corresponded to the previous SEM observations. High-resolution (HR) TEM images of different regions of the TC200 nanorod are shown in Figure 3b,c. However, due to the repeated stacking of TiO2 and Cu2O, the lattice fringe arrangements in the inner region of the images cannot be easily distinguished. In contrast, clear lattice fringe arrangements can be observed in the outer regions of the HRTEM images, indicating the crystalline features of the decorated Cu2O layer. The spacing between these lattice fringes was measured to be approximately 0.24 nm and 0.3 nm in different orientations, and these lattice spacings corresponded to the interplanar spacings of the (111) and (110) planes of cuprite Cu2O, respectively [30]. Figure 3d shows selected area electron diffraction (SAED) patterns of several TC200 composite nanorods. It shows several diffraction spots arranged in concentric circles with different radii. These concentric circles correspond to rutile TiO2 ((110), (101), and (200)) and cuprite Cu2O ((111), (211), (110), and (200)). This confirms the formation of a crystalline Cu2O layer on the TiO2 nanorod. Figure 3e shows the cross-sectional EDS line-scan profiling spectra, in which the signal of Ti is distributed across the inner region of the nanorod, the signal of O is uniformly distributed over the entire nanorod region, and the signal of Cu is concentrated in the outer region of the nanorod. This indicates that the main core of the nanorod was TiO2 and the surface was covered with a layer of copper oxide. A TiO2 composite nanorod well-decorated with a Cu2O layer is visibly displayed. rangements can be observed in the outer regions of the HRTEM images, indicating the crystalline features of the decorated Cu2O layer. The spacing between these lattice fringes was measured to be approximately 0.24 nm and 0.3 nm in different orientations, and these lattice spacings corresponded to the interplanar spacings of the (111) and (110) planes of cuprite Cu2O, respectively [30]. Figure 3d shows selected area electron diffraction (SAED) patterns of several TC200 composite nanorods. It shows several diffraction spots arranged in concentric circles with different radii. These concentric circles correspond to rutile TiO2 ((110), (101), and (200)) and cuprite Cu2O ((111), (211), (110), and (200)). This confirms the formation of a crystalline Cu2O layer on the TiO2 nanorod. Figure 3e shows the cross-sectional EDS line-scan profiling spectra, in which the signal of Ti is distributed across the inner region of the nanorod, the signal of O is uniformly distributed over the entire nanorod region, and the signal of Cu is concentrated in the outer region of the nanorod. This indicates that the main core of the nanorod was TiO2 and the surface was covered with a layer of copper oxide. A TiO2 composite nanorod well-decorated with a Cu2O layer is visibly displayed.

Figure 3a shows a low-magnification TEM image of a single TC200 nanorod. The

entire TiO2 nanorod was uniformly covered by a continuous Cu2O film. Rough and irregular surface features can be observed on the sidewalls of the nanorod. The decorated copper oxide layer thickness was estimated to be approximately 32 nm. The feature that appeared corresponded to the previous SEM observations. High-resolution (HR) TEM

region of the images cannot be easily distinguished. In contrast, clear lattice fringe ar-

**Figure 3.** TEM analysis of TC200: (**a**) low-magnification TEM image; (**b**,**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC200 nanorods; (**e**) EDS line scanning profiles across the composite rod. **Figure 3.** TEM analysis of TC200: (**a**) low-magnification TEM image; (**b**,**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC200 nanorods; (**e**) EDS line scanning profiles across the composite rod.

Figure 4a shows a low magnification TEM image of a single TC350 nanorod. Compared with the TEM image of TC200 (Figure 3a), the copper film originally coated onto the surface of the TiO2 nanorod was transformed into a discontinuously decorated layer after the postannealing procedure. The discontinuous decorated layer consisted of numerous granular crystallites with a particle size of approximately 37 nm. Figure 4b,c present HRTEM images of different peripheral regions from Figure 4a. The decorated particles were further analyzed using HRTEM. The lattice fringes arranged with spacings of 0.25 nm, 0.27 nm, and 0.24 nm could be measured in the different orientations. They Figure 4a shows a low magnification TEM image of a single TC350 nanorod. Compared with the TEM image of TC200 (Figure 3a), the copper film originally coated onto the surface of the TiO2 nanorod was transformed into a discontinuously decorated layer after the postannealing procedure. The discontinuous decorated layer consisted of numerous granular crystallites with a particle size of approximately 37 nm. Figure 4b,c present HRTEM images of different peripheral regions from Figure 4a. The decorated particles were further analyzed using HRTEM. The lattice fringes arranged with spacings of 0.25 nm, 0.27 nm, and 0.24 nm could be measured in the different orientations. They corresponded to the (002) and (110) crystal planes of the CuO phase and the (111) crystal plane of the Cu2O phase, respectively. The HRTEM images showed that the crystalline features of the Cu2O

corresponded to the (002) and (110) crystal planes of the CuO phase and the (111) crystal

*Nanomaterials* **2022**, *12*, 2634 7 of 19

plane of the Cu2O phase, respectively. The HRTEM images showed that the crystalline features of the Cu2O and CuO phases coexisted in the decorated discontinuous layer. This further confirmed the results for the previous XRD patterns. When the annealing

two different oxides, which coexisted on the surface of the TiO2 nanorods. Figure 4d

and CuO phases coexisted in the decorated discontinuous layer. This further confirmed the results for the previous XRD patterns. When the annealing temperature was raised above 300 ◦C, the original pure copper film was transformed into two different oxides, which coexisted on the surface of the TiO2 nanorods. Figure 4d shows the SAED patterns of multiple TC350 composite nanorods. The contributions of the crystallographic planes of cuprite Cu2O, tenorite CuO, and rutile TiO2 are visibly exhibited, proving that the ternary phases of Cu2O, CuO, and TiO2 coexisted in the TC350 nanorods. Figure 4e presents the cross-sectional EDS line-scan profiling spectra of the TC350 nanorods. The Cu signal was very strong in the outer region, and the Ti signal was mainly distributed in the inner region of the composite nanorod. The O signal was evenly distributed over the composite nanorod. A TiO2 composite nanorod well-shelled with copper oxide is demonstrated here, and the EDS analysis revealed that the Cu/Ti had an atomic ratio of 0.24. When the annealing temperature was further increased to 400 ◦C, as the low-magnification TEM image (Figure 5a) of the TC400 nanorod shows, the size of the particles wrapped over the sidewall surface of the TiO2 nanorod changed significantly compared to TC350 (Figure 4a). The size of the particles wrapped over the surface of the TiO2 nanorod was further increased to 55–70 nm. These seriously agglomerated particles on the TiO2 with relative large sizes can be attributed to the marked increase in the annealing temperature, which led to a substantially increased rate of nucleation and accelerated crystal size growth under the given annealing condition. During particle coalescence, the initially formed copper oxide particles could migrate to the TiO2 nanorod template surface and coalesce if motion yielded a reduction in overall system energy. Evidence for such a thermal annealing-induced Ostwald ripening process has been provided in other heterogeneous catalyst systems [31]. shows the SAED patterns of multiple TC350 composite nanorods. The contributions of the crystallographic planes of cuprite Cu2O, tenorite CuO, and rutile TiO2 are visibly exhibited, proving that the ternary phases of Cu2O, CuO, and TiO2 coexisted in the TC350 nanorods. Figure 4e presents the cross-sectional EDS line-scan profiling spectra of the TC350 nanorods. The Cu signal was very strong in the outer region, and the Ti signal was mainly distributed in the inner region of the composite nanorod. The O signal was evenly distributed over the composite nanorod. A TiO2 composite nanorod well-shelled with copper oxide is demonstrated here, and the EDS analysis revealed that the Cu/Ti had an atomic ratio of 0.24. When the annealing temperature was further increased to 400 °C, as the low-magnification TEM image (Figure 5a) of the TC400 nanorod shows, the size of the particles wrapped over the sidewall surface of the TiO2 nanorod changed significantly compared to TC350 (Figure 4a). The size of the particles wrapped over the surface of the TiO2 nanorod was further increased to 55–70 nm. These seriously agglomerated particles on the TiO2 with relative large sizes can be attributed to the marked increase in the annealing temperature, which led to a substantially increased rate of nucleation and accelerated crystal size growth under the given annealing condition. During particle coalescence, the initially formed copper oxide particles could migrate to the TiO2 nanorod template surface and coalesce if motion yielded a reduction in overall system energy. Evidence for such a thermal annealing-induced Ostwald ripening process has been provided in other heterogeneous catalyst systems [31].

**Figure 4.** TEM analysis of TC350: (**a**) low-magnification TEM image; (**b**–**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC350 nanorods; (**e**) EDS line scanning profiles across the composite rod. **Figure 4.** TEM analysis of TC350: (**a**) low-magnification TEM image; (**b**,**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC350 nanorods; (**e**) EDS line scanning profiles across the composite rod.

temperature, which led to a substantially increased rate of nucleation and accelerated

When the annealing temperature was further increased to 400 °C, as the

low-magnification TEM image (Figure 5a) of TC400 nanorod shows, the size of the particles wrapped over the sidewall surface of the TiO2 nanorod changed significantly compared to TC350 (Figure 4a). The size of the particles wrapped over the surface of TiO2 nanorod was further increased to 55–70 nm. These seriously agglomerated particles with

phase and crystallite size on the TiO2 nanorod template.

crystal size growth under the given annealing condition. During particle coalescence, the initially formed copper oxide particles can migrate to the TiO2 nanorod template surface and coalesce if motion yields an overall system-energy reduction. Evidence for such a thermal annealing-induced Ostwald ripening process has been found in other heterogeneous catalyst systems [31]. Figure 5b,c show HRTEM images of the periphery of TC400 nanorod. Lattice spacings of 0.23 nm, 0.25 nm, and 0.27 nm in different orientations can be measured in Figure 5b,c, which corresponded to the interplanar distances of CuO (111), (002), and (110), respectively. These results confirm that the large-sized particles attached to the surface of TiO2 nanorod after annealing at 400 °C were CuO crystallites. Figure 5d shows the SAED patterns obtained from multiple TC400 composite nanorods. Obvious diffraction spots are arranged in concentric circles with different radii. Several crystallographic planes of CuO (111), (002), (11-2), (110), and (021) can be indexed in the SAED pattern. No other copper oxide phases were identified, indicating a TiO2-CuO composite structure for the TC400 nanorods. The TEM structural analyses showed the same results as revealed in XRD patterns. In addition, the cross-sectional elemental profiling spectra shown in Figure 5e also demonstrated a good compositional distribution for the copper oxide-decorated TiO2 nanorod composite structure. The TEM analysis re-

**Figure 5.** TEM analysis of TC400: (**a**) low-magnification TEM image; (**b**–**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC400 nanorods; (**e**) EDS line scanning profiles across the composite rod. **Figure 5.** TEM analysis of TC400: (**a**) low-magnification TEM image; (**b**,**c**) HRTEM images of various regions of the composite rod as marked in (**a**); (**d**) SAED patterns of several TC400 nanorods; (**e**) EDS line scanning profiles across the composite rod.

Figure 6a displays the high-resolution Cu 2p XPS spectra for TC200. The distinct peaks centered at approximately 932.4 eV and 952.4 eV can be attributed to Cu 2p3/2 and Cu 2p1/2, respectively. The Cu binding energies matched the Cu+1 binding state in the Cu2O phase, and this was consistent with the results from a report on the XPS analysis of a sol-gel-derived thin Cu2O film [32]. Figure 6b shows the high-resolution Cu 2p XPS spectra for TC 350. The appearance of the XPS spectra is similar to that observed in a study on CuO@Cu2O heterostructures derived using the solvothermal method [33]. In contrast to the Cu 2p spectra for TC200, oscillating satellite peaks could be detected for TC350 at the binding energies of approximately 942.2 eV and 961.8 eV, which further indicated the existence of a CuO phase in TC350. This has also been demonstrated in the Cu 2p spectra analysis of pristine Cu2O and CuO thin films, in which pure Cu2O and CuO could easily be observed without and with the appearance of satellite peaks from the XPS When the annealing temperature was further increased to 400 ◦C, as the low-magnification TEM image (Figure 5a) of TC400 nanorod shows, the size of the particles wrapped over the sidewall surface of the TiO2 nanorod changed significantly compared to TC350 (Figure 4a). The size of the particles wrapped over the surface of TiO2 nanorod was further increased to 55–70 nm. These seriously agglomerated particles with a relative large size on the TiO2 can be attributed to the marked increase in the annealing temperature, which led to a substantially increased rate of nucleation and accelerated crystal size growth under the given annealing condition. During particle coalescence, the initially formed copper oxide particles can migrate to the TiO2 nanorod template surface and coalesce if motion yields an overall system-energy reduction. Evidence for such a thermal annealing-induced Ostwald ripening process has been found in other heterogeneous catalyst systems [31]. Figure 5b,c show HRTEM images of the periphery of TC400 nanorod. Lattice spacings of 0.23 nm, 0.25 nm, and 0.27 nm in different orientations can be measured in Figure 5b,c, which corresponded to the interplanar distances of CuO (111), (002), and (110), respectively. These results confirm that the large-sized particles attached to the surface of TiO2 nanorod after annealing at 400 ◦C were CuO crystallites. Figure 5d shows the SAED patterns obtained from multiple TC400 composite nanorods. Obvious diffraction spots are arranged in concentric circles with different radii. Several crystallographic planes of CuO (111), (002), (11-2), (110), and (021) can be indexed in the SAED pattern. No other copper oxide phases were identified, indicating a TiO2-CuO composite structure for the TC400 nanorods. The TEM structural analyses showed the same results as revealed in XRD patterns. In addition, the cross-sectional elemental profiling spectra shown in Figure 5e also demonstrated a good compositional distribution for the copper oxide-decorated TiO2 nanorod composite structure. The TEM analysis results demonstrate that the annealing temperature effectively dominated the copper oxide phase and crystallite size on the TiO2 nanorod template.

> Figure 6a displays the high-resolution Cu 2p XPS spectra for TC200. The distinct peaks centered at approximately 932.4 eV and 952.4 eV can be attributed to Cu 2p3/2 and Cu 2p1/2, respectively. The Cu binding energies matched the Cu+1 binding state in the Cu2O phase, and this was consistent with the results from a report on the XPS analysis

of a sol-gel-derived thin Cu2O film [32]. Figure 6b shows the high-resolution Cu 2p XPS spectra for TC 350. The appearance of the XPS spectra is similar to that observed in a study on CuO@Cu2O heterostructures derived using the solvothermal method [33]. In contrast to the Cu 2p spectra for TC200, oscillating satellite peaks could be detected for TC350 at the binding energies of approximately 942.2 eV and 961.8 eV, which further indicated the existence of a CuO phase in TC350. This has also been demonstrated in the Cu 2p spectra analysis of pristine Cu2O and CuO thin films, in which pure Cu2O and CuO could easily be observed without and with the appearance of satellite peaks from the XPS spectra, respectively [32]. The spectra detected herein were further separated into several contributions. The intense fitted peaks located at 933.5 eV and 953.4 eV (blue line) were attributed to Cu 2p3/2 and Cu 2p1/2 of the CuO phase, respectively. There was a difference of approximately 20 eV between the Cu 2p3/2 and Cu 2p1/2 peaks of CuO, which matches well with the reported results for hydrothermally derived CuO nanoflowers [34]. In contrast, two relatively weak peaks (green line) appeared at 932.6 eV and 952.2 eV, corresponding to Cu 2p3/2 and Cu 2p1/2 of the Cu2O phase, respectively [35]. These results verify the coexistence of Cu2O and CuO phases in the decorated copper oxide layer in TC350. Figure 6c presents the high-resolution XPS spectra for Cu 2p in TC400. The distinct appearance of the satellite peaks (located at 942.2 and 961.8 eV) was observed (Figure 6c). The characteristic peaks centered at the binding energies of 933.3 eV and 953.3 eV corresponded to Cu 2p3/2 and Cu 2p1/2 of the CuO phase, respectively. These XPS results demonstrate that an adjustable copper oxide phase was obtained in the decorated copper film layer by varying the annealing temperature. *Nanomaterials* **2022**, *12*, 2634 9 of 19 spectra, respectively [32]. The spectra detected herein were further separated into several contributions. The intense fitted peaks located at 933.5 eV and 953.4 eV (blue line) were attributed to Cu 2p3/2 and Cu 2p1/2 of the CuO phase, respectively. There was a difference of approximately 20 eV between the Cu 2p3/2 and Cu 2p1/2 peaks of CuO, which matches well with the reported results for hydrothermally derived CuO nanoflowers [34]. In contrast, two relatively weak peaks (green line) appeared at 932.6 eV and 952.2 eV, corresponding to Cu 2p3/2 and Cu 2p1/2 of the Cu2O phase, respectively [35]. These results verify the coexistence of Cu2O and CuO phases in the decorated copper oxide layer in TC350. Figure 6c presents the high-resolution XPS spectra for Cu 2p in TC400. The distinct appearance of the satellite peaks (located at 942.2 and 961.8 eV) was observed (Figure 6c). The characteristic peaks centered at the binding energies of 933.3 eV and 953.3 eV corresponded to Cu 2p3/2 and Cu 2p1/2 of the CuO phase, respectively. These XPS results demonstrate that an adjustable copper oxide phase was obtained in the decorated copper film layer by varying the annealing temperature.

**Figure 6.** High-resolution XPS Cu 2p spectra: (**a**) TC200, (**b**) TC350, and (**c**) TC400. **Figure 6.** High-resolution XPS Cu 2p spectra: (**a**) TC200, (**b**) TC350, and (**c**) TC400.

Figure 7a shows the referenced Ti 2p core-level doublet spectra for the TiO2 nanorod template. The high-resolution XPS spectra were deconvoluted into four subpeaks. The more intense subpeaks at 458.3 eV and 463.9 eV corresponded to Ti 2p3/2 and Ti 2p1/2 for the Ti4+ valence state in TiO2, respectively. Furthermore, the subpeaks with weaker intensities and smaller binding energies of 457.2 eV and 462.9 eV corresponded to Ti 2p3/2 and Ti 2p1/2 in the Ti+3 valence state [36,37]. The presence of the mixed Ti4+/Ti3+ valance state indicates the possible presence of oxygen vacancies on the surfaces of the TiO2 nanorod template. Figure 7b shows a comparison of the Ti 2p core-level doublet spectra for TiO2, TC200, TC350, and TC400. It can be seen that the Ti 2p XPS spectra of the TC composite nanorods demonstrated positive shifts in binding energy positions in comparison with the binding energy position of pristine TiO2. The modification of TiO2 nanorods with copper oxides described herein might have changed the electronic state of Ti in Ti-O because of the formation of heterojunctions between the n-type TiO2 and p-type copper oxides. This has been demonstrated with CuO@TiO2 powders and core–shell N-TiO2@CuOx heterojunction composites formed using ball milling [38,39]. Notably, the Cu/Ti atomic ratio of TC350 was evaluated to be 3.6. The investigation depth of XPS is usually below 10 nm. This Cu/Ti atomic ratio substantially differs from the Cu/Ti atomic Figure 7a shows the referenced Ti 2p core-level doublet spectra for the TiO2 nanorod template. The high-resolution XPS spectra were deconvoluted into four subpeaks. The more intense subpeaks at 458.3 eV and 463.9 eV corresponded to Ti 2p3/2 and Ti 2p1/2 for the Ti4+ valence state in TiO2, respectively. Furthermore, the subpeaks with weaker intensities and smaller binding energies of 457.2 eV and 462.9 eV corresponded to Ti 2p3/2 and Ti 2p1/2 in the Ti+3 valence state [36,37]. The presence of the mixed Ti4+/Ti3+ valance state indicates the possible presence of oxygen vacancies on the surfaces of the TiO2 nanorod template. Figure 7b shows a comparison of the Ti 2p core-level doublet spectra for TiO2, TC200, TC350, and TC400. It can be seen that the Ti 2p XPS spectra of the TC composite nanorods demonstrated positive shifts in binding energy positions in comparison with the binding energy position of pristine TiO2. The modification of TiO2 nanorods with copper oxides described herein might have changed the electronic state of Ti in Ti-O because of the formation of heterojunctions between the n-type TiO2 and p-type copper oxides. This has been demonstrated with CuO@TiO2 powders and core–shell N-TiO2@CuOx heterojunction composites formed using ball milling [38,39]. Notably, the Cu/Ti atomic ratio of TC350 was evaluated to be 3.6. The investigation depth of XPS is usually below 10 nm. This Cu/Ti atomic ratio substantially differs from the Cu/Ti atomic ratios calculated from the EDS

ratios calculated from the EDS spectra of electron microscopes because of the different

measurement depths of the various analysis methods.

spectra of electron microscopes because of the different measurement depths of the various analysis methods. *Nanomaterials* **2022**, *12*, 2634 10 of 19

**Figure 7.** High-resolution XPS Ti 2p spectra: (**a**) TiO2 and (**b**) comparison of Ti 2p spectra of various samples. **Figure 7.** High-resolution XPS Ti 2p spectra: (**a**) TiO2 and (**b**) comparison of Ti 2p spectra of various samples.

Figure 8a shows the optical absorption characteristics of pristine TiO2 nanorods and TC composite nanorods. A sharp absorption drop appeared at approximately 410 nm for TiO2 nanorod template, and this absorption edge was consistent with the inherent band-gap absorption of rutile TiO2 [1]. Notably, the TC composite nanorods demonstrated a significant red-shift extension of the absorption edge in comparison with that of the pristine TiO2. This can be attributed to the decoration of the TiO2 nanorods with Cu2O and CuO visible-light sensitizers. These visible-light sensitizers helped to absorb the longer wavelength spectra, making up for the inability of TiO2 to absorb visible light, and enhanced the absorption in the visible light region. The higher annealing temperature resulted in a larger size for the red-shift of the absorption edge of the TC composite samples; this was associated with the fact that the CuO formed at the higher annealing temperature had a narrower band-gap energy than that of Cu2O [40,41]. Figure 8b shows the Kubelka–Munk function (F(R)) vs. energy plots for various nanorod samples [42]. Notably, the TiO2 and copper oxides used herein were expected to exhibit a direct transition in the band-gap measurements. Therefore, the band-gap energy of the TiO2 nanorods and TC composite nanorods could be deduced from the (F(R)*hv*)2 vs. *hv* plots by extrapolating the straight portion of the curves to the energy axis. The TiO2 nanorod template was estimated to have an energy gap of approximately 3.03 eV. The energy gap values of TC 200, TC300, TC350, and TC400 were estimated to be approximately 2.59 eV, 2.43 eV, 2.34 eV, and 2.27 eV, respectively. The phase evolution of the decorated layer from Cu2O to CuO with increased annealing temperature visibly demonstrated a decreased energy gap in the TC composite nanorods. The band-gap energy variation in the copper oxides due to the phase evolution was consistent with a report on electrodeposited Cu2O/CuO powder oxides [43]. The UV-vis analysis demonstrated that the energy gap size of the TC composite nanorods could be effectively tuned by varying the postannealing temperature. In addition, the energy gaps of single CuO and Cu2O films were also estimated from the Tauc plot (Figure 8c). The energy gap values for CuO and Cu2O were estimated to be 1.76 eV and 2.04 eV, respectively, by extrapolating the curve tangent to the energy axis in Figure 8c. These values are similar to those from previous work on Cu2O formed with copper foil annealing and sputtering CuO [44,45]. Figure 8a shows the optical absorption characteristics of pristine TiO2 nanorods and TC composite nanorods. A sharp absorption drop appeared at approximately 410 nm for TiO2 nanorod template, and this absorption edge was consistent with the inherent band-gap absorption of rutile TiO2 [1]. Notably, the TC composite nanorods demonstrated a significant red-shift extension of the absorption edge in comparison with that of the pristine TiO2. This can be attributed to the decoration of the TiO2 nanorods with Cu2O and CuO visible-light sensitizers. These visible-light sensitizers helped to absorb the longer wavelength spectra, making up for the inability of TiO2 to absorb visible light, and enhanced the absorption in the visible light region. The higher annealing temperature resulted in a larger size for the red-shift of the absorption edge of the TC composite samples; this was associated with the fact that the CuO formed at the higher annealing temperature had a narrower band-gap energy than that of Cu2O [40,41]. Figure 8b shows the Kubelka–Munk function (F(R)) vs. energy plots for various nanorod samples [42]. Notably, the TiO2 and copper oxides used herein were expected to exhibit a direct transition in the band-gap measurements. Therefore, the band-gap energy of the TiO2 nanorods and TC composite nanorods could be deduced from the (F(R)*hv*) 2 vs. *hv* plots by extrapolating the straight portion of the curves to the energy axis. The TiO2 nanorod template was estimated to have an energy gap of approximately 3.03 eV. The energy gap values of TC 200, TC300, TC350, and TC400 were estimated to be approximately 2.59 eV, 2.43 eV, 2.34 eV, and 2.27 eV, respectively. The phase evolution of the decorated layer from Cu2O to CuO with increased annealing temperature visibly demonstrated a decreased energy gap in the TC composite nanorods. The band-gap energy variation in the copper oxides due to the phase evolution was consistent with a report on electrodeposited Cu2O/CuO powder oxides [43]. The UV-vis analysis demonstrated that the energy gap size of the TC composite nanorods could be effectively tuned by varying the postannealing temperature. In addition, the energy gaps of single CuO and Cu2O films were also estimated from the Tauc plot (Figure 8c). The energy gap values for CuO and Cu2O were estimated to be 1.76 eV and 2.04 eV, respectively, by extrapolating the curve tangent to the energy axis in Figure 8c. These values are similar to those from previous work on Cu2O formed with copper foil annealing and sputtering CuO [44,45].

samples.

**Figure 8.** (**a**) Absorption spectra of various samples and Band-gap evaluations of (**b**) various nanorod samples, (**c**) pristine Cu2O and CuO films.

work on Cu2O formed with copper foil annealing and sputtering CuO [44,45].

**Figure 7.** High-resolution XPS Ti 2p spectra: (**a**) TiO2 and (**b**) comparison of Ti 2p spectra of various

Figure 8a shows the optical absorption characteristics of pristine TiO2 nanorods and TC composite nanorods. A sharp absorption drop appeared at approximately 410 nm for TiO2 nanorod template, and this absorption edge was consistent with the inherent band-gap absorption of rutile TiO2 [1]. Notably, the TC composite nanorods demonstrated a significant red-shift extension of the absorption edge in comparison with that of the pristine TiO2. This can be attributed to the decoration of the TiO2 nanorods with Cu2O and CuO visible-light sensitizers. These visible-light sensitizers helped to absorb the longer wavelength spectra, making up for the inability of TiO2 to absorb visible light, and enhanced the absorption in the visible light region. The higher annealing temperature resulted in a larger size for the red-shift of the absorption edge of the TC composite samples; this was associated with the fact that the CuO formed at the higher annealing temperature had a narrower band-gap energy than that of Cu2O [40,41]. Figure 8b shows the Kubelka–Munk function (F(R)) vs. energy plots for various nanorod samples [42]. Notably, the TiO2 and copper oxides used herein were expected to exhibit a direct transition in the band-gap measurements. Therefore, the band-gap energy of the TiO2 nanorods and TC composite nanorods could be deduced from the (F(R)*hv*)2 vs. *hv* plots by extrapolating the straight portion of the curves to the energy axis. The TiO2 nanorod template was estimated to have an energy gap of approximately 3.03 eV. The energy gap values of TC 200, TC300, TC350, and TC400 were estimated to be approximately 2.59 eV, 2.43 eV, 2.34 eV, and 2.27 eV, respectively. The phase evolution of the decorated layer from Cu2O to CuO with increased annealing temperature visibly demonstrated a decreased energy gap in the TC composite nanorods. The band-gap energy variation in the copper oxides due to the phase evolution was consistent with a report on electrodeposited Cu2O/CuO powder oxides [43]. The UV-vis analysis demonstrated that the energy gap size of the TC composite nanorods could be effectively tuned by varying the postannealing temperature. In addition, the energy gaps of single CuO and Cu2O films were also estimated from the Tauc plot (Figure 8c). The energy gap values for CuO and Cu2O were estimated to be 1.76 eV and 2.04 eV, respectively, by extrapolating the curve tangent to the energy axis in Figure 8c. These values are similar to those from previous

Figure 9a shows the transient photoresponses of various samples. Irradiation was applied with the full-band spectrum, and a bias potential of 1.2 V was used to measure the photocurrent. Photocurrent generation occurred entirely as a result of the on and off responses to the irradiation. Seven cycles of on/off irradiation were repeated, as shown in Figure 9a, and all samples could obtain a stable photogenerated current when the irradiation was turned on, indicating that the samples were stable under cycling chopping irradiation. A higher photocurrent indicates better efficiency for the separation of photogenerated charges and better photocatalytic activity for the photoelectrode [46]. Comparatively, all the TC composite nanorods exhibited improved photoresponses compared to that of the pristine TiO2. This was attributed to the fact that decoration with Cu2O and CuO visiblelight sensitizers enhanced the light-harvesting ability of the TiO2 nanorod template, and the formation of heterojunctions in the composite system resulted in improved photogenerated carrier separation efficiency. Furthermore, compared to TC200, which had a single-phase Cu2O decoration, TC400 (with single-phase CuO decoration) had a higher photocurrent, which can be attributed to the narrower energy gap in CuO compared to Cu2O. This led to TC400 absorbing across a longer wavelength range than TC200, as revealed in the previous absorption analysis, thereby increasing light absorption and promoting the photoexcited carrier density. This has also been demonstrated in previous work on the photoactive performance of a Cu2O/CuO system [32,47]. Notably, TC300 and TC350 displayed the best photoresponse abilities among the various nanorod samples, revealing that the composite nanorod system decorated with dual Cu2O and CuO phases was a more efficient material combination for enhancing the photoactivity of the copper oxide–TiO2 composite nanorods. Figure 9b presents the Nyquist plots of various samples measured at the frequency range from 100 kHz to 0.1 Hz and a potential amplitude of 10 mV. The radius of the semicircles in Nyquist plots is associated with the interfacial charge transfer resistance [48]. Notably, TC350 had the smallest semicircular radius, and the pure TiO2 nanorod template exhibited the largest semicircular radius, indicating that TC350 had the smallest charge transfer resistance and TiO2 the largest. The sizes of the semicircle radii from the Nyquist plots for various samples were ordered in the following trend: TiO2 > TC200 > TC400 > TC300 > TC350. This result was also found with the transient photoresponse measurements. The multi-interface heterostructures consisting of TiO2, CuO, and Cu2O in TC350 and TC300 effectively helped to enhance the separation and transfer abilities of electron–hole pairs, as revealed in the previous I-t curves (Figure 9a). Similar coexistence of ternary phases leading to substantial improvements in PEC properties has also been demonstrated in BiVO4/CdS/CoOx core-shell composites [49]. These improvements can provide an opportunity to induce electron redistribution and synergistic effects at the interfaces for heterogeneous catalysis consisting of two or more components connected by well-defined interfaces [50,51]. The existence of multiple heterointerfaces in ternary phase composites improves their PEC properties. The charge transfer resistance can be estimated by fitting the arc radius of the Nyquist curves according to the proposed equivalent circuits in Figure 9c. Rs, CPE, and Rct represent the series resistance, constant phase element, and charge transfer resistance, respectively. Similar equivalent circuits have also been used in a

ternary Fe2O3–MoS2–Cu2O nanofilm system to determine the Rct [52]. The representative fitting parameters for TC350 were Rct = 582 Ohm and Rs = 43.89 Ohm. After fitting the Nyquist plots using the proposed equivalent circuits, the Rct values for the other samples, TiO2, TC200, TC300, and TC400, were 3653, 1832, 702, and 1284 Ohm, respectively. Notably, although the TC300 and TC350 were both ternary-phase composite nanorods, lower interfacial charge transfer resistance in TC350 was observed in comparison to that of TC300. This might have been associated with the fact that, as the annealing temperature increases, the crystallite size of the decorated copper oxides increased, and this could have reduced the grain boundaries in the decoration layer. Therefore, TC350 had a better charge transport ability than TC300, and this was evidenced in the Rct. further calculated to be 0.56 and 0.81 eV (vs. NHE), respectively. The VB positions assessed herein are close to previously reported results for Cu2O and CuO [58,59]. Figure 10d shows the M-S curves for various TC composite nanorods. Inverted V-shaped M-S curves were observed for the all composite nanorods, demonstrating that the composites had both n-type and p-type electronic properties and confirming the formation of p-n junctions in the TC composite nanorods [60]. Construction of p-n junctions in composite systems has been posited to be a sensible strategy to enhance photocatalytic activity. The formation of a p-n junction with space charge regions at the heterointerface could induce the electric field-driven diffusion of electrons and holes and further inhibit the recombination of photogenerated charges [54,61].

grain boundaries in the decoration layer. Therefore, TC350 had a better charge transport

In order to further analyze and construct the energy-band structure of the composite nanorods, measurements of the flat-band potential of the TiO2 nanorod template, Cu2O film, and CuO film were carried out and presented in Figure 10a–c. The M-S curves exhibited a positive slope for the TiO2 and negative slopes for the Cu2O and CuO, revealing the n-type nature of the TiO2 and the p-type nature of the Cu2O and CuO. According to the M-S equation [53], when 1/C2 is extrapolated to a value equal to 0, the X-axis intercept is equal to the flat-band potential of the material [54]. The flat-band potential of pure TiO2 was estimated to be about −0.11 eV (vs. NHE). The flat-band potential in n-type semiconductors is closer to the conduction band (CB) and the CB position of an n-type semiconductor is generally more negative (0.1 eV) than the flat-band potential [55]. After calculation, it was deduced that the CB of TiO2 was −0.21 eV. In contrast, the flat-band potential of the p-type semiconductor is closer to its valence band (VB) [56,57]. The flat-band potentials of Cu2O and CuO were estimated to be approximately 0.46 eV and 0.71 eV, respectively, as shown in Figure 10b,c. The VB positions of Cu2O and CuO were

*Nanomaterials* **2022**, *12*, 2634 12 of 19

ability than TC300, and this was evidenced in the Rct.

**Figure 9.** (**a**) Photocurrent density versus time curves for various samples at 1.2 V (vs. Ag/AgCl) under chopping illumination. (**b**) Nyquist plots for various samples under irradiation. (**c**) Possible equivalent circuits for Rct evaluation. **Figure 9.** (**a**) Photocurrent density versus time curves for various samples at 1.2 V (vs. Ag/AgCl) under chopping illumination. (**b**) Nyquist plots for various samples under irradiation. (**c**) Possible equivalent circuits for Rct evaluation.

In order to further analyze and construct the energy-band structure of the composite nanorods, measurements of the flat-band potential of the TiO2 nanorod template, Cu2O film, and CuO film were carried out and presented in Figure 10a–c. The M-S curves exhibited a positive slope for the TiO2 and negative slopes for the Cu2O and CuO, revealing the n-type nature of the TiO2 and the p-type nature of the Cu2O and CuO. According to the M-S equation [53], when 1/C2 is extrapolated to a value equal to 0, the X-axis intercept is equal to the flat-band potential of the material [54]. The flat-band potential of pure TiO2 was estimated to be about −0.11 eV (vs. NHE). The flat-band potential in n-type semiconductors is closer to the conduction band (CB) and the CB position of an n-type semiconductor is generally more negative (0.1 eV) than the flat-band potential [55]. After calculation, it was deduced that the CB of TiO2 was −0.21 eV. In contrast, the flat-band potential of the p-type semiconductor is closer to its valence band (VB) [56,57]. The flat-band potentials of Cu2O and CuO were estimated to be approximately 0.46 eV and 0.71 eV, respectively, as shown in Figure 10b,c. The VB positions of Cu2O and CuO were further calculated to be 0.56 and 0.81 eV (vs. NHE), respectively. The VB positions assessed herein are close to previously reported results for Cu2O and CuO [58,59]. Figure 10d shows the M-S curves for various TC composite nanorods. Inverted V-shaped M-S curves were observed for the all composite nanorods, demonstrating that the composites had both n-type and p-type electronic properties and confirming the formation of p-n junctions in the TC composite nanorods [60]. Construction of p-n junctions in composite systems has been posited to be a sensible strategy to enhance photocatalytic activity. The formation of a p-n junction with space charge regions at the heterointerface could induce the electric field-driven

diffusion of electrons and holes and further inhibit the recombination of photogenerated charges [54,61]. *Nanomaterials* **2022**, *12*, 2634 13 of 19

**Figure 10.** Mott–Schottky plots for various samples: (**a**) TiO2, (**b**) Cu2O, and (**c**) CuO. (**d**) A series of M−S plots for various composite nanorods. **Figure 10.** Mott–Schottky plots for various samples: (**a**) TiO2 , (**b**) Cu2O, and (**c**) CuO. (**d**) A series of M−S plots for various composite nanorods.

The photocatalytic ability of the samples was further estimated by using the formula whereby the percentage degradation = C/Co, where Co is the initial concentration of RhB solution and C is the time-dependent concentration of RhB solution upon irradiation. Figure 11a presents C/Co vs. irradiation duration plots for RhB solution with different samples. Adsorption–desorption equilibrium was reached by placing the photocatalysts in the RhB solution for 45 min in the dark before starting the photodegradation experiments. Under dark equilibration conditions for 45 min, the C/Co values for TiO2, TC200, TC300, TC350, and TC400 were approximately 3.1%, 4.2%, 6.8%, 7.9%, and 5%, respectively. This indicated that the TiO2 nanorod template decorated with copper oxide had an improved surface dye absorption capacity. After offsetting with a dark adsorption contribution, the degradation rates of TC200, TC300, TC350, and TC400 were approximately 59%, 83%, 90%, and 70%, respectively, with 60 min irradiation. The TC composite nanorods exhibited improved photodegradation abilities towards RhB solution in comparison to the pristine TiO2 nanorod template. Furthermore, among the various TC composites, TC350 had the highest photodegradation ability towards RhB solution under the given test conditions. In addition, the photodegradation kinetics of the RhB solution with all samples were also investigated and presented in Figure 11b. The pseudo-first-order kinetic equation is expressed as: kt = ln Co/C, where k represents the pseudo-first-order rate constant (min−1) for the initial degradation [61]. All the TC composite samples displayed larger k values than that of the pristine TiO2. Furthermore, TC350 had the highest k value of 0.04578 min−1. The photodegradation abilities of the photocatalysts towards organic pollutants were significantly related to the separation efficiency for electrons and holes. The magnitude trends for the k values for the various samples investigated herein were consistent with the previously measured PEC and EIS experimental results. In addition, the photocatalytic reaction was closely related to the active species produced in the process. The role of these species in the degradation reaction was investigated by measuring the variation in the degradation performance of the RhB solution with TC350 through the addition of various radical scavengers after 60 min irradiation. The radical The photocatalytic ability of the samples was further estimated by using the formula whereby the percentage degradation = C/Co, where Co is the initial concentration of RhB solution and C is the time-dependent concentration of RhB solution upon irradiation. Figure 11a presents C/Co vs. irradiation duration plots for RhB solution with different samples. Adsorption–desorption equilibrium was reached by placing the photocatalysts in the RhB solution for 45 min in the dark before starting the photodegradation experiments. Under dark equilibration conditions for 45 min, the C/Co values for TiO2, TC200, TC300, TC350, and TC400 were approximately 3.1%, 4.2%, 6.8%, 7.9%, and 5%, respectively. This indicated that the TiO2 nanorod template decorated with copper oxide had an improved surface dye absorption capacity. After offsetting with a dark adsorption contribution, the degradation rates of TC200, TC300, TC350, and TC400 were approximately 59%, 83%, 90%, and 70%, respectively, with 60 min irradiation. The TC composite nanorods exhibited improved photodegradation abilities towards RhB solution in comparison to the pristine TiO2 nanorod template. Furthermore, among the various TC composites, TC350 had the highest photodegradation ability towards RhB solution under the given test conditions. In addition, the photodegradation kinetics of the RhB solution with all samples were also investigated and presented in Figure 11b. The pseudo-first-order kinetic equation is expressed as: kt = ln Co/C, where k represents the pseudo-first-order rate constant (min−1 ) for the initial degradation [61]. All the TC composite samples displayed larger k values than that of the pristine TiO2. Furthermore, TC350 had the highest k value of 0.04578 min−1 . The photodegradation abilities of the photocatalysts towards organic pollutants were significantly related to the separation efficiency for electrons and holes. The magnitude trends for the k values for the various samples investigated herein were consistent with the previously measured PEC and EIS experimental results. In addition, the photocatalytic reaction was closely related to the active species produced in the process. The role of these species in the degradation reaction was investigated by measuring the variation in the degradation performance of the RhB solution with TC350 through the addition of various radical scavengers after 60 min irradiation. The radical capture experiments were

capture experiments were performed using tert-butanol (TBA) as a hydroxyl radical (·OH) scavenger, ammonium oxalate (AO) as a hole quencher, and benzoquinone (BQ) as

performed using tert-butanol (TBA) as a hydroxyl radical (·OH) scavenger, ammonium oxalate (AO) as a hole quencher, and benzoquinone (BQ) as a superoxide radical (·O2 −) scavenger. As shown in Figure 11c, when 1 mM AO was added, the RhB degradation efficiency slightly decreased to 69%, indicating that holes played a minor role in the degradation process. In contrast, adding TBA or BQ scavengers resulted in a more intense decrease in the photodegradation level of the RhB solution. This shows that ·O2 − and ·OH were the main radicals involved in the photodegradation process of the RhB solution with TC350. Comparatively, the removal of the ·O2 − active species resulted in the most significant decrease in the degradation efficiency. TC300. This can be attributed to optical absorption and the microstructural differences between the TC300 and TC350. The TC350 exhibited a better light absorption ability than TC300, as revealed by the previous UV-vis absorption analysis, which enhanced the generation efficiency of photoexcited charges in TC350. Moreover, TC350 also had a larger surface particle size in the decorated copper oxide layer in comparison to that of TC300. A larger grain size reduces grain boundaries in the decorated copper oxide layer, resulting in enhanced charge transport [47]. The ternary phase and suitable microstructural and optical properties of TC350 mean that it has excellent photoactivity compared to the other TC composite samples.

This is supported by work on introducing a protective layer of TiO2 in Cu2O–CuO heterojunction films to prevent the photocorrosion effect [69]. The multiple charge transfer routes shown in Figure 12 explain the superior photoactive performance of TiO2–Cu2O– CuO composite systems (TC350 and TC300) among the various TC composite nanorods. Finally, it should be mentioned that TC350 exhibited better photoactivity than that of

*Nanomaterials* **2022**, *12*, 2634 15 of 19

**Figure 11.** (**a**) C/Co vs. irradiation duration plots; (**b**) ln (Co/C) vs. irradiation duration plots; (**c**) degradation percentages of RhB solution with TC350 in the presence of various scavengers. **Figure 11.** (**a**) C/Co vs. irradiation duration plots; (**b**) ln (Co/C) vs. irradiation duration plots; (**c**) degradation percentages of RhB solution with TC350 in the presence of various scavengers.

The band structures of pristine TiO2, Cu2O, and CuO were constructed according to the M-S measurements and the UV-vis analysis results, as shown in Figure 12. As shown in the previous scavenger experiments, the main active species involved in the TC350 photodegradation process with RhB solution were superoxide and hydroxyl radicals. Moreover, superoxide radicals demonstrated a greater contribution than hydroxyl radicals, as seen from the constructed band alignment in the ternary TC350 composite nanorods. If the electron–hole transfer route in the TC350 composite nanorods had followed the type II transfer mode, superoxide and hydroxyl radicals would not have been formed according to the relative band positions of the CB, VB, and redox potentials [62]. Therefore, none of the electrons/holes at the CB/VB positions would reach the required redox potential, so superoxide and hydroxyl active species would not have been produced with this mechanism. This contradicts the previous scavenger experiments. The Z-scheme mechanism shown in Figure 12 is more appropriate to explain the movement of photogenerated electrons/holes and the generation of active species for photodegradation. Under irradiation, photoinduced carriers form in the composite system (reaction 3). Through the movement of photogenerated carriers in the Z-scheme mechanism, the holes finally accumulated in the VB of TiO2 (2.82 eV vs. NHE), which was significantly higher than the oxidation potential of water or (–OH) molecules, which is 2.4 eV. Therefore, the holes were able to react with water (or –OH) molecules and generate hydroxyl radicals (reaction 4) [63]. In contrast, electrons accumulated in the CB of Cu2O (−1.48 eV vs. NHE). The electrons were located significantly lower than the reduction potential of oxygen (−0.33 eV), and electrons could react with oxygen to form superoxide radicals (reaction 5) [64]. These main reactive species

**References**

could further react with RhB dye molecules and decompose into carbon dioxide and water (reaction 6) [65]:

$$\text{TiO}\_2/\text{CuO}/\text{Cu}\_2\text{O} + hv\_{(\text{UV}-visible)} \rightarrow \text{TiO}\_2\left(e\_{\text{CB}}^-h\_{VB}^+\right)/\text{CuO}\left(e\_{\text{CB}}^-h\_{VB}^+\right)/\text{Cu}\_2\text{O}\left(e\_{\text{CB}}^-h\_{VB}^+\right) \tag{3}$$

$$\text{TiO}\_2\left(h\_{VB}^+\right) + \text{H}\_2\text{O} \rightarrow \cdot \text{OH} + \text{H}^+ + \text{TiO}\_2\tag{4}$$

$$\text{Cu}\_2\text{O}(\text{e}\_\text{CB}^-) + \text{O}\_2 \rightarrow \cdot \text{O}\_2^- + \text{Cu}\_2\text{O}\tag{5}$$

$$\cdot \cdot \text{OH} + \cdot \text{O}\_2^- + \text{RhB} \rightarrow \cdot \text{product} \,\left(\text{ex} : \text{H}\_2\text{O} + \text{CO}\_2\right) \tag{6}$$

**Figure 12.** Possible mechanisms of charge transfer in TC350 under irradiation. **Figure 12.** Possible mechanisms of charge transfer in TC350 under irradiation.

**4. Conclusions** The morphology of copper oxide decorated on a TiO2 nanorod template changed from a continuous layer morphology to granular aggregates when the postannealing temperature was varied from 200 to 400 °C. The composite nanorods formed at 350 °C (TC350) exhibited superior photoactive performance compared to the other composite nanorods. The larger particle size resulting from the copper oxide modification in TC350 reduced the grain boundaries in the decorated layer, thereby increasing the charge transport ability. Moreover, the surface-modified Cu2O–CuO mixed crystallites on the TiO2 template could absorb sunlight more efficiently. These factors enhanced the photoactive performance of the TC350 composite nanorods studied herein. The scavenger tests demonstrated that the Z scheme was the possible carrier movement mechanism in TC350 under irradiation, and that result explains the high photocatalytic degradation ability of TC350 towards organic pollutants. The experimental results obtained herein demonstrate that regulation of the composition phase and microstructure of the modified copper oxide layer through control of the thermal annealing budget for the thin copper layer on TiO2 nanorod templates is a promising approach to design copper oxide–TiO2 composite nanorods with satisfactory photoactive performance. **Author Contributions:** Methodology, Y.-C.L. and T.-H.L.; formal analysis, T.-H.L.; investigation, T.-H.L.; writing—original draft preparation, Y.-C.L. and T.-H.L.; supervision, Y.-C.L. All authors have read and agreed to the published version of the manuscript. **Funding:** This research was funded by Ministry of Science and Technology of Taiwan, grant no. Notably, from among the band alignments proposed for the TiO2 , Cu2O, and CuO, multiple photoinduced charger transfer routes could occur in the ternary TiO2–Cu2O–CuO composite system. The p-n junctions formed between the n-type TiO2 and p-type copper oxides induced an internal electric field at the heterointerfaces, promoting charge separation under irradiation. The stepped-band edge arrangement in the composite system caused multiple Z-scheme transfer routes for the photoinduced charges. This finally resulted in the accumulation of holes at the VB of TiO2 and of electrons at the CB of Cu2O. A similar Z-scheme carrier movement was also exhibited in a ternary ZnO–Cu2O–CuO photocatalyst system [66]. In the report by Wei et al., the composite material TiO2–Cu2O showed carrier movement with a Z scheme under irradiation [67]. These examples echo the carrier movement mechanism proposed in this work. In addition, the Z-scheme charge transfer in the composite nanorods had an important contribution in preventing the photocorrosion of Cu2O. Photocorrosion has been demonstrated in previous work on single-phase Cu2O photocatalysts [68]. The Cu2O phase coupling with TiO2 (TC200) or TiO2–CuO (TC300 and TC350) in the composite system effectively guided the photoexcited electrons and holes accumulated in the Cu2O and inhibited photocorrosion. Therefore, a stable photocurrent curve could be observed in the previous photoresponse plots. This is supported by work on introducing a protective layer of TiO2 in Cu2O–CuO heterojunction films to prevent the photocorrosion effect [69]. The multiple charge transfer routes shown in Figure 12 explain the superior photoactive performance of TiO2–Cu2O–CuO composite systems (TC350 and TC300) among the various TC composite nanorods. Finally, it should be mentioned that TC350 exhibited better photoactivity than that of TC300. This can be attributed to optical absorption and the microstructural differences between the TC300 and TC350. The TC350 exhibited a better light absorption ability than TC300, as revealed by the previous UV-vis absorption analysis, which enhanced the generation efficiency of photoexcited charges in TC350. Moreover, TC350 also had a larger surface particle size in the decorated copper oxide layer in comparison to that of TC300. A larger grain size reduces grain boundaries in the decorated copper oxide layer, resulting in enhanced charge

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

1. Liang, Y.-C.; Zhao, W.-C. Morphology-dependent photocatalytic and gas-sensing functions of three-dimensional TiO2–ZnO

2. Wei, N.; Cui, H.; Wang, C.; Zhang, G.; Song, Q.; Sun, W.; Song, X.; Sun, M.; Tian, J. Bi2O3 nanoparticles incorporated porous TiO2 films as an effective p-n junction with enhanced photocatalytic activity. *J. Am. Ceram. Soc.* **2017**, *100*, 1339–1349.

MOST 108-2221-E-019-034-MY3.

nanoarchitectures. *CrystEngComm* **2020**, *22*, 7575–7589.

transport [47]. The ternary phase and suitable microstructural and optical properties of TC350 mean that it has excellent photoactivity compared to the other TC composite samples.

#### **4. Conclusions**

The morphology of copper oxide decorated on a TiO2 nanorod template changed from a continuous layer morphology to granular aggregates when the postannealing temperature was varied from 200 to 400 ◦C. The composite nanorods formed at 350 ◦C (TC350) exhibited superior photoactive performance compared to the other composite nanorods. The larger particle size resulting from the copper oxide modification in TC350 reduced the grain boundaries in the decorated layer, thereby increasing the charge transport ability. Moreover, the surface-modified Cu2O–CuO mixed crystallites on the TiO2 template could absorb sunlight more efficiently. These factors enhanced the photoactive performance of the TC350 composite nanorods studied herein. The scavenger tests demonstrated that the Z scheme was the possible carrier movement mechanism in TC350 under irradiation, and that result explains the high photocatalytic degradation ability of TC350 towards organic pollutants. The experimental results obtained herein demonstrate that regulation of the composition phase and microstructure of the modified copper oxide layer through control of the thermal annealing budget for the thin copper layer on TiO2 nanorod templates is a promising approach to design copper oxide–TiO2 composite nanorods with satisfactory photoactive performance.

**Author Contributions:** Methodology, Y.-C.L. and T.-H.L.; formal analysis, T.-H.L.; investigation, T.-H.L.; writing—original draft preparation, Y.-C.L. and T.-H.L.; supervision, Y.-C.L. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by Ministry of Science and Technology of Taiwan, grant no. MOST 108-2221-E-019-034-MY3.

**Data Availability Statement:** Not applicable.

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

#### **References**

## *Article* **Experimental Study on the Stability of a Novel Nanocomposite-Enhanced Viscoelastic Surfactant Solution as a Fracturing Fluid under Unconventional Reservoir Stimulation**

**Xiaodong Si 1,2 , Mingliang Luo 1,2,\* , Mingzhong Li 1,2, Yuben Ma 3 , Yige Huang 1,2 and Jingyang Pu 1,2**

**Abstract:** Fe3O4@ZnO nanocomposites (NCs) were synthesized to improve the stability of the wormlike micelle (WLM) network structure of viscoelastic surfactant (VES) fracturing fluid and were characterized by Fourier transform infrared spectrometry (FT-IR), scanning electron microscopy (SEM), energy dispersive spectrometry (EDS), X-ray diffraction (XRD) and vibrating sample magnetometry (VSM). Then, an NC-enhanced viscoelastic surfactant solution as a fracturing fluid (NC-VES) was prepared, and its properties, including settlement stability, interactions between NCs and WLMs, proppant-transporting performance and gel-breaking properties, were systematically studied. More importantly, the influences of the NC concentration, shear rate, temperature and pH level on the stability of NC-VES were systematically investigated. The experimental results show that the NC-VES with a suitable content of NCs (0.1 wt.%) shows superior stability at 95 ◦C or at a high shear rate. Meanwhile, the NC-VES has an acceptable wide pH stability range of 6–9. In addition, the NC-VES possesses good sand-carrying performance and gel-breaking properties, while the NCs can be easily separated and recycled by applying a magnetic field. The temperature-resistant, stable and environmentally friendly fracturing fluid opens an opportunity for the future hydraulic fracturing of unconventional reservoirs.

**Keywords:** nanocomposite; worm-like micelle; viscoelastic surfactant; fracturing fluid; stability

#### **1. Introduction**

In recent years, unconventional oil and gas resources, especially shale gas and tight oil, have become an important part of the world energy landscape and a substitute resource for conventional oil and gas [1–5]. Hydraulic fracturing technology is one of the key technologies for the efficient development of unconventional reservoirs [6–8]. During the fracturing operation, fracturing fluid (FF) is injected into the formation through the wellbore using a high-pressure pump, forming high well pressure, and resulting in the fracture of the reservoir rock, which opens high-conductivity fracture channels for hydrocarbons migrating from the reservoir to the wellbore [9,10]. Another function of the FF is to carry proppant (such as quartz sand) into these channels to maintain their opening [11,12]. Therefore, FF is a key component of the fracturing operation, and its performance directly affects the success of the fracturing operation. Traditional FFs are mainly water-based fracturing fluids based on polymer thickening compounds such as polyacrylamide (PAM), guar gum and their derivatives [13–15]. However, these polymer thickeners cannot be completely broken and degraded, and insoluble residues are preserved in the formation [16]. The residues block rock pores, reducing reservoir porosity and permeability, resulting in extremely

**Citation:** Si, X.; Luo, M.; Li, M.; Ma, Y.; Huang, Y.; Pu, J. Experimental Study on the Stability of a Novel Nanocomposite-Enhanced Viscoelastic Surfactant Solution as a Fracturing Fluid under Unconventional Reservoir Stimulation. *Nanomaterials* **2022**, *12*, 812. https://doi.org/10.3390/ nano12050812

Academic Editor: Sam Lofland

Received: 26 January 2022 Accepted: 25 February 2022 Published: 28 February 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

unfavourable conditions for subsequent exploitation of oil and gas resources [17,18]. To effectively reduce damage to reservoirs, a viscoelastic surfactant (VES) fracturing fluid called clean fracturing fluid was developed [19].

The structure of VES is a worm-like micelle (WLM) network structure with hydrophilic groups facing outward and hydrophobic groups facing inward, which is self-assembled by small molecular surfactants under the action of some organic or inorganic salts (such as sodium salicylate, potassium chloride and sodium chloride) [20]. The viscoelastic surfactant solution as a fracturing fluid (VESFF) shows excellent viscoelasticity and has the advantages of good proppant-transporting performance, a strong drag reduction effect and low impact on the environment [21]. Under the action of oil, gas or formation fluids, the WLM network structure of VESFF can completely break without damaging the formation, which is ideal compared to traditional FFs [22]. However, in harsh in situ conditions of unconventional reservoirs such as high temperature and shear, the WLM network structure of VES is easily destroyed, which greatly weakens the performance of the clean FF [23,24]. Encouragingly, the rapid development and application of nanoparticles (NPs) provide new ideas to meet these challenges. Some NPs have been used to improve the structural stability of VESFF to withstand the complex environment during fracturing. Nettesheim et al. [25] found that SiO2 NPs incorporated into a NaNO3/cetyltrimethylammonium bromide (CTAB) WLM system could promote the entanglement of micelles to effectively improve the viscosity of the fracturing fluid. Similarly, Zhang et al. [26] proposed a SiO2 NP-enhanced CTAB/NaSal (sodium salicylate) WLM structure, while SiO2 NPs act as junctions of micelles to produce a cross-linked micellar system. Philippova et al. [27] further found that NP-WLM with a pseudo-crosslinking structure could increase the viscosity of the fracturing fluid by three orders of magnitude. García et al. [28] studied the influence of Al2O3 NPs on the rheological behaviour of nanoparticle-enhanced VES and found that the Al2O3 NPs increased the viscosity but not the elastic properties of the VES. In addition, some pyroelectric NPs (such as BaTiO3 and ZnO NPs) were investigated to improve the rheological properties and temperature resistance of VESFF, and they show good proppant-transporting performance at high temperature [29,30]. However, the abovementioned NPs do not have external response characteristics and are easily detained in the formation pores, which will damage the formation. Moreover, these NPs mixed with oil and water are difficult to separate and reuse, which not only pollutes the environment but also increases the application cost and wastes nanomaterials [31–33]. Fortunately, magnetic NPs can respond to an external magnetic field, which is easily separated from oil and water for reuse [34]. Magnetic NPs (such as Fe3O4 NPs) have been reported in fracture and oil/water monitoring [35,36], fracturing drag reduction [37] and enhanced oil recovery [38]. What is more, the rheological properties of magnetic NPs /WLM systems can be improved with a magnetic field [39,40]. Therefore, a nanomaterial with both magnetic and pyroelectric effects may be more suitable to improve the stability, proppant-transporting capacity, environmental protection and reusability of VESFF with NPs in complex conditions for hydraulically fracturing unconventional reservoirs.

Accordingly, this study aims to prepare a novel Fe3O4@ZnO nanocomposite-enhanced viscoelastic surfactant solution as a fracturing fluid (NC-VES) using the interactions between nanoparticles and micelles, and gain further insights into the influence and mechanism of NC concentration, shear rate, temperature and pH on the stability of the NC-VES. First, Fe3O4@ZnO nanocomposites were synthesized and characterized by Fourier transform infrared spectrometry (FT-IR), scanning electron microscopy (SEM), energy dispersive spectrometry (EDS) and vibrating sample magnetometry (VSM). Then, the NC-VES was prepared, and its properties, including settlement stability, interactions between NCs and WLMs, proppant-transporting performance and gel-breaking property, were systematically studied.

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

#### *2.1. Materials*

Ferric chloride (FeCl3) and ferrous sulfate heptahydrate (FeSO4·7H2O) with purities of ≥99.5 wt.% were purchased from Shanghai Aladdin Biochemical Technology Co., Ltd. (Shanghai, China). Acetoxyzinc dihydrate (Zn(Ac)2·2H2O), triethanolamine (TEOA) and sodium hydroxide (NaOH) with purities of ≥99.5 wt.% were manufactured by Shanghai Macklin Biochemical Co., Ltd. (Shanghai, China). Octadecyl trimethyl ammonium chloride (OTAC), ethyl alcohol and sodium salicylate (NaSal) with purities of ≥99 wt.% were provided by Sinopharm Chemical Reagent Co., Ltd. (Shanghai, China). The Krafft points of OTAC surfactant before and after addition of NaSal are 40 ◦C and 25 ◦C, respectively. Polyacrylamide (PAM) and guar gum were provided by PetroChina Changqing Downhole Technical Operation Company (Xi'an, China).

#### *2.2. Synthesis of Fe3O4@ZnO NCs*

First, Fe3O4 NPs were prepared by a chemical coprecipitation method, where 1.62 g FeCl3 and 2.08 g FeSO4·7H2O were dissolved in 150 mL deionized water. Then, 0.1 mol/L NaOH solution was slowly added to the mixed solution until the pH reached 10 and stirred at constant speed mixer at 300 r/min for 3 h. The Fe3O4 precipitate was then washed with ethanol and deionized water and dried at 60 ◦C for 10 h.

Second, ZnO NPs were synthesized by a hydrothermal method where 1.09 g Zn(Ac)2·2H2O was dissolved in 100 mL deionized water. A total of 50 mL of 0.2 mol/L NaOH solution was mixed in zinc acetate solution. Then, the mixed solution was placed in a high-pressure reactor at 160 ◦C with a mixing rate of 300 r/min for 8 h. The ZnO precipitate was subsequently washed with ethanol and deionized water and dried at 60 ◦C for 10 h.

Finally, 0.25 g Fe3O4 NPs were dispersed in 50 mL deionized water with ultrasonic vibration for 15 min. Then, 20 mL TEOA solution (1.6 mol/L) and 30 mL zinc acetate solution (0.02 mol/L) were added into the previous solution at 90 ◦C with stirring at 300 r/min for 10 h. Thus, the composite nanoparticles were obtained after repeated purification more than 5 times through centrifugation (Centrifuge ST16, Thermo Fisher Scientific Inc., Osterode, Lower Saxony, Germany), washing, drying and magnetic adsorption separation.

#### *2.3. Preparation of the NC-VES*

First, 2.87 mmol OTAC was dissolved in 80 mL deionized water at 40 ◦C. Fe3O4@ZnO NCs were dispersed in the OTAC solution then ultrasonically vibrated for 15 min. Then, 20 mL NaSal solution (0.156 mol/L) was added into the previous suspension solution at a mixing rate of 180 r/min for 10 min. NC-VES was obtained after standing for 12 h at room temperature. Meanwhile, VES solution with Fe3O4 nanoparticles and without nanoparticles were prepared with the same concentration (OTAC: 0.287 mmol/L; NaSal: 0.312 mmol/L) using the same methods.

#### *2.4. Characterization of Nanoparticles*

The main functional groups of the samples were examined by FT-IR with a Nicolet 6700 FT-IR instrument (Thermo Fisher Scientific Inc., Waltham, MA, USA). The surface morphology of the nanoparticles was observed by scanning electron microscopy (JSM-5800, JEOL Ltd., Toyoshima, Tokyo, Japan). The elemental content of the nanoparticles was examined with an energy dispersive spectrometer (JEOL Ltd., Toyoshima, Tokyo, Japan). The size of the synthesized nanoparticles was measured by a Malvern laser particle size analyzer (ZS90, Malvern Panalytical, Malvern, UK). The crystal structure of the synthesized samples was measured by X-ray diffractometer (XRD, PANalytical B.V., Almelo, The Netherlands). The magnetizing ability of the samples was evaluated through a vibrating sample magnetometer (Lake Shore 7400, Lake Shore Cryotronics Inc., Westerville, OH, USA).

#### *2.5. Property Tests of the NC-VES* some or all of the NCs in the system will settle out due to the density difference between

*2.5. Property Tests of the NC-VES*

If the network structure of the NC-VES system is unstable or seriously damaged, some or all of the NCs in the system will settle out due to the density difference between the NCs and liquid phase. In this test, the settlement rate of NCs was used to characterize the stability of the NC-VES system. The slower the settlement rate is, the better the stability of the NC-VES. The settlement rates (Equation (1)) were measured by the weighing method with a precision balance (the accuracy was 0.0001 g; Shanghai Fangrui Instrument CO., LTD., Shanghai, China).  Additionally, the effect of various factors on the settlement stability of NC-VES were studied by changing the nanoparticle concentration, shear rate, temperature and pH. the NCs and liquid phase. In this test, the settlement rate of NCs was used to characterize the stability of the NC-VES system. The slower the settlement rate is, the better the stability of the NC-VES. The settlement rates (Equation (1)) were measured by the weighing method with a precision balance (the accuracy was 0.0001 g; Shanghai Fangrui Instrument CO., LTD., Shanghai, China).  Additionally, the effect of various factors on the settlement stability of NC-VES were studied by changing the nanoparticle concentration, shear rate, temperature and pH. *WW*

If the network structure of the NC-VES system is unstable or seriously damaged,

$$set!element\ rate = \frac{W\_2 - W\_1}{W\_0} \times 100\% \tag{1}$$

(1)

2

−

where *W*0 is the total weight of the NCs added to the NC-VES, g; *W*1 is the weight of the empty sample bottle, g; and is the total weight of the sample bottle and settled NCs after being dried, g. where *W*0 is the total weight of the NCs added to the NC-VES, g; *W*1 is the weight of the empty sample bottle, g; and is the total weight of the sample bottle and settled NCs after being dried, g.

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**Figure 1.** Test of the settlement stability (**a**) and proppant-transporting performance (**b**). **Figure 1.** Test of the settlement stability (**a**) and proppant-transporting performance (**b**).

Proppant-transporting performance is one of the key properties of fracturing fluid. In the test, the settling velocity of proppant (quartz sand) in the NC-VES is used to characterize the proppant-transporting performance. As shown in Figure 1b, the settling velocity was calculated by the Formula *v* = h/t, where h and t were recorded as the vertical distance (cm) and time (s) taken for quartz sand as proppant to settle from the surface to Proppant-transporting performance is one of the key properties of fracturing fluid. In the test, the settling velocity of proppant (quartz sand) in the NC-VES is used to characterize the proppant-transporting performance. As shown in Figure 1b, the settling velocity was calculated by the Formula *v* = *h*/*t*, where *h* and *t* were recorded as the vertical distance (cm) and time (s) taken for quartz sand as proppant to settle from the surface to the bottom of 40 mL NC-VES.

#### the bottom of 40 mL NC-VES. **3. Results and Discussion**

#### **3. Results and Discussion** *3.1. Characterization of Fe3O4@ZnO Nanocomposites*

*3.1. Characterization of Fe3O4@ZnO Nanocomposites* Figure 2 shows the comparison of the chemical structures of Fe3O4, ZnO and Fe3O4@ZnO based on the FT-IR spectra obtained over the wavenumber range of 400–4000 cm−1. The stretching vibration peak at approximately 3410 cm−1 in all the IR spectra is attributed to hydroxyl groups (O–H) due to the adsorbed water molecules on the surface of the samples. The peak at approximately 570 cm−1 represents Fe-O bonds in spectra Figure 2a,b. The stretching peak at approximately 450 cm−1 is related to Zn-O bonds in spectra Figure 2 shows the comparison of the chemical structures of Fe3O4, ZnO and Fe3O4@ZnO based on the FT-IR spectra obtained over the wavenumber range of 400–4000 cm−1 . The stretching vibration peak at approximately 3410 cm−1 in all the IR spectra is attributed to hydroxyl groups (O–H) due to the adsorbed water molecules on the surface of the samples. The peak at approximately 570 cm−1 represents Fe-O bonds in spectra Figure 2a,b. The stretching peak at approximately 450 cm−1 is related to Zn-O bonds in spectra Figure 2a,c. Both the Fe-O bond and Zn-O bond existing in the synthesized Fe3O4@ZnO composite indicate that ZnO was adsorbed on the surface of the Fe3O4 particles.

Figure 2a,c. Both the Fe-O bond and Zn-O bond existing in the synthesized Fe3O4@ZnO

composite indicate that ZnO was adsorbed on the surface of the Fe3O4 particles.

60

**Figure 2.** FT-IR spectra of the samples, (**a**) Fe3O4@ZnO, **(b**) Fe3O4, (**c**) ZnO. **Figure 2.** FT-IR spectra of the samples, (**a**) Fe3O4@ZnO, **(b**) Fe3O4 , (**c**) ZnO.

The SEM images with EDS analysis of the Fe3O4 (a), ZnO (b) and Fe3O4@ZnO samples (c) are depicted in Figure 3. As seen from the SEM images, the average sizes of the synthesized Fe3O4 NPs, ZnO NPs and Fe3O4@ZnO NCs are approximately 25 nm, 10 nm and 60 nm, respectively. As shown in Figure 4, the medium diameters of the three nanoparticles with a narrow size distribution, were consistent with the observation results of the SEM images. The pure Fe3O4 NPs appear to exhibit an octahedral structure (Figure 3a), while the ZnO NPs display a spherical-like structure with smaller dimensions (Figure 3b). Figure 3c clearly shows that ZnO NPs are successfully adsorbed on the surface of Fe3O4 NPs. Additionally, the comparison results of the EDS analysis further showed that the composite nanoparticles were successfully synthesized. Meanwhile, the mass ratio of The SEM images with EDS analysis of the Fe3O4 (a), ZnO (b) and Fe3O4@ZnO samples (c) are depicted in Figure 3. As seen from the SEM images, the average sizes of the synthesized Fe3O4 NPs, ZnO NPs and Fe3O4@ZnO NCs are approximately 25 nm, 10 nm and 60 nm, respectively. As shown in Figure 4, the medium diameters of the three nanoparticles with a narrow size distribution, were consistent with the observation results of the SEM images. The pure Fe3O4 NPs appear to exhibit an octahedral structure (Figure 3a), while the ZnO NPs display a spherical-like structure with smaller dimensions (Figure 3b). Figure 3c clearly shows that ZnO NPs are successfully adsorbed on the surface of Fe3O4 NPs. Additionally, the comparison results of the EDS analysis further showed that the composite nanoparticles were successfully synthesized. Meanwhile, the mass ratio of Fe3O4 and ZnO (5.14:1) in the Fe3O4@ZnO NCs was obtained from the EDS analysis. Besides, in all EDS diagrams, the oxygen content higher than the theoretical value was mainly attributed to the adsorption of water molecules.

Fe3O4 and ZnO (5.14:1) in the Fe3O4@ZnO NCs was obtained from the EDS analysis. Besides, in all EDS diagrams, the oxygen content higher than the theoretical value was mainly attributed to the adsorption of water molecules. XRD analysis of synthesized Fe3O4, ZnO and Fe3O4@ZnO nanoparticles and Rietveld refinement of the XRD data of the Fe3O4@ZnO composite were shown in Figure 5. As shown in Figure 5a, the characteristic peaks and Bragg lattice planes reported at 30.08◦ (220), 35.42◦ (311), 43.08◦ (400), 53.56◦ (422), 56.98◦ (511) and 62.63◦ (440) are related to Fe3O4 structure, while the characteristic peaks and Bragg lattice planes reported at 31.72◦ (100), 34.44◦ (002), 36.21◦ (101), 47.49◦ (102), 56.51◦ (110), 62.81◦ (103) and 67.86◦ (112) are related to ZnO structure, which are in accordance with the patterns of standard pure Fe3O4 and ZnO, respectively [41,42]. Almost no obvious impurity peak was observed, which confirmed the high purity for the synthesized samples. As shown in Figure 5b, Rietveld refinement with good fit (goodness of fit *x* 2 = 1.19) was carried out to obtain more crystal phase information of the composite. Parameters for Fe3O4 (crystal structure: cubic; lattice constant: a = 8.372 Å; space group: Fd3m) and parameters for ZnO (lattice constant: a = 3.249 Å, c = 5.208 Å; crystal structure: hexagonal; space group: P63mc) were obtained. Meanwhile, the mass ratio of Fe3O4 and ZnO (4.91:1) in the Fe3O4@ZnO NCs from the Rietveld refinement is consistent with the EDS test within the range of allowable error. The results revealed that Fe3O4@ZnO NCs with a high purity were successfully synthesized.

*Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 6 of 15

**Figure 3.** SEM (**left**) and EDS analysis (**right**) of Fe3O4 (**a**), ZnO (**b**) and Fe3O4@ZnO (**c**).**Figure 3.** SEM (**left**) and EDS analysis (**right**) of Fe3O4 (**a**), ZnO (**b**) and Fe3O4@ZnO (**c**). **Figure 3.** SEM (**left**) and EDS analysis (**right**) of Fe3O4 (**a**), ZnO (**b**) and Fe3O4@ZnO (**c**).

**Figure 4.** Particle size distribution of Fe3O4, ZnO and Fe3O4@ZnO nanoparticles. **Figure 4.** Particle size distribution of Fe3O4, ZnO and Fe3O4@ZnO nanoparticles. **Figure 4.** Particle size distribution of Fe3O4 , ZnO and Fe3O4@ZnO nanoparticles.

XRD analysis of synthesized Fe3O4, ZnO and Fe3O4@ZnO nanoparticles and Rietveld refinement of the XRD data of the Fe3O4@ZnO composite were shown in Figure 5. As

(220), 35.42° (311), 43.08° (400), 53.56° (422), 56.98° (511) and 62.63° (440) are related to Fe3O4 structure, while the characteristic peaks and Bragg lattice planes reported at 31.72° (100), 34.44° (002), 36.21° (101), 47.49° (102), 56.51° (110), 62.81° (103) and 67.86° (112) are

XRD analysis of synthesized Fe3O4, ZnO and Fe3O4@ZnO nanoparticles and Rietveld

Fe3O4 structure, while the characteristic peaks and Bragg lattice planes reported at 31.72° (100), 34.44° (002), 36.21° (101), 47.49° (102), 56.51° (110), 62.81° (103) and 67.86° (112) are sized.

refinement with good fit (goodness of fit *x*

**Figure 5.** XRD patterns for synthesized nanoparticles (**a**) and Rietveld refinement of the XRD data of the Fe3O4@ZnO composite (**b**). **Figure 5.** XRD patterns for synthesized nanoparticles (**a**) and Rietveld refinement of the XRD data of the Fe3O4@ZnO composite (**b**).

related to ZnO structure, which are in accordance with the patterns of standard pure Fe3O4 and ZnO, respectively [41,42]. Almost no obvious impurity peak was observed, which confirmed the high purity for the synthesized samples. As shown in Figure 5b, Rietveld

phase information of the composite. Parameters for Fe3O4 (crystal structure: cubic; lattice constant: a = 8.372 Å ; space group: Fd3̅m) and parameters for ZnO (lattice constant: a = 3.249 Å , c = 5.208 Å ; crystal structure: hexagonal; space group: P63mc) were obtained. Meanwhile, the mass ratio of Fe3O4 and ZnO (4.91:1) in the Fe3O4@ZnO NCs from the Rietveld refinement is consistent with the EDS test within the range of allowable error. The results revealed that Fe3O4@ZnO NCs with a high purity were successfully synthe-

2= 1.19) was carried out to obtain more crystal

Figure 6 shows the magnetization of the Fe3O4 and Fe3O4@ZnO samples. It is clearly observed that the magnetization curves for the measured samples are S-shaped. When the magnetic field is zero, the remanence and coercive force are close to zero, indicating that the synthesized Fe3O4 NPs and Fe3O4@ZnO NCs possess good soft magnetic and practically superparamagnetic properties. The saturation magnetization measured for pure Fe3O4 NPs is 76.03 emu/g, while the saturation magnetization of Fe3O4@ZnO NCs decreases slightly but still maintains a high value of 62.26 emu/g. The high magnetization ensures that the Fe3O4@ZnO NCs have strong magnetic response ability and are easily separated from the liquid phase by magnetic field during quantitatively measuring the settlement rate of NCs (seen in Section 2.5). On the other hand, Fe3O4@ZnO NCs is simply recycled from the dispersed solution by using a magnet (seen in Section 3.3). What is more, the mass ratio of Fe3O4 and ZnO calculated from the comparative saturation magnetization is 5.02:1, consistent with the above tests (EDS analysis and XRD analysis), which further proves the purity of the synthetic Fe3O4@ZnO NCs. Figure 6 shows the magnetization of the Fe3O4 and Fe3O4@ZnO samples. It is clearly observed that the magnetization curves for the measured samples are S-shaped. When the magnetic field is zero, the remanence and coercive force are close to zero, indicating that the synthesized Fe3O4 NPs and Fe3O4@ZnO NCs possess good soft magnetic and practically superparamagnetic properties. The saturation magnetization measured for pure Fe3O4 NPs is 76.03 emu/g, while the saturation magnetization of Fe3O4@ZnO NCs decreases slightly but still maintains a high value of 62.26 emu/g. The high magnetization ensures that the Fe3O4@ZnO NCs have strong magnetic response ability and are easily separated from the liquid phase by magnetic field during quantitatively measuring the settlement rate of NCs (seen in Section 2.5). On the other hand, Fe3O4@ZnO NCs is simply recycled from the dispersed solution by using a magnet (seen in Section 3.3). What is more, the mass ratio of Fe3O4 and ZnO calculated from the comparative saturation magnetization is 5.02:1, consistent with the above tests (EDS analysis and XRD analysis), which further proves the purity of the synthetic Fe3O4@ZnO NCs. *Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 8 of 15

**Figure 6.** Magnetization of the Fe3O4 NPs and Fe3O4@ZnO NCs. **Figure 6.** Magnetization of the Fe3O4 NPs and Fe3O4@ZnO NCs.

*3.2. Stability of the NC-VES*

the viscosity of NC-VES decreases sharply at high shear rates.

focuses on the NC-VES system with a 0.1 wt.% NC concentration.

was studied with shear rates ranging from 0.1 to 1000 s−1 at 25 °C (tested with an Anton Paar rheometer, Physica MCR 302, Anton Paar GmbH, Graz, Austria). Figure 7 shows that the viscosity remains unchanged at low shear rates, while a notably reduced slope is observed under high shear rates, and the shear-thinning phenomenon occurs in all samples. The dependence of viscosity on shear rate is usually explained by shear banding behaviour, which has been reported in previous research results [43–46]. It is found that the NCs can improve the viscosity of the VES system. When the concentration of the NCs is low (0.01 wt.%), the viscosity increases but is not clear compared with the VES without NCs. When the NC concentration reaches 0.1 wt.%, the viscosity at 170 s−1 of the NC-VES is higher than VES without NCs. However, when the NC concentration is higher (0.3 wt.%),

Dynamic modulus (storage modulus G′ and loss modulus G″) as a function of frequency is shown in Figure 8. Within the measured test frequency range, the storage modulus remains almost unchanged, while the loss modulus increases gradually. For NC-VES systems, the storage modulus and loss modulus are greater than those of the VES system, suggesting that the WLM network in VES is strengthened by NCs. Here, the nanoparticles are incorporated into worm-like micellar systems to form nanoparticle-micelle junctions as the connection point of WLMs, which improves the structural stability and viscoelasticity of the WLM systems [47–49]. Combined with Figures 7 and 8, this indicates that the NC concentration (0.1 wt.%) enables the NC-VES system to possess sufficient stability and high viscosity, maintaining good proppant-transporting performance during fracturing operations. Therefore, subsequent research on the influencing factors of fracturing fluid

3.2.1. Effects of Shear Rate and NC Concentration

#### *3.2. Stability of the NC-VES*

3.2.1. Effects of Shear Rate and NC Concentration

The shear rate-dependent viscosities of the VESFF for different NC concentrations was studied with shear rates ranging from 0.1 to 1000 s−1 at 25 ◦C (tested with an Anton Paar rheometer, Physica MCR 302, Anton Paar GmbH, Graz, Austria). Figure 7 shows that the viscosity remains unchanged at low shear rates, while a notably reduced slope is observed under high shear rates, and the shear-thinning phenomenon occurs in all samples. The dependence of viscosity on shear rate is usually explained by shear banding behaviour, which has been reported in previous research results [43–46]. It is found that the NCs can improve the viscosity of the VES system. When the concentration of the NCs is low (0.01 wt.%), the viscosity increases but is not clear compared with the VES without NCs. When the NC concentration reaches 0.1 wt.%, the viscosity at 170 s−1 of the NC-VES is higher than VES without NCs. However, when the NC concentration is higher (0.3 wt.%), the viscosity of NC-VES decreases sharply at high shear rates. *Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 9 of 15

**Figure 7.** Viscosity as a function of shear rate at different NC concentrations.**Figure 7.** Viscosity as a function of shear rate at different NC concentrations.

10 Modulus (Pa) G'(0% NCs) G〞(0% NCs) G'(0.01wt.% NCs) G〞(0.01wt.% NCs) G'(0.1wt.% NCs) G〞(0.1wt.% NCs) G'(0.3wt.% NCs) G〞(0.3wt.% NCs) Dynamic modulus (storage modulus G0 and loss modulus G") as a function of frequency is shown in Figure 8. Within the measured test frequency range, the storage modulus remains almost unchanged, while the loss modulus increases gradually. For NC-VES systems, the storage modulus and loss modulus are greater than those of the VES system, suggesting that the WLM network in VES is strengthened by NCs. Here, the nanoparticles are incorporated into worm-like micellar systems to form nanoparticle-micelle junctions as the connection point of WLMs, which improves the structural stability and viscoelasticity of the WLM systems [47–49]. Combined with Figures 7 and 8, this indicates that the NC concentration (0.1 wt.%) enables the NC-VES system to possess sufficient stability and high viscosity, maintaining good proppant-transporting performance during fracturing operations. Therefore, subsequent research on the influencing factors of fracturing fluid focuses on the NC-VES system with a 0.1 wt.% NC concentration.

#### 3.2.2. Effect of Temperature

3.2.2. Effect of Temperature

1 10 1 Frequency (rad/s) All samples were heated to the corresponding experimental temperature by a water bath and held for 10 h, then the settlement rate of NCs in each sample was obtained by weighing and calculation, as shown in Figure 9. The results show that when the temperature increased from 25 to 65 ◦C, the settlement rate of NCs increased from 0 to 0.1%, while almost no sedimentation phenomenon occurred, indicating that the NC-WLM network structure of the system was very stable in the low to medium temperature range. As the

**Figure 8.** Variations of storage modulus (G′) and loss modulus (G″) with frequency at 25 °C.

All samples were heated to the corresponding experimental temperature by a water bath and held for 10 h, then the settlement rate of NCs in each sample was obtained by weighing and calculation, as shown in Figure 9. The results show that when the temperature increased from 25 to 65 °C, the settlement rate of NCs increased from 0 to 0.1%, while almost no sedimentation phenomenon occurred, indicating that the NC-WLM network structure of the system was very stable in the low to medium temperature range. As the temperature continued to rise, the sedimentation rate increased slightly, but even at a high temperature of 95 °C, the settlement rate was only about 2%, indicating that the temperature resistance of the system was very good. Conventional VESFF is a WLM network structure formed by the spontaneous aggregation and self-assembly of high-concentration surfactant molecules, which has a specific viscoelasticity and can meet the needs of fracturing in medium- and low–temperature reservoirs. However, the WLM structure in traditional VESFF is easily broken at high temperatures, because the interaction between molecules such as van der Waals forces and hydrogen bonding forces weakens and the

temperature continued to rise, the sedimentation rate increased slightly, but even at a high temperature of 95 ◦C, the settlement rate was only about 2%, indicating that the temperature resistance of the system was very good. Conventional VESFF is a WLM network structure formed by the spontaneous aggregation and self-assembly of high-concentration surfactant molecules, which has a specific viscoelasticity and can meet the needs of fracturing in medium- and low–temperature reservoirs. However, the WLM structure in traditional VESFF is easily broken at high temperatures, because the interaction between molecules such as van der Waals forces and hydrogen bonding forces weakens and the stability of aggregates decreases [23], leading to a sharp decrease in the viscosity of the system, which severely weakens the performance of the fracturing fluid. 0.1 1 10 100 1000 100 Viscosity (mPa·s)Shear rate (s-1) **Figure 7.** Viscosity as a function of shear rate at different NC concentrations. *Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 10 of 15 stability of aggregates decreases [23], leading to a sharp decrease in the viscosity of the

VES

 NC-VES with 0.01wt.% NCs NC-VES with 0.1wt.% NCs NC-VES with 0.3wt.% NCs

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1000

10,000

**Figure 8.** Variations of storage modulus (G′) and loss modulus (G″) with frequency at 25 °C. **Figure 8.** Variations of storage modulus (G0 ) and loss modulus (G00) with frequency at 25 ◦C. ity of the NC-WLM system at high temperature.

**Figure 9.** Settlement rate as a function of temperature, in which the insertions from left to right refer to the stable state of the NC-WLM system at 25 °C, 65 °C and 95 °C respectively.(OTAC: 0.287 mmol/L; NaSal: 0.312 mmol/L).**Figure 9.** Settlement rate as a function of temperature, in which the insertions from left to right refer to the stable state of the NC-WLM system at 25 ◦C, 65 ◦C and 95 ◦C respectively.(OTAC: 0.287 mmol/L; NaSal: 0.312 mmol/L).

NC-VES (0.1wt.% NCs)

VES with 0.1wt.% Fe3O4 NPs

20

30

40

50

60

Temperature (°C)

70

80

90

100

VES

Temperature

0 10 20 30 40 50 60

**Figure 10.** Viscosity as a function of time at high temperature (heated to 95 °C).

Time (min)

20

40

60

80

100

Viscosity (mPa·s)

120

140

160

180

The viscosities as a function of time at 95 ◦C with a constant heating rate of 3 ◦C/min and a shear rate of 170 s−1 are displayed in Figure 10. When the temperature is lower than 60 ◦C, a slight increase in viscosity was observed in all three curves. However, as the temperature continues to increase to 95 ◦C, the viscosity of the samples decreases sharply. The viscosity of VES with Fe3O4 NPs is lower than that without NPs at 95 ◦C. The reason is that a certain amount of surfactant molecules are adsorbed on the surface of Fe3O4 NPs, resulting in a decrease in the surfactant concentration involved in WLM structure in the liquid phase [50]. Compared with the other two systems, the high-temperature resistance of the NC-VES system is mainly attributed to the ZnO NPs on the surface of the NCs. As shown in Figure 11, the ZnO nanoparticles adsorbed on the surface of NCs possess a pyroelectric effect, releasing charges with increasing temperature. The charged NCs easily adsorb micelles and play a role in the junction of the WLM network structure. The electrostatic screening of charged WLMs promotes further growth of wormlike micelles, and maintains good stability of the WLM network structure [29]. Therefore, the NCs incorporated into the WLM network act as a skeleton-like structure, greatly improving the stability of the NC-WLM system at high temperature. 30 40 50 60 70 80 90 100 0.0 0.5 1.0 1.5 Settlement rate (%)Temperature (℃) **Figure 9.** Settlement rate as a function of temperature, in which the insertions from left to right refer to the stable state of the NC-WLM system at 25 °C, 65 °C and 95 °C respectively.(OTAC: 0.287 mmol/L; NaSal: 0.312 mmol/L).

*Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 10 of 15

ity of the NC-WLM system at high temperature.

2.0

2.5

3.0

system, which severely weakens the performance of the fracturing fluid.

stability of aggregates decreases [23], leading to a sharp decrease in the viscosity of the

and a shear rate of 170 s−1 are displayed in Figure 10. When the temperature is lower than 60 °C, a slight increase in viscosity was observed in all three curves. However, as the temperature continues to increase to 95 °C, the viscosity of the samples decreases sharply. The viscosity of VES with Fe3O4 NPs is lower than that without NPs at 95 °C. The reason is that a certain amount of surfactant molecules are adsorbed on the surface of Fe3O4 NPs, resulting in a decrease in the surfactant concentration involved in WLM structure in the liquid phase [50]. Compared with the other two systems, the high-temperature resistance of the NC-VES system is mainly attributed to the ZnO NPs on the surface of the NCs. As shown in Figure 11, the ZnO nanoparticles adsorbed on the surface of NCs possess a pyroelectric effect, releasing charges with increasing temperature. The charged NCs easily adsorb micelles and play a role in the junction of the WLM network structure. The electrostatic screening of charged WLMs promotes further growth of wormlike micelles, and maintains good stability of the WLM network structure [29]. Therefore, the NCs incorporated into the WLM network act as a skeleton-like structure, greatly improving the stabil-

The viscosities as a function of time at 95 °C with a constant heating rate of 3 °C/min

**Figure 10. Figure 10.** Viscosity as a function of time at high temperature (heated to 95 Viscosity as a function of time at high temperature (heated to 95 °C). ◦C).

**Figure 11.** Schematic diagram of the stability of the NC-WLM network structure at high tempera-**Figure 11.** Schematic diagram of the stability of the NC-WLM network structure at high temperatures.

#### tures. 3.2.3. Effect of pH

other hand, excessive H+

sity (about 5.26 g/cm3

of WLMs (**b**).

3.2.3. Effect of pH As shown in Figure 12a, in a highly acidic environment (pH = 1–5), the settlement phenomenon of the system is very clear, while the settlement rate is as high as 94.9%. This shows that the system cannot exist stably in strong acids, which can be attributed to two aspects. On the one hand, in the presence of a strong acid, some NCs react with H+ (Fe3O4 As shown in Figure 12a, in a highly acidic environment (pH = 1–5), the settlement phenomenon of the system is very clear, while the settlement rate is as high as 94.9%. This shows that the system cannot exist stably in strong acids, which can be attributed to two aspects. On the one hand, in the presence of a strong acid, some NCs react with H+ (Fe3O4 + 8H+ = Fe2+ + 2Fe3+ + 4H2O, ZnO + 2H+ = Zn2+ + H2O) and are dissolved in the

+ 8H+ = Fe2+ + 2Fe3+ + 4H2O, ZnO + 2H+ = Zn2+ + H2O) and are dissolved in the solution. The solution will then exhibit the crimson colour of the Fe2+ and Fe3+ aqueous solutions. On the

large settlement of NCs [51,52]. From the results, when pH = 1, the settlement rate decreases, which does not mean that the settlement stability of the NC-VES becomes better, but that more NCs have been dissolved in the solution. Therefore, it can be considered that the stronger the acidity, the worse the settlement stability of the system. When the pH is weakly acidic–neutral–weakly alkaline (pH = 6–9), the settlement rate of the system is very small. Meanwhile, Figure 12b shows that it is difficult for NCs to form stable suspension in water without WLMs even under neutral conditions, because the particle den-

environment (pH = 10–14), the settlement rate of NCs increases rapidly with increasing pH. When pH ≥ 13, a large number of NCs settle out. On the one hand, ZnO is dissolved

OH− destroys the self-assembly mechanism of surfactant molecules, leading to micelles unsuccessful for connecting with each other and assembling into a WLM network structure [53]. Therefore, the pH environment has a strong impact on the stability of the NC-VES, with the system being able to maintain a good stability in a wide range of pH = 6–9.

**Figure 12.** Settlement rate as a function of pH (**a**) and behaviour of NCs in deionized water without

) is much greater than that of the liquid. In the strongly alkaline

). On the other hand, an excessive

2−

by alkaline solution (ZnO + 2OH− + H2O = [Zn(OH)4]

of WLMs (**b**).

tures.

3.2.3. Effect of pH

other hand, excessive H+

*Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 11 of 15

solution. The solution will then exhibit the crimson colour of the Fe2+ and Fe3+ aqueous solutions. On the other hand, excessive H+ leads to an increase in the repulsive force between the head groups of the OTAC molecule and destroys the WLM network structure, which results in large settlement of NCs [51,52]. From the results, when pH = 1, the settlement rate decreases, which does not mean that the settlement stability of the NC-VES becomes better, but that more NCs have been dissolved in the solution. Therefore, it can be considered that the stronger the acidity, the worse the settlement stability of the system. When the pH is weakly acidic–neutral–weakly alkaline (pH = 6–9), the settlement rate of the system is very small. Meanwhile, Figure 12b shows that it is difficult for NCs to form stable suspension in water without WLMs even under neutral conditions, because the particle density (about 5.26 g/cm3 ) is much greater than that of the liquid. In the strongly alkaline environment (pH = 10–14), the settlement rate of NCs increases rapidly with increasing pH. When pH ≥ 13, a large number of NCs settle out. On the one hand, ZnO is dissolved by alkaline solution (ZnO + 2OH− + H2O = [Zn(OH)4] 2−). On the other hand, an excessive OH− destroys the self-assembly mechanism of surfactant molecules, leading to micelles unsuccessful for connecting with each other and assembling into a WLM network structure [53]. Therefore, the pH environment has a strong impact on the stability of the NC-VES, with the system being able to maintain a good stability in a wide range of pH = 6–9. large settlement of NCs [51,52]. From the results, when pH = 1, the settlement rate decreases, which does not mean that the settlement stability of the NC-VES becomes better, but that more NCs have been dissolved in the solution. Therefore, it can be considered that the stronger the acidity, the worse the settlement stability of the system. When the pH is weakly acidic–neutral–weakly alkaline (pH = 6–9), the settlement rate of the system is very small. Meanwhile, Figure 12b shows that it is difficult for NCs to form stable suspension in water without WLMs even under neutral conditions, because the particle density (about 5.26 g/cm3 ) is much greater than that of the liquid. In the strongly alkaline environment (pH = 10–14), the settlement rate of NCs increases rapidly with increasing pH. When pH ≥ 13, a large number of NCs settle out. On the one hand, ZnO is dissolved by alkaline solution (ZnO + 2OH− + H2O = [Zn(OH)4] 2− ). On the other hand, an excessive OH− destroys the self-assembly mechanism of surfactant molecules, leading to micelles unsuccessful for connecting with each other and assembling into a WLM network structure [53]. Therefore, the pH environment has a strong impact on the stability of the NC-VES, with the system being able to maintain a good stability in a wide range of pH = 6–9.

leads to an increase in the repulsive force between the head

(Fe3O4

**Figure 11.** Schematic diagram of the stability of the NC-WLM network structure at high tempera-

As shown in Figure 12a, in a highly acidic environment (pH = 1–5), the settlement

phenomenon of the system is very clear, while the settlement rate is as high as 94.9%. This shows that the system cannot exist stably in strong acids, which can be attributed to two

+ 8H+ = Fe2+ + 2Fe3+ + 4H2O, ZnO + 2H+ = Zn2+ + H2O) and are dissolved in the solution. The

groups of the OTAC molecule and destroys the WLM network structure, which results in

aspects. On the one hand, in the presence of a strong acid, some NCs react with H+

**Figure 12.** Settlement rate as a function of pH (**a**) and behaviour of NCs in deionized water without **Figure 12.** Settlement rate as a function of pH (**a**) and behaviour of NCs in deionized water without of WLMs (**b**).

#### *3.3. Proppant-Transporting Performance and Gel-Breaking Property*

In the process of a hydraulic fracturing operation, fracturing fluid is pumped into the formation, cracking the rock formation and forming a fracture channel to improve the oil and gas production efficiency. To prevent fracture closure, proppants (such as quartz sand) are transported by fracturing fluid into the formation to prop fractures, keeping the fracture open for a long time under formation pressure. Therefore, it is very important that the fracturing fluid possess good proppant-transporting performance, which is one of the key factors for the success of fracturing operations. In this work, the proppant-transporting performance of four fracturing fluids was compared, and the results are shown in Table 1. Compared with conventional fracturing fluids (Guar FF and PAM FF), the VES systems (NC-VES and VESFF) show better proppant-transporting performance. The settlement velocity of quartz sand in the NC-VES is only 0.52 × 10−3 cm/s, far less than 0.08 cm/s, which displays superior proppant-transporting performance [54]. In addition, by adding a small amount of kerosene to NC-VES, gel breaking occurs rapidly with no residue, and the NCs can be easily recycled using a magnet to apply a magnetic field, as shown in Figure 13. The results imply that the NC-VES system is a low-damage, environmentally friendly and cost-saving fracturing fluid. mentally friendly and cost-saving fracturing fluid. **Table 1.** Settlement velocity of different FF samples at room temperature.

**Table 1.** Settlement velocity of different FF samples at room temperature. **Samples Constituents (Per 100 mL Water) Viscosity (mPa·s) Settlement Velocity (cm/s)**

*Nanomaterials* **2022**, *11*, x FOR PEER REVIEW 12 of 15

*3.3. Proppant-Transporting Performance and Gel-Breaking Property*

In the process of a hydraulic fracturing operation, fracturing fluid is pumped into the

formation, cracking the rock formation and forming a fracture channel to improve the oil and gas production efficiency. To prevent fracture closure, proppants (such as quartz sand) are transported by fracturing fluid into the formation to prop fractures, keeping the fracture open for a long time under formation pressure. Therefore, it is very important that the fracturing fluid possess good proppant-transporting performance, which is one of the key factors for the success of fracturing operations. In this work, the proppanttransporting performance of four fracturing fluids was compared, and the results are shown in Table 1. Compared with conventional fracturing fluids (Guar FF and PAM FF), the VES systems (NC-VES and VESFF) show better proppant-transporting performance. The settlement velocity of quartz sand in the NC-VES is only 0.52 × 10−3 cm/s, far less than 0.08 cm/s, which displays superior proppant-transporting performance [54]. In addition, by adding a small amount of kerosene to NC-VES, gel breaking occurs rapidly with no residue, and the NCs can be easily recycled using a magnet to apply a magnetic field, as shown in Figure 13. The results imply that the NC-VES system is a low-damage, environ-

**Figure 13.** Gel-breaking test and the separation of NCs by a magnet. **Figure 13.** Gel-breaking test and the separation of NCs by a magnet.

#### **4. Conclusions**

**4. Conclusions** In this study, the influence and mechanism of NC concentration, shear rate, temperature and pH on the settlement stability of the NC-VES fracturing fluid was systematically investigated. First, NCs with good magnetic response ability were synthesized and acted as junctions of micelles to improve the stability of the WLM network structure. The results showed that the NC-VES system with the optimal concentration of 0.1 wt.% possesses good shear and temperature resistance. At high temperatures (such as 95 °C), the ZnO NPs with a pyroelectric effect on the surface of NCs can effectively reduce the decomposition of the WLM network structure to prevent the NCs from settling out of the fracturing fluid. Strong acid and strong alkaline solutions seriously damage the NC-VES network In this study, the influence and mechanism of NC concentration, shear rate, temperature and pH on the settlement stability of the NC-VES fracturing fluid was systematically investigated. First, NCs with good magnetic response ability were synthesized and acted as junctions of micelles to improve the stability of the WLM network structure. The results showed that the NC-VES system with the optimal concentration of 0.1 wt.% possesses good shear and temperature resistance. At high temperatures (such as 95 ◦C), the ZnO NPs with a pyroelectric effect on the surface of NCs can effectively reduce the decomposition of the WLM network structure to prevent the NCs from settling out of the fracturing fluid. Strong acid and strong alkaline solutions seriously damage the NC-VES network structure by dissolving NCs and hindering the self-assembly behaviour of surfactant molecules, resulting in an acceptable pH range of 6~9. In addition, the settling velocity of quartz sand in the system at room temperature was only 0.52 × 10−3 cm/s, indicating that the NC-VES has good proppant-transporting performance. Finally, no residue was found after gel breaking of the NC-VES and the recovery of NCs by magnetic adsorption, implying no damage to the formation and the environment. A new type of temperature-resistant, stable and environmentally friendly fracturing fluid was thus provided for hydraulic fracturing of unconventional reservoirs.

**Author Contributions:** Conceptualization, M.L. (Mingliang Luo) and M.L. (Mingzhong Li); methodology, X.S., M.L. (Mingzhong Li) and M.L. (Mingliang Luo); experimental, X.S., Y.H. and Y.M.; writing, X.S. and M.L. (Mingliang Luo); analytical studies, X.S. and M.L. (Mingliang Luo); visualization, X.S.; funding acquisition, M.L. (Mingliang Luo) and J.P. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (grant numbers 51874334 and 52104057) and the China Postdoctoral Science Foundation (grant number 2021M693506).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Data are contained within the article.

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

#### **References**

## *Article* **Wettability Improvement in Oil–Water Separation by Nano-Pillar ZnO Texturing**

**Xiaoyan Liu 1 , Shaotong Feng 1 , Caihua Wang 1 , Dayun Yan 2,\* , Lei Chen 3 and Bao Wang 1,3,\***

**Abstract:** The nanostructure-based surface texturing can be used to improve the materials wettability. Regarding oil–water separation, designing a surface with special wettability is as an important approach to improve the separation efficiency. Herein, a ZnO nanostructure was prepared by a two-step process for sol–gel process and crystal growth from the liquid phase to achieve both a superhydrophobicity in oil and a superoleophobic property in water. It is found that the filter material with nanostructures presented an excellent wettability. ZnO-coated stainless-steel metal fiber felt had a static underwater oil contact angle of 151.4◦ ± 0.8◦ and an underoil water contact angle of 152.7◦ ± 0.6◦ . Furthermore, to achieve water/oil separation, the emulsified impurities in both water-in-oil and oil-in-water emulsion were effectively intercepted. Our filter materials with a small pore (~5 µm diameter) could separate diverse water-in-oil and oil-in-water emulsions with a high efficiency (>98%). Finally, the efficacy of filtering quantity on separation performance was also investigated. Our preliminary results showed that the filtration flux decreased with the collection of emulsified impurities. However, the filtration flux could restore after cleaning and drying, suggesting the recyclable nature of our method. Our nanostructured filter material is a promising candidate for both water-in-oil and oil-in-water separation in industry.

**Keywords:** nanostructured surface; wettability; metal fiber felt; oil–water separation; emulsion

#### **1. Introduction**

Surface wettability is an important material property, and it is determined by surface structure and surface chemical composition [1,2]. Based on existing theoretical methods, a special material wettability can be achieved by modification of surface texture and surface chemical composition [3,4]. During the last decades, there has been increasing interest in the materials with unique functions due to their promising application in various fields such as self-cleaning, antimicrobial, anti-icing, biocompatible materials, or oil transportation [4–9]. Particularly, depending on the different interfacial energy between organic compounds and water in immiscible oil–water mixtures, the materials with special wettability have gained widespread attention in oil–water separation [10,11]. As a simple preparation method, the meshed membrane owned many advantages, such as high permeability of pore space, good oil–water separation efficacy, and promising corrosion resistance [12,13]. Membrane emulsification method attracted plenty of attention. The factors to affect the demulsification performance have been systematically investigated, which included membrane pore size, membrane thickness, transmembrane pressure, as well as emulsion composition [14–17]. In contrast, traditional methods such as gravity-driven separation, biological treatment, adsorption, and electrochemical techniques have many flaws, such as low separation efficiency, low separation capacity, long time cost, and noticeable secondary pollution [18–20].

**Citation:** Liu, X.; Feng, S.; Wang, C.; Yan, D.; Chen, L.; Wang, B. Wettability Improvement in Oil–Water Separation by Nano-Pillar ZnO Texturing. *Nanomaterials* **2022**, *12*, 740. https://doi.org/10.3390/ nano12050740

Academic Editors: Henrich Frielinghaus and Fabien Grasset

Received: 30 December 2021 Accepted: 17 February 2022 Published: 22 February 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 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 (https:// creativecommons.org/licenses/by/ 4.0/).

Furthermore, the structural parameters of filter membranes may play pivotal roles in the performance of oil–water separation, such as pore size, porosity, and membrane thickness [21–23]. To date, some simulation and experimental studies have provided clues for filter membrane design. Some representative examples were introduced here. You et al. utilized electrospinning techniques to prepare two fibrous membranes with different nanopore diameters. They found that the physical cutting effect dominated the coalescence process when oil droplets crossed a smaller pore on the membrane [24]. Zhu et al. found that the oil droplets continuously separated emulsion via using squeezing coalescence demulsification (SCD) within a narrow pore size. Their hydrogel nanofiber membranes had a sustained separation capacity. More importantly, these results demonstrated that the pore size, the membrane wettability, and the interfacial tension coefficient of oil–water could control the SCD's efficiency [25]. Recently, Wei et al. fabricated highly hydrophilic and underwater oleophobic polytetrafluoroethylene membranes, which showed a high oil–water separation efficiency and achieved a high oil–water separation flux. They found that the emulsion separation may be an interfacial problem and may be related to the pore structure of a filtration membrane [26]. Furthermore, Xi et al. manufactured underwater superoleophobic paper-based materials with excellent wet intensity through green papermaking techniques. They also demonstrated that the water flux could be increased by controlling the average membrane pores' size [27]. These preliminary studies could provide an important basis for the design of filter membranes.

Among these structural parameters, the diameter of the liquid distributor and the hole shape dominated the oil–water separation. To achieve a high demulsification efficiency, the filtering precision should be improved accordingly [28,29]. For the highly precise filtering materials at the submicron scale, a nanostructured membrane is an ideal candidate to improve separation efficacy [30]. In addition, surface energy, a key material characteristic, could only be modified by the coating on that material. Due to the limitation of coating technology, these treatment approaches have been used to promote surface wettability without achieving a satisfied performance in terms of coating combinations [31].

In addition, membrane filtration modified by organic polymer could achieve effective separation of oil–water mixtures. Due to its poor stability in the aqueous medium, however, it is hard to use the common polymer membranes with a special wettability in practice [32]. Therefore, the stable materials such as the inorganic membranes are promising alternatives. To date, all previous studies just focused on the fabrication of hierarchical surface structures without giving adequate attention to the substrate, which determines the material strength and stability.

ZnO nanostructured pillars could be used in complex industrial oil–water mixtures [33–36]. In this study, the stable ZnO nano-pillars were prepared on a high-intensity metal fiber filter by crystal growth, which had a good combination with the substrate. Therefore, the modified filter material possessed stability of mineral coating and strength. Our method promoted the surface wettability and achieved both a superhydrophobicity in oil and a superoleophobic property in water. The filter membrane modified by nano-pillars could be stably used in both water-in-oil and oil-in-water conditions. The filtration flux variation with the collection of emulsified impurities was also investigated. This study provided a novel strategy to achieve stability and strength of a mineral coating simultaneously and might have wide application in separating water-in-oil and oil-in-water emulsions.

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

#### *2.1. Preparation of ZnO Seed Layers*

The stainless-steel fiber felts with the filtration precision of 5 µm were used as a substrate in this study. To eliminate surface contamination, all substrate samples were ultrasonically cleaned with absolute ethanol and ultra-pure water for 10 min. (CH3COO)2Zn (0.75 M (CH3COO)2Zn) (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) and Monoethanolamine ((CH3COO)2Zn and Monoethanolamine in the proportion 1:1) (Aladdin Biochemical Technology Co., Ltd., Shanghai, China) were dissolved in a 2-methoxyethanol

solution (Aladdin Biochemical Technology Co., Ltd., Shanghai, China). This solution was magnetically stirred and heated in a temperature-controlled water bath (70 ◦C, 60 min). The seed solutions were obtained by standing at room temperature for 24 h. Substrates are submerged in the seed solution and lifted repeatedly. Finally, the substrates with the ZnO nanostructured coating were sintered (350 ◦C, 20 min) to stabilize the coating.

#### *2.2. Preparation of ZnO Nano-Pillar Coated Substrates*

As shown in Figure 1, the seed layer coated substrates were first immersed in Zn(NO3)2 (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China)/hexamethylenetetramine (HMTA) (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China)/DAP (Macklin Inc., Shanghai, China) (0.05 M Zn(NO3)2, Zn(NO3)2, and DAP in the proportion 1:6) aqueous solution (10 min) and then was heated in a constant temperature water bath (90 ◦C, 3 h). After the reaction, the substrates were separated from the solution and washed with ultra-pure water to obtain the substrates with ZnO nanostructures.

**Figure 1.** Schematic illustration of manufacturing ZnO modified metal fiber felt.

#### *2.3. Characteristic Analysis*

A field emission scanning electron microscope was used to examine the surface morphology of the samples as they were developed (SEM, HITACHI-SU8220, Hitachi, Ltd., Tokyo, Japan). Chemical component was characterized by energy dispersive spectrometry (EDS, Bruker, Billerica, MA, USA). Data-physics was used to characterize surface wettability, including water contact angle (WCA), oil contact angle (OCA), underoil–water contact angle (UWCA), and underwater–oil contact angle (UOCA) (Data-Physics, DataPhysics Instrumente GmbH., Filderstadt, Germany). The liquid droplets utilized for WCA, UWCA, CA and UOCA analyses had a volume of 4 µL. The amount of water remained in the collected filtrates was determined by Karl Fischer Titrator (TP653, TimePower Measure and Control Equipment Co. Ltd., Beijing, China). Total organic carbon (TOC) equipment was used to measure the oil content of the filtrates (TOC-L, Shimadzu (Shanghai) Global Laboratory Consumables Co. Ltd., Shanghai, China). The Zetasizer Nano ZS (Malvern Instruments Ltd., Malvern, UK) was used to investigate the dispersion of water droplets in oil and oil droplets in water. The microscopy pictures of W/O and O/W emulsions were acquired by using an inverted microscope (Keyence Corporation, Osaka, Japan).

#### *2.4. Oil–Water Separation*

Here, 99.6 g of infused oils (Diesel, Decane, Dodecane, and N-tetradecane) were combined with 0.4 g of water and ultrasonically agitated for 1 min to create water-in-oil(W/O) emulsions. To prepare oil-in-water (O/W) emulsions, 0.4 g of oil (Diesel, Decane, Dodecane, and N-tetradecane) was combined with 99.6 g of ultra-pure water and ultrasonically agitated for 10 min. The experiment of separation was performed by using a home-made filter device that sandwiched the created membrane between two glass tubes. Before

assembling the separation device, the prepared filter membranes were placed in ultrapure water or oil solutions for 30 s, dependent on the types of emulsions. The pre-wetted filter membranes were used in the separation experiments. A glass cylinder tube containing metal fiber felt as a filter was placed at the position shown in Figure 2 to realize filtration of oil–water emulsions under gravity. The filtrate was collected at the bottom of the vessel after the emulsion was poured through the open end of the glass cylinder. The oil–water emulsion was quickly poured into the home-made separation device. The liquid column was sustained at a depth of 10 cm during the separation. The separation efficiency was estimated by the following equation:

$$\eta = 1 - \frac{\mathcal{C}\_1}{\mathcal{C}\_0} \tag{1}$$

where *C*0 represented the dispersed phase content in the initial W/O and O/W emulsions (ppm); *C*1 represented the water/oil content in the collected filtrate (ppm); and *η* represented the separation efficiency. Meanwhile, the filtration flux (*L*) was calculated by:

$$L = \frac{m}{\rho \pi r^2 t} \tag{2}$$

where *m* was the mass of collected filtrate measured with a balance scale; *ρ* was the density of water/oil (*ρ* = 103 kg/m3 ); *r* was the radius of a glass tube (the radius of the glass tube in this study was 8 mm); *t* was the filtration duration, with the filtration flux computed by measuring the quantity of filtrate 1 min per time.

**Figure 2.** Schematic diagram of experimental setup.

#### **3. Results and Discussion**

#### *3.1. Surface Morphology and Chemical Composition of ZnO Coated Metal Fiber Felt*

The seed layer was coated on a metal fiber felt substrate using the sol-gel method. We fabricated ZnO-coated metal fiber felts by immersing the metal fiber felts with the seed layer into the precursor solution for the liquid phase growth process (Figure 1). At the initial growth stage, HMTA tardily decomposed to liberate OH− ions. The generated OH− ions form insoluble ZnO precipitated at supersaturated Zn(OH)4 2− ions or Zn2+ ions. As a polar crystal, ZnO (0001) exhibited a positive charge on the crystalline surface and attracted OH− ions and Zn(OH)4 2− ions in the solution. As a result, ZnO had the highest growth rate along the C-axis compared to other crystal planes and eventually forms ZnO nanostructures [37–40]. The overall chemical reactions were presented as follows:

$$\mathrm{C\_6H\_{12}N\_4 + 6H\_2O \to 6HCHO + 4NH\_3}$$

$$\mathrm{NH\_3 + H\_2O \to NH\_4^+ + OH^-}$$

$$\mathrm{Zn^{2+} + 2OH^- \to ZnO(s) + H\_2O}$$

SEM images indicated that the raw stainless-steel fiber felt was composed of metal fibers with an intricate arrangement and the surface of the metal fibers was smooth (Figure 3a,b). The average metal fiber diameter of the fabricated substrates increased, and the average pore size decreased after being coated with ZnO nanostructures (Figure 3c,d). The zoomed-in imaging showed that ZnO had nano pillar-like structures. Furthermore, the ZnO nanopillars in our growth experiments had hexagonal cross section with an average top diameter of 185 ± 81 nm. Finally, ZnO nanostructures wrapped metal fibers' surface without blocking pores.

**Figure 3.** SEM images of original filter membrane (**a**,**b**). SEM images of ZnO-coated filter membrane (**c**,**d**).

The liquid-phase ZnO growth had two stages: dissolution and crystallization. The growth solution concentration was one of the critical parameters for ZnO's nucleation and growth. The effect of growth solution concentration on the formation of ZnO coating on filter membrane was shown in Figure 4. We found that the surface was accumulated by dense ZnO structures at low growth solution concentrations (Figure 4a). The nanostructures had a slight pitch and low ratio of length to neck diameter. During crystal growth, the concentration of growth solution determined the solutions' concentration at the center of crystal growth face. As the concentration of growth solution increased, the supersaturation of Zn-ions in the solution was enhanced. It led to apparent orientation growth behavior and eventually formed nanoarray structures (Figure 4b). When the concentration increased to 0.05 mol/L, the filter membrane surface presented arranged pillar-like structures. A large gap between nanostructures facilitated the droplets' immersion into nanostructures. The average diameter of nano-pillars decreased from 314 ± 52 nm to 91 ± 24 nm, and the average length of the nano-pillars just increased slightly. ZnO nanostructure transformed from nano-pillars to nano-needles (Figure 4c). The decrease in the diameter of the nanostructures increased the size of structural interstices, resulting in a coating deficiency on metal fibers and weakened stability of ZnO coatings. Energy-dispersive X-ray spectroscopy (EDS) data were shown in Figure 4d–f. Zn and O element distribution and measurements in coated metal fiber confirmed that the final coating was constituted of ZnO with an excellent coverage density.

**Figure 4.**  Surface Wettability*