In these applications, crystalline materials are usually preferred for their superior electrical transport characteristics as well as spatial homogeneity and low surface roughness over the device area for reduced scattering. Here, we show applied electrical currents can induce a permanent reconfiguration of polycrystalline VO2 nanowires into crystalline nanowires, resulting in a dramatically reduced hysteresis across the phase transition and reduced resistivity. Low currents below 3 mA were sufficient to cause the local temperature in the VO2 to reach about 1780 K to activate the irreversible polycrystalline-to-crystalline transformation. The crystallinity was confirmed by electron microscopy and diffraction analyses. This simple yet localized post-processing of insulator-metal phase transition materials may enable new methods of studying and fabricating nanoscale structures and devices formed from these materials.Vanadium dioxide (VO Thermal modelling reveals that the applied current causes heating localized to the VO2 nanowire, which likely led to the rearrangement of grain boundaries in VO2. Our discovery is a new way to selectively and rapidly post-process nanostructures of polycrystalline VO2, and potentially other phase transition materials, deposited on amorphous substrates into a crystalline form.Materials capable of insulator-metal phase transitions (IMTs) enable nanoelectronic and nanophotonic devices, such as transistors124578102 on amorphous substrates, we drew inspiration from the recrystallization of metals and amorphous Si in which thermal annealing is often used to rearrange existing grain boundaries and promote the growth in size of the newly nucleated grains24252 micro/nanostructures, VO2 could crystallize upon cooling without the need to heat the entire substrate. Here in our manuscript we define the term “polycrystalline” to be a form when the grain size in the film is less than 100 nm in width and the “crystalline” form is when the grain boundaries are no longer observable through electron microscopy and diffraction analyses. Moreover, if the VO2 micro/nanostructure is smaller than the final grain size, the resultant structure would effectively be “single crystalline”. A way to deliver heating that is local and specific to the VO2 micro/nanostructure is by applying an electrical current. For sufficiently small cross-section areas, large current densities, hence sufficient heating, may be achieved at low (<3 mA) applied currents and the crystallization process can be extremely rapid. This process is in contrast to the typical thermal annealing of VO2, in which the film is heated gradually and uniformly to reduce V2O5 into VO2 or create oxygen vacancies2728To crystallize VO2 nanowire devices fabricated to demonstrate and investigate this phenomenon. A 120 nm thick VO2 film was deposited onto a 2 μm thick thermally grown SiO2 on a Si substrate using radio-frequency (RF) magnetron sputtering2 nanowires with varying widths and lengths using electron beam lithography and reactive ion etching. The two lateral electrodes were formed by electron-beam lithography and lift-off of a 100 nm thick palladium (Pd) film. The contacts were separated by a gap of length L varying between 0.75 μm and 7.5 μm, and the VO2 wire width, W, varied between 0.40 μm and 50 μm. We chose to use the two-contact geometry, rather than the four-point probe geometry in ref. VI) characteristics before and after the crystallization using the same devices. Details of the fabrication are described in the Methods section.2 film would not have a significant contribution to the total measured resistance, the surface area of the VO2 film under each Pd contact was kept to a large area of 4 μm × 50 μm. This surface area was identical for all the devices such that their contact resistances would be nominally identical. The H shaped VO2 region is indicated by the long dashed line in L = 7.5 μm and W = 5.0 μm. The X-ray diffraction (XRD) scan of the VO2 film in 2 but also had some contribution of V2O5. The calculated weight percentage of VO2 and V2O5 was 68.9 wt.% and 31.1 wt.%, respectively. The broad low amplitude peak between 15 °2θ and 47 °2θ with a maximum around 22 °2θ in the scan is due to the amorphous SiO2 under the VO2 film.To ensure that the contact resistance between the Pd contact pads and VOVI relations of the polycrystalline VO2 by sweeping the applied current between 0 mA and 0.3 mA with 1 μA increments, which is applied for 2.7 μs. We ensure that the devices undergo the nominal current-induced two-step phase transition in VO2311215VI plots are shown in blue in 2 nanowires with two different dimensions. The “before” VI curve has a large hysteresis in the second step due to the Joule heating assisted phase transition2 wiresThe fabricated devices were measured by contacting tungsten probes to the two Pd pads. Using a precision sourcemeter with a 1 kΩ resistor in series, current was applied to the devices. This resistor protected the device from being overdriven during the IMT. First, we obtained the “before” IC . This means the change in the VI characteristics was not due to the contacts but rather the VO2 wire. Furthermore, this suggests the crystallinity of the VO2 film under the Pd contacts experienced minimal change.After the “before” measurements, for each sample, we increased the applied current beyond 0.3 mA in increments of 0.2 mA until we observed a rapid, abrupt drop in the voltage at a critical current, IC . The volρ, as a function of current for the device of ρ vs. temperature plot in 2. This is consistent with an increased carrier scattering time or mean free path in the material by at least 2.5 times. However, the overall change in resistivity across the phase transition was reduced. The ratio between the metallic and insulator phase resistivity in the “before” case is 2.3 × 10−2, and the resistivity ratio in the “after” case is 3.6 × 10−2. The “before” and “after” phase transition temperatures were about the same at approximately 343 K, but in the “after” case, the hysteresis width was reduced to about 3 K, identical to that of single crystal VO2192023VI and ρ vs. temperature measurements were repeatable in both the “before” and “after” case. Device failure occurred when a current beyond IC was injected into the device.2 since the phase transition temperature in the “before” and “after” cases are similar . However, the “after” state should be dominantly VO similar . It is uely352O5  or chang2 into crystalline VO2. 2 nanowires in the “before” and “after” states. L = 7.5 μm and W = 0.91 μm. The top surface of the VO2 in 2 film in the “before” state in 2 surface is smooth with no visible grain boundaries. The expanded width on the right end of the wire is suggestive of material reflow. The XTEM of the “after” state in 2 wire width and length and we were not able to confirm whether the wire was entirely single crystal. For devices with W > 5 μm, only a narrow region in the VO2 wire appeared smooth, suggesting that the current was concentrated through a filament that had a minimum resistance between the two contact pads  from previous studies2 sample relative to the incident electron beam. To determine the incident direction of the beam, in c-axis lattice constant of VO2 (c = 2.85 Å), so we identified the diffraction spot to be from the (002) plane and varied the electron beam incident direction until the rest of the diffraction pattern was matched. A confirmation of the fitting is that with the correct view direction of the diffraction pattern, [210], the simulated diffraction pattern aligns with the SAED pattern. The faint spots in the diffraction pattern that do not match with the VO2 are likely due to the residual V2O5. The complete fittings of the diffraction patterns for both the monoclinic and rutile phases and the contribution from V2O5 are included in the W and L combined with heat transport simulations, we can ascribe the observed polycrystalline-to-crystalline transformation to a thermally activated process, akin to recrystallization in metals and amorphous semiconductors2425k is the thermal conductivity, and T is the temperature) to the VO2 nanowire, in the one-dimensional steady-state, if P) over cross-section area (Wt), thenThrough measurements of devices with varying T is the change in temperature. Therefore, the dissipated power is expected to be linearly proportional to the geometric parameter, W/L. In our experiment, R = Rtot − 2Rcon,m where Rtot is the total measured resistance of the device. P, at the critical current for the structural change is indeed a linear function of W/L. Using the thermal conductivity of VO2 in ref. T = 1470 K as shown in 2 wire was around 1770 K for the polycrystalline-to-crystalline transformation to occur. This temperature is about 80% of the melting point of VO2 of 2240 K, which is consistent with the requirements of recrystallization processes242425where Δ2 nanowires, including the Pd contacts and substrate, at the critical current. The calculations were performed using Joule heating module of COMSOL Multiphysics a stationary study in three dimensions using a non-uniform mesh (minimum mesh size of 5 nm) for different VO2 wire dimensions. The top surfaces of the structure are modelled as convective cooling surfaces allowing thermal transport into the air. The side and bottom boundaries were set to room temperature and were sufficiently far away from the VO2 wire such that they did not influence the results of the heat transport in the wire. The heat capacity of VO2 was modelled as ref. 2 nanowire with the experimental critical current density required for crystallization. 2 wire with L = 0.78 μm and W = 0.41 μm. The current density is the highest near the interface between the Pd contacts and VO2 wire. The centre of the VO2 wire reaches the highest temperature of 1780 K, similar to the value extrapolated from the one dimensional Fourier model and 2 to crystallize. Moreover, at these high temperatures it is unlikely that VO2 will further oxidize with the oxygen in the air to form V2O5. This is because the boiling temperature of V2O5 is 1750 KTo obtain a more accurate estimate of the temperature for the polycrystalline-to-crystalline transformation and to confirm whether such temperatures could have been generated in our experiment, we simulated the temperature profile in VO2 into crystalline VO2 in nanowire geometries by injecting electrical currents of the order of 1 mA to reach local temperatures near 1780 K. The reconfiguration is likely through a process akin to recrystallization in metals and semiconductors by thermal annealing, with a key difference that the transformation here was rapid, occurring within microseconds. The final crystal orientation of the VO2 nanowires was not controlled in these experiments, but it may be possible to induce a preferential orientation by introducing strain or patterning the substrate, for example. The crystallized devices exhibited reduced hysteresis across the phase transition and reduced resistivity, essential features for common electronic and photonic devices. Because polycrystalline VO2 films can be deposited at low temperatures (less than around 450–500 °C) on amorphous dielectrics or silicon, compatible with back-end of line complementary metal oxide semiconductor (CMOS) fabrication processing172 and other IMT materials with CMOS devices. Further, the simplicity of the current induced crystallization can broaden the investigation of crystalline VO2 and IMT nanostructures on unconventional substrates, including those that are lattice-mismatched, amorphous, or flexible.In summary, we have demonstrated the transformation of polycrystalline VO2 film was deposited using reactive RF magnetron sputtering of a 50 mm vanadium target (99.7% purity). A 4 cm × 4 cm piece of Si with a 2 μm thick of thermally grown SiO2 was placed on a 150 mm diameter Si wafer, which served as the substrate holder. During the deposition, the substrate holder was held at 500 °C and the chamber was filled with high purity argon (Ar) and oxygen (O2) at flow rates of 80.0 sccm and 2.0 sccm, respectively at a total pressure of 10 mT. The RF power was ramped to 200 W and the O2 flow rate was closely monitored to keep the DC bias of the sample constant. The deposition rate was approximately 0.3 Å/s.The VOμm × 20 μm tungsten alignment markers for the electron-beam lithography were formed onto the VO2 film. Then, the patterns for the VO2 nanowires were written using an electron-beam writer  in a 450 nm thick layer of ZEP-520A resist. The chip was then developed in ZED-N50 for 60 s and then in methyl isobutyl ketone (MIBK) for 30 s. Prior to etching, the resist was baked for 4 min at 100 °C. The sample was then etched using an inductively coupled plasma reactive ion etcher (ICP-RIE). The sample was placed on an aluminum chuck and pre-heated on a hotplate to 80 °C prior to insertion into the ICP-RIE, so the VO2 was in its metallic phase during the etch. A gas mixture of Cl2 (6 sccm), H2 (6 sccm), and Ar (9 sccm) at a total pressure of 5 mT was used for the etch. At a RF power of 125 W and ICP power of 500 W, 50 s was required for a full etch of the VO2. After etching, the ZEP resist was removed by immersing the sample in ZDMAC for 10 min at 70 °C followed by 1 min of sonication in the same solution.With this fabrication process, the minimal feature size of the VO2 nanowire obtained was 400 nm in width and 750 nm in length. Finally, contact pads were formed on the VO2 nanowires by aligned electron-beam lithography (with ZEP-520A resist), thermal evaporation of a 3 nm thick chromium adhesion layer and a 100 nm thick Pd layer (deposition rate: 1.5 Å/s), and lift-off.The devices were fabricated using aligned electron-beam lithography, dry etching, and metal deposition. First, 20 How to cite this article: Jeong, J. et al. Current induced polycrystalline-to-crystalline transformation in vanadium dioxide nanowires. Sci. Rep.6, 37296; doi: 10.1038/srep37296 (2016).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


T1 and T2 mapping to detect intramyocardial hemorrhage (IMH) in ST‐segment elevation myocardial infarction (STEMI) patients treated by primary percutaneous coronary intervention (PPCI).To investigate the performance of T1, T2, and T2* maps at 4 ± 2 days. Receiver operating characteristic (ROC) analyses were performed to assess the performance of T1 and T2 to detect IMH.Fifty STEMI patients were prospectively recruited between August 2013 and July 2014 following informed consent. Forty‐eight patients completed a 1.5T cardiac magnetic resonance imaging (MRI) with native T2* maps and 26/39 (67%) of the patients had IMH (T2* <20 msec on T2* maps). Both T1 and T2 values of the hypointense core within the area‐at‐risk (AAR) performed equally well to detect IMH . Using the binary assessment of presence or absence of a hypointense core on the maps, the diagnostic performance of T1 and T2 remained equally good  with good sensitivity and specificity .The mean age was 59 ± 13 years old and 88% (24/48) were male. In all, 39 patients had interpretable T1 and T2 maps can detect IMH equally well and with good sensitivity and specificity in reperfused STEMI patients and could be used as an alternative when T2* images are not acquired or are not interpretable.The presence of a hypointense core on the Level of Evidence: 2Technical Efficacy: Stage 2J. MAGN. RESON. IMAGING 2017;46:877–886 T2* imaging has been shown to detect IMH in animal models of acute MIT2* imaging is currently the reference standard for the detection of IMH, it is prone to motion, flow and off‐resonance artifacts.T2* data. Furthermore they also showed that MVO and IMH are dynamic and follow distinct time courses and had different prognostic significance.Primary percutaneous coronary intervention (PPCI) is the reperfusion strategy of choice in ST‐segment elevation myocardial infarction (STEMI). However, paradoxically, the process of reperfusion itself can lead to microvascular obstruction (MVO) and intramyocardial hemorrhage (IMH).T2‐weighted short tau inversion recovery (STIR) imaging, which has been used to delineate the area‐at‐risk (AAR) in reperfused STEMI patients, has also been used to detect IMH as the presence of a hypointense core within the AAR.T2* imaging.T1 and T2 maps have been explored as alternative methods for quantifying the AAR in reperfused STEMI patientsT1T2 maps.T1 and T2 maps to detect the presence of IMH and MVO following STEMI has not been directly compared. Therefore, the main aim of this study was to investigate the performance of hypointense core on the native T1 (referred to as T1 throughout the article for simplicity) and T2 maps to detect IMH within the first week in reperfused STEMI patients, using T2* mapping as the reference standard for IMH.T1 and T2 maps to detect early and late MVO, using gadolinium enhancement as the reference standard for MVO.The local Ethics Committee approved this study. The patients included in this study have been previously reported.T1 mapping, T2 mapping, three  LV short axis slices of T2* mapping and full LV short axis coverage of early and late gadolinium enhancement (EGE and LGE). All the short axis images were aligned with the cine short axis slice position. Colored T1/T/T2* maps using the default look up table from the scanner were generated.Cardiac MR was performed on a 1.5T scanner  using a 32‐channel phased‐array cardiac coil. The imaging protocol included full left ventricular (LV) short axis coverage of cines, native T1 maps were acquired with a steady‐state free precession (SSFP)‐based modified Look–Locker inversion recovery (MOLLI) sequence  using a 5s(3s)3s sampling protocol.T1 map.Native T2 maps consisting of pixel‐wise T2 values were generated inline following motion correction and fitting to estimate T2 relaxation timesT2 preparation times .Colored T2* maps were obtained (bandwidth 814(×8) Hz/pixel; echo times × 8: 2.7, 5, 7.3, 9.6, 11.9, 14.2, 16.5, 18.8 msec; flip angle = 18°; matrix = 256 × 115; slice thickness = 8 mm; uninterpolated resolution = 1.6 × 1.6 mm) and a colored pixel‐wise T2* map was generated inline by the scanner.T1‐weighted sequence at a fixed high inversion time of 440 msecEGE images were acquired 1–2 minutes after the injection of 0.1 mmol/kg of gadoterate meglumine  using inversion recovery single‐shot SSFP Imaging analysis was performed using CVI42 software . The endocardial and epicardial borders were manually drawn and MI size was quantified in grams and as a percentage of the LV volume (%LV) using a signal intensity threshold of 5 standard deviations (SD)T1 and T2 maps were performed with breath‐hold and motion‐correction, the raw images were visually assessed for any misalignment due to significant motion between single‐shot raw images, mistriggering, partial volume effects, and artifacts as previously described by von Knobelsdorff‐Brenkenhoff et al.T2* maps and their raw images were visually assessed for breathing, motion, and off‐resonance artifacts.Although the T2 maps and the AAR was obtained using a threshold of 2SDT2 relaxation time and expressed as %LV. Areas of hypointense core within the areas of hyperenhancement were manually included as part of the MI zone.The endocardial and epicardial borders were manually drawn on the T1 and T2 maps matching the T2* maps, the 2SD threshold was also used to identify the hypointense core  on both maps to identify the AAR and the hypointense core , the salvaged myocardium within the AAR , and the remote myocardium (expressed as T1Remote and T2Remote) on both maps. In cases when no hypointense core was identified, the ROI was drawn in the area of infarct. The T1 and T2 maps were also graded in a binary fashion to detect the presence or absence of a hypointense core (with the help of the semiautomated method as above).On the Ten randomly selected patients were separately analyzed by two investigators, 2 months apart for intraobserver and interobserver variability.T2* maps and the remote myocardium  and representative values were recorded.Corresponding ROIs were drawn in the hypointense core of the T2* maps with a T2* <20 msec as previously validatedIMH was defined as a hypointense core on the Early MVO was defined as areas of dark core within the infarcted territory visually detected  analyses were used to assess the diagnostic performance for T1 and T2 maps for detecting IMH on the acute scans and were compared using the method described by Delong et al.T1 and T2 values of the hypointense core were assessed in 10 patients and expressed as an intraclass correlation coefficient (ICC) . Cohen's kappa was used to assess interobserver and intraobserver agreement for the binary assessment of the maps. All statistical tests were two‐tailed, and P < 0.05 was considered statistically significant.Statistical analysis was performed using SPSS v. 22  and MedCalc for Windows v. 15.6.1 . Continuous data were expressed as mean ± SD or median (interquartile range) and compared using paired or unpaired tests where appropriate. Categorical data were reported as frequencies and percentages. Patients with no interpretable Figure T1 and T2 maps were not interpretable and were predominantly the basal or apical slices. All patients had a least one short axis T1 and T2 map matching the T2* maps. T2* maps were not interpretable in 19% (9/48) of the patients (due to motion flow and off‐resonance artifacts) but their corresponding T1 and T2 maps were interpretable.Ten percent of the T1 and T2 mapping acquisition (full LV coverage) and an additional 3 minutes for three short‐axis T2* mapping acquisitions.The average scanning time was 59 ± 4 minutes, longer than a clinical cardiac MR scan, as on average an additional 15 minutes was required for Figure T2 maps in 60% (29/48) of the patients and on the T1 maps in 63% (30/48) of the patients. Sixty‐seven percent (26/39) of the patients with interpretable T2* maps had IMH.A hypointense core was present on the T2* maps, the mean T2* value of the core was 13 ± 3 msec compared to 33 ± 4 msec in the remote myocardium, P < 0.001. IMH occurred in 67%(26/39) of patients. As expected, patients with IMH were more likely to have larger MI size  and worse ejection fraction . Further details on the cardiac MR findings are summarized in Table In patients with a hypointense core on the T1Core was lower than T1Remote  whereas T2Core was similar to T2Remote . In patients without IMH, T1Core and T2Core were higher than T1Remote and T2Remote as shown in Fig. T1 and T2 values of the hypointense core showed that both mapping techniques performed equally well at detecting IMH on the acute scans (T1 maps: area under the curve (AUC) 0.86 [95% CI 0.72‐0.99], cutoff value for T1Core: <1080 msec; T2 maps: AUC 0.86 [95% CI 0.74‐0.99]; P = 0.94, cutoff value for T2Core: <54 msec) . The presence of a hypointense core had a sensitivity of 88% and a specificity of 85% to detect IMH on the T1 maps and a sensitivity of 85% and specificity of 85% on the T2 maps. The accuracy by T1 mapping was 87% and 85% by T2 mapping. The positive predictive value by T1 and T2 mapping were both 92%. The negative predictive value was highest by T1 mapping at 87% and 85% by T2 mapping.In patients with IMH, T1  and 0.82 (95% CI 0.69–0.97) for T2 . For the detection of late MVO the AUC was 0.79 (95% CI 0.64–0.95) for T1  0.80 (95% CI 0.62–0.93) for T2 .The AUC for the presence of a hypointense core on the maps to detect early MVO was 0.83 (95% CI 0.70–0.97) for T1 and T2 maps to detect IMH, early MVO, and late MVO.Table T1 and T2 mapping to detect IMH, early MVO, and late MVO There was no significant difference in the diagnostic performance for T2* maps and early MVO. Eighty‐one percent (26/32) of the patients had early MVO with IMH and 19% (6/32) of the patients had early MVO without IMH. Both T1 and T2 values were significantly lower in patients with early MVO with IMH compared to those with early MVO without IMH . The same comparison was not performed for late MVO as only one patient had late MVO without IMH. All patients with early MVO with IMH developed late MVO compared to 2/6 (33%) of those with early MVO without IMH had late MVO, P < 0.0001.Thirty‐two patients had analyzable T1 maps was 0.944 (0.792–0.986) and for T2 maps was 0.903 (0.637–0.976). For interobserver variability, the ICC for T1 maps was 0.935 (0.746–0.984) and for T2 maps was 0.887 (0.528–0.972). On the qualitative mapping analysis for the hypointense core detection using the semiautomatic method, the interobserver and intraobserver agreement was 100%, Cohen's kappa = 1.0, P < 0.0001.On the quantitative mapping analysis of the hypointense core, for intraobserver variability the ICC for the quantification of the hypointense core for the T1 and T2 maps obtained at day 4 following a reperfused STEMI detected the presence of IMH with good specificity and sensitivity compared to the reference standard method using T2* maps. The binary assessment of presence or absence of a hypointense core performed as well as the quantitative assessment of the actual T1 and T2 values of the hypointense core. However, T2* maps are still the modality of choice when available, as the accuracy for T1 and T2 maps to detect IMH was only 85 to 87%, respectively. T1 and T2 maps provide an alternative option for the detection of IMH when T2* imaging is not available.We have shown that the presence of a hypointense core within the area if hyperenhancement (AAR) on the T1 and T2 mapping performed equally well to detect early and late MVO. This is not surprising for late MVO, as most patients with late MVO also had IMH and therefore this approach performs well to detect hemorrhagic MVO in our cohort. Although numerically the AUCs were higher for the detection of early MVO than late MVO, ROC curves comparison did not show a statistical difference. T1 and T2 mapping could differentiate between early MVO with and without IMH. Furthermore, those with IMH were more likely to display late MVO. Early MVO likely represents a spectrum of etiologies for microvascular injury and late MVO represents the more severe form as IMH. A recent studyT2‐STIR imaging could not discriminate between IMH and MVO but imaging was performed at 8 days and no T2* data were acquired. Given the recent evidence of the dynamic nature for the detection of the paramagnetic properties of IMH,The presence of a hypointense core on the T1 of the infarct core was recently shown to be more prognostic than LV ejection fraction, infarct core T2, and IMH in a large cohort of STEMI patients.T2* for detection of IMH, CMR were performed at 2–3 days  and also shown in a recent large study with 14% of T2* maps being not analyzable.T2* were not available and this was not the aim of this study, and some patients with small areas of IMH may have been missed. Histological validation for the low T1 and T2 of the hypointense core was not possible in this study and warrants further investigation. We used two LGE readouts to accommodate for patients who preferred not to breath‐hold for the LGE acquisition given the long duration of the scan, and the difference in signal‐to‐noise ratio between the two may have affected the detection of MVO. The large majority of our patients with MVO also had IMH. Therefore, we could not assess whether the hypointense core on the T1 and T2 maps could differentiate between late MVO with IMH and late MVO without IMH. However, we did find a difference in the T1Core and T2Core between those with early MVO with IMH and early MVO without IMH and this needs to be confirmed in future studies. We did not acquire data on black blood T2‐STIR images or susceptibility‐weighted cardiac MRI (which has been shown to improve the detection of IMH at 1.5 TT1 and T2 using SSFP readouts could have been due to a combination of both the on‐resonance and off‐resonance signals from the paramagnetic components of the IMH and this was not elucidated in this study.This was a small study of 39 patients and no formal power calculation was performed, but was similar in size to previous studiesT1 and T2 maps within the first week following a STEMI can detect IMH equally well and with good sensitivity and specificity. The T1 and T2 mapping techniques provide an alternative approach for the detection of IMH in situations when T2* maps are not interpretable or not available. However, T2* mapping currently remains the reference standard in the clinical setting and T1 and T2 mapping therefore may play a complementary role in future studies targeting IMH.In conclusion, the presence of a hypointense core on


We report the addition of M–H bonds  to a Rh(iii) intermediate generated from the reductive elimination of triethylsilane from [Cp*Rh(H)2(SiEt3)2]. iii) intermediate generated from the reductive elimination of triethylsilane from [Cp*Rh(H)2(SiEt3)2]. A series of new heterobimetallic complexes possessing Rh–M bonds have been isolated and characterised by a number of spectroscopic  and computational techniques . Experimental and computational data are consistent with cleavage of the M–H bond upon addition to rhodium with formation of new Rh–M and Rh–H bonds. Upon photolysis the Al analogue of this series undergoes a further elimination reaction producing triethylsilane and a highly unusual Rh2Al2H4 containing cluster proposed to contain an Al(i) bridging ligand.We report the addition of M–H bonds  to a Rh( We demonstrate that in combination the transformations may be defined by a spectrum of reactivity that lies between two extreme definitions: oxidative addition and hydride transfer. We show that as electronegativity difference between the M and H atoms increases, the description of hydride transfer becomes more accurate than oxidative addition.This paper concerns the addition of main group hydrides to a transition metal complex. We have studied the addition of M–H  bonds to a 16-electron Rh8].8 Despite these findings, little is known about the activation of heavier main group hydrides at transition metal centres, and less still about the heterobimetallic complexes that would result from hydride transfer from the main group to the transition metal. Herein we report the addition of M–H bonds  to a Rh(iii) intermediate generated from the reductive elimination of triethylsilane from [Cp*Rh(H)2(SiEt3)2].The oxidative addition of element–hydrogen bonds to transition metal centres is a reaction of fundamental importance to catalysis. Upon coordination of H2(SiEt3)2] under thermal or photochemical conditions to give the corresponding metal boryl complex by B–H activation.10 In line with these findings, thermolysis of [Cp*Rh(H)2(SiEt3)2] in the presence of β-diketiminate stabilised aluminium, magnesium or zinc hydrides in C6D6 solution led to the formation of the corresponding heterobimetallic complexes 1–3 along with a single equiv. of triethylsilane (12Hartwig and co-workers have shown that H–Bpin reacts with  by neutron diffraction experiments and assigned it as a Rh(v) complex.b While Hartree–Fock calculations support this formulation,11 subsequent calculations suggest that assignment of the +5 formal oxidation state may be misleading due to the fact that the Si···H distance may be compressed with very little energetic cost.11103Rh–1H HMBC experiments show that this latter species (δ = –1709 ppm) has a similar magnetic environment at rhodium to the heterobimetallic complexes we have isolated . The absence of H···H bonding was supported by long T1 relaxation times of the Rh–H resonances of 1–3 (T1 = 0.9–1.0 s). The lack of a strong M–H interaction was further evidenced by infrared data on 1–3. The Rh–H stretches  are shifted to slightly lower energy relative to that in [Cp*Rh(H)2(SiEt3)2] (2019 cm–1).11Maitlis and co-workers have characterised , and is similar to that found in a Rh/Al heterobimetallic recently isolated by our group.18 To the best of our knowledge, no examples of crystallographically characterised complexes containing a Mg–Rh bond exist. While the element–hydrogen bond lengths should be treated with caution it is noteworthy that in all cases the Rh–H bond lengths are short (1.42(3)–1.56(3) Å) and the M–H distances are long (2.06(3)–2.17(3) Å). For comparison the Al–H bond length of the terminal hydride in 1 is 1.65(2) Å  Å .A) and a rhodiate complex in which a cationic main group fragment is stabilised by coordination to a rhodium anion (B). Both bonding descriptions are likely to be augmented by additional non-classical M···H interactions. As a result the assignment of a formal oxidation state has limited meaning. In order to verify the position of the hydride atoms and to gain a deeper insight into the bonding within the ground-state structures of 1–3 a series of gas-phase DFT calculations were conducted. For comparison, [Cp*Rh(H)2(Bpin)(SiEt3)] and [Cp*Rh(H)2(SiEt3)2] were also analysed by DFT methods 2(SiEt3)2] and [Cp*Rh(H)2(Bpin)(SiEt3)], the Rh–H WBIs are consistent across the series (2(SiEt3)2] which demonstrate a low energy barrier to hydride translation and formation of a weak Si···H bond. It has also been concluded that [Cp*Rh(H)2(Bpin)(SiEt3)] contains a weak B···H interaction as evidenced by the line-broadening of the hydride resonances due to 1JB–H coupling.10 It is worth noting that a series of related complexes, including [Cp*Rh(PMe3)(η2-HSiR3)(SiR3)]+  and [CpRh(SiMe3)2(η2-HSiEt3)] have been formulated as σ-silane adducts based on the observation of significant 1JSi–H coupling constants > 20 Hz.fNBO calculations revealed that, while the Rh–M Wiberg bond indices of e series , ESI†. S19 The NPA charges on the bridging hydride atoms are small throughout the series. While the charge on Rh is also small for the B, Si and Al complexes it becomes significant for the Zn and Mg analogues taking values of –1.01 and –0.99 respectively , 0.30 (Si), 0.55 (Zn), 0.59 (Al), 0.89 (Mg).Our findings were further underscored by quantum theory atoms-in-molecules (QTAIM) calculations, which in all cases reveal bond critical points (BCPs) between Rh and M, Rh and H and not between M and H , ESI†. Tectively  and ESI†2(SiEt3)2] to form [Cp*Rh(H)(SiEt3)] was found to be energetically accessible. A transition state could not be located for this transformation. Scanning the potential energy surface by gradually increasing a Rh···Si distance within [Cp*Rh(H)2(SiEt3)2] revealed that reductive elimination of H–SiEt3 to form the reactive intermediate is energetically uphill but without an easily identifiable maxima corresponding to a transition state. The addition of the M–H bond to the resulting 16-electron intermediate was calculated to be increasingly exergonic across the series Si ∼ B < Al ∼ Zn < Mg (2(SiEt3)2].To gain insight into the mechanism of bond activation and to provide support for an elimination-addition process, additional DFT calculations were conducted. The reductive elimination of triethylsilane from [Cp*Rh(H) Zn < Mg . While t3)] may be non-reversible, while the addition of Si–H and B–H bonds is reversible (2(SiEt3)(Bpin)] undergoes competitive reductive elimination of both the silane, H–SiEt3, and the borane, H–Bpin upon heating in the presence of P(p-tol)3.10 While in the current case extended thermolysis of 1–3 led only to slow decomposition, photolysis of 1 for 6 h using a 400 W Hg lamp gave clean formation of 24, a product resulting from the selective elimination of H–SiEt3 from the heterobimetallic precursor 2(SiEt3)2].The calculations suggest the addition of Al–H, Zn–H and Mg–H bonds to  (4), which could be described as a Rh(iii) complex supported by an Al(i) ligand, or a Rh(i) complex with an η2;η2-bound aluminium dihydride  orbital would be disrupted. Hence, we tentatively propose that the Al fragment in 24 acts as an LZ-type ligand rather than an X2-type ligand and that the monomeric unit 4 could be formulated with a terminal Al(i) ligand (see ESI6Triethylsilane elimination from d see ESI.6i)/Co(iii) formulation in the cationic fragment of the cobalt complex [(OC)3Co(μ-H)2AlBDI2][Co(CO)4]. This latter species also contains an additional donor–acceptor interaction in the form of coordination of an oxygen based lone pair from the anionic [Co(CO)4] unit to Al. Furthermore, the generation of a gallium(i) ligand from the spontaneous transition metal mediated dehydrogenation of a β-diketiminate supported gallium dihydride has been reported by the same group.8e Two pathways could explain the generation of 4 from 1, either a 1,2-elimination of H–SiEt3 from across the Rh–Al bond or reductive elimination of H–SiEt3 from the Rh metal centre followed by α-migration of the terminal hydride on Al to Rh. Based on the current data we cannot discriminate these mechanisms.Aldridge and co-workers have provided computational data to support a similar Al(iii) complex [Cp*Rh(H)(SiEt3)]. The reaction results in the cleavage of the M–H bond and the formation of a new Rh–M and Rh–H bond. While parallels may be drawn with an oxidative addition mechanism, only in the case of Al is this description valid. For Zn–H and Mg–H calculations reveal considerable charge localisation on rhodium and hence hydride transfer to generate an ate complex is a more apt description of the reaction. DFT calculations suggest that the addition of M–H  bonds to the Rh(iii) intermediate is non-reversible. This latter hypothesis has been supported by initial reactivity studies, which demonstrate the selective elimination of triethylsilane from a rhodium complex containing ligands derived from both a silane and an alane. We are continuing to study the unusual heterobimetallics reported herein as catalysts and reagents for difficult bond transformations.We have demonstrated the addition of aluminium, zinc and magnesium hydrides to a coordinatively unsaturated, 16-electron, Rh(


A reference to the metabolic cost in Patagonia sea lions is missing in the first paragraph of the discussion. The following sentence should be corrected: The estimated resting metabolic rates were similar to those measured in Steller sea lions and California sea lions in water  were similar to those measured in Steller sea lions and California sea lions in water  were similar to those measured in Patagonia sea lions, Steller sea lions and California sea lions in water  and the reported value for s 2 should be 53 mL O2 • min−1 • kg−0.6 or 8.6 ± 1.9 mL O2 • min−1 • kg−1.4. A typo was detected in the reported mass-specific metabolic rate (s The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


The obtained HM‐TiO2‐NB superstructure demonstrates high reversible capacity, excellent cycling performance, and superior rate capability. Most importantly, a self‐improving phenomenon of Li+ insertion capability based on two simultaneous effects, the crystallization of amorphous TiO2 and the formation of Li2Ti2O4 crystalline dots on the surface of TiO2 nanowires, has been clearly revealed through ex situ transmission electron microcopy (TEM), high‐resolution transmission electron microscopy (HRTEM), X‐ray diffraction (XRD), Raman, and X‐ray photoelectron spectroscopy (XPS) techniques during the Li+ insertion process. When discharged for 100 cycles at 1 C, the HM‐TiO2‐NB exhibits a reversible capacity of 174 mA h g−1. Even when the current density is increased to 50 C, a very stable and extraordinarily high reversible capacity of 96 mA h g−1 can be delivered after 50 cycles. Compared to the previously reported results, both the lithium storage capacity and rate capability of our pure TiO2 material without any additives are among the highest values reported. The advanced electrochemical performance of these HM‐TiO2‐NB superstructures is the result of the synergistic effect of hybriding of amorphous and crystalline (anatase/rutile) phases and hierarchically structuring of TiO2 nanowire bundles. Our material could be a very promising anodic material for lithium‐ion batteries.A hierarchical mesoporous TiO Such super hierarchical structures used as anode materials in lithium‐ion batteries exhibit high reversible capacity, excellent cycling performance, and superior rate capability. More interestingly, a self‐improving phenomenon of the Li+ insertion capability based on the crystallization of amorphous TiO2 and the formation of new Li2Ti2O4 crystal islands in the HM‐TiO2‐NB super structure has been revealed during the discharge–charge processes. The synergy of the amorphous/crystalline (anatase–rutile) phases hybriding and the hierarchical structuring of mesoporous TiO2 nanowire bundles leads to the large enhancement of the electrochemical performances and the self‐improving phenomenon.Herein, we demonstrate the design and synthesis of hierarchically structured mesoporous TiO2Figure2‐NB sample. The diffraction peaks at 25.3°, 37.9°, 48.1°, 54.1°, 62.8°, and 75.3° can be indexed to the (101), (103), (200), (105), (213), and (107) planes of anatase , respectively. The other peaks at 27.3°, 36.1°, and 41.3° can be attributed to the (110), (101), and (111) planes of rutile , respectively. The XRD patterns indicate that the obtained HM‐TiO2‐NB structure is a mixture of anatase and rutile phases. In addition, all the reflections have a large half‐peak width and the XRD pattern is slightly rough, indicating the low crystallinity with the possible existence of the amorphous region or the small crystallites in the HM‐TiO2‐NB nanowire‐bundled structure. It can be calculated that the weight ratio of anatase and rutile is ≈43:57.2 phases in the HM‐TiO2‐NB structure was further proven by the Raman spectroscopic analysis  and 514 (A1g) cm−1 are attributed to the anatase phase,Eg) and 605 (A1g) cm−1 can be indexed to the rutile phase, respectively.2‐NB sample, which is composed of numerous well‐defined flower‐like superstructures of about 500–600 nm in diameter, which are constructed by the hierarchical assembly of several nanowire bundles of a diameter of 50–100 nm , (110), (103), and (200) planes of the anatase phase and (101) and (200) planes of the rutile phase, respectively, being in agreement with the XRD data and Raman spectroscopic study . Figure 2‐NB sample in our intensive observations . Such a superstructure with hybrid nanowires can provide, on one hand, higher specific surface areas and more active sites for the Li+ insertion/extraction and promote the Li+ storage properties. On the other hand, each core of nanowire is formed by crystalline structures (mixture of anatase and rutile) of TiO2, which are a stable host for the Li+ insertion/extraction reaction and facilitate charge across the electrode/electrolyte interface.Transmission electron microcopy (TEM) observations were carried out to present more detailed information of the as‐prepared HM‐TiOe Figure e–g. The nding to 01, (110)2‐NB superstructure has been illustrated in Figure2 precursor occurs accompanied by the decrease of the charge density of TiO2 at the interface. This process is caused by the increase of the interfacial curvature of the surfactant matrix,2 nanoparticles . It gives a type‐IV isotherm with a type H3 hysteresis loop, indicating a mesoporous structure.2‐NB has a high specific BET surface area of 117.8 m2 g−1 with a pore size centered at 6 nm , being in well consistent with the TEM observation . Since without the addition of hydrochloric acid, only amorphous nanoparticles with different sizes appear in the final product . Under our reaction system, the copolymer P123 and humidity can control the hydrolysis and condensation of the titanium isopropoxide precursor.2/Li half‐cell is thought to be4+ to Ti3+ and a theoretical capacity of anatase TiO2 of 335 mA h g−1.0.5TiO2 phase with orthohombic symmetry. The lithium ions are randomly distributed over half of the available interstitial octahedral sites. Therefore, from a practical point of view, the maximum reversible insertion level is limited to an insertion coefficient of 0.5 Li per formula unit of TiO2, corresponding to the two phases equilibrium between the Li‐poor Li0.05TiO2 and the Li‐rich Li0.5TiO2 phases. The associated reversible capacity is generally 167 mA h g−1.2. It has been observed that microcrystalline rutile TiO2 does not show an excellent Li storage property, due to the peculiar tetragonal structure, which does not facilitate Li+ transport along the ab‐plane in comparison to that along the c‐axis.2 with various morphologies demonstrated an insertion coefficient of ≈0.5 mol of Li per TiO2,2 can enhance the Li ions storage capacities owing to a pseudocapacitance effect.+ diffusion paths and the electron transport lengths.Generally, the principle reaction that governs the electrochemical processes in nanostructured TiOFigure2‐NB electrode at 0.5 mV s−1 for the first few cycles at 1–3 V (vs Li/Li+), displaying redox reaction at 1.3 V and 1.8 V (vs Li/Li+) during discharge and at 2.0 V during charge. The peaks at 1.75 V (cathodic sweep) and 1.9 V (anodic sweep) correspond to the Ti4+/Ti3+ redox reaction in the anatase phase2‐NB . The results show that the HM‐TiO2‐NB can accommodate Li up to Li0.88TiO2 (293 mA h g−1) during the first discharge process. Subsequently, 0.7 mol of Li per 1 mol TiO2  can be cycled reversibly, leading to a high Coulombic efficiency of 81%. The discharge curve in the first cycle can be divided into two stages. In the first stage, it is observed that the voltage decreases monotonically from the open circuit sites of ≈2.9 V to ≈1.8 V, corresponding to a Li+ insertion capability of 34 mA h g−1. In this region, lithium is stored at the surface or inserted into the amorphous surface layer. In the second stage, it can be observed that the potential decreases slowly from 1.8 V to 1.0 V, leading to a Li+ capability of 259 mA h g−1. The lithium capacity in the second stage corresponds to lithium insertion into the TiO2 lattice and further insertion of Li+ into the surface layer of the electrode material. Remarkably, there is only a very short plateau (not obvious) compared with that previously reported in the literature.2−1 with a corresponding charge capacity of 222 mA h g−1, leading to a Coulombic efficiency of 88%. In the sixth discharge–charge cycle, the Coulombic efficiency even increased to 95%.Figure  C 1 C = 7 mA g−1.2‐NB electrode at various current densities. The discharge‐specific capacities of 292, 199, 175, 162, 134, and 112 mA h g−1 are obtained when the current rates are 0.2, 0.5, 1, 2, 5, and 10 C, respectively. Figure 2‐NB electrode at various current densities. It is important to note that an initial discharge capacity of 292 mA h g−1 could be obtained at a current rate of 0.2 C, leading to a lithium insertion coefficient of 0.88, which is much higher than the theoretical value of 0.5.2‐NB electrode to have a discharge capacity of 112 mA h g−1 even at a rate of 10 C. Although the sluggish kinetic of additional surface capacity, with its increase in current rate due to the tardiness of lithium storage at the electrode surface, decreases the total storage capacity of the electrode, the extraordinarily high capacity observed at the rate of 10 C should be assigned to the very high specific surface area and the amorphous phase of our HM‐TiO2‐NB superstructure providing more active sites for lithium insertion. The straight nanochannels in each nanowire bundle can provide shorter lengths for Li+ diffusion and electron transport, as well as quick liquid electrolyte penetration. A capacity of 199 mA h g−1 was resumed while the current density was reduced back to 0.2 C again, indicating the very high stability of the HM‐TiO2‐NB electrode. Thus, the second round cycle performance was performed on the same LIB cell. Obviously, the HM‐TiO2‐NB superstructure demonstrates excellent capacity retention at various current densities. When charged at 1 C, the HM‐TiO2‐NB has an initial discharge capacity of 189 mA h g−1 and a subsequent charge capacity of 188 mA h g−1, leading to an irreversible capacity loss as low as 0.4%. Very interestingly, a self‐improving phenomenon of the Li+ insertion capability of the HM‐TiO2‐NB material during the ultra‐long discharge–charge cycles was observed. The initial discharge capacity of 189 mA h g−1 in the cycling region is higher than that of 175 mA h g−1 in the rating region at 1 C. Figure S6 (Supporting Information) shows the discharge–charge profiles of HM‐TiO2‐NB anode material at 1 C in the rating and cycling processes, respectively. When discharged in the rating process, the HM‐TiO2‐NB anode material demonstrates capacities of 34 mA h g−1 for the first stage and 141 mA h g−1 for the second stage, respectively. However, when discharged at 1 C in the cycling process, the HM‐TiO2‐NB anode material possesses capacities of 34 mA h g−1 for the first stage and 155 mA h g−1 for the second stage, respectively. It clearly displays that the capacities at the second stage in the cycling process is higher than that in the rating process. According to previous research, this self‐improving phenomenon may be generated by the crystallization of amorphous TiO22Ti2O4 crystal particles.2‐NB electrode was discharged after 100 cycles, a very high reversible capacity of 174 mA h g−1 was retained, indicating the excellent electrochemical and structural stabilities of the HM‐TiO2‐NB anode material. This corresponds to a lithium coefficient of 0.52, which is slightly higher than the previously determined maximum value of 0.5. Its reversible volumetric capacity at 1 C after 100 cycles is ≈283.6 mA h cm−3, according to the measured tap density value of 1.63 g cm−3. Figure S7 (Supporting Information) demonstrates the Coulombic efficiency of HM‐TiO2‐NB anode material at 1 C for 100 cycles. It clearly shows that the Coulombic efficiency almost retains at ≈100% for 100 cycles. The discharge capacity decreased to 163 mA h g−1 when the current density was increased to 2 C. When the electrode was charged/discharged at 2 C after 50 cycles, a reversible discharge capacity of 159 mA h g−1 was retained. The cycle performance shows that while charged at high current rates of 5, 10, 20, and 30 C after 50 cycles, the HM‐TiO2‐NB electrode is still able to deliver reversible capacities of 147, 131, 112, and 104 mA h g−1, respectively, which are much higher than many other TiO2 structures.−1 could still be delivered after 50 cycles. These extraordinarily high capacities and rate capabilities are among the best performances of TiO2 in the literature. These extraordinary performances can be ascribed to particular hierarchically structured mesoporous TiO2 amorphous/crystalline hybrid nanowire bundles. This special superstructure can provide enough spaces for the conductive phase to Li ions storage and shorten the Li+ diffusion paths and the electron transport lengths. Furthermore, the porous straight channel structure of the HM‐TiO2‐NB can further facilitate the electrolyte penetration and the lithium diffusion, as well as the electron transportation in the electrode.2‐NB sample. Thus, the amorphous, crystalline phases, and the hierarchical mesoporous nanowire bundled structure can synergistically enhance the Li+ insertion capability, high rate capacity, and cycle life.Figure 2‐NB material as an anode, the as‐prepared HM‐TiO2‐NB was heat treated at 350 °C (HM‐TiO2‐NB‐350) in the air. From the XRD , SEM , and TEM images , it can be observed that the hierarchical structure morphology can be generally retained, while the amorphous phase is completely transformed into crystalline phases accompanied by the broadening of the diameter of nanowires to 15 nm, leading to the decrease in number of straight nanochannels. Thus, the HM‐TiO2‐NB‐350 sample has a lower surface area of 72.4 m2 g−1 with a narrower pore size of 4–5 nm compared to the as‐prepared HM‐TiO2‐NB sample . The N2 adsorption–desorption isotherms are still of type IV with type H3 hysteresis . In the CV curves  and 2.1 V (anodic sweep) than the HM‐TiO2‐NB electrode, indicating that the HM‐TiO2‐NB‐350 electrode has more obvious and stable redox potentials2‐NB‐350 electrode  phases in each nanowire, the very high surface area of HM‐TiO2‐NB superstructure, and the presence of a large number of straight nanochannels giving an easy accessibility of electrolyte and Li ions to nanowires are critical for advanced lithium storage.To better understand the excellent electrochemical performance of HM‐TiOes Figure4a, the2‐NB electrode, and the self‐improvement phenomenon of Li+ insertion capability, the post‐mortem studies were carried out after 100 discharge–charge cycles at 1 C rate. As shown in Figure2 nanowire bundles is well retained , respectively. HRTEM observation was performed on the lithiated HM‐TiO2‐NB sample. The lattice spacing of the particle in Figure 2Ti2O4, confirming the formation of crystalline Li2Ti2O4 phase.2Ti2O4 according toIn order to further understand the advanced lithium storage property, the structure stability of the HM‐TiOd Figure a,b. A lad Figure c, inset.d Figure c. These e Figure c,e. FiguxTiO2 will lead to an atomic rearrangement in the HM‐TiO2‐NB structures  to form the new cubic crystals of Li2Ti2O4 corresponding to 1 Li inserted per formula unit of TiO2.2Ti2O4 crystal can further facilitate the Li+ insertion capability of the HM‐TiO2‐NB anode material during the electrochemical reaction, leading to the self‐improvement of the Li+ insertion capability at the same current density  during the cycling stage (after the rating stage), as shown in Figure 2‐NB sample has a higher capacity of 189 mA h g−1 at 1 C at the cycling process than that at the rating process (175 mA h g−1). Moreover, when the current density was increased to 50 C for 50 cycles, a high capacity of 96 mA h g−1 is still retained.Based on the above TEM, HRTEM, and SAED results shown in Figure 2‐NB electrode.Moreover, the distances of lattice spacing were measured to be 0.35 and 0.22 nm , respectively. Compared with the as‐prepared HM‐TiO2‐NB material, all the diffraction peaks of the reacted HM‐TiO2‐NB anode material from the initial anatase and rutile phases are weaker, especially the (101) plane of anatase and the (110) plane of rutile, indicating the Li+ insertion during the discharge process. Moreover, it can be observed that the peak of the anatase (101) plane decreases in intensity more rapidly than the peak of the rutile (110) plane. The appearance of diffraction peaks at 39.3° and 43.2°  and Ti3+ (456.7 eV) characteristics. This means that Li+ inserts into the “nanohybrid” regions of HM‐TiO2‐NB, leading to LixTiO2 formation. The formed LixTiO2 in the HM‐TiO2‐NB anode material has an enhanced electrical conductivity due to Ti3+ existence and further improves the Li+ storage performance.+ insertion mechanism into the hierarchical HM‐TiO2‐NB structure inducing the crystallization of amorphous phase and the formation of Li2Ti2O4 phase is proposed and illustrated in Figure+ ions can quickly intercalate into the amorphous layer and nanodomains of nanowires in the HM‐TiO2‐NB material, leading to the formation of LixTiO2 matrix and excellent Li+ insertion capability. Then, with the further Li+ insertion into the LixTiO2 matrix, the disordered atoms of Li, Ti, and O rearrange into a new ordered Li2Ti2O4 crystalline structure to reduce the free energy of the system. This newly obtained Li2Ti2O4 crystal can further facilitate the Li+ insertion capability of the HM‐TiO2‐NB anode material during the discharge–charge process, leading to a self‐improving phenomenon of Li+ insertion capability.Then, the ex situ XRD, Raman, and X‐ray photoelectron spectroscopy (XPS) techniques were utilized to detect the crystalline structure and the surface electronic state of the HM‐TiO° Figure b and the° Figure c indicat32‐NB) has been synthesized. This superstructure can provide more active sites for the Li+ insertion/extraction and facilitate the Li+ diffusion and the Li+ insertion/extraction process to promote the Li+ storage and electrochemical stability of the anode. The crystalline phases inside the hierarchical structure are a stable host and can facilitate charge across the electrode/electrolyte interface resulting in long cycle life and high rate performances. Compared with that presented in the literature, our material demonstrates superior electrochemical performances. The synergy of amorphous layer, mesoporous straight channels, and the crystalline nanodomains make a great enhancement of the electrochemical performances of TiO2. Through the nanowire amorphous surface, Li+ can continuously intercalate into the HM‐TiO2‐NB structure to form an amorphous LixTiO2 matrix and a new Li2Ti2O4 crystal with a size range of 5–10 nm. This newly obtained Li2Ti2O4 crystal can further facilitate the Li+ insertion capability of the HM‐TiO2‐NB anode material during the discharge–charge process, leading to a self‐improving phenomenon of Li+ insertion capability. Our results indicate that HM‐TiO2‐NB with amorphous surface and straight channels is a good candidate for advanced lithium storage.By designing on the basis of hybriding of amorphous/crystalline phases in nanowires and hierarchical structuring of mesoporous bundled nanowires, a very promising anode material , 0.5 g of P123 and 1 mL of concentrated HCl were added to 20 mL of anhydrous ethanol under vigorous stirring. Then, 2.96 mL of TTIP was added to the above solution. The final mixture was further stirred in a 50 mL beaker for 12 h at room temperature. The beaker was then placed in an oven at 40 °C with humidity over 60 RH% for 12 h and subsequently heated at 80 °C for 12 h. Finally, the product was washed using ethanol reflux method for 12 h and dried at 60 °C.Characterizations: XRD patterns were obtained on a Bruker D8 system with Cu Κα radiation (λ = 0.15405 nm) with 40 mA and 40 kV. SEM observation was carried out using an S‐4800 field‐emission SEM . TEM and HRTEM images and SAED patterns of the samples were recorded on carbon‐coated copper grids by using a JEM‐2100F microscope at an acceleration voltage of 200 kV . Nitrogen adsorption–desorption isotherms were obtained using a TriStar II 3020 surface area and porosity analyzer at 77 K . The specific surface area was calculated by the Brunauer–Emmett–Teller (BET) method. The pore size distribution was calculated by the Barret–Joyner–Halenda (BJH) method. Raman measurements were carried out at room temperature, and the signals were recorded by an Invia Raman Microscope  using the 632.8 nm line of a Nd YAG laser as the excitation source. The surface electronic states of the Ti element were analyzed by XPS (VG Multilab 2000). Tap density was measured by the BT‐301 tap denser. The as‐prepared powders were poured into 25 mL glass cylinder, and this cylinder was tapped for 10 min with 300 taps min−1 frequency.Electrochemical Characterization: Electrochemical experiments were performed with Swagelok‐type cells with pure lithium metal as both the counter electrode and the reference electrode at room temperature. The working electrode consisted of active material (HM‐TiO2‐NB), a conductive agent , and a polymer binder (poly vinylidenedifluoride) (PVDF) in a 70:20:10 weight ratio. After these materials had been thoroughly mixed in an N‐methyl‐2‐pyrrolidone (NMP) solution, the prepared slurry was coated on Cu foil. The coated electrode was dried at 120 °C in a vacuum oven for 12 h. Circular disk electrodes were punched from the foil and a lithium metal foil was used as the counter electrode. The electrolyte used was 1.0 m LiPF6 in a 50:50 (w/w) mixture of ethylene carbonate and diethyl carbonate. Cell assembly was carried out in an argon‐filled glove box. Cyclic voltammetry (1–3 V) was performed using an electrochemical workstation (CHI 660C). The charge–discharge tests were performed using a battery tester (LAND) with a voltage window of 1–3 V at various current rates.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The nano-sized spinel crystallites connected together to form mesoporous structure by annealing homogeneous complex oxalates precursor at a low rate of heating. It is found that the low anneal rate plays an important role for the formation of mesoporous ZnCo2O4 rods. The effects of the heat temperature on the phase, morphology and catalytic properties of the products were studied. The XRD, SEM TEM, and N2 absorption/desorption have been done to obtain compositional and morphological information as well as BET surface area of the as-prepared sample. Catalytic activities of mesoporous ZnCo2O4 rods toward the thermal decomposition of ammonium perchlorate (AP) were investigated with differential scanning calorimetry (DSC) and thermogravimetry (TG) techniques. The results show that the addition of ZnCo2O4 rods to AP dramatically reduces the decomposition temperature. The ZnCo2O4 rods annealed at 250 °C possesses much larger specific area and exhibits excellent catalytic activity (decrease the high decomposition temperature of AP by 162.2 °C). The obtained mesoporous ZnCo2O4 rods are promising as excellent catalyst for the thermal decomposition of AP.Mesoporous ZnCo Ammonium perchlorate (AP), as the main high energy constituent of solid propellants in the national defense field, accounts for 60–90% of the total weight of composite propellants4. Hence, the burning velocity and energy features of the propellants are significantly affected by the decomposition of AP. A comprehensive research about the thermal decomposition of AP was carried out by researchers5. Results indicate that a small quantity of catalysts can reduce the thermal decomposition temperature of AP, especially that corresponding to the high temperature decomposition (HTD), boosting up the apparent decomposition heat of AP, and increasing the burning velocity and efficiency of propellant accordingly. Recently, various kinds of transition metal oxides have been explored as catalyst for use in the thermal decomposition of AP, primarily including Fe2O3, CuO, Co3O4, NiO, ZnO and other metal oxide powders9. Compared to the single metal oxide, the complex oxides (containing two or more types of cations) with spinel structure have allured a lot of attention in material research because that the stabilization of active phases and synergistic interactions between two different oxides may improve its catalytic performance12. The presence of a partially filled-3d orbital in CuCo2O4 structure, so easy to accept electrons and improve the transferring electrons from perchlorate ions to the ammonium ions, which exhibits its stronger catalytic activity than probably CuO and Co3O4 as reported by Gheshlaghi et al.13. On account of the synergistic effect between Cu2+ and Cr3+ in the CuCr2O4 nanoparticles, the CuCr2O4 exhibits better catalytic effect than CuO15.Due to their various applications including propulsion for large space vehicles and tactical missiles, gas generators for airbags, composite solid propellants are being intensively pursued2O4 is an attractive material with the bivalent Zn-ions occupying the tetrahedral sites in the cubic spinel structure and the trivalent Co-ions occupying the octahedral sites. ZnCo2O4 has recently attracted more and more attention for applying in catalyst in thermal decomposition of AP, owing to its excellent physicochemical performance and its abundant resources, low cost and environmental friendliness17. Porous structured materials have gained great attention for their superior physicochemical characters, including large specific area, high porosity, low density, high permeability and high adsorption performance. Hence, reports are available on the application of this type of materials in catalyst, for the fact that it favors the exposure of active sites and offer rapid mass transfer processes19. So it is an important part of this article to make further improvement on the catalytic activities of AP by introducing porous structure into spinel ZnCo2O4. Up to now, most approaches towards porous materials focus on template-assisted processes, including hard templates (porous silicon and polystyrene sphere) and soft templates (surfactants and block copolymers). Tomboc et al. synthesized hierarchical mesoporous ZnCo2O4 via using PVP as the soft template20. Highly ordered mesoporous spinel ZnCo2O4 with high surface area and narrow pore size was synthesized by using SBA-15 as the hard templates21. However, this synthetic route involves multi-step process, which is time-consuming and relatively complicated. So, it is very necessary to look for a simple and effective method to synthesize porous structure ZnCo2O4. A facial oxalate co-precipitation method to fabricate mesoporous spinel Co3O4 and MCo2O4(4.5)  with high surface areas by a controlled pyrolytic of metal oxalate precursor as reported in22. This strategy is simpler and more efficient compared with other porous-casting method.As a binary oxide, ZnCo2O4 rod has been successfully synthesized via oxalate co-precipitation method without any template. The nano-sized spinel crystallites connected together to form mesoporous structure by annealing homogeneous complex oxalates precursor at a low rate of heating. The effects of the heat temperature on the phase, morphology and catalytic properties of the products were studied. The composition, morphology, porous structure, surface area and the catalytic activities to AP’s thermal decomposition of as-prepared mesoporous ZnCo2O4 rod are investigated in detail. The preliminary analysis of catalytic mechanism is discussed.In this paper, the mesoporous ZnCoAll the reagents used in the experiments were purchased from commercial sources of analytical grade and used without further purification.2O4 rod is as follows: 16 mM Co(NO3)2·6H2O and 8 mM Zn(NO3)2·6H2O were successively dissolved in 80 mL deionized water followed by magnetic stirring for 30 min to obtain a homogeneous solution. Subsequently, the 24 mL saturated sodium oxalate solutions was added slowly. After continuous stirring for 2 h, the precipitate was collected by centrifugation and washed with deionized water and absolute ethanol several times and dried in air at 60 °C overnight.The typical procedure adopted for the preparation of mesoporous ZnCoAfterward, the pink precursor was annealed at various temperature  for 2 h at a lower heating rate (1 °C/min) in temperature programming furnace. After that, the product was cooled inside the furnace to room temperature. The resultant black powder was collected and directly subjected to the various characterizations.α X-ray beam was adopted as the excitation source and power was set to 250 W. Vacuum pressure of the instrument chamber was 1 × 10−7 Pa as read on the panel. Measured spectra were decomposed into Gaussian components by a least-square fitting method. Bonding energy was calibrated with reference to C1s peak (285.0 eV). The pore size distributions and the BET surface areas were measured by nitrogen adsorption/desorption using a NOVA2200e gas sorption analyzer (Quantachrome Corp.). Prior to the measurements, the sample was degassed at 300 °C in vacuum for 3 h. The functional groups and coordination of the samples were studied via the FT-IR analysis performed on the Nicolet iS 10 (Semerel technology Corp.), the frequency is range from 4000 to 400 cm−1. Thermogravimetric and differential thermal analysis (TG-DTA) were performed on a HCT-3 thermal analyzer (Bei-jing) at a heating rate of 10 °C/min from 25 °C to 400 °C.X-ray powder diffraction (XRD) patterns of the product was carried out on a Rigaku D/max-3B diffractometer with an incident X-ray wavelength of 1.540 Å (Cu Kα line), operated at 40 kV, 100 mA. The morphology was observed using field emission scanning electron microscopy (FESEM) taken on FEI nova nanosem 450 with microscope operating at 30 kV. Detailed studies of the microstructure were also carried out by transmission electron microscopy (TEM) (JEOL JEM-2100) at an acceleration voltage of 200 kV. XPS was carried out at room temperature in a PHI 5500 spectrometer with polychromatic Al/Mg-Kα X-ray source. During XPS analysis, Al K2O4 rod in the thermal decomposition of AP were performed using a HCT-3 thermal analyzer (Bei-jing) at a heating rate of 20 °C/min in nitrogen atmosphere over the range of 25–500 °C.The as-prepared sample was mixed with AP to reach certain mass ratios of 2%, 5%, 7% and 10%, respectively. The mixture was fully grinded in the presence of a certain amount of anhydrous ethanol until the ethanol volatilize. Afterwards, the catalytic activities of mesoporous ZnCo23. Therefore, the pyrolytic behavior of homogeneous complex oxalates precursor was first explored by TG-DTA analysis in air atmosphere, the results of which is showed in Fig. 2O4. This transitional process is accompanied by a strong exothermic peak at approximately 285 °C in the DTA curve. The single-phase feature of cobalt-zinc complex oxalate is further confirmed by the phenomenon that only one dehydration and one decomposition stage can be observed, as the endothermic and exothermic peaks of zinc oxalate do not accord with those of cobalt. Thus, it can be concluded that the zinc and cobalt atoms distribute on a molecular level in the resulting oxide lattice19. The decomposition temperature has huge effect on the crystalline size and the surface area of the resulting sample, and the anneal temperature can be reduced finely with the heating rate decreasing23. Hence, in this paper, in order to unveil the effects of the heat temperature on the phase, morphology and catalytic properties of the products, the pyrolysis temperature was set at various values from 250 °C to 400 °C on the basis of the above TG-DTA analysis.Researches show that the pyrolytic process of precipitated precursors has a strong effect on the crystallite phase and the morphology features  of the product2O4 (JCPDS card no. 80-1543). It also can be seen that the diffraction peaks become increasingly sharp as the calcination temperature elevating, which confirms better crystallization and bigger grain size of the decomposed products with the temperature elevating. But there are diffraction peaks originating from the other phases besides the pure ZnCo2O4 phase  in Fig. XRD patterns of final products obtained by sintering the precursor at 250 °C, 300 °C, 350 °C and 400 °C are shown in Fig. 2O4 calcined at 250 °C and 400 °C are studied with FESEM, and the obtained results are shown in Fig. Surface morphologies of the as-synthesized ZnCo2O4 nano-crystals interconnected together to form porous structure. The low and high magnification TEM images of the ZnCo2O4 annealed at 400 °C are exhibited in Fig. 2O4. The porous structure might be formed by large amounts of gases slowly released from the micrometer sized oxalate particles leaving over plenty of space during the pyrolysis process19.The more information about the shape and crystallite size of the calcined products were further analysed by TEM Fig. . Figure 2O4 obtained at different calcination temperature were measured by nitrogen adsorption/desorption method at 77 K. Figure 2 adsorption-desorption isotherms and corresponding pore size distribution plots  of the as-synthesized ZnCo2O4. As can be seen in Fig. 2O4 calcined at different temperatures exhibit a type IV isotherms and H3 hysteresis loop according to the IUPAC classification, which suggest that the mesoporous structure is formed by slit-like mesopore. Those pores were produced form the decomposition of cobalt - zinc oxalate hydrate crystallization during the pyrolysis process. The structural parameters and BET specific surface areas of the as-prepared ZnCo2O4 are derived from the isotherms and tabulated in Table 2O4 calcined at 250 °C, 300 °C, 350 °C, 400 °C are determined to be 102.34, 101.36, 68.02 and 43.25 m2·g−1, respectively. The above data show that the calcination temperature has a significant impact on the specific surface areas of the products, i.e., the high calcination temperature will induce extensive growth of the spinel ZnCo2O4 and the collapse of the pore network. Thus, the specific surface area of ZnCo2O4 calcined at 400 °C dropped to 43.25 m2·g−1. From the pore size distribution curve , as shown in Fig. 2O4 rod. Characteristic peaks for Zn, Co, O and C elements can be observed in the obtained curve. The binding energy values of the major peaks are 780.5 and 795.5 eV in the Co 2p spectrum  at 1390 cm−1 were observed29. The bands at 666 and 566 cm−1 can be assigned to the metal-oxygen vibration frequency of the metal at tetrahedral clearance (Zn-O) and octahedral clearance (Co-O), indicating the formation of ZnCo2O4 spinel structure30.The FT-IR spectrum of the as-synthesized ZnCo32, the particle size of AP is studied through SEM before the thermal decomposition analysis. From the SEM images of pure AP  peak. After that, complete decomposition at 471.15 °C is revealed by the high temperature decomposition (HTD) peak13. Compared with the pure AP, an obvious difference for AP decomposition in the presence of mesoporous ZnCo2O4 rod with a mass ratio of 2% can be seen in Fig. 2O4 rod, while significant declines in the value of temperature can be seen from LTD and HTD peaks , indicating that the final decomposition temperature of AP had been significantly decreased by the additive mesoporous ZnCo2O4 rod calcined at different temperature : 329.62 °C, AP + 2%ZnCo2O4 (350 °C): 328.89 °C, AP + 2%ZnCo2O4 (300 °C): 296.71 °C, AP + 2%ZnCo2O4 (250 °C): 292.71 °C). This phenomenon considerably concurs with the exothermic peaks of the DSC curves. It can be seen that the mesoporous ZnCo2O4 rod calcined at 250 °C manifests highest catalytic activity. In addition, the catalytic activity of the mesoporous ZnCo2O4 rod are the highest than those reported in literatures39, as shown in Table The catalytic performance of as-prepared ZnCoaks Fig. . All of 2O4 rod. DSC curves of AP with a different mass ratio of mesoporous ZnCo2O4 rod calcined at 250 °C and pure AP are shown in Fig. 2O4 rod  of different weight ratios. The HTD peaks of AP are lowered with the increasing mass ratio of the mesoporous ZnCo2O4 rod additive (from 308.93 °C at 2% mass ratio to 295.98 °C at 10% mass ratio). The results reveal that better catalytic performance is achieved with mesoporous ZnCo2O4 within the mixture. TG curves of pure AP and mixtures of ZnCo2O4 and AP with different mass ratios are given in Fig. The thermal decomposition of AP is also influenced by the blend ratio of mesoporous ZnCo41. At low temperatures decomposition, the AP decomposes leading to formation of a small number of intermediate products, the main pivotal step is that the electrons transfer from perchlorate ion to ammonium ion, which would transfer to NH3 and HClO4 by dissociation and sublimation and as follows:Up to now, the thermal decomposition mechanism of AP is not yet fully understood because the decomposition process of AP is a complex solid-gas multiphase reaction process involving reactions in the solid, absorbed and gaseous phases. Several unsolved issues remain till nowet al. assumed the electrons transfer to happen on the surface not interior of crystal42. After the NH4+ accepts electron to become activated NH40, the activated ammonium radicals can decompose to ammonia and hydrogen atom:Boldyrev 40 can react with hydrogen atom and form HClO4, the HClO4 can further reacted with the H atom:The activated ClO3 can converted into ClO3− which then react with NH3 in the adsorbed gas to produce various such as NO, N2O, H2O15:As an electron absorption body, ClO3 and HClO4 by dissociation and sublimation of AP are not only absorbed on the surface of perchlorate crystal to react, but also desorbed and sublimed into the gas phase35. Because the adsorbed NH3 can not be completely oxidized by the decomposition products of HClO4 at low temperature, it overlays continually on the surface of AP. Hence, the NH3 adsorbed on the surface gets saturated, which causes cessation of the reaction and incomplete transformation of perchlorate. As the temperature rising sequentially, the reaction between NH3 and HClO4 absorbed on the surface of AP will be triggered again to produce final volatile products including HCl, H2O, NO, N2O and so on43. In the high-temperature decomposition process, the controlling step is the transformation from O2 to superoxide ion O2−, which can further react with NH3 to form N2O, NO2 and H2O44. Therefore, the high electron transfer capacity and large specific surface area of the catalyst have an important effect on the thermal decomposition of AP. According to the electron transfer mechanism, the ZnCo2O4 as p-type semiconductor materials, it has effective sites  to accept released electron from perchlorate, accompanied by the abstraction of atomic oxygen from the perchlorate ion45.−oxide represents a positive hole in the valence band of the oxide and Ooxide is an abstracted oxygen atom from oxide. The mechanism of catalytic action is based on the presence of superoxide ion O2− on the surface of the catalysts, so the ZnCo2O4 catalyst may promoted the dissociation of ClO4− species into ClO3− and O2. In addition, the transition metal ion Co3+ in the spinel structure of ZnCo2O4 has outermost d orbitals with 3d6 electronic configurations, and the d orbitals are not filled with electrons and have hole conductivity. It can easily accept the released electron from ClO4− to form Co2+ (3d7) cations, the electron was transferred to the surface of the catalyst and reacted with NH4+ to decompose into ammonia and hydrogen atom46.The intermediate products NH2O4 serves as a bridge for transferred electrons from perchlorate ions to the ammonium ions and the other transformation from O2 to superoxide ion O2−, as depicted in Fig. 13. Due to the high specific surface area and great adsorption of the mesoporous ZnCo2O4 rods, the decomposed intermediate products in the gaseous phase of AP can be facilely adsorbed on the surface of mesoporous ZnCo2O4 rods shown in Fig. 2O4 to AP can increase the contact area of the catalytic reaction, augment the number of active sites, which promotes the thermal decomposition of AP. So that the mesoporous ZnCo2O4 rod calcined at 250 °C manifests highest catalytic activity than other products, which is mainly because of its relatively highest specific surface area.In the thermal decomposition process of AP, ZnCo2O4 rod has been successfully synthesized via a controlled thermal decomposition of homogeneous complex oxalates precursor, which is no need of the assistance of soft/hard template. XRD, SEM, TEM, XPS and nitrogen adsorption/desorption have been done to systematically characterize the structural and morphological features of the as-prepared products. After calcined at given temperature with a low rate of heating, the nano-sized ZnCo2O4 crystallites connected together to form mesoporous rod. The as-prepared material ZnCo2O4 calcined at 250 °C showed much larger surface area (102.34 m2·g−1) and high catalytic activity, shifting the AP high thermal decomposition temperature downwardly to about 162.1 °C. The results suggest that the as-prepared mesoporous ZnCo2O4 rod has great catalytic properties on thermal decomposition of AP, which can be used as a promising additive in the future.In summary, mesoporous ZnCoSupplementary Information


Although indications are that a single chiral quantum anomalous Hall(QAH) edge mode might have been experimentally detected. There have been very many recent experiments which conjecture that a chiral QAH edge mode always materializes along with a pair of quasi-helical quantum spin Hall (QSH) edge modes. In this work we deal with a substantial ‘What If?’ question- in case the QSH edge modes, from which these QAH edge modes evolve, are not topologically-protected then the QAH edge modes wont be topologically-protected too and thus unfit for use in any applications. Further, as a corollary one can also ask if the topological-protection of QSH edge modes does not carry over during the evolution process to QAH edge modes then again our ‘What if?’ scenario becomes apparent. The ‘how’ of the resolution of this ‘What if?’ conundrum is the main objective of our work. We show in similar set-ups affected by disorder and inelastic scattering, transport via trivial QAH edge mode leads to quantization of Hall resistance and not that via topological QAH edge modes. This perhaps begs a substantial reinterpretation of those experiments which purported to find signatures of chiral QAH edge modes albeit in conjunction with quasi helical QSH edge modes. There have been some other quite recent experiments4 where it has been reported that QAH edge modes occur in conjunction with quasi helical quantum spin Hall(QSH) edge modes5. The latter are prone to backscattering and are nothing but QSH edge modes which occur in a trivial insulator. These experiments which “see” QAH edge modes are in fact designed out of QSH edge mode setups in a topological insulator. By applying an extra Ferromagnetic layer or otherwise, an energy gap is sought to be created between the pair of helical edge modes in a QSH sample splitting these modes away from each other and suppressing one of these leads to a single chiral QAH edge mode in a sample. However, contrary to expectation it is not just a chiral QAH mode which was seen in those experiments4 but it always comes with the additional baggage of quasi-helical QSH edge modes in the trivial phase5.Although, the experiment depicted in ref.Helical QSH edge modes from which these chiral QAH edge modes evolve not only occur in topological insulators but they also do occur in a trivial insulator. Now applying a similar technique as before or attaching a ferromagnetic layer to a trivial insulator, we can again make the trivial helical edge modes evolve into chiral QAH edge modes. But in the latter case, the chiral QAH edge mode so produced wont have a topological character and therefore this chiral QAH edge mode won’t be protected against backscattering. Now this begs the question how can one be sure of the topological character of QAH edge modes.4 is via addition of magnetic impurities or a ferromagnetic layer. This may destroy their topological character since helical QSH edge modes are susceptible to spin flip scattering in presence of magnetic impurities. In this context our work becomes relevant, since in those QAH experiments what is quite evident is that the quantization of Hall resistance is attributed to chiral topological QAH edge modes which exist in combination with quasi helical QSH edge modes. What our work reveals is that a chiral trivial QAH edge mode which exists in combination with quasi helical QSH edge modes gives the quantization of Hall resistance and not the chiral topological QAH edge mode when combined with trivial QSH edge modes. Thus a shadow of doubt creeps up regarding the interpretation of those experiments5.Another question which can crop up is, does the topological nature of the QAH edge modes which evolve from helical QSH edge modes in a topological insulator survive the evolution. This “evolution” from helical QSH to chiral QAH edge mode as has been described in refs5) and finally the case wherein a trivial QAH edge mode exists with a pair of trivial QSH edge modes (our ‘What If?’ scenario). Both the 4 terminal and 6 terminal samples are analyzed in three distinct regimes: 1. where there is no disorder and inelastic scattering- the ideal case, 2. when there is disorder but no inelastic scattering and finally 3. when both disorder and inelastic scattering are present in the sample. The disorder we consider in our sample is restricted to terminal/contacts while inelastic scattering is present inside the sample and leads to the energy equilibration of the edge modes, see refs9 for further details on energy equilibration as applied in different contexts in quantum Hall and quantum spin Hall samples.We focus specifically on 4 and 6 terminal quantum anomalous Hall samples. We distinguish three cases one in which there is just a single chiral QAH edge mode which is topological in character , the second wherein the chiral topological QAH edge mode exists alongwith a pair of trivial QSH edge modes , two terminal local resistance RT2 and finally the non-local resistance RNL for a four terminal sample with disorder and inelastic scattering and distinguish between the aforesaid three cases. In section IV we consider a six terminal sample and calculate the longitudinal resistance RL for afore-mentioned three different cases. We end our manuscript with a conclusion wherein we tabulate all the results of the 4 terminal and 6 terminal samples.The rest of the paper is organized as follows- in section II we address the situation of two terminal and three terminal samples and distinguish between three cases of chiral topological QAH edge mode, chiral topological QAH edge mode with trivial QSH edge modes and finally chiral trivial QAH edge mode with trivial QSH edge modes. In section III we calculate the Hall resistance  and Ii is the current passing through the same terminal. Tij is the transmission probability from the jth to ith probe and Gij is the conductance. N denotes the number of terminals in the system.The Landauer-Buttiker formalismV2 = 0, we derive the local (two probe) resistance The case of a single chiral QAH edge mode is represented in Fig. f0 between QAH chiral and trivial helical edge modes, will decide whether the QAH edge mode is topological (if f0 = 0) or trivial (if f0 ≠ 0). While f defines the spin-flip scattering between the trivial helical edge modes. This situation of topological chiral plus trivial helical QSH is shown in Fig. V2 = 0, we derive local (two probe) resistance f0 = 0 will give the two terminal resistance for topological QAH with trivial helical edge modes. The local resistance for topological QAH with trivial edge modes is f = f0 = 0.5 we see that trivial Here we have considered the general case, where the spin-flip scattering parameter I2 is zero. We further choose reference potential V3 = 0 which gives V2 = V1. So, local  resistance The single chiral QAH edge mode case is represented in Fig. V3 = 0 and I2 = 0 (as probe 2 is a voltage probe), we get f0 = 0) case the local (two probe) resistance f = f0 = 0.5 we see that trivial The case of a single chiral QAH edge mode with a pair of quasi-helical QSH edge modes is shown in Fig. RH = R13,24, the local (two probe) resistance RT2 = R13,13 and the non-local resistance RNL = R14,23 for various cases starting with just a single chiral QAH edge mode, then the chiral QAH edge mode with trivial quasi-helical QSH edge modes and finally the case of chiral QAH edge mode with trivial quasi-helical QSH edge modes.The four terminal sample is shown in Fig. V3 = 0, further since 2 and 4 are voltage probes, I2 = I4 = 0, we thus have V2 = V1 and V3 = V4 = 0. So, local (two probe) resistance ... Hall resistance- RNL we consider 2, 3 as voltage probes and 1, 4 as current probes, we get V2 = V3 which gives RNL = 0. Thus disorder has no effect on a single chiral QAH edge mode.Herein we consider two of the contacts  to be disordered, see Fig. Similar to before, we consider two of the contacts  are disordered, see Fig. V3 = 0 and I2 = I4 = 0, since 2 and 4 are voltage probes, we thus derive V2 = V1 and V3 = V4 = 0. So, local (two probe) resistance The currents and voltages at the contacts from 1 to 4 are related by the equations-R2 and R4 in the above expressions for T11, T12, …, T23 rest of the transmission probabilities T31 to T44 can be deduced. The transmission probabilities are calculated in this way- say T23, the transmission probability of electron from terminal 3 to 2 can be explained as the sum of paths available from 3 to 2 for one chiral topological edge mode and one pair of trivial helical edge modes. An electron in the topological edge mode coming out of probe 3 has probability zero to reach probe 2. But an electron in the trivial helical edge mode has finite probability to reach probe 2 from 3. An electron coming out of probe 3 can reach probe 2 with probability T2(1 − f), but that is just one path, it can also reach 2 with probability fR2T2(1 − f) following a second path due to spin flip scattering, similarly a third path is Herein, as before we consider two of the contacts 2 and 4 to be disordered, see Fig. V3 = 0, and since 2 and 4 are voltage probes, we derive the local (two probe) resistance in absence of disorder (R2 = R4 = 0) as RH, RT2 and RNL are too large to be reproduced here, so we will analyze them via plots, see Fig. Choosing reference potential i is the contact index from 1 to 4. The trivial helical edge modes equilibrate with other trivial helical edge modes and these equilibrate their energy to Vi by-V3 = 0, and the contact 2 and 4 to be voltage probe as before, i.e., I2 = I4 = 0, we derive local (two probe) resistance in absence of disorder as R2 = R4 = 0). The Hall resistance RH, RT2 and RNL are again large, so we will analyze them via plots, see Fig. Herein we consider the effect of both disorder and inelastic scattering on topological QAH edge modes as shown in Fig. R2 and R4 in the above equation , but that is just one path, it can also reach 2 with probability (f + f0)R2T2(1 − f − f0) following a second path due to spin flip scattering, similarly probability for a third path is T23 as in Eq.  we derive local (two probe) resistance in absence of any disorder as RH, RT2 and RNL are again too large to be reproduced here, so we will examine them via plots, see Fig. Choosing reference potential I’ is from 1 to 4). The contacts 2 and 4 are disordered as in the previous case. The currents and voltages at the contacts from 1 to 4 are related by the equations-Vi are written in the supplementary material (section I. A).Herein, we consider the trivial QAH edge modes with both disorder and inelastic scattering as shown in Fig. V3 = 0, and as before I2 = I4 = 0 (these are voltage probes), we derive the local (two probe) resistance RH, RT2 and RNL are large, so again we will analyze them via plots as in Fig. Choosing reference potential RH(↑) = −RH(↓) and RNL = 0. We see that RNL(↑) ≠ RNL(↓) for topological QAH  while this isn’t case for trivial QAH  edge modes, see Fig. 2 have been interpreted as in ref.5.In Table RL. For a single chiral QAH edge mode RL = 0, but the experiments4 revealed a finite longitudinal resistance. This result prompted the interpretation of the experiments4 as seeing not just a chiral QAH edge mode but in addition also a pair of quasi-helical QSH edge modes5. Since a non zero RL is the hallmark of helical QSH edge modes. Here we probe further by comparing as in sections before the three cases and try to find out if a topological QAH edge mode or a trivial QAH edge mode occurring with quasi-helical edge modes results in a non-zero RL.In this section we analyze a six terminal QAH bar, we especially focus on the longitudinal resistance V4 = 0 and I2 = I3 = I5 = I6 = 0 (as these are voltage probes), we get Herein we consider two of the contacts 1 and 4 to be disordered. The relations between the currents and voltages at the various terminals can be derived from the conductance matrix below:Herein we consider two of the contacts  to be disordered. The electrons incoming from probe 6 with energy The currents and voltages at the contacts from 1 to 4 are related by the equations-V4 = 0 and I2 = I4 = I5 = I6 = 0 (as these are voltage probes), we thus derive Choosing the reference potential As before we consider two of the contacts 1 and 4 to be disordered, see Fig. R1 with R4 in the above equation rest of the transmission probabilities T31 to T56 can be deduced. Choosing reference potential V4 = 0, and I2 = I3 = I5 = I6 = 0 (as these are voltage probes), we derive longitudinal resistance Replacing A = 1 − R1f and C = 1 − R4f, where the potential Vi by-Vi are mentioned in the supplementary material (section I. B) .Herein we consider the effect of disorder and inelastic scattering on the various resistances for the sample as shown in Fig. V4 = 0, and since 2, 3, 5, 6 are voltage probes, I2 = I3 = I5 = I6 = 0 we derive longitudinal resistance RL in presence of both disorder and inelastic scattering is large so we analyze them via plots as in Fig. Choosing reference potential R1 with R4 in the above equation rest of the transmission probabilities T31 to T66 can be deduced.Herein again we consider two of the contacts 1 and 4 to be disordered, see Fig. V4 = 0, and I2 = I4 = I5 = I6 = 0  we derive longitudinal resistance RL are too large to reproduce here, so we will analyze them via plots as in Fig. Choosing reference potential V′i (`i′ is from 1 to 6). The contacts 1 and 4 are disordered as in the previous case. The currents and voltages at the contacts from 1 to 6 are related by the equations-Herein we consider the trivial QAH edge modes with disorder and inelastic scattering, as shown in Fig. Vi are given in the supplementary material (section I. C) . Choosing reference potential V4 = 0, and as before I2 = I3 = I5 = I6 = 0 (these are voltage probes), we derive the longitudinal resistance The relation between potentials 5. In Table RL(↑) = RL(↓) for trivial QAH edge mode which does not hold for a topological QAH edge mode with quasi helical QSH edge modes. This response of the trivial QAH edge mode is again in line with what was experimentally seen.The expression for longitudinal resistance in the general case of arbitrary disorder are quite large so we examine it via plots as in Fig. 5 than the topological QAH edge mode with trivial quasi-helical QSH edge modes is. This implies a reevaluation of the consensus regarding those quantum anomalous Hall experiments4. Perhaps, something else is happening and maybe these are not true chiral quantum anomalous Hall edge modes which were seen, but rather what could be described as chiral QAH edge modes.We conclude by analyzing the Tables As this is an analytical work, all the data needed to plot the Figs Supplementary Information