varies upon experimental conditions.49 Phonon modes were calculated using density functional perturbation theory (DFPT) with GGA. For this purpose, the structures were reoptimized using more strict convergence criteria of 10−7 eV for total energy and 10−6 eV/Å for forces, and the VASP-reported dynamical (Hessian) matrices were used to obtain the phonon modes by Phonopy.50, 51 The Cu migration barriers were calculated at 0 K temperature based on Ta S (TaS ) and Ta S (Ta S ) structures (Supplementary Table 1) using the climbing-image nudged elastic band (CI-NEB) method with force convergence <0.02 eV/Å, and three intermediate images were inserted along each path between neighboring local minima.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors investigated Turing-patterned Ta2S3 ultrathin films as a Cu diffusion barrier. While the adaptation of Ta2S3 layers for this purpose is intriguing, the manuscript lacks sufficient experimental evidence to convincingly demonstrate superior barrier properties.\n1. What is the thickness of a monolayer of TaS2? How many TaS2 layers correspond to the 1.2 nm thickness of the Ta2S3 film? It appears that Cu can easily diffuse between adjacent TaS₂ layers, as there are potential diffusion paths at the interfaces between discrete layers.\n2. The authors argued that Turing pattern periodicities are approximately to be 19 nm and 25 nm for 9 and 13 ALD cycles, respectively. However, such large periodicities may not be suitable for advanced interconnect applications, as they result in increased surface area, which is undesirable in scaled device nodes.\n3. In Figure 3k, high-resolution STEM imaging is necessary to clearly resolve the Turing structure. Also, the authors should provide additional evidence to quantify and validate the differences in surface area where these periodic patterns are observed as pointed out in question #2.\n4. TDDB measurements were used to evaluate the barrier performance of 1.5 nm TaS2, 2.5 nm TaN, and 1.2 nm Ta2S3 films using capacitor structures. However, further experimental data are required. The electric field applied across the SiO2 layer may vary due to differences in the resistance of each barrier layer. What is the resistivity or resistance of the Ta2S3 layer? If its resistance is significantly higher, the voltage drop across the SiO2 layer would be lower than in the other samples, potentially skewing the TDDB results.\n5. The authors should provide diffusion barrier performances on different surface of Ta2S3 because of the existence of periodicities. Also, Cu diffusion profiles obtained via appropriate imaging or analytical techniques should be compared.\n6. Finally, the ALD process using TBTDET and H₂S is not a unique approach. The authors should discuss what distinguishes their process or structure from prior works using similar chemistry.", "reviewer_2": "In this submission the authors present an interesting study of a novel copper diffusion barrier for CMOS back end of line interconnects, using a modified structure of TaS2, with composition Ta2S3.\nThe work is comprehensive and generally well described and covers all the analysis needed to demonstrate the functioning of Ta2S3 as a copper diffusion barrier, including high aspect ratio structure testing and some selective ALD to deposit the barrier on the sidewalls.\nI would be happy to recommend publishing this work as it is a good advance in an important area that is a potential roadblock for future CMOS scaling.\nSome comments\n- Page 3, Line 60/61 are these % values calculated by the authors or from other work?\n- Page 4, it would be useful to define a \"Turing mechanism\" for the reader not familiar with the concept.\n- In many figures, the text is very small and hard to read. Figure 1 caption should have \"TP\" instead of \"Tp\" for Turing Patterned\n- Is this Turing Pattern structure and hence the performance of the diffusion barrier reproducible overy multiple samples? It appears that this is stochastic so does the author's ALD approach ensure that the TP appears consistently over different samples\n- I am not sure what Figure 2(d) is saying to help understand the ALD mechanism - it could probably be removed\n- Figure 3(c) is a good example of the small font size. Can authors indicate in this figure the approximate conditions for the ALD material?\n- There are a number of places in the work where comparison with TaS2 is made but others where a comparison could be made but is lacking. For example, page 10 on copper wettability.\n- Can the ALD process be used to control the extent of Ta incorporation between the TaS2 layers and assess the role of Ta-content\n- Figure 4 has Gibbs free energies - what temperature is used? What species are used in the vibrational entropy calculation?\n- Figure 5 caption - replace \"penal\" with \"panel\"\n- How stable is the Ta2S3 material over time?\n- Could a comparison of Co or Ru with Cu be made?\n- Experimental\n(1) what is the ALD purge time?\n(2) why does the palsma ALD promote the TP and thermal ALD does not?\n(3) DFT part: valence electron configurations are missing\n(4) what are the lattice dimensions and supercell expansion of TaS2 and Ta2S3\n(5) How is formation energy of adding Ta calculated?\n(6) how many images are in the CI-NEB pathways\n(7) what smearing is used?\n(8) what is convergence criterion and displacement for the DFPT\n- Page 11, since authors did DFT calculations, it would be little effort to explore the interaction strength of Cu on the TaSx materials and compare how strong the proposed Cu-S bond formation is. This could also be compared with Cu-MoS2"}, "author_response": "Responses to Comments raised by Reviewers\nFor Reviewer #1:\nThe authors investigated Turing-patterned Ta S ultrathin films as a Cu diffusion\n2 3\nbarrier. While the adaptation of Ta S layers for this purpose is intriguing, the\n2 3\nmanuscript lacks sufficient experimental evidence to convincingly demonstrate\nsuperior barrier properties.\nAuthor response: We sincerely thank you for reviewing our manuscript and providing\nthe constructive comments. According to your comments, we have carefully revised\nour manuscript, and the following are our point-to-point responses to your comments.\nIn particular, the superior barrier properties of the Turing-patterned Ta S are best\n2 3\nreflected in Figure 4 for the excellent wetting, adhesion properties, and, more\nimportantly, the much-improved TDDB barrier performance. Additionally, using our\nfabrication approach, the Turing-patterned Ta S can also be conformally deposited into\n2 3\nhigh aspect ratio trench structures and selectively deposited only on dielectrics, as\nshown in Figure 5, which are essential in terms of industrial manufacturability to\nresolve the important bottleneck for future CMOS interconnect scaling. Therefore, we\nbelieve the data are convincing and important.\n1) What is the thickness of a monolayer of TaS ? How many TaS layers correspond to\n2 2\nthe 1.2 nm thickness of the Ta S film? It appears that Cu can easily diffuse between\n2 3\nadjacent TaS layers, as there are potential diffusion paths at the interfaces between\n2\ndiscrete layers.\nAuthor response: The thickness of a monolayer of TaS is approximately 0.70 nm. The\n2\n1.2 nm thickness of the Ta S film corresponds to approximately two TaS layers on\n2 3 2\naverage (Figure 2e,f). Indeed, for the conventional TaS , which is a 2D material with\n2\nvan der Waals (vdW) interlayer interaction, Cu can easily diffuse between adjacent TaS\n2\nlayers; however, for our Ta S , of which the structure (as schematically illustrated in\n2 3\nFigure 1c or 3b) has a bridging Ta layer to bond together the adjacent TaS layers, the\n2\ndiffusion energy barrier for Cu is extremely high, so there is essentially no diffusion\npath at the interface between discrete layers. Figure 4e,f compares the DFT-calculated\nCu diffusion barriers for TaS and Ta S , where the Cu interlayer diffusion barrier is\n2 2 3\nextremely high to be at least 11.98 eV for Ta S (and can even be 22.53 eV, if transition\n2 3\nstates are considered, see Supplementary Fig. 10), as compared to only 0.17 eV for the\nvan der Waals TaS . Also, experimentally, we found that after the TDDB stress, the\n2 3 2 2 2\npronounced Cu signals (see Supplementary Figs. 6 and 7 in Author response to\nComment 5). These results are in good consistence with the above mechanism, because\nfor the conventional vdW 2D materials (such as TaS ), once Cu diffuses one layer step\n2\ndownward, it can diffuse laterally in a very fast pace to fill the entire vdW interlayer\nspace (see Author response to Comment 5). Therefore, it is important to have the Ta-\nbridged Ta S material to serve as the barrier layer.\n2 3\nTo address your comment, in the revised manuscript, we revised Figure 1c by\nadding “×” symbols to denote the interlayer diffusion paths that are of very high\nmigration barriers (these paths are essentially blocked in Tp-Ta S ).\n2 3\nRevised Figure 1(c). Schematic illustration of the possible Cu ion diffusion paths for\nthe conventional 2D barrier materials and Tp-Ta S barrier. Red “×” symbols denote\n2 3\nthe diffusion paths that are of very high migration barriers.\n2) The authors argued that Turing pattern periodicities are approximately to be 19 nm\nand 25 nm for 9 and 13 ALD cycles, respectively. However, such large periodicities may\nnot be suitable for advanced interconnect applications, as they result in increased\nsurface area, which is undesirable in scaled device nodes.\nAuthor response: We realized that the description in the original manuscript may have\ncaused some confusion, and we are sorry for the confusion. Here in this work, the\nTuring pattern periodicities (i.e., approximately 19 nm and 25 nm for 9 and 13 ALD\ncycles) refer to the in-plane periodicities (i.e., in the film lateral direction); whereas the\nactual film height (i.e., in the out-of-plane direction) is only of at most one monolayer\nTaS thickness (ca. 0.7 nm) in difference over a period. Experimentally, as shown in\nq 2 3\nonly 0.453 nm, which is rather negligible as compared to the CD size of the advanced\ninterconnect structures (e.g., CD = 15 nm for the 7Å node, as shown in Figure 5).\nTherefore, the Turing pattern would not appreciably increase the surface area.\nExperimentally, the Tp-Ta S film coated on high AR trench sidewall also support in\n2 3\nthis regard (see Figure 5f).\nTo better illustrate the concept of using the Tp-Ta S barrier for advanced\n2 3\ninterconnect structures, we added the following scheme as Supplementary Fig. 2 in the\nrevised submission. In this illustrative example, an advanced interconnect trench\nstructure (CD = 15 nm, AR = 1.5, corresponding to the 7Å node) is first covered by the\n1.2 nm Tp-Ta S barrier layer and then filled with Cu. The benefit of this scheme (for\n2 3\nreduced line resistance) essentially comes from the enlarged cross-sectional area of the\nfilled Cu, owing to the thinned down barrier layer (see Figure 1b and Figure 5a). The\nin-plane Turing periodicity is not directly relevant to the barrier layer coverage, nor is\nthe barrier layer roughness a large portion of the CD size – therefore, we consider this\nTuring pattern concept is compatible with the scaled device nodes.\nTo address your comment, in the revised manuscript, we added Supplementary Fig.\n2, revised Figure 2e,f to display R , and also added the following content:\nq\n“It is worth noting that both Tp-TaS and Tp-Ta S films were quite smooth, with\n2 2 3\nroot-mean-square (rms) roughnesses (R ) of 0.509 and 0.453 nm, respectively. These\nq\nsmooth films are fairly compatible with the advanced interconnect structures\n(Supplementary Fig. 2).”\nSupplementary Fig. 2 Schematic illustration of an advanced interconnect trench\nstructure (CD = 15 nm, AR = 1.5, corresponding to the 7Å node) filled with Cu and the\nTp-Ta S barrier. The “×” symbol denotes the possible Cu diffusion path that is of very\n2 3\nhigh migration barrier (see Figure 4f).\n3) In Figure 3k, high-resolution STEM imaging is necessary to clearly resolve the\nTuring structure. Also, the authors should provide additional evidence to quantify and\nvalidate the differences in surface area where these periodic patterns are observed as\npointed out in question #2.\nAuthor response: As per your comment, we revised the high-resolution STEM image\n(Figure 3k) as the following. We removed the overlaid eye-guiding lines, leaving the\noriginal high-resolution STEM image for better clarity (Figure 3k upper); below this\nSTEM image, we also added a comparative scheme (Figure 3k lower) to better illustrate\nthe Turing structure. Additionally, we also added another cross-sectional STEM image\nas Supplementary Fig. 3 showing a larger view of the Tp-Ta\n3\nfilm. In this image, the\nperiods of the Turing structure (1/f) are labeled, which are in good consistence with the\ntop-view high-resolution TEM images shown in Figures 2c and 3e.\nAs for your comment regarding the surface area, we have mostly responded it in\nour Author response to Comment 2. In brief, the rms roughness (R ) of the 13-cycle Tp-\nq\nTa S film is only 0.453 nm, which is negligible as compared to the CD size of the\n2 3\nadvanced interconnect structures (typically, CD = 15~17 nm for the advanced nodes, as\nshown in Figure 5). Therefore, the Turing pattern would not appreciably increase the\nsurface area (see Supplementary Fig. 2 for illustration). Experimental STEM results\n(Figure 5f) also support in this regard. To further quantify the increase of the surface\narea, we adopted a model (ACS Appl. Polym. Mater. 2025, 7, 19, 13276–13285) based\non the sinusoidal approximation of the surface height modulation. Assuming that the\nlateral modulation wavelength of the surface height is equal to the Turing periodicity\n(1/f = 25 nm), given R = 0.453 nm, the percentage increase of the surface area (ΔA/A)\nq\ncan be calculated to be ΔA/A = 2π²R ²f² = 0.6%, which is indeed negligible. Therefore,\nq\nwe are confident that the negligible increase of surface area is not a concern for adopting\nthe Turing-patterned barrier concept for advanced interconnect technology.\nTo address your comment, in the revised manuscript, we revised Figure 3k, added\nSupplementary Fig. 3, and revised the relevant text as the following:\n“As shown in Figure 3k (see also Supplementary Fig. 3), a clear intertwined layer\nstructure can be observed, which confirmed the successful synthesis of the Turing-\npatterned Ta S film.”\n2 3\nRevised Figure 3(k). Cross-sectional STEM image of the Tp-Ta S film (upper) and the\n2 3\ncorresponding schematic illustration (lower).\nSupplementary Fig. 3 A larger-view cross-sectional STEM image of the Tp-Ta S film.\n2 3\nThe Turing periodicities (1/f) are marked.\n4) TDDB measurements were used to evaluate the barrier performance of 1.5 nm TaS ,\n2\n2.5 nm TaN, and 1.2 nm Ta S films using capacitor structures. However, further\n2 3\nexperimental data are required. The electric field applied across the SiO layer may\n2\nvary due to differences in the resistance of each barrier layer. What is the resistivity or\nresistance of the Ta S layer? If its resistance is significantly higher, the voltage drop\n2 3\nacross the SiO layer would be lower than in the other samples, potentially skewing the\nAuthor response: Actually, Ta S is metallic, so the voltage drop on the Ta S layer is\n2 3 2 3\nnegligible. As shown in the Experimental section (Interconnect resistance estimation),\nthe resistivity of Tp-Ta\n3\nwas measured to be 5.0 mΩ cm by the standard four-point\nprobe method. Likewise, the resistivities of s-TaS and TaN were 6.1 and 0.9 mΩ cm,\n2\nrespectively. In contrast, the resistivity of SiO was measured to be 3.3×1011 Ω cm from\n2\nthe leakage current of the Cu/SiO /n++-Si capacitor. Therefore, given the TDDB\n2\ncapacitor area (A) of 7.1×10−4 cm2 (300 μm in diameter), the resistances of the barrier\nlayers (R = ρt/A) are R = 1.3×10−6 Ω, R = 3.2×10−7 Ω, and R = 8.5×10−7\ns-TaS2 TaN Tp-Ta2S3\nΩ. In contrast, the resistance of the 50 nm SiO layer is R = 2.3×109 Ω, which is 15\norders of magnitude larger. Therefore, the fractional voltage drop on these barrier layers\nare all on the order of 10−16 (calculated from R /R ), so the voltage drop across\nbarrier total\nthe SiO layer is essentially identical for all these cases.\nfigure as Supplementary Fig. 5 to illustrate the voltage profile across the\nCu/barrier/SiO /n⁺⁺-Si stack during the TDDB stress.\n2+\nSupplementary Fig. 5 Schematic illustration of the voltage profile across the\nCu/barrier/SiO /n⁺⁺-Si stack during the TDDB stress. For the thin metallic barrier layer\n2\nused in this work (i.e., 1.5 nm s-TaS , 2.5 nm TaN, or 1.2 nm Tp-Ta S ), the voltage drop\n2 2 3\nis almost fully on the insulating SiO layer (the fractional voltage drop on the barrier\n2\nlayer is only on the order of 10−16).\n5) The authors should provide diffusion barrier performances on different surface of\nTa S because of the existence of periodicities. Also, Cu diffusion profiles obtained via\n2 3\nappropriate imaging or analytical techniques should be compared.\nAuthor response: We felt that there might be some misunderstanding – the periodicity\nitself is not of the paramount key factor to its superior barrier performance. As our\nTuring-patterned barrier concept described in Introduction and Figure 1, the 2D layers\nthemselves are already good barriers, but the real problem is the grain boundary\ndiffusion. Forming the Turing pattern can well resolve this issue by creating an\nintertwined layer structure (Figure 1c) that forces the Cu atoms to first migrate laterally\nbetween the 2D TaS layers; together with the special attribute of the Ta S structure,\n2 2 3\nwhich has the interlayer bridging Ta atoms that can strongly bond together the\nneighboring TaS layers and substantially increase the interlayer Cu diffusion barrier\n2\n(extremely high barrier of at least 11.98 eV), the Cu lateral diffusion path is essentially\nblocked. Therefore, it is the intertwined layer structure of Ta S that contributes most\n2 3\nto its superior barrier performance. The size of the periodicity contributes second to the\nbarrier performance, but as long as there is enough intertwined overlay, the superior\nbarrier performance should be well retained.\nRegarding the surface area, as we pointed out in Author responses to Comments 2\nand 3, the rms roughness (R ) of our Tp-Ta S film is only 0.453 nm, which is quite\nq 2 3\nsmall, and thus, the Turing pattern would not appreciably increase the surface area.\nOn a separate note, our TDDB measurement uses a circular-shaped capacitor\nelectrode of 300 μm in diameter, which is 4 orders of magnitude larger than the Turing\nperiodicity (e.g., 25 nm); therefore, any effect from the periodicity has already been\nconsidered and averaged. Also, it is worth noting that TDDB quantifies the time for the\nweakest point to break down. Equivalently speaking, TDDB finds out the weakest point\n(worst case) of all possible structural scenarios (including different local stacks,\ndifferent boundary structures, different local periodicities, etc.). Therefore, our TDDB\nresults are highly reliable to gauge the barrier performance for practical applications.\nAs per your request for additional imaging for the Cu diffusion profile, we further\ncharacterized the cross-sectional STEM images of the Cu/Tp-Ta S /SiO /n++-Si and\n2 3 2\nCu/s-TaS /SiO /n++-Si stacks after the TDDB measurements. Additional data are added\n2 2\nas Supplementary Figs. 6 and 7 in the revised submission. As shown in Supplementary\nFigs. 6a-d and 7a, the Tp-Ta S barrier layer was well preserved after the TDDB stress,\n2 3\nand the upper Cu layer was completely blocked to reach the underneath SiO . (Notably,\n2\nthe TDDB stress always stops at the first breakdown event, so the single breakdown\nlocation can hardly be caught by a random cut for the cross-sectional STEM imaging.)\nIn contrast, as shown in Supplementary Figs. 6e-h and 7b, the s-TaS layer failed to\n2\nblock Cu after TDDB, as clear Cu EDS signals can be observed in between the s-TaS\n2\nand SiO layers. In addition, it is worth noting that Cu seemed to form a complete layer\n2\nunderneath s-TaS , which is consistent with the proposed lateral interlayer diffusion\n2\nmechanism for vdW 2D materials. (Once Cu diffuses one layer step downward, it can\ndiffuse laterally in a very fast pace to fill the entire vdW interlayer space.)\nTherefore, these additional STEM results can well support our proposed\nmechanisms.\nTo address your comment, in the revised manuscript, we added Supplementary\nFigs. 5 and 6, along with the following content:\n“We further took the cross-sectional STEM/EDS characterizations of the Cu/Tp-\nTa S /SiO /n++-Si and Cu/s-TaS /SiO /n++-Si stacks after the above TDDB stress. As\n2 3 2 2 2\nshown in Supplementary Figs. 6 and 7, the Tp-Ta S barrier layer could well block Cu\n2 3\nfrom diffusing through; whereas the s-TaS layer failed to block Cu, showing\n2\npronounced Cu signals at the s-TaS /SiO interface, possibly owing to the facile Cu\n2 2\ndiffusion in between the vdW 2D TaS layers, as will be discussed later.”\nCu/barrier/SiO /n++-Si stacks after the TDDB measurements. (a) STEM image and\n2 3 2\nstack. (e) STEM image and EDS maps of (f) Cu, (g) Ta, and (h) S for the TDDB-stressed\nCu/s-TaS /SiO /n++-Si stack. “1” and “2” mark the locations for the line scans shown\n2 2\nin Supplementary Fig. 7.\nSupplementary Fig. 7 EDS line-scan profiles of the atomic concentrations across the\nTDDB-stressed (a) Cu/Tp-Ta S /SiO /n++-Si and (b) Cu/s-TaS /SiO /n++-Si stacks\n2 3 2 2 2\nshown in Supplementary Fig. 6. Particularly in (b), a clear shoulder of Cu signal can\nbe seen in the dashed box 2, highlighting the accumulation of Cu at the s-TaS /SiO\n2 2\ninterface.\n6) Finally, the ALD process using TBTDET and H₂S is not a unique approach. The\nauthors should discuss what distinguishes their process or structure from prior works\nusing similar chemistry.\nAuthor response: Actually, to the best of our knowledge, there is no prior report on\nthe ALD process of using TBTDET and H₂S for TaS (see ALD process database\nx\nhttps://www.atomiclimits.com/alddatabase/). A similar (but not the same) process was\nreported in 2024 by Deijkers et al.,41 who used a different Ta precursor of\nTa[N(CH ) ] [NC(CH ) ] together with H S plasma to deposit TaS . In their work, aside\n3 2 3 3 3 2 x\nfrom using a different Ta precursor, their H S plasma conditions were also rather\n2\ndifferent from ours (they used much higher H S concentration and much higher plasma\n2\npower), so their deposited TaS films were quite different from ours (for instance, their\nx\n150°C-deposited films had S-to-Ta ratio of 4.1) – certainly, no Turing pattern was\nreported in their work.\nTo address your comment, in the revised manuscript, we added the following\ncontent in the Experimental Section (Fabrication of Tp-Ta S ):\n2 3\n“also, the Turing pattern did not show in another recent ALD TaS report using\nx\nTa[N(CH ) ] [NC(CH ) ] precursor and H S plasma with different deposition\n3 2 3 3 3 2\nconditions,41 which suggests the importance of the ALD conditions to form the Turing\npattern.”\nAdded Reference:\n41. Deijkers JH, et al. Atomic layer deposition of transition metal chalcogenide\nTaS using Ta[N(CH ) ] [NC(CH ) ] precursor and H S plasma. J. Vac. Sci.\nx 3 2 3 3 3 2\nTechnol. A 42, 062404 (2024).\nFor Reviewer #2:\nIn this submission the authors present an interesting study of a novel copper diffusion\nbarrier for CMOS back end of line interconnects, using a modified structure of TaS ,\n2\nwith composition Ta S .\n2 3\nThe work is comprehensive and generally well described and covers all the analysis\nneeded to demonstrate the functioning of Ta S as a copper diffusion barrier, including\n2 3\nhigh aspect ratio structure testing and some selective ALD to deposit the barrier on the\nsidewalls.\nI would be happy to recommend publishing this work as it is a good advance in an\nimportant area that is a potential roadblock for future CMOS scaling.\nAuthor response: We sincerely thank you for reviewing our manuscript and providing\nthe constructive comments. According to your comments, we have carefully revised\nour manuscript, and the following are our point-to-point responses to your comments.\n1) Page 3, Line 60/61 are these % values calculated by the authors or from other work?\nAuthor response: These values were calculated by us. The detailed calculations are\nprovided in the Experimental section (Interconnect resistance estimation).\nTo make it clear, in the revised manuscript, we added the calculation description\nin the caption of Figure 1b, as the following:\n“CD = 16 nm and AR = 1.5 are used for the calculation, as per the IRDS™\nRoadmap1 for the M1 lines for the 15-Å node.”\n2) Page 4, it would be useful to define a \"Turing mechanism\" for the reader not familiar\nwith the concept.\nAuthor response: As per your suggestion, in the revised manuscript, we added a note\nof “(see next section for details)” after “Turing mechanism” on Page 4, and explain the\nTuring mechanism concept in detail on Page 5, as the following:\n“In general, the formation of Turing pattern (i.e., Turing mechanism) relies on\nsimultaneous presence of an activator and an inhibitor species in a reaction-diffusion\nsystem and the inhibitor diffuses significantly faster than the activator.20, 21”\n3) In many figures, the text is very small and hard to read. Figure 1 caption should have\n\"TP\" instead of \"Tp\" for Turing Patterned.\nAuthor response: As per your comment, in the revised manuscript, we enlarged the\ntext font size in Figures 3c, 4a-d, and 5a,c-f. As for the abbreviation for “Turing\npatterned”, we choose to use “Tp” for consistency throughout the manuscript.\n4) Is this Turing Pattern structure and hence the performance of the diffusion barrier\nreproducible over multiple samples? It appears that this is stochastic so does the\nauthor's ALD approach ensure that the TP appears consistently over different samples.\nAuthor response: Yes, this Turing pattern structure is highly reproducible, thanks to\nthe high reproducibility of ALD. The excellent diffusion barrier performance is also\nhighly reproducible. In fact, each set of the TDDB data (Figure 4d) were measured on\n15 separate capacitor devices, so the reproducibility of the excellent barrier\nperformance has already been by guaranteed to some extent. In this revised manuscript,\nwe added Supplementary Figs. 8 and 9 (mainly to address your Comments 11 and 12),\nwhich provide more solid data to support the reproducibility of the excellent barrier\nperformance.\n5) I am not sure what Figure 2(d) is saying to help understand the ALD mechanism - it\ncould probably be removed.\nAuthor response: Figure 2(d) serves two roles: (1) it is an illustration of our proposed\nlayer-by-layer Frank-van der Merwe growth mode; (2) we marked “1/f” in this figure\nto better illustrate the characteristic periodicity of the Turing pattern. Given these two\nimportant roles, we would like to keep Figure 2(d) in the manuscript.\n6) Figure 3(c) is a good example of the small font size. Can authors indicate in this\nfigure the approximate conditions for the ALD material?\nAuthor response: As per your comment, we increased the font size in Figure 3(c) in\nthe revised manuscript. Also, the experimental ALD film growth condition\ncorresponded to the sulfur-rich condition as indicated in the figure. Therefore, in the\nrevised manuscript, we also added the relevant statement as the following for\nclarification:\n“As shown in Figure 3c, under the S-rich condition, which corresponds to the\nexperimental ALD condition, TaS\n2\nis the stable stoichiometry; whereas under the Ta-\nrich condition, which can be achieved by vacuum or H /Ar gas annealing, Ta S\n2 2 3\nbecomes more stable.”\nRevised Figure 3. (c) DFT-calculated formation energies for a series of Ta S\nx y\nstoichiometries.\n7) There are a number of places in the work where comparison with TaS is made but\n2\nothers where a comparison could be made but is lacking. For example, page 10 on\ncopper wettability.\nAuthor response: As per your comment, we added the copper wettability results for\nTaS as Supplementary Fig. 4 in the revised manuscript. As shown below, the Cu film\n2\non s-TaS was also of very good wettability, probably owing to a similar strong Cu−S\n2\nbond at the Cu/s-TaS interface. In the revised manuscript, we added the following\n2\ncontent:\n“Note that the Cu film on s-TaS was also of good wettability (Supplementary Fig.\n2\n4).”\nSupplementary Fig. 4 Top-view SEM image of Cu on the s-TaS surface, showing good\n2\nwettability of s-TaS .\n2\n8) Can the ALD process be used to control the extent of Ta incorporation between the\nTaS layers and assess the role of Ta-content?\nannealing in H /Ar to tune the Ta:S ratio – we found that this two-step approach is\n2\nhighly reproducible to fabricate high-performance uniform films. In theory, tuning the\nALD condition should be able to tune the Ta:S ratio in the as-deposited films, but the\nfollow-up annealing step may change the composition in the final films.\n9) Figure 4 has Gibbs free energies - what temperature is used? What species are used\nin the vibrational entropy calculation?\nAuthor response: Actually, we used 0 K temperature for data in Figure 4 (thus, no\nvibrational entropy was included). As the migration barriers compared in Figure 4e,f,\nthe discrepancy is so large (0.11 vs. 11.98 eV for intermediate states) that the entropy\ncorrection is not necessary.\nTo clarify this issue, we revised Figure 4(e,f) by replacing “ΔG” with “Migration\nbarrier”, and added the description of 0 K calculation in the Experimental section (DFT\ncalculation) of the revised manuscript, as the following:\n“The Cu migration barriers were calculated 0 K temperature...”\nRevised Figure 4 (e,f). DFT-calculated relative energies of the transition/intermediate\nstates for the Cu atom migration in the layered structures of (e) TaS and (f) Ta S .\n2 2 3\n10) Figure 5 caption - replace \"penal\" with \"panel\"\nAuthor response: We corrected this typo in the revised manuscript.\n11) How stable is the Ta S material over time?\n2 3\nAuthor response: Our Ta S is very stable over time. We measured the TDDB\n2 3\nperformance of the Cu/Tp-Ta S /SiO /n++-Si stack after one-month storage. As shown\n2 3 2\nin Supplementary Fig. 8, no significant change was observed. Therefore, the Ta S\n2 3\nmaterial is very stable over time. To address your comment, in the revised manuscript,\nwe added Supplementary Fig. 8, together with the following content:\n“It is worth noting that the Tp-Ta S barrier was also fairly stable. We measured\n2 3\nthe TDDB performance of the Cu/Tp-Ta S /SiO /n++-Si stack after one-month storage,\n2 3 2\nand as shown in Supplementary Fig. 8, no significant change was observed.”\nSupplementary Fig. 8 TDDB results for the as-prepared Cu/Tp-Ta S /SiO /n++-Si stack\n2 3 2\nand the same stack after one month storage.\n12) Could a comparison of Co or Ru with Cu be made?\nAuthor response: As per your comment, we added the TDDB data for Co and Ru as\nSupplementary Fig. 9 in the revised manuscript. As shown below, Tp-Ta S is also a\n2 3\ngood barrier for Co and Ru. After inserting the 1.2 nm Tp-Ta S barrier layer, the\n2 3\ncharacteristic TTF for Co increased from 16.1 to 186.2 s, and the TTF for Ru\n63.2% 63.2%\nincreased from 78.6 to 504.9 s. These results well demonstrate that the Turing-patterned\nbarrier strategy is of very good generality, as the Tp-Ta S barrier not only worked for\n2 3\nCu but also worked for other two important BEOL metals of Co and Ru.\nTo address your comment, in the revised manuscript, we added Supplementary Fig.\n9, together with the following content:\n“In addition, we found that this Turing-patterned barrier layer strategy not only\nworked for Cu but also worked for other two important BEOL metals of Co and Ru. As\nshown in Supplementary Fig. 9, the characteristic TTF values were also found to\n63.2%\nsignificantly increase from 16.1 to 186.2 s for Co and from 78.6 to 504.9 s for Ru, after\nthe 1.2 nm Tp-Ta S barrier layer was inserted. These results well demonstrate that the\n2 3\nTuring-patterned barrier strategy is of very good generality.”\nSupplementary Fig. 9 Barrier performances of 1.2 nm Tp-Ta S for Co and Ru. (a)\n2 3\nTDDB results for Co/SiO /n++-Si and Co/Tp-Ta S /SiO /n++-Si stacks. (b) TDDB\n2 2 3 2\nresults for Ru/SiO /n++-Si and Ru/Tp-Ta S /SiO /n++-Si stacks.\n2 2 3 2\n13) what is the ALD purge time?\nAuthor response: The ALD purge time was 30 s for each half-cycle. We added this\ninformation in the Experimental section (Fabrication of Tp-Ta S ) of the revised\n2 3\nmanuscript, as the following:\n“…and the purge time was 30 s for each half-cycle”\n14) why does the plasma ALD promote the TP and thermal ALD does not?\nAuthor response: As we explain in Results and Discussion (Fabrication of Turing-\npatterned Ta S ), to realize the characteristic “short-range activation and long-range\n2 3\ninhibition” Turing mechanism, the chemical reaction system requires an inhibitor\nspecies that diffuses significantly faster than the counterpart activator species. In our\nALD reaction system, the H S plasma is critically important in order to generate the\n2\ninhibitor species on surface (i.e., inert sulfur clusters, e.g., S ) and accelerate the\n8\ninhibitor surface diffusion by kinetic energy transfer (via collision with energetic\nplasma species). In contrast, thermal ALD does not have such diffusion-accelerating\nmechanism, and therefore, the critical Turing conditions cannot be met under thermal\nALD.\n15) DFT part: valence electron configurations are missing.\nAuthor response: 5p65d36s2 (Ta), 3s23p4 (S), and 3d104s1 (Cu) were treated as the\nvalence electrons for the DFT calculation. We added this information in the\nExperimental section (DFT calculation) of the revised manuscript, as the following:\n“5p65d36s2 (Ta), 3s23p4 (S), and 3d104s1 (Cu) were treated as valence electrons.”\n16) what are the lattice dimensions and supercell expansion of TaS and Ta S ?\n2 2 3\nAuthor response: As per your comment, we added the following content in the\nExperimental section (DFT calculation) of the revised manuscript:\n“The Cu migration barriers were calculated … based on Ta S (TaS ) and Ta S\n18 36 2 24 36\n(Ta S ) structures (Supplementary Table 1)”\n2 3\nSupplementary Table 1 Lattice parameters of Ta S (TaS ) and Ta S (Ta S ) used\n18 36 2 24 36 2 3\nfor DFT.\nStructure a (Å) b (Å) c (Å)\nTa S (TaS ) 9.973 9.973 33.760\n18 36 2\nTa S (Ta S ) 9.840 9.840 33.760\n24 36 2 3\n17) How is formation energy of adding Ta calculated?\nAuthor response: As per your comment, we added the following content in the\nExperimental section (DFT calculation) of the revised manuscript:\n“The formation energies (E) of Ta S (can be rewritten as Ta S ) were\nf x y m+n 2m\ncalculated by E = E (Ta S ) − E (Ta S ) − nμ , where E (Ta S ) and\nf tot m+n 2m tot m 2m Ta tot m+n 2m\nE (Ta S ) are the total energies of Ta S and Ta S , respectively, n is the number\ntot m 2m m+n 2m m 2m\nof Ta atoms intercalated into the TaS bilayer structure, and μ is the chemical potential\nof Ta atom, which varies upon experimental conditions.49”\nAdded Reference:\n49. Wang M, et al. Enhanced magnetic anisotropy in two-dimensional 2H−TaS by\n2\nself-intercalation: A DFT study. J. Magn. Magn. Mater. 553, 168988 (2022).\n18) how many images are in the CI-NEB pathways\nAuthor response: Three images are used in the CI-NEB pathways. We added this\ninformation in the Experimental section (DFT calculation) of the revised manuscript,\nas the following:\n“…and three intermediate images were inserted along each path between\nneighboring local minima.”\n19) what smearing is used?\nAuthor response: Gaussian smearing with 0.02 eV smearing width was used. We\nadded this information in the Experimental section (DFT calculation) of the revised\nmanuscript, as the following:\n“Gaussian smearing with 0.02 eV smearing width was used.”\n20) what is convergence criterion and displacement for the DFPT\nAuthor response: For the DFPT calculation, more strict convergence criteria of 10−7\neV for total energy and 10−6 eV/Å for forces were used for the structural optimization.\nThen, the VASP-reported dynamical (Hessian) matrices were used to obtain the phonon\nmodes by Phonopy.50 In this approach, the phonon frequencies were calculated directly\nfrom the dynamical (Hessian) matrix, rather than from a finite-displacement supercell\nscheme, which avoids the need to construct explicit displaced configurations. Therefore,\nno explicit atomic displacement is set in our DFPT calculation.\nTo address your comment, in the revised manuscript, we added the following\ncontent in the Experimental section (DFT calculation):\n“For this purpose, the structures were reoptimized using more strict convergence\ncriteria of 10−7 eV for total energy and 10−6 eV/Å for forces, and the VASP-reported\ndynamical (Hessian) matrices were used to obtain the phonon modes by Phonopy.50,51”\nAdded References:\n50. Togo A, Chaput L, Tadano T, Tanaka I. Implementation strategies in phonopy\nand phono3py. J. Phys.: Condens. Matter 35, 353001 (2023).\n51. Togo A. First-principles Phonon Calculations with Phonopy and Phono3py. J.\nPhys. Soc. Jpn. 92, 012001 (2022).\n21) Page 11, since authors did DFT calculations, it would be little effort to explore the\ninteraction strength of Cu on the TaS materials and compare how strong the proposed\nx\nCu-S bond formation is. This could also be compared with Cu-MoS .\nx\nrather complicated. One has to consider different crystal orientations of Cu and their\ninterface lattice mismatches with respect to TaS – the latter is particularly complicated\nx\nto treat under the periodic DFT modeling framework, because the lattice mismatch can\npossibly impose an artificial strain that substantially affects the calculated interface\nenergy. In reality, this lattice strain can be largely relaxed via grain boundaries (because\nof the polycrystalline nature of Cu), which can hardly be accurately modeled by the\ncurrent DFT framework.\nNevertheless, Cu is known to bond strongly with S, so a strong Cu–TaS (or Cu–\nx\nMoS ) interaction strength is expected. On the other hand, in the real case, the adhesion\n2\nbottleneck is actually located at the underneath TaS –TaS (or MoS –MoS ) interfaces,\n2 2 2 2\nbecause of the very weak van der Waals interlayer interactions for the conventional 2D\nmaterials. In this work, Ta S is proposed to solve this issue, as its structure (see Figure\n2 3\n1c or 3b) contains a bridging Ta layer that can strongly bond together the adjacent TaS\n2\nlayers, thereby significantly enhancing the adhesion energy as we experimentally\nobtained (10.33 J/m²).\nTo address your comment, we construct simplified DFT models (Figure R1) by\nputting one Cu atom on TaS , Ta S , and MoS surfaces to calculate the Cu atom\n2 2 3 2\nadsorption energies (reflecting the Cu–S bond strengths). As the results shown below\n(Table R1), the calculated Cu adsorption energies are −2.65, −2.06, and −1.15 eV for\nTaS , Ta S , and MoS , respectively. These numbers are all substantially negative,\n2 2 3 2\nwhich suggest strong Cu–S bonds for all cases; and comparatively, the Cu–S bonds for\nTaS are stronger than that for MoS , suggesting a stronger Cu–TaS interaction than\nx 2 x\nCu–MoS .\n2 2 3 2\nSubstrate Cu adsorption energy (eV)\nTaS (Ta S ) −2.65\n2 18 36\nTa S (Ta S ) −2.06\n2 3 24 36\nMoS (Mo S ) −1.15\n2 18 36\nFigure R1. Schematic illustration of the DFT modeling structures for the adsorption\nenergy calculations of Cu on (a) TaS\n2\n, (b) Ta\n3\n, and (c) MoS\n2\n."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors revised the manuscript according to the reviewers’ suggestions, which strengthened the manuscript. I recommend this manuscript for publication.", "reviewer_2": "The authors have answered all reviewer comments and improved the manuscript by including extra information and clarifications.\nI am happy to recommend accepting this submission for publication."}, "author_response": ""}]

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