**Supercritical Antisolvent Processing of Nitrocellulose: Downscaling to Nanosize, Reducing Friction Sensitivity and Introducing Burning Rate Catalyst**

**Oleg S. Dobrynin 1, Mikhail N. Zharkov 1, Ilya V. Kuchurov 1, Igor V. Fomenkov 1, Sergey G. Zlotin 1, Konstantin A. Monogarov 2, Dmitry B. Meerov 2, Alla N. Pivkina 2 and Nikita V. Muravyev 2,\***

Received: 28 August 2019; Accepted: 26 September 2019; Published: 27 September 2019

**Abstract:** A supercritical antisolvent process has been applied to obtain the nitrocellulose nanoparticles with an average size of 190 nm from the nitrocellulose fibers of 20 μm in diameter. Compared to the micron-sized powder, nano-nitrocellulose is characterized with a slightly lower decomposition onset, however, the friction sensitivity has been improved substantially along with the burning rate increasing from 3.8 to 4.7 mm·s−1 at 2 MPa. Also, the proposed approach allows the production of stable nitrocellulose composites. Thus, the addition of 1 wt.% carbon nanotubes further improves the sensitivity of the nano-nitrocellulose up to the friction-insensitive level. Moreover, the simultaneous introduction of carbon nanotubes and nanosized iron oxide catalyzes the combustion process evidenced by a high-speed filming and resulting in the 20% burning rate increasing at 12 MPa. The presented approach to the processing of energetic nanomaterials based on the supercritical fluid technology opens the way to the production of nitrocellulose-based nanopowders with improved performance.

**Keywords:** nitrocellulose; supercritical antisolvent process; nanoparticles; combustion

### **1. Introduction**

Nitrocellulose (NC) is widely used as a base substance for conservation, adhesives, membranes and coatings [1–5] with an important application for solid rocket propellants and gunpowder production [6–8]. Long-term experience with nitrocellulose reveals some of its limitations, including certain thermal and mechanical hazards [9], inhomogeneity problems resulting from the organic nature of the cellulose precursor [10], and poor combustion performance. These challenges can be addressed by downscaling nitrocellulose to nanosized powder that should improve the burning rate, as was reported for another energetic substance, 1,3,5-trinitro-1,3,5-triazinane (RDX), whose nanoparticles demonstrate a twofold burning rate increase [11]. Also, the available data for the nanosized energetic materials show that their mechanical sensitivity is often beneficially reduced [12–14]. To date, several examples of NC with nanosized morphology have been reported. Electrostatic spinning of NC solution results in the formation of fibers with diameters from 80 to 300 nm [15,16]. The samples demonstrate the slight lowering of the decomposition onset from 198 ◦C to 190 ◦C [16], while the heat of reaction was increased [15] for nanosized NC as compared to as-received material. A complementary use of the sol-gel approach and sc-CO2 drying allows obtaining the NC aerogel with 30–50 nm particles that also

show a lower decomposition onset temperature compared to raw NC [17]. Besides, the thermolysis mechanism of this porous material is altered by its increased adsorption capacity. Zhang and Weeks explored the low temperature solvent evaporation technique to prepare nearly 500 nm-sized NC spheres (5 ◦C, dimethylformamide) and thin films of submicron NC (5–45 ◦C, MeOH or EtOH) [18]. Notably, the obtained films exhibited a more complete combustion with the rate 350% higher than that for films of ordinary nitrocellulose. From the chemical engineering prospective, there are definite complications related to the reported preparative techniques, i.e., physicochemical methods (e.g., anti-solvent [19,20], suspension dispersion [21], recrystallization [18]) are inevitably accompanied by the issues of wasted organic solvents utilization and removal of their traces trapped inside the product, whereas mechanical methods (e.g., milling [22]) are incompatible with the polymeric nature of NC. Thereby, the challenge to develop a 'greener' and more efficient approach for a NC morphology modification has not been addressed yet.

Tailoring the burning rate of the double-based propellants, i.e., nitrocellulose-based formulations, is usually performed by adding of 1–10 wt.% of catalysts or inhibitors [23,24]. The general conclusion drawn on the basis of the experience with catalysis of double-based propellants is that carbonaceous skeleton formed above the burning surface promotes the burning rate via the prolongated residence of catalyst particles and increase of heat flux to the condensed phase [24–26]. The desired carbonaceous layer can be formed either during combustion of some components, e.g., dinitrotoluene, dibutyl phthalate plasticizers, or by introduced prior to combustion additives (e.g., 0.5–3 wt.% of carbon black in [24]). Interestingly, the addition of carbon in the form of carbon nanotubes (CNTs) can increase the burning rate per se, by improving the thermal conductivity [27–29]. Thus, in the recent studies [30,31] the optimal level of multiwall carbon nanotubes additive was found to be about 1 wt.%, whereas at higher loadings the combustion is depressed. Nano-sized metal oxides render the catalytic action in diverse processes, including the decomposition and combustion of energetic materials [32–34]. Specifically, nano-Fe2O3 has been found to catalyze the decomposition of NC, with the governing effect of the contact surface between two components [35]. Therefore, improvement of the nitrocellulose combustion requires the composites designed to deliver the nano-sized additives to the reaction zone.

Recently, sub- and supercritical fluids media have proved valuable for modification and synthetic applications in chemical engineering [36,37]. Particularly, supercritical carbon dioxide (sc-CO2) is most attractive due to its low critical point (Pc = 7.38 MPa and Tc = 31.0 ◦C [38]), non-toxicity and environmental friendliness [39,40]. Supercritical Anti-Solvent (SAS) applications of CO2 are based on its very good miscibility with the majority of organic solvents and insolubility of the target substrate in resulting binary solvent systems (sc-CO2/solvent) that causes the precipitation (recrystallisation) of the final material [41,42]. The SAS technology has already been applied for micronization of nitramine explosives [43,44]. Herein, we report the first application of the supercritical antisolvent process to prepare the uniform nano-sized nitrocellulose powder. Furthermore, NC-based composites with carbon nanotubes and nano-Fe2O3 as the burning rate modifiers have been also fabricated via the SAS processing. Thermal behavior, mechanical sensitivity, and combustion performance were evaluated showing the enhanced burning rate and reduced sensitivity for prepared nano-nitrocellulose and its composites.

### **2. Materials and Methods**

### *2.1. Materials*

Commercial nitrocellulose was used as a precursor for synthesis and as a reference material. The powder is constituted by fibers with a characteristic diameter of about 20 μm, as shown in Figure 1b. The elemental composition of the raw NC was found to be 26.7 ± 0.2 wt.% C, 2.9 ± 0.1% H, and 11.5 ± 0.1% N. Multi-walled carbon nanotubes ("Taunit-M" trademark by NanoTechCenter Ltd., Tambov, Russia) were used as-received. In line with the manufacturer specification [45] the inner diameter of CNTs was ca.10 nm, the outer diameter near 20 nm, and the length no less than 2 μm

(Figure S1c,d, see Supporting information for details). The laboratory sample of nanosized α-iron oxide was produced by the plasma condensation technique and appeared as spherical particles 0.05–0.3 μm in diameter (Figure S1e,f, Supporting information). High purity carbon dioxide (Russian standard TU 2114-006-05015259-2014, Linde Gas Rus, Ltd., Balashikha, Russia) and acetone (Russian standard TU 2633-012-29483781-2009, Chimmed Ltd., Moscow, Russia) were used during the study.

**Figure 1.** Supercritical anti-solvent (SAS) fabrication of nitrocellulose (NC)-based composites: (**a**) scheme of the experimental setup, (**b**) scanning electron microscopy (SEM) of raw NC fibers, (**c**) product obtained under non-optimized conditions, (**d**) pure NC after SAS treatment in optimized conditions.

### *2.2. Fabrication of Composites*

Figure 1a illustrates the scheme of the experimental setup for the composite fabrication with the key units indicated. Cooled CO2 was transferred from the cylinder through the high-pressure pump into the 500 mL autoclave reactor with temperature-control. The pressure inside the autoclave was maintained with the automated back pressure regulator (ABPR). The NC acetone solution (2 wt.%) or suspension of selected fillers in this solution was fed with another pump to the filled with sc-CO2 autoclave through the capillary in the lid. The mutual diffusion of CO2 and acetone results in the precipitation of NC or its composite particles which were collected on the metallic filter at the bottom. The binary mixture CO2-acetone left the autoclave, went through ABPR and was divided inside the separator at the lower pressure controlled with the manual back pressure regulator (MBPR). The main units of the setup were connected with 1/8" steel pipes.

The general procedure implies filling with CO2 and pre-heating the system up to 40 ◦C. Once the desired temperature was achieved, the carbon dioxide feeding rate was set at 50 g·min−1. NC solution or suspension, prepared and maintained by ultrasonic vibrations, was fed at 1–2 mL·min−1 rate. After the solution/suspension was fully transferred to an autoclave, and CO2 flow pumped for an extra 30–40 min to flash off the residual acetone. Finally, the system was decompressed, and ~1 g of the target solid was taken out.

### **3. Results**

### *3.1. Optimization of the Process Conditions*

Preliminary runs resulted in the "herring-bone" structures instead of the desired uniform powder, as shown in Figure 1c. Several parameters of SAS process were tuned to improve the product morphology, i.e., the autoclave pressure, the ratio between the feeding rates of the solution and that of CO2 (f/F), NC concentration in the solution, and the additive material type (Section S3, Supporting information). Pressure variation within the 9–15 MPa range specifies the optimal value to be 10–12 MPa. At lower pressures the "herring-bone" was observed, whereas at higher pressures the uniform powder was obtained (cf. Figure 1c,d). The factor f/F controlled both the production time and CO2 consumption; however, its increase was shown to worsen the composite morphology. Thus, the solution flow rate appeared to be no higher than 2 mL·min−1 at the maximum CO2 feeding rate of 50 g·min−1 given by the experimental setup. The nitrocellulose concentration in the solution apparently determines the acetone and CO2 consumption and has to be maximized. On the other hand, its increase above 2 wt.% had adverse effects on the the NC dissolution time and the pump performance (due to an increase in the mixture viscosity). When the filler, i.e., and/or nano-Fe2O3, was introduced into the solution, the ultrasonic processing (30% of 70 W, 20 kHz) was applied continuously to prevent settling and to maintain the suspension shelf life.

### *3.2. Nano-Nitrocellulose and Composites*

Once the SAS process parameters were optimized, four samples were obtained: *NC(sc)* pure nitrocellulose after SAS, *NC*/*CNT(sc)* nitrocellulose with 1 wt.% of carbon nanotubes, *NC*/*Fe2O3(sc)* nitrocellulose with 5 wt.% of nano-iron oxide, *NC*/*CNT*/*Fe2O3(sc)* nitrocellulose with 1 wt.% of CNTs and 5 wt.% of nano-Fe2O3. The elemental composition of *NC(sc)* was unaffected by the processing: 26.7 ± 0.1 wt.% C, 2.9 ± 0.4% H, and 11.6 ± 0.1% N. In contrast to the large fibers of initial NC (Figure 1b), the product after the SAS processing with no additives appeared as nanopowder. Electron microscopy revealed that the microstructure of the powder was formed by the nano-sized spheres of 50–450 nm in diameter that were located on thin polymeric fibers (Figure 2a,b). Numerical particle size distribution gave the average diameter of 190 nm (Figure 2c, Section S4, Supporting information). Note, that transmission electron microscope (TEM) images have captured a sole fraction that is below 100 nm. Apparently, the bigger fractions were lost during the sample preparation.

**Figure 2.** Morphology of *NC (sc)*: scanning (**a**), transmission (**b**) electron microscopy, and numerical particle size distribution (**c**). Note that image (b) shows the nanoparticles localized on transmission electron microscope (TEM) grid.

With the introduction of nanotubes the skeletal structure becomes more pronounced (Figure 3a). However, the further addition of nano-Fe2O3 resulted in the material with uniformly distributed nanoparticles and with no evidence of webbing (Figure 3b). These must have been the nanoparticles of iron oxide that promote the formation of NC nanospheres instead of fibers. Notably, the catalyst content within the fabricated composites was close to the target value, i.e., no material is lost during preparation (Section S5, Supporting information).

**Figure 3.** Electron microscopy of (**a**) NC/carbon nanotube (CNT)(sc) and (**b**) *NC*/*CNT*/*Fe2O3(sc)*.

### *3.3. Thermal Behavior*

The linear heating of all samples shows the general pattern consistent with the typical nitrocellulose behavior [46,47]: a single peak of the heat release started at about 190 ◦C with the corresponding mass loss of nearly 80 wt.%. The heat released during thermolysis was within the range of 1800–2000 J·g−1 and agreed closely for all analyzed species within the experimental error. A closer examination of the four obtained nano-NC-based samples showed a lower decomposition onset temperature (see Figure 4 and Table 1), previously observed for nitrocellulose with nano-sized features [16–18]. The isoconversional Friedman analysis showed that the activation energy remained the same as for micron-NC at low- and high-conversion degrees, while at the medium range the barrier was notably lower for nano-NC (Figure 4c). A similar conclusion is drawn from the formal kinetic fit with a single reaction model in flexible reduced Sestak–Berggren form [48,49]. Kinetic parameters for nano-NC, *E*a = 186 ± 1 kJ·mol−1 and ln *A* = 43.0 ± 0.2 s−1 were lower than that for raw NC, *E*a = 196 ± 1 kJ·mol−1 and ln *A* = 45.7 ± 0.3 s−1 (Section S6, Supporting Information). Both pairs of the kinetic parameters were consistent with the global autocatalytic reaction proposed by Brill and Gongwer [50]. The shift of the kinetic parameters for nano-NC was apparently caused by the higher surface area available for autocatalytic agents and its increased contribution to the overall decomposition process.

**Figure 4.** Thermal analysis under linear heating: mass loss (**a**), heat flow (**b**) data, and isoconversional kinetic analysis results (**c**).

1 Measured at 5 K·min−1 rate. 2 For pressed pellets burned at 12 MPa.

### *3.4. Friction Sensitivity*

The full set of experiments according to standard STANAG 4487 [51] have been performed to derive the friction force corresponding to the 50% probability of explosion. The SAS processing decreased the pure nitrocellulose *NC(sc)* sensitivity, and the required initiation stimulus accordingly rose from 234 to 342 N. The effect must have been caused by the particle size reduction since the

chemical composition of the sample was unaltered as shown above. The addition of the iron oxide sensitized the *NC*/*Fe2O3(sc)* sample down to the level of untreated raw compound. Introduction of carbon nanotubes ranked the composite *NC*/*CNT(sc)* as friction insensitive (10% of explosions at highest load of 360 N) and with the addition of both carbon nanotubes and iron oxide *NC*/*CNT*/*Fe2O3(sc)* the effect was still observed (40% of explosions at 360 N).

### *3.5. Combustion Tests*

The energetic performance of the samples was estimated by combustion in air with the high-speed filming of the burning surface (2000 fps) and directly measured as the burning rate–pressure dependency in a Crawford bomb at elevated pressures. For the raw nitrocellulose, the observation revealed active bubble formation on the surface during combustion (Videos S1,S2, Supporting information). In this foam zone, small black particles were formed that moved inside the liquid layer to aggregate (see Figure 5a,b for sequential frames with 0.075 s time difference). Thus, the formed framework remained above the surface until burning out in a flame of the final reactions between the semiproducts of NC decomposition or until blowing away by the gas flow. Similar observations were drawn for the *NC(sc)* sample with the only exception of higher carbonaceous framework formation that remained after tests as the solid residue (Figure 5c and Figure S6, Video S2, Supporting information). The finer structure of unreacted nano-nitrocellulose seemed to either result in a more distributed localization of the seeds or in the structural modification of the formed framework to keep it intact in gas flow. Considering the burning rate values, the nano-NC exhibited a higher velocity as compared to raw NC, e.g., 3.8 against 4.7 mm·s−1 at 2 MPa (Figure 5g). This enhancement was apparently linked with the lowering of the activation barrier for condensed-phase decomposition as revealed by thermal analysis. The introduction of carbon nanotubes gave more seeds for carbonaceous framework (Figure 5d), and in fact a considerable porous layer grew above the surface (Figure S7, Video S3, Supporting information). However, it easily broke down into fragments, which were blown away as the combustion proceeds. The burning rate for the *NC*/*CNT(sc)* composite was the same as for nano-NC monopropellant showing no significant effect of the increased heat conductivity on the combustion front propagation.

**Figure 5.** Combustion of the composites: still images from the high-speed filming (**a–f**) and the burning rate-pressure dependencies (**g**). Images shows the typical shots of NC (**a**,**b**), *NC(sc)* (**c**), *NC*/*CNT(sc)* (**d**), *NC*/*Fe2O3(sc)* (**e**), and *NC*/*CNT*/*Fe2O3(sc)* (**f**) burning. Frame width is 1.22 mm.

Figure 5e illustrates the burning process for the NC nanocomposite modified with nano-Fe2O3 addition. Seeds of the size bigger than that in previous samples stuck together on the surface and stayed in gas phase as a skeleton (Video S4, Supporting information). Recently, Shin et al. have proposed a facile method of nano-Fe2O3 transformation to Fe3O4 nanoparticles covered by a carbonaceous layer using a porous film made of nano-Fe2O3 particles filled with nitrocellulose [52]. After the burning of the NC matrix with the registered temperatures of 500–900 ◦C, the desired product was recovered. In line with these results, Denisyuk and Demidova [24] noticed the reduction of SnO2 to a mixture of SnO and Sn during the combustion of an NC-based propellant. Here, we could assume the reduction of nano-Fe2O3 in our experiments; however the necessary high-temperature flame was not registered. As for the burning rate, we noticed the acceleration of the combustion by iron oxide addition under 8 MPa pressure (blue curve, Figure 5g).

When carbon nanotubes and nano-iron oxide are introduced together to the composite, we observe a different picture of combustion (Video S5, Supporting information). Figure 5f shows a typical shot of the burning surface covered by a crumbly carbonaceous layer with luminous spots that appeared occasionally. These luminous regions give strong evidence of the catalyzed reactions and the increased heat flux seems to be the main driver for the combustion velocity increase. Figure 5g summarizes the burning rate-pressure data showing the highest catalytic effect at high pressures, achieving a maximum increase in 20% at 12 MPa (Table 1). The simplified combustion models [53,54] assumed the pressure exponent *n* of the burning rate dependency *U* = *B*·*P*n to be an indicator of the importance of gas-phase reactions, and the observed increase in it is consistent with the action of the catalyst in gas-phase zone.

To estimate the value of the SAS process in the arrangement of nanocatalyst and CNTs, we prepare the sample of the same composition *NC*/*CNT*/*Fe2O3* by means of conventional dry mixing. Testing for sensitivity and combustion reveal its inferior performance as compared to the composite produced via the SAS (Table 1).

### **4. Conclusions**

Supercritical CO2 has been applied for the first time as anti-solvent for obtaining nano-sized nitrocellulose powder and various nitrocellulose-based composites. Thus prepared materials exhibit a significantly lower sensitivity to friction and a higher burning rate than the raw fibrous nitrocellulose, while maintaining the chemical composition and thermal stability. For the benchmark dry mixture of *NC*/*CNT*/*Fe2O3* the results are notably lower than those for the SAS-produced composite. Overall, the proposed approach for introducing the burning rate modifiers into the matrices of energetic materials allows increasing the performance, satisfies the high technological and ecological requirements and may be useful for modern energetic composites production.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/2079-4991/9/10/1386/s1: Section S1: Experimental Methods, Section S2: Starting Materials, Figure S1: Microstructure of (a, b) untreated micron-sized NC fibers, (c,d) carbon nanotubes, (e,f) nano-sized iron oxide, Section S3: Details of the SAS process, Table S1: SAS process parameters for NC modification, Section S4: Morphology of Nano-NC, Figure S2: Atomic force microscopy (AFM) image of nitrocellulose powder produced via SAS process, Table S2: Parameters of the particle-size distribution of nano-NC powder, Section S5: Determination of catalyst content in fabricated composites, Figure S3: Temperature and mass profiles during annealing of NC-based composites with Fe2O3 addition, Section S6: Thermokinetic analysis, Figure S4: Results of formal kinetic analysis for micron-(a) and nano-(b) sized nitrocellulose, Figure S5: Data from review and the results of the current study, Section S7: Combustion tests, Figure S6: Combustion of untreated (a–c) and processed (d–f) nitrocellulose: (a,d) view of pellet before test, (b,e) still image of the burning surface, (c,f) residue after experiment, Figure S7: Combustion nitrocellulose with additives of nano-iron oxide (a–c), carbon nanotubes (d–f) and both Fe2O3 and CNTs (g–i): (a,d,g) view of pellet before test, (b,e,h) still image of the burning surface, (c,f,i) residue after experiment, Table S3: Results of combustion experiments, Figure S8: Combustion velocity against pressure for all involved compositions, Videos S1–S6.

**Author Contributions:** Conceptualization, I.V.K. and S.G.Z.; methodology, M.N.Z., I.V.K. and K.A.M.; software, N.V.M.; validation, M.N.Z. and N.V.M.; formal analysis, I.V.K.; investigation, O.S.D., K.A.M. and D.B.M., resources, I.V.F., S.G.Z. and A.N.P.; data curation, D.B.M.; writing – original draft, N.V.M.; writing – review and editing, M.N.Z., A.N.P. and N.V.M., visualization, D.B.M. and K.A.M., project administration, I.V.K. and N.V.M.; supervision, I.V.F., S.G.Z. and A.N.P.

**Funding:** The authors are grateful for financial support of synthetic part and combustion behavior studies by Russian Foundation for Basic Research (Grant No. 18-29-06023) K.A.M., A.N.P., and N.V.M. acknowledge the State research project #0082-2018-0002 (CITIS #AAAA-A18-118031490034-6) to Semenov Institute of Chemical Physics for a financial support of thermal analysis.

**Acknowledgments:** Authors thank Ekaterina K. Kosareva for performing the atomic force microscopy and Lyubov N. Mitireva for thorough English language editing.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **Nomenclature**

### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Preparation and Characterization of Silicon-Metal Fluoride Reactive Composites**

**Siva Kumar Valluri 1, Mirko Schoenitz 1 and Edward Dreizin 1,2,\***

Received: 16 October 2020; Accepted: 25 November 2020; Published: 28 November 2020

**Abstract:** Fuel-rich composite powders combining elemental Si with the metal fluoride oxidizers BiF3 and CoF2 were prepared by arrested reactive milling. Reactivity of the composite powders was assessed using thermoanalytical measurements in both inert (Ar) and oxidizing (Ar/O2) environments. Powders were ignited using an electrically heated filament; particle combustion experiments were performed in room air using a CO2 laser as an ignition source. Both composites showed accelerated oxidation of Si when heated in oxidizing environments and ignited readily using the heated filament. Elemental Si, used as a reference, did not exhibit appreciable oxidation when heated under the same conditions and could not be ignited using either a heated filament or laser. Lower-temperature Si fluoride formation and oxidation were observed for the composites with BiF3; respectively, the ignition temperature for these composite powders was also lower. Particle combustion experiments were successful with the Si/BiF3 composite. The statistical distribution of the measured particle burn times was correlated with the measured particle size distribution to establish the effect of particle sizes on their burn times. The measured burn times were close to those measured for similar composites with Al and B serving as fuels.

**Keywords:** reactive materials; nanocomposite; metal combustion; ignition; thermal analysis

### **1. Introduction**

Silicon serves as a fuel in several pyrotechnic compositions [1]. Its high energy density, abundance in the lithosphere [2] and cheap manufacturing make it an attractive fuel for a broad range of applications, including propellants and explosives. However, it is known to be relatively difficult to burn [3,4]. Silicon has a high boiling point (3538 K) compared to its oxide (2503 K) and hence burns heterogeneously. Further, the nascent refractory oxide layer on silicon particles limits the inward diffusion of oxygen at temperatures below 1273 K [5]. The direct oxidation of silicon from gaseous oxygen [6] and nitrogen [7] is only accelerated at higher temperatures, when diffusion rates increase.

To make silicon more reactive, previous work focused on employing aggressive condensed phase oxidizers. Transition metal oxides [4,8–11], alkali metal nitrates [4,9–11], perchlorates [9] and manganates [9] were explored as oxidizers for micron-sized and nanometric silicon powders. These compositions, though favorably reactive, do not fully exploit the potential of silicon as they yield largely condensed phase combustion products. Respective silicon-condensed oxidizer composites have reduced energy content compared to elemental silicon reacting with oxygen gas; thus, the reactions have reduced adiabatic flame temperatures. Use of high surface area, porous nanosilicon has attracted attention [9,12–14]; however, such materials may be difficult to handle and their applications are mostly limited to pyrotechnics.

In comparison to oxidation, fluorination of silicon is preferred in fuel-limited systems because the negative heat of formation of silicon tetrafluoride, SiF4, (−1615 kJ/mol) is substantially lower than that of SiO2 (−911 kJ/mol), making the reaction more exothermic. Silicon fluorides (and most oxyfluorides [15]) are gases under standard conditions.

Combustion-related work on fluorination of silicon was mostly limited to the use of fluoropolymers, e.g., polytetrafluoroethylene (PTFE) as oxidizers. Silicon-PTFE compositions are predicted to have a higher adiabatic combustion temperature and produce more gaseous products compared to combinations of silicon with oxide-based oxidizers [9]. Silicon-PTFE milled composites with a range of compositions (10–40 wt.% silicon) have been explored [16–18].

In recent studies it was shown that metal fluorides, such as CoF2, NiF2 and BiF3, could be effective oxidizers with metal fuels, e.g., aluminum. Reduced ignition temperatures and higher burn rates than for pure aluminum (in air) were reported for respective composites [19,20]. Further, metal fluorides were combined with boron, known to be more difficult to ignite than aluminum. A marked improvement in the ignition and combustion characteristics for boron/metal fluoride composites was observed with as little as 10 wt.% of BiF3 [21,22]. The addition of fluorides altered combustion of boron particles: it occurred in a single stage with shortened burn times [21–23]. In constant volume explosion experiments, the 10 wt.% BiF3 coated boron powders exhibited a higher rate of pressurization and higher maximum pressure compared to fine aluminum [22]. Experiments with aluminum and boron-based reactive composites containing metal fluorides as oxidizers show consistently significant improvements in kinetics of reactions leading to ignition [24,25].

Despite the increased reactivity, the fluoride composites made with aluminum and boron were found to be safer to handle, as they were insensitive to electrostatic stimulus unlike oxide based compositions [24,25]. It should also be noted that the gas-phase fluorinated reaction products could be aggressive; thus care must be taken when such products are released.

The use of metal fluorides as oxidizers for silicon is explored in the current work. The fluorides used in this study, BiF3 and CoF2 were selected because they are less hygroscopic, easy to handle [25], and contain metals that interact differently with metal fuels: Bi does not alloy with most metals or metalloids while Co forms alloys readily. The phase of the reduced metal (pure element or alloy) could affect further oxidation and combustion reactions. Based only on gravimetric and volumetric heats of reaction shown in Table 1, replacing oxides of Co and Bi with respective fluorides does not lead to tangible advantages. The heats of reactions are nearly matching for the metal-rich composites (with 50 wt.% of Si) reacting in presence of the external oxidizer, so that excess of Si and the metals obtained from the reduced fluorides can further be oxidized. Thus, the anticipated advantages of using metal fluorides are not based on the expected theoretical energy release but are based on the expected accelerated reaction rates and formation of gas phase products.

**Table 1.** Estimated heats of reactions involving Si with oxides and fluorides of Co and Bi for stoichiometric and fuel-rich (50/50 wt.%) composites reacting in inert and oxidizing environments.

**Table 1.** *Cont.*

\* For clarity of the concept, reactions are written without balancing reactants and products.

Both BiF3 and CoF2 were incorporated into silicon-metal fluoride composites prepared by arrested reactive milling [26].

### **2. Experimental**

### *2.1. Materials Preparation*

The silicon-metal fluoride composite powders were prepared in a SPEX Certiprep 8000 series shaker mill (by SPEX CertiPrep, Metuchen, NJ, USA). The constituent silicon (−325 mesh crystalline Si, 99%) and anhydrous fluoride powders, bismuth (III) fluoride (BiF3, 98%) and cobalt (II) fluoride (CoF2, 99%) were sourced from Alfa Aesar, Ward Hill, MA, USA. Preliminary inspection using electron microscopy showed that BiF3 contained crystalline particle agglomerates, mostly in the 10–15 μm range. For CoF2, the powder appeared as finer powder, with most particles in the single micron range. The compositions targeted 30 and 50 wt.% of silicon with the remaining mass of bismuth or cobalt fluoride. The materials were prepared in 5 g batches in steel milling vials. The powders were milled using steel milling media with hexane as the process control agent. The ball to powder mass ratio was fixed at 10.

 This material is referred to as 1-stage 50Si·50BiF3. Examination of 1-stage 50Si·50BiF3 samples showed the presence of large Si particles that were not mixed with BiF3. To improve homogeneity, milling was therefore carried out it two steps for all other prepared composites. Silicon was milled separately in the first milling step to refine it. This step lasted 60 min and used 14 mL of hexane and 5-mm steel balls. This size-reduced silicon was also used as reference material for thermal analysis and other experiments.

In the second step, the size-reduced (premilled) silicon was milled with bismuth fluoride or cobalt fluoride for 60 min using 4.76-mm steel balls with 7 mL of hexane. These composites are referred to as 50Si·50BiF3, 50Si·50CoF2, 30Si·70BiF3 and 30Si·70CoF2, with respective mass percentages of components used as identifiers. Table 2 presents characteristics of the prepared materials. The equivalence ratios describe reactions.

$$
\text{Si} + \frac{4}{\text{x}} \text{MeF}\_{\text{x}} \rightarrow \text{SiF}\_{4} \uparrow + \frac{4}{\text{x}} \text{Me} \tag{1}
$$

where Me stands for Bi or Co. Reactions (1) reduce the respective metal fluorides and yield SiF4. All prepared composites are fuel-rich, making them potential fuel candidates for applications where oxygenated gas environments are available.

**Table 2.** Compositions of prepared silicon-metal fluoride reactive composites.

The prepared materials were passivated in an argon-filled glovebox, where oxygen was present at a low partial pressure. The passivation lasted 48 h. Passivated materials were covered by hexane and stored in the laboratory for the duration of this project.

Reactivity was initially assessed by taking 100–200 mg of powder on a ventilated filter paper and igniting the paper by a butane lighter inside a fume hood. The 30 wt.% silicon powders burned quickly with a bright noiseless flash; the 50 wt.% silicon powders were comparatively slower and generated particle streaks jetting from the burning powder bulk. For the 50 wt.% silicon compositions, despite a coarser scale of mixing and poor homogeneity, 1-stage 50Si·50BiF3 was visibly more reactive than two-stage milled 50Si·50BiF3. There was no apparent difference between composites with BiF3 and CoF2 serving as oxidizers.

### *2.2. Material Characterization*

The surface morphology and homogeneity of mixing between components in the prepared powders was examined using a JSM-7900 field emission scanning electron microscope (SEM) by JEOL, Tokyo, Japan. The imaging was performed by back-scattered electrons to achieve compositional contrast. To identify material composition and impurities incorporated by milling, energy dispersive X-ray spectroscopy (EDX) was used. The prepared powders were also analyzed using X-ray diffraction (XRD) on a PANalytical Empyrean multipurpose research diffractometer by Malvern Panalytical, Malvern, UK, operated at 45 kV and 40 mA using filtered Cu Kα radiation.

Reactivity of powders containing 50 wt.% silicon was tested by thermogravimetric analysis (TG) coupled with differential thermal analysis (DTA)/differential scanning calorimetry (DSC) on a STA409PG by Netzsch, Selb, Germany. The powders with a composition of 30 wt.% silicon were not tested to avoid releasing large amounts of aggressive fluorinated gases in the instrument.

Samples were heated in open alumina DTA crucibles only up to 973 K (700 ◦C) to limit release of volatile fluorinated species. Samples were tested in argon (99.998%) flowing at 50 mL/min as an inert environment and in an oxidizing environment, with an additional 50 mL/min flow of oxygen (99.994%). Both gases were sourced from Airgas. For each case, a baseline was obtained by heating an empty crucible under identical conditions. The sample mass was fixed at 5 mg, except for a few runs in argon performed with larger sample masses of 10 and 24 mg. For the oxidizing environment, heating rates 1, 2 and 5 K/min were used. Experiments in argon were performed with heating rates of 2, 5 and 10 K/min. To probe the progress of reactions, samples were recovered from specific intermediate temperatures and analyzed using XRD.

### *2.3. Heated Filament Ignition*

The ignition temperatures for the prepared powders were obtained by coating a sample onto an electrically heated 24 gauge (0.5105 mm) nickel-chromium wire. The time of ignition was determined from a high-speed video and the wire temperature was measured optically, using a calibrated photo diode. The basic experimental setup and procedure have been presented earlier [27]. The photodiode-based pyrometer has been described in Ref. [24]. Briefly, the wire was coated with a suspension of the powder in hexane. Once the hexane dried, the wire was electrically heated using a variable set of rechargeable large cell batteries (7238K57 McMaster Carr. Elmhurst, IL, USA) and a 1 Ω rheostat connected in series to achieve targeted heating rates in the range of 103–104 K/s. A fiber optics cable was focused at a section of the wire that was uncoated and fed into germanium photodiode (PDA30B2 by Thorlabs, Newton, NJ, USA). The photodiode was calibrated against a black body emission source (BB4A by Omega Engineering, Bridgeport, NJ, USA) to function as a high-speed pyrometer. To correlate the time of ignition to the temperature of the wire, the ignition was recorded using a MotionPro 500 camera by Redlake, San Diego, CA, USA, synchronized with the ignition circuit. At least 5 runs were performed for each targeted heating rate.

### *2.4. Particle Combustion*

The particle combustion was studied by aerosolizing the materials in a room-temperature air stream fed into the focal spot of a sealed CO2 laser (Evolution 125 by Synrad, Mukilteo, WA, USA). The optical emission of the ignited particles was captured by filtered photomultiplier tubes (PMT) to obtain burn times. Each burning particle produced an emission pulse; the pulse width taken at 10% of its peak value was interpreted as the burn time.

The powder feeder described elsewhere [28] consisted of a screw driven by a DC motor, with 30 threads coated by 0.12 g of powder. The powder lodged between the threads was blown off the screw and aerosolized by a carrier gas, air flown at 0.68 L/min. The aerosolized particles were fed via a 2.39 mm internal diameter brass tube. The upper end of the tube was placed 2 mm below the focal spot of the laser. The laser beam was focused by a ZnS lens to a spot of about 250 μm diameter. The laser was operated at 37.5 W, or ca. 30% of its maximum power, which is much higher than the ignition threshold for the materials tested. Thus, particles passing through the laser beam ignited consistently. Further experimental details are available elsewhere [29].

A fiber optics bundle trained to capture the particle emissions was placed 2 cm above and 16 cm away from the brass feeder tube. The fiber was split to feed to two R3896-03 PMTs by Hamamatsu, Hamamatsu City, Japan, filtered at 700 and 800 nm. The signal from the PMTs was acquired through a 16 bit PCI-6123 board by National Instruments, Austin, TX, USA, at 100,000 samples per second and processed using LabView software (Labview 16 by National Instruments, August 2016). Signal was collected for periods of 8 s at a time. Several such 8 s runs were required to acquire over 800 pulses produced by individual burning particles for each material. Overlapping, or closely spaced pulses were dismissed. The widths of the collected individual particle pulses, interpreted as burn times, were developed into respective statistical distributions employing a smoothing kernel density function [21].

### *2.5. Aerosolized Particle Collection and Sizing*

The aerosolization of the composite powders may cause size classification. Thus, the size distributions for the particles observed to burn in the present experiments could differ from those of the prepared powders. To account for this, samples of powders passed through the feeder were collected and analyzed. A flat aluminum substrate covered with a double-sided adhesive carbon tape was placed 2.5 cm above the brass feeder. Particles were trapped onto the tape for 10–15 s using the feeder operated at the same gas flowrate as in the combustion tests; however, with the laser beam turned off. The collected particles were imaged by the SEM and their sizes were obtained using ImageJ software [30]. From a table of all particles of one material, size distributions were generated using a kernel density function. The particle size distributions were correlated with the burn time distributions to identify the effect of particle size on the burn time. Such correlations broadly assume that larger particles burn longer.

### **3. Results**

### *3.1. Particle Morphology and Composition*

SEM images of the prepared 50 wt.% silicon- 50 wt.% metal fluoride composites are presented in Figure 1. The grey crystalline phase was silicon, and the brightest phase was the fluoride. The carbon tape in the background appeared darkest. The single stage milled 1-stage 50Si·50BiF3 powder is shown in Figure 1A,B. It is apparent from Figure 1A that the fluoride was not uniformly distributed in the material. Many relatively large, ca. 20–40 μm, Si particles and some unattached BiF3 clusters were observed. Part of the image in Figure 1A with a relatively well-homogenized composite particle was magnified and shown as Figure 1B. Individual Si and BiF3 phases were interspersed on a submicron scale. Such particles were few with a majority of silicon particles having little to no fluoride attached to them.

**Figure 1.** Backscattered electron images of as-milled 50 wt.% silicon-50 wt.% fluoride composites: (**A**,**B**) 1-stage 50Si·50BiF3, (**C**,**D**) 50Si·50BiF3 and (**E**,**F**) 50Si·50CoF2.

A composite prepared using the two-stage protocol, 50Si·50BiF3 is shown in Figure 1C,D. All particles show good mixing of the fluoride and silicon and reduced particle sizes in Figure 1C. A magnified image of a representative particle in Figure 1D shows angular silicon crystals decorated by bright particles of BiF3.

The images of the two-stage milled 50Si·50CoF2 powder are presented in Figure 1E,F. In Figure 1E, a large, agglomerated particle of 60 μm, with no discernable difference in phase contrast is seen. Such agglomerates were common. The inset in Figure 1E shows that the particle surface had distinct angular platelets and smaller particles. Through EDX, the larger platelets were confirmed to be Si and the smaller particles were found to be CoF2. A smaller particle of 50Si·50CoF2 is shown in Figure 1F. Distinct morphology enables one to identify both silicon and CoF2 phases mixed rather uniformly on the nanoscale.

SEM images of powders with 30 wt.% silicon and 70 wt.% metal fluoride are presented in Figure 2. These images were obtained using secondary electrons to highlight the surface morphology, while losing on the phase contrast. General particle morphology for 30Si·70BiF3 is illustrated in Figure 2A,B. Similarly, the morphology for 30Si·70CoF2 is illustrated by Figure 2C,D.

**Figure 2.** Secondary electron images of as milled 30 wt.% silicon-70 wt.% fluoride composites: (**A**,**B**) 30Si·70BiF3 and (**C**,**D**) 30Si·70CoF2.

A large amount of unattached BiF3 particles and small crystalline silicon particles were observed in Figure 2A. Some relatively large BiF3 particles decorated agglomerates of silicon crystals (Figure 2B); although the coating did not appear to be uniform.

In Figure 2C, cobalt fluoride was observed as small fuzzy particles covering jagged crystals of silicon. No unattached CoF2 was observed. Generally, 30Si·70CoF2 composite contained silicon particles with a broad size distribution and all Si surfaces were rather uniformly coated by cobalt fluoride, unlike 30Si·70BiF3. Occasional large particles (>10 μm) were observed in both materials, as shown in Figure 2B,D.  SEM images show that such large particles consisted of fine silicon particles embedded in a fluoride matrix.

The XRD patterns of milled composites are presented in Figure 3. Peaks corresponding to BiF3 and Bi7F11O5 [31] were observed in all Bi-containing composites. Detectable amounts of reduced Bi were also observed. In the two-stage milled composites, the fluoride and reduced Bi peak intensities scale with the loaded BiF3 concentration; stronger fluoride peaks were observed for 30Si·70BiF3 as compared to 50Si·50BiF3. Results of quantitative composition analysis for BiF3-containing composites based on the whole pattern refinement performed using the X-pert Highscore software package [32] are shown in Table 3. The data suggest that 50Si·50BiF3 retained only ca. 64% of the initial fluorine. On the other hand, the poorly refined 1-stage 50Si·50BiF3 exhibited the weakest bismuth peaks and strongest BiF3/Bi7F11O5 peaks among the three Si·BiF3 composites; it retained approximately 71% if the initial fluorine.

**Figure 3.** XRD patterns of the prepared composite powders.

The XRD pattern of the composites with CoF2 shows peaks corresponding to silicon and cobalt fluoride. No reduced cobalt is observed. The scaling of the fluoride peak intensity as a function of the loaded fluoride concentration was less apparent.

**Table 3.** Compositions of the prepared Si·BiF3 composites obtained from the whole pattern refinement of the XRD traces.

### *3.2. Thermal Analysis*

The TG and their corresponding DTA traces for 50Si·50CoF2 heated in both oxidizing and inert environments are shown in Figure 4. The DTA traces for the oxidizing and inert runs used separate vertical scales because of difference in the heat effects observed. The TG trace measured in Ar (Figure 4A) shows a single mass loss event with the onset below 673 K (400 ◦C); the total observed mass loss was close to 23% of the initial mass. The respective DTA trace (Figure 4B) shows weak features, which can be interpreted at two broad exothermic humps (peaks at 498 and 893 K (or 225 and 620 ◦C)) or a very broad exothermic feature overlaid with an endothermic event. To support the latter interpretation, much of the endothermic feature correlated with the observed mass loss.

**Figure 4.** (**A**) Thermogravimetric analysis (TG) and corresponding (**B**) differential thermal analysis (DTA) traces for 50Si·50CoF2 heated at 5 K/min in different environments.

For TG in Ar/O2, the onset for a gradual mass gain was at 548 K (275 ◦C). The mass gain was reversed at 723 K (450 ◦C), and the maximum mass loss of ca. 8.3% was observed by 798 K (525 ◦C). After the minimum, the mass began increasing again. The corresponding DTA trace exhibited a rising

baseline (despite the correction applied using an empty crucible). A small exothermic peak was noted at 280 ◦C, corresponding to the onset of the mass gain. A stronger exotherm was observed at 773 K (500 ◦C) correlating with the mass loss.

The TG and DTA plots for 50Si·50BiF3 heated in different environments are presented in Figure 5. In Figure 5A, the oxidative TG runs for the milled/size-reduced Si and 1-stage 50Si·50BiF3 were presented for comparison. No mass change is observed for the Si reference.

**Figure 5.** (**A**) TG and corresponding (**B**) DTA traces for 50Si·50BiF3 in heated at 5 K/min in oxidizing and inert environments. The oxidative TG runs for 1-stage 50Si·50BiF3 and milled Si are also shown.

The TG trace for 50Si·50BiF3 in Ar shows a gradual mass loss of 8.2%. It begins around 423 K (150 ◦C) and ends by 873 K (600 ◦C). Weak but reproducible step-like features are noted in the TG trace; the onset for the mass loss are marked by a filled circle as labeled in Figure 5A. The corresponding DTA plot in Figure 5B, shows multiple weak peaks. It is possible to interpret the trace as showing a broad endothermic peak, similar to that in Figure 4. Additionally, several exothermic features are noted. The first three of these features, labeled in Figure 5B, correlated with the mass loss steps. The second exotherm, at ca. 543 K (270 ◦C; marked Peak II), had a small but reproducible superimposed endothermic feature.

For 50Si·50BiF3 experiments in both Ar and Ar/O2 were performed at different heating rates. The shift of the respective mass loss and mass gain points to higher temperatures at greater heating rates was used to consider relevance of the respective reactions to ignition, as discussed below.

In Ar/O2, 1-stage milled 50Si·50BiF3, exhibited a small mass loss (ca. 1.5%, onset at 583 K (310 ◦C)) and then mass gain, which was not complete by the end of the run. For the two-stage milled composite, 50Si·50BiF3, heated in Ar/O2, the mass began to increase at ca. 423 K (150 ◦C) as marked in Figure 5A. The mass was stabilized by about 593 K (320 ◦C). It began increasing again at 658 K (385 ◦C); the rate

of mass gain increased around 873 K (600 ◦C). The final measured mass gain was 15.5%, it was not complete by the end of the run. The corresponding DTA curve for 50Si·50BiF3 seen in Figure 5B exhibited a rising baseline indicating a broad exothermic process initiated at least from 393 K (120 ◦C). The exothermic feature ended around 923 K (650 ◦C).

For both oxidizing and inert environments runs for 50Si·50BiF3, partially reacted samples were recovered from 473, 673 and 973 K (200, 400 and 700 ◦C).

To better resolve weak processes occurring in Ar, runs with the powder mass increasing from 5 to 10 and to 23 mg were performed for 50Si·50BiF3. The respective TG traces are presented in Figure 6. The step-like features observed during mass loss were prominent for the lowest sample loading of 5 mg, for which a mass gain was also noted around 543 K (270 ◦C), which was close to the melting temperature of bismuth (marked in the Figure 6). The mass gain could have been caused by oxidation of the sample reacting with trace oxygen remaining in the furnace despite it being flushed with Ar. Melting of Bi could have accelerated such oxidation. When the sample was heated to a higher temperature of 1473 K (1200 ◦C), cf. 23 mg trace, a strong mass loss of 33% was observed with an onset at 1173 K (900 ◦C), which could only be associated with the loss of Bi.

**Figure 6.** TG traces for different mass samples of 50Si·50BiF3 heated in Ar at 5 K/min.

SEM images of partially reacted 50Si·50BiF3 recovered from different temperatures in inert and oxidizing environments are presented in Figure 7. The powder heated to 700 ◦C in Ar (Figure 7A) shows bright spherical bismuth particles studded on darker silicon particles. The dispersion of bismuth across silicon particles was largely preserved. For the sample heated to 700 ◦C in Ar/O2 (Figure 7B), elongated crystals, appearing to be Bi2SiO5 from EDX analysis, are observed; see the inset for details. Minor bright spheroidal particles rich in Bi were found decorating the surface of the darker Si-rich matrix. EDX shows oxygen present in all phases.

Figure 7C shows a characteristic particle heated to 1473 K (1200 ◦C) in Ar. It is porous and consists of several fused parts. Figure 7D provides a higher magnification image, presenting the surface morphology with multiple small spherical submicron bismuth-rich particles and lamellae of silicon crystal. In materials heated in both Ar and Ar/O2, EDX detects iron and chromium; both contaminants came from the steel milling balls.

**Figure 7.** Backscattered electron images of the two-stage milled composite 50Si·50BiF3 heated at 5 K/min and recovered from 973 K (700 ◦C): (**A**) in Ar, (**B**) in Ar/O2 and recovered from 1473 K (1200 ◦C) in Ar (**C**,**D**).

XRD patterns for 50Si·50BiF3 and 50Si·50CoF2 heated to different temperatures in Ar and recovered for analysis are presented in Figure 8, along with the reference pattern for the as-milled material. The specific peaks of oxyfluoride Bi7F11O5 in the as milled sample have not been presented in the Figure 8 due to space constraints.

**Figure 8.** The XRD traces of the two-stage milled composites: (**A**) 50Si·50BiF3 and (**B**) 50Si·50CoF2 heated at 5 K/min in Ar and quenched at different temperatures.

Despite Ar purge, traces of oxygen were expected in these runs. For 50Si·50BiF3 (Figure 8A) peaks associated to BiF3 are depressed by 473 K (200 ◦C). Simultaneously, BiF1.9O0.55 peaks [33] became dominant and Bi7F11O5 peaks from starting powder disappeared. At 673 K (400 ◦C), strong peaks of Bi, Si and the fluorine containing BiOF became predominant. By 973 K (700 ◦C), peaks corresponding to fluorinated species disappeared, remaining were peaks of Bi and Si and minor peaks attributed to Bi2SiO5 and Fe2Si. At 1473 K (1200 ◦C), Bi peaks diminished, and peaks of Si became stronger. Patterns attributed to iron and chromium impurities, such as Fe2Si and Si1.07Fe0.97Cr0.01 were also noted.

For 50Si·50CoF2, the XRD patterns of the starting powder and of the material heated to 700 ◦C in Ar are shown in Figure 8B. CoF2 peaks disappeared for the heated material; aside from elemental Co and Si, a pattern for CoSi was observed.

XRD patterns for the powders heated to different temperatures in Ar/O2 for 50Si·50BiF3 and 50Si·50CoF2 are presented in Figure 9. As seen in Figure 9A, the as milled 50Si·50BiF3 had prominent peaks of BiF3, Si, Bi and weaker peaks of Bi7F11O5. The material recovered from 473 K (200 ◦C) exhibited weaker BiF3 peaks while Bi and the oxyfluoride BiF1.9O0.55 peaks became stronger. By 673 K (400 ◦C), the BiF3 peaks were absent, Si peaks diminished appreciably, while a new oxyfluoride species, BiOF was observed. By 973 K (700 ◦C), oxidized products, Bi2SiO5 and Bi6Cr2O15 were the primary species. Minor peaks corresponding to Si were observed as well.

**Figure 9.** XRD patterns of the two-stage milled composites (**A**) 50Si·50BiF3 and (**B**) 50Si·50CoF2 heated at 5 K/min in Ar/O2 and recovered from different temperatures.

The XRD patterns for as-milled 50Si·50CoF2 and that quenched at 973 K (700 ◦C) in Ar/O2 are presented in Figure 9B. Strong peaks of the oxide of cobalt, Co3O4 and unreacted Si were observed.

### *3.3. Heated Filament Ignition*

High-speed video frames for all materials igniting on a heated filament are presented in Figure 10. The frames are labeled with the temperature of the filament.

All prepared materials exhibited ignition at low temperatures well before the filament turns incandescent. The incandescence of particles marks the ignition in fuel-rich 50 wt.% silicon compositions as indicated by arrows in the respective first frames. The 1-stage 50Si·50BiF3 composite exhibited less bright, short-lived particle streaks within an incandescent plume. At higher temperatures (see 857 K) a weak secondary ignition event was observed as marked in the frame.

**Figure 10.** High-speed video frames of the ignition of 1-stage 50Si·50BiF3 (heated at 2100 K/s) and two-stage milled materials; 50Si·50BiF3 (heated at 3800 K/s), 50Si·50CoF2 (heated at 2100 K/s), 30Si·70BiF3 (heated at 2200 K/s) and 30Si·70CoF2 (heated at 2200 K/s), on an electrically heated nichrome wire. The wire temperatures are indicated.

For 50Si·50BiF3, bright smoke enveloped the burning particles flying off the filament, as seen in frames captured above 796 K. The 50Si·50CoF2 material ignited producing a bright plume with few particle streaks, which became prominent at higher wire temperatures (>1123 K).

For materials with higher fluoride content, 30Si·70BiF3 and 30Si·70CoF2, ignition caused luminescent plume flares. For 30Si·70BiF3, particle streaks and smoke were observable above 806 K. The 30Si·70CoF2 composite generated a bright plume and no detectable burning particles.

The measured ignition temperatures of all prepared materials are shown as a function of heating rates in Figure 11. The error bars represent the scatter in 5 runs performed for each targeted heating rate. All ignition temperatures increase with the heating rate.

Lower ignition temperatures were measured for the materials with BiF3 compared to those with CoF2. The composites with 50 wt.% Si ignited consistently at lower temperatures compared to the composites with 30 wt.% Si. The 1-stage 50Si·50BiF3 had a higher ignition temperature comparable to its two-stage milled analog, 50Si·50BiF3. The secondary ignition for 1-stage 50Si·50BiF3 could only be clearly detected at lower heating rates. At higher heating rates, this secondary ignition could have occurred at much higher temperatures and was not discernable due to the bright filament emission.

**Figure 11.** Ignition temperature in air of all prepared materials as a function of the heating rate.

### *3.4. Particle Combustion*

In experiments employing the laser ignition of composite particles, 50Si·50CoF2 powder could not be fed consistently. Conversely, 50Si·50BiF3 powder was easy to feed; thus, only the experimental results with this powder are presented. Attempts to ignite milled Si powder, which could be fed into the laser beam, were not successful. Additionally, condensed combustion products could not be collected for analysis, likely because of the substantial amounts of gaseous combustion products.

A representative emission pulse of a laser-ignited 50Si·50BiF3 particle burning in air is shown in Figure 12. All pulses observed had single peaks with minor oscillatory features. A distribution of the particle burn times is shown in Figure 13. It peaked at 0.43 ms with a tail trailing towards longer times.

**Figure 12.** A representative emission pulse produced by a 50Si·50BiF3 particle burning in air.

**Figure 13.** The burn time distribution for particles of 50Si·50BiF3 burning in air.

The size distribution of the aerosolized 50Si·50BiF3 particles passed through the feeder is shown in Figure 14. The distribution was nearly symmetrical, with a mode of 3.09 μm.

**Figure 14.** The particle size distribution for aerosolized 50Si·50BiF3 powder passed through the feeder.

The burn time distribution was correlated with the particle size distribution for 50Si·50BiF3; the correlation plot is presented in Figure 15 along with earlier results for other metal-metal fluoride composites obtained using the same data processing method [20,23]. Like Al and B-based composites using BiF3 as an oxidizer, 50Si·50BiF3 composite powder particles burn rapidly, with most burn times under 1 ms. These times were shorter than reported burn times for Al particles of the same sizes.

**Figure 15.** Burn time–particle size correlation for 50Si·50BiF3 powder ignited in air by a laser. For comparison, similar correlations are shown for powders of 50Al·50BiF3 [20] and 50B·50BiF3 [23] obtained by arrested reactive milling.

The adiabatic flame temperature and predicted mole fractions of the combustion products of 50Si·50CoF2 burning in air calculated by NASA CEA code [34] are presented in the Supplementary Section Figure S1. The particle surface temperature during combustion is expected to be close to 2750 K (predicted to be the adiabatic flame temperature), sufficiently high for gasifying a significant fraction of the main combustion product SiO2.

### **4. Discussion**

### *4.1. Material Preparation*

A single-stage milling protocol does not effectively mix the physically different fluoride and silicon powders. Conversely, using premilled silicon for the second step milling with metal fluorides yields well-homogenized nanocomposite powders.

Materials with CoF2 have distinct morphological features, where the fluoride formed a coating on the larger silicon particles. Even with increased fluoride content, no loose fluoride or uncoated silicon particles were observed in 30Si·70CoF2. It is surprising that such a fine scale of mixing was achieved without detectable mechanochemical reaction, e.g., lack of reduced Co detected by XRD of as-milled powders (see Figure 3). The XRD results were consistent with the thermal analysis, Figure 4A. The complete reduction of CoF2 by silicon to form gaseous SiF4 resulted in a mass loss of 26%. This was close to the mass loss of 23 wt.% measured by TG.

Based on XRD of the as-milled 50Si·50BiF3 showing the presence of reduced bismuth and bismuth oxyfluorides, it was estimated that BiF3 made up only 28.5–30 wt.%, with the remaining 20–21.5 wt.% of BiF3 lost before or during milling. Accounting for this initial amount of BiF3 available in the as-milled powder, it is expected that the complete reaction of Si with the remaining F in the material would result in an observable mass loss of close to 8.9%. This agreed reasonably with the mass loss of 8.3% observed in TG (Figure 7A). While a smaller fraction of the fluoride reacted and was lost for the 1-stage 50Si·50BiF3, the mixing scale and homogeneity for that composite material were inadequate.

### *4.2. Low-Temperature Reactions*

The TG for 50Si·50BiF3 and 50Si·50CoF2 in Ar shows a mass loss due to the evolution of SiF4, described by reaction (1). As seen in Figure 4, SiF4 formation in 50Si·50CoF2 occurred in a single step between 661 and 873 K (388 and 600 ◦C). Conversely, for 50Si·50BiF3, the mass decreased gradually, with discernable stages starting from ca. 423 K (150 ◦C) and extending to 873 K (600 ◦C).

In Ar/O2 oxidation of the reduced metal and oxidation of silicon occur parallel to fluorination of silicon. Both oxide formation reactions form condensed products and thus increase the mass. Assuming for simplicity that the fluorination occurs similarly in both Ar and Ar/O2 environments, the mass gain due to oxidation, Δ*mox*, can be estimated from the mass difference between measurements in Ar/O2 and in Ar:

$$
\Delta m\_{0\overline{x}} = \Delta m\_{Ar/O\_2} - \Delta m\_{Ar} \tag{2}
$$

This change of mass is shown in Figure 16. The mass gain due to oxidation consists of two components, caused by oxidation of silicon, Δ*mSi ox* and that of the reduced metal, Δ*mMe ox* . The value of Δ*mMe ox* can be estimated from the mass measured in Ar, assuming Δ*mAr* is entirely due to evolution of SiF4 and assuming that all the reduced metal oxidizes immediately when the reaction occurs in Ar/O2:

$$
\Delta m\_{\rm ox}^{\rm Mc} = \Delta m\_{Ar}^{\rm Mc} = \frac{M\_{\rm Mc}}{M\_{\rm SiF\_4}} \cdot \frac{\chi}{4} \Delta m\_{Ar} \tag{3}
$$

Subtracting this mass from Δ*mox* gives:

$$
\Delta m\_{\rm ox}^{Si} = \Delta m\_{\rm ox} - \Delta m\_{\rm ox}^{Mc} \tag{4}
$$

Finally, the rate of Si oxidation can be obtained dividing the derivative *d* Δ*mSi ox* /*dt* by the mass of Si remaining in the material accounting for SiF4 leaving as gas (Equation (1)), also based on the TG measured in Ar.

**Figure 16.** The (**A**) oxidative mass gain and corresponding (**B**) silicon oxidation rate of two-stage milled materials 50Si·50BiF3 and 50Si·50CoF2.

Processing experimental traces using Equations (2) and (4) yields traces for Δ*mox*, Δ*mSi ox*, and the rate, 1 *mSi d*(Δ*mSi ox*) *dt* shown in Figure 16. Figure 16B shows that the rate of silicon oxidation parallels the observed rate of oxidation in the Ar/O2 environment. Since only as much Me oxide can form as MeFx has been reduced, the bulk of the observed mass gain represents the oxidation of Si. The apparent negative Si oxidation rate seen for 50Si·50CoF2 at between 450 and 500 ◦C suggests that the assumption that SiF4 formation is identical in Ar and Ar/O2 environments is not strictly true.

For 50Si·50CoF2, Figure 16 suggests that oxidation of Si began at a higher temperature than the corresponding SiF4 formation (cf. Figure 4); it also began at a higher temperature than for 50Si·50BiF3, for which oxidation began simultaneously with fluorination (cf. Figure 5). For both composites, the oxidation rate of Si increased initially and then became relatively stable over a range of temperatures. For both composites the low-temperature rates of heat release associated with the observed initial rate of oxidation of Si could be added to the rates of heat release caused by the Si fluorination occurring simultaneously (e.g., recovered from data in Figures 4 and 5 for 50Si·50CoF2 and 50Si·50BiF3, respectively) to describe the exothermic reactions leading to ignition of such composite materials.

Aside of relatively small changes in the reaction rate for 50Si·50BiF3, a stepwise increase in the oxidation rate is observed between 803 and 973 K (530 and 600 ◦C). As discussed below, this accelerated oxidation at relatively high temperatures was unlikely to affect ignition of the prepared composite. For 50Si·50CoF2 the apparent oxidation rate of Si dropped around 723 K (450 ◦C). However, based on Figure 5, it can be concluded that this effect was superficial and simply means that the fluorination of Si accelerated rapidly at such temperatures, with the associated mass loss becoming stronger than mass gain due to oxidation of both Si and reduced Co. Indeed, after a while, the oxidation rate of Si returned to about the same level it was at just below 723 K.

### *4.3. Reactions Leading to Ignition*

The onsets of TG features observed in both inert and oxidizing environment runs for both 50Si·50BiF3 and 50Si·50CoF2 were plotted along with their respective heated filament ignition temperatures in the Kissinger plot shown in Figure 17. The onsets of the mass loss observed in Ar and mass gain in Ar/O2 for 50Si·50BiF3 nearly coincided with each other for different heating rates. Extrapolating the kinetic trend from the ignition data points down to the range of low heating rates

used in thermal analysis pointed clearly to the onsets of the observed mass gain (in Ar/O2) and loss (in Ar) as reactions associated with ignition. As discussed above, for 50Si·50BiF3, both fluorination and oxidation of Si occurred nearly simultaneously with both exothermic reactions contributing to igniting the material.

**Figure 17.** Kissinger plot combining TG and DTA features and heated filament ignition data for two-stage milled 50Si·50BiF3 and 50Si·50CoF2.

The onset of mass loss observed in Ar for 50Si·50CoF2 (for which the measurement was only made at one heating rate) was shifted from that of the mass gain observed for this material heated in Ar/O2 at a lower temperature. The kinetic trend implied by the respective ignition temperatures lined up better with the mass loss observed in Ar, and thus it was likely that the ignition was mostly associated with SiF4 formation for this composite. It is nonetheless expected that the low-temperature oxidation of Si, once started, assisted ignition. The apparent activation energy that can describe ignition was roughly evaluated from the slopes of line fits shown in Figure 17. These activation energies were around 62 and 140 kJ/mol for 50Si·50BiF3 and 50Si·50CoF2, respectively. For 50Si·50BiF3, the inferred activation energy was comparable to that reported for fluorination of silicon wafers by elemental fluorine at reduced fluorine partial pressures [35].

The higher temperature reactions are not shown but would be shifted far left in Figure 17 and thus would be clearly irrelevant for ignition.

### *4.4. Particle Combustion*

The micron-sized silicon particles could not be initiated in the laser-assisted combustion experiment. Conversely, the prepared 50Si·50BiF3 composite could be readily initiated and burned in the time frame comparable to other reactive metal-fluoride composite powders. Most likely, initial fluorination of Si made it easier to initiate the reaction; however, it is expected that the laser energy was sufficient to rapidly heat Si particles, which did not sustain combustion once exposed to the room temperature air. It is thus suggested that the presence of reduced or partially reduced Bi accelerates both lowand high-temperature heterogeneous oxidation of Si, assisting in establishing and sustaining the particle flame.

As estimated by the CEA code (Figure S1), the high adiabatic flame temperatures volatilize a significant mass fraction of refractory silicon oxides formed during combustion of 50Si·50CoF2. In the case of 50Si·50BiF3, the mass fraction of volatile species might be higher because both bismuth and bismuth oxide vaporize easily. However, once bismuth is oxidized in the presence of a more reactive metal fuel, it is likely to be immediately reduced, serving as an oxygen shuttle. Thus, presence of Bi can effectively accelerate oxidation of Si, similarly to how it was suggested to accelerate oxidation of Al [20] and B [23]. When BiF3 serves as an oxidizer with different fuels, including Al, B and Si, gas phase products (including fluorides and oxyfluorides of the respective fuels) are favored a the high combustion temperatures. At the same time, the surface reaction may be governed by oxidation of Bi. Respectively, similar rates of combustion, with heterogeneous surface reactions yielding volatile products, are expected for all such composite materials. The rate of such reactions is expected to be controlled by diffusion of the oxidizer (oxygen) to the particle surface in all cases. This reasoning can explain comparable burn times observed for different composites, 50Al·50BiF3, 50B·50BiF3 and 50Si·50BiF3 for similarly sized particles (Figure 15)

### **5. Conclusions**

Preparing Si-based composites with BiF3 and CoF2 as oxidizers required a premilling step to refine the starting commercial Si powder in order to achieve nanoscale mixing between the material components. Preparation of well homogenized composites of Si with BiF3 could not be achieved without partial reaction of BiF3; however, no such reaction of CoF2 was observed. Upon heating, oxidation of Si occurs at lower temperatures than its fluorination when the added oxidizer is CoF2. Conversely, both oxidation and fluorination occurred nearly simultaneously and at a lower temperature with BiF3 as an oxidizer. The presence of reduced Co and Bi accelerated oxidation of Si upon its heating in an oxygen-containing environment.

Unlike elemental Si, prepared composite powders ignited readily when placed as a coating on an electrically heated filament. Ignition of the composite with BiF3 as an oxidizer occurred at consistently lower temperatures than that of the composite with CoF2. Results suggest that low-temperature fluorination and oxidation caused ignition for the composites with BiF3 as an oxidizer; apparent activation energy for reactions leading to ignition was close to 62 kJ/mol. For the composites with CoF2 as an oxidizer, it is likely that ignition is associated with Si fluorination.

Prepared composite powder with the BiF3 oxidizer ignited readily when passed through a CO2 laser beam in room temperature air. The measured particle burn times were nearly the same as those of similar composites combining BiF3 with such fuels as Al and B. Combustion products were mostly gaseous making it impossible to collect and analyze produced condensed species.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/2079-4991/10/12/2367/s1, Figure S1: The calculated adiabatic flame temperature and the combustion products presented as mole fractions at a range of equivalence ratios for 50Si·50CoF2.

**Author Contributions:** S.K.V. performed interpreted experiments, prepared first draft of the manuscript; M.S. discussed experimental results and correlated material properties with the reactions observed; E.D. set objectives of the effort, analyzed results, and edited the manuscript. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded in parts by the Office of Naval Research, Grant N00014-19-1-2048 and by the Defense Threat Reduction Agency, Grant HDTRA1-17-1-0044.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**

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### *Communication* **Layered Al**/**CuO Thin Films for Tunable Ignition and Actuations**

### **Ludovic Salvagnac, Sandrine Assie-Souleille and Carole Rossi \***

CNRS, LAAS, University of Toulouse, 7 Avenue du Colonel Roche, F-31400 Toulouse, France<EMAIL_ADDRESS>(L.S<EMAIL_ADDRESS>(S.A.-S.)

**\*** Correspondence<EMAIL_ADDRESS>
Received: 20 August 2020; Accepted: 8 October 2020; Published: 12 October 2020

**Abstract:** Sputter-deposited Al/CuO multilayers are used to manufacture tunable igniters and actuators, with applications in various fields such as defense, space and infrastructure safety. This paper describes the technology of deposition and the characteristics of Al/CuO multilayers, followed by some examples of the applications of these energetic layers.

**Keywords:** nanothermite; pyroMEMS; nanoenergetics; reactive thin film

### **1. Introduction**

Energetic materials are widely used to suddenly generate high amounts of thermal or mechanical energy under an electrical, mechanical or thermal stimulus. As typical energetic materials, nanothermites, which contain Al and oxide, have attracted much attention as they exhibit not only better combustion efficiencies and better ignitability compared to traditional explosives, but also their reaction outputs can be tuned, thanks to the selections of fuels, oxidizers, architectures and reactant size, allowing multiple actions. Among nanothermites, sputter-deposited Al/CuO multilayers represent interesting energetic thin film nanomaterials for tunable ignition, to replace old hot-wire ignitors.

In that context, our research group demonstrated several miniature pyrotechnical ignition devices integrating Al/CuO multilayers within microelectromechanical systems (MEMS) based microheaters, called pyroMEMS, for civilian and military applications such as triggering the inflation of airbags, micro-propulsion systems, and arm and fire devices.

This paper briefly reviews the technological process, focusing on the sputter-deposition technique, presents the properties of Al/CuO multilayered films, details the fabrication process flow of a micro-igniter, and present examples of applications of pyroMEMS.

### **2. Al**/**CuO Multilayered Nanothermites**

Two decades ago, the introduction of nanotechnologies enabled the emergence of a new class of energetic materials called nanothermites that could play a great role in future society. Nanothermites use raw materials that can be found in abundance, are low-cost and non-polluting (green materials). They also exhibit better combustion efficiencies and better ignitability compared to typical CHNO energetic mixtures, while being safer. Importantly, the reactions and ability to trigger specific actions can be tailored by varying the size and composition of the oxide, allowing multiple applications. Focusing on thermite multilayers, the overwhelming majority of works concern Al/CuO systems, as they feature an exceptionally high-energy release with gaseous production.

### *2.1. Multilayers Deposition*

Al/CuO thermite multilayers are mainly produced by the sputter-deposition technique [1–5], as this provides excellent control over the layering and stoichiometry. Cupric oxide thin film is synthesized using the direct current (DC) reactive magnetron sputtering technique in an oxygen-enriched environment. To produce Al/CuO multilayered films, layers of Al and CuO are deposited on top of each other in an alternating fashion (see Figure 1) without venting the chamber. However, after each layer of deposition, the sample stage is cooled at ambient temperature for 600 s. At each interface between the cupric oxide and Al layers, an intermixing layer of 4–8 nm is present that forms during the deposition itself.

**Figure 1.** Cross-sectional transmission electron micrographs of the Al/CuO multilayer obtained by magnetron sputtering.

Depositions parameters have been published several times and can be found in [6,7].

### *2.2. Ignition and Reaction Properties*

Applying an external source of energy locally on the multilayer results in an increase in the temperature of the multilayer section directly in contact with the heated surface. This can be done by electrostatic discharge [8], mechanical impact [9], laser irradiation [10], electrical heating (spark) [11] or thermal hot points [12], the latter of which is the widest used. After being ignited, the multilayers react by a self-sustained propagating reaction following the chemistry of Equation (1).

$$2\text{Al} + \text{ЭCuO} \rightarrow \text{Al}\_2\text{O}\_3 + \text{ЭCu} \text{ (+ } \Delta \text{H)}\tag{1}$$

In other words, after the temperature is raised locally and rapidly to a characteristic onset reaction temperature, the reaction enthalpy (ΔH) spreads into neighboring unreacted portions of the film, leading to a self-sustained reaction wave, which is highly luminous.

As the exothermic chemical reactions are controlled by the outward migration of oxygen atoms from the CuO matrix towards the aluminum layers, the reaction rate strongly depends on the reactant spacing (bilayer thickness). The sustained combustion rate is commonly characterized with a standard high-speed camera under ambient pressure [5,7].

The combustion velocity increases for thinner bilayers, with a maximum at ~90 m/s for free-standing foils (see Figure 2a). However, the reaction stops suddenly for extremely small bilayer thicknesses, which is due to the presence of interfacial layers, as described earlier, which reduces the amount of stored energy, i.e., reaction enthalpy, ΔH.

**Figure 2.** (**a**) Sustained combustion velocity in air of free-standing multilayers as a function of bilayer thickness (total thickness t = 2 μm); (**b**) ignition time as a function of the electrical power sent through the micro-heater for 5 bilayers of Al/CuO.

### **3. PyroMEMS Integration and Characterization**

When integrated onto a device/substrate, the widest and simplest ignition method for igniting an Al/CuO multilayered thin film is using hot-wire, which functions via local metallic thin film resistance. A pyroMEMS is mainly composed of a substrate, thin metallic resistance and the Al/CuO nanothermites (see Figure 3a).

**Figure 3.** (**a**) 3D schematic representation of a pyroMEMS; (**b**) photo of the pyroMEMS; (**c**) photo during the nanothermite reaction.

To minimize the ignition energy, the substrate must be a good thermal insulator. We chose 500 μm thick 4-inches glass substrates AF32 (Schott AG, Mainz, Germany). After being cleaned in oxygen plasma to remove surface contaminants, 300 nm thick titanium followed by 300 nm thick gold layers are evaporated onto the glass substrate, and the filament is patterned using the lift-off technique. Then the gold is chemically etched and patterned to define the Ti resistance and Au electrical pads. Finally, Al/CuO multilayers are sputter-deposited through a silicon shadow mask. A photo of a pyroMEMS thus manufactured is presented in Figure 3b.

We characterized that, for any heating surface areas (micro-heater sizes), the ignition time (delay between the time of application of the electrical power and the emission of the light) rapidly decreases when the ignition power density increases, until an asymptotic value, defining the minimum response ignition time, which is a characteristic of the multilayered film itself.

As an illustration, Figure 2b provides the ignition time versus the ignition power for a thermite made of 15 300 nm thick 1:1 Al/CuO bilayers. Below a certain electrical power, no ignition occurs, regardless of the duration of application of the power. This is simply explained by the fact that, below this ignition power threshold, the electrical energy supplied to the micro-heater is not sufficient to compensate for the energy lost by conduction through the substrate, by radiation from the exposed surface, or by convection in the air. The multilayer cannot reach its ignition temperature, also known as the onset reaction temperature.

This threshold highly depends on the heating surface area, and detailed results can be found in [12].

### **4. Examples of Applications**

To date, igniters are the widest investigated applications and are widely used to trigger the inflation of airbags, micro-propulsion systems, and the arm and fire devices used in missiles, rockets and any other ordnance systems. Traditionally, igniter technology consists of a metallic hot-wire (Ni/Cr, et al.) or bridge wire in contact with a secondary or primary explosive. The fabrication and integration of explosives in igniters requires extreme precaution due to their high sensitivity to electrostatic discharge, friction or shock.

Al/CuO multilayers can efficiently replace the hot-wire and primary explosive. Other advantages of replacing traditional hot-wire with pyroMEMS are as follows: (1) The overall integration is easier, requiring only the electrical connections of the pyroMEMS, whereas traditional pyrotechnical igniters require at least two steps (hot-wire and primary pyrotechnic composition deposition). (2) The Al/CuO multilayers are safe and less sensitive to the environment. (3) Ignition threshold and reaction output (flame temperature and gaseous products) can be easily tuned by changing the Al/CuO bilayer thickness and stoichiometry (bilayer thickness ratio).

Taton et al. [13] first reported the design, realization and characterization of hot-wire ignition integrating Al/CuO multilayers (see Figure 4) for use in pyrotechnical systems for space applications. The reactive Al/CuO multilayered thin film resides on a 100 μm thick epoxy/polyethyleneterephtalate (PET) membrane to insulate the reactive layer from the bulk substrate. When current is supplied, Al/CuO reacts and the products of the reaction produce sparks that can ignite any secondary energetic composition, such as RDX (hexogen). The authors demonstrated a 100% success of ignition over a 0.25–4 A firing current range, corresponding to 80–244 μJ, and with response times ranging from 2 to 260 μs. Then, Glavier et al. [14] used this igniter to cut and propel a thin metallic foil and ignite RDX in a detonation. Authors showed that a stainless-steel flyer of 40 mg can be properly cut and propelled at velocities calculated from 665 to 1083 m/s, as a function of the RDX's extent of compaction and ignition charge. The impact of the flyer can directly initiate the detonation of an RDX explosive, which is very promising in terms of removing primary explosives from detonators. A schematic view and photo of one miniature detonator is given in Figure 4.

**Figure 4.** (**a**) Schematic and (**b**) photo of the reactive electro-thermal initiator on Epoxy/PET membrane. Chip dimension is 4 mm × 2 mm. (**c**) Schematic and (**d**) photo of one miniature detonator integrating Al/CuO multilayers as the initiator.

More recently, Nicollet et al. [12] developed the Al/CuO multilayer igniter concept for several substrates, and simplified the fabrication process to adapt it to airbag initiation. Additionally, our team has demonstrated miniature one-shot circuit breakers [15], based on the combustion of a nanothermite. Each device is simply made from two assembled printed board circuits (PCBs) to define a hermetic cavity in which an Al/CuO multilayer initiator chip ignites—in less than 100 μs—a few milligrams of

nanothermite, to cut a thick copper connection (see Figure 5). The authors demonstrated the good operation (100% success rate) with a response time of 0.57 ms, which is much lower than the response time of classical mechanical circuit breakers (>ms).

**Figure 5.** (**a**) 3D schematic of the nanothermite-based circuit breaker with exploded views and (**b**) snapshots of high speed images taken for the tested circuit breaker. The time between each picture is 100 μs. Taken from [15].

### **5. Conclusions**

Sputter-deposited Al/CuO multilayers, composed of alternating Al and CuO nanolayers with total thicknesses ranging from 1 to tens of μm, present an opportunity for tunable ignition and actuation. The interest in and integration of igniter applications have been demonstrated for different applications, such as in inflators, aerospace or actuators.

**Author Contributions:** L.S. developed the sputter-deposition process and assisted in pyroMEMS development. S.A.-S. developed all the experimental benches and assisted in the characterization part. C.R. supervised the research and wrote the paper. All authors have read and agreed to the published version of the manuscript.

**Funding:** C.R. received funding from the European Research Council (ERC) under the European Union's Horizon 2020 research and innovation program (grant agreement No 832889—PyroSafe). L.S. received funding for sputter-deposition equipment from FEDER/Region Occitanie program Grant THERMIE.

**Acknowledgments:** The authors acknowledge support from the European Research Council (H2020 Excellent Science) Researcher Award (grant 832889—PyroSafe) and the Occitanie Region/European Union for their FEDER support (THERMIE grant). The authors thank Séverine Vivies from LAAS-CNRS and Teresa Hungria-Hernandez from "Centre de Micro Caractérisation Raymond Castaing (UMS 3623)" for their great support with the sample preparation and TEM experiments.

**Conflicts of Interest:** The authors declare no conflict of interests.

### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Article*