text:
\section{Introduction}
Monolayer (1L) MoS$_2$\xspace, a two dimensional (2D) transition metal dichalcogenide (TMD), has become one of the most studied materials in the last decade for its unique electronic and optical properties\cite{mak1}, among which spin-valley coupling\cite{mak2,VP,VP1} and direct band gap at the K-points of the Brillouin zone\cite{mak1,bs} are the most important ones. Because of its direct band gap of $\sim$1.9 eV and large exciton binding energy, 1L-MoS$_2$\xspace offers high luminescence yield in the visible range\cite{mak1,bs,VP1}. Atomically thin size and flexible nature make the material a potential candidate for 3D integrated and flexible electronics of the future \cite{3D_flex}. Furthermore, there are multiple theoretical proposals that predict ferromagnetic behaviour even above room temperature in 1L-MoS$_2$\xspace, when Mo is substituted by certain transition metal ions. Such a two dimensional dilute magnetic semiconductor (2D-DMS) has potential for injecting spin polarized excitons and carriers for future 2D spintronic and valleytronic devices\cite{magnet1,magnet2,magnet3,VP2}. Achieving controllable doping is one of the important requirements for any technology application of the material. However, doping of 1L-MoS$_2$\xspace is always onerous because of the strong bonding between Mo and S, which makes it almost inert for any substitutional reaction to take place\cite{XPS1,bs}. This has become even more challenging as the unambiguous detection of doping is extremely difficult in an atomically thin layer. Even though, there are efforts by various groups to incorporate different elements in 1L-MoS$_2$\xspace using varied  means, such as chemical doping\cite{ch1,ch2,ch3},  physisorption\cite{therm,ad2,ad3,ad4}, metal evaporation\cite{evap1,raman2} and ion irradiation\cite{ion1,raman1}, a method for stable and sustainable doping of the layer is yet to be established.  Previously, Zhang \textit{et al.} have reported up to 2 atomic percentage of  Manganese (Mn) doping in 1L-MoS$_2$\xspace by vapor phase deposition technique\cite{XPS2}. Mn incorporation has also been tried using Mn ion irradiation technique\cite{raman1}, which has been found to result in a large amount of structural defects in the lattice. More studies are needed to understand the mechanism of incorporation of Mn and other metals in 1L-MoS$_2$\xspace. It is also important  to investigate how the incorporation can influence the electrical, optical and magnetic properties of the layer.

Here, we present a detailed study of manganese doping in 1L-MoS$_2$\xspace through pulsed laser deposition (PLD) technique and its influence on the optical properties of the films, which are grown on sapphire substrates by microcavity based chemical vapor deposition (CVD) technique\cite{pkm,tailoring}. Our investigation using high resolution X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), Raman and PL spectroscopy reveals two clearly distinct regimes of Mn-incorporation in 1L-MoS$_2$\xspace. When the deposition amount is below a certain critical value, Mn islands covering large areas are found to be formed on MoS$_2$\xspace layers and substitution of Mo ions by Mn are observed. As the deposition amount goes beyond the critical value, dewetting leads to the formation of Mn-droplets on MoS$_2$\xspace layers. In this regime, substitutional incorporation of Mn is decreased, while the disorder in the lattice increases with the deposition time.  PL study shows that Mn-deposition strongly influence the excitonic recombination dynamics by opening up non-radiative recombination channels. Both the density and the barrier height for such recombinations are found to increase  with the amount of Mn-deposited.  These non-radiative transitions are attributed to certain Mo related defects.

Strictly monolayer (1L) MoS$_2$\xspace films were grown on \textit{c}-sapphire substrates by the microcavity based CVD route where molybdenum trioxide (MoO$_3$) and sulfur (S) powders were used as precursors and argon (Ar) was used as the carrier gas.  More details about the growth can be found elsewhere\cite{pkm,tailoring}. Manganese (Mn) was deposited on these layers using pulsed laser deposition (PLD) technique where a polycrystalline Mn pellet was used as target. The pellet was prepared from 99.99\% pure Mn powder by arc melting technique in Ar (5N pure) atmosphere. Substrate (1L-MoS$_2$\xspace/sapphire) temperature was maintained at 300$$^{\small{\text{o}}}$$C during the deposition. Distance between the Mn-target and the sample was kept at 5 cm for all depositions. A KrF excimer laser of wavelength 248 nm and pulse width of 25 ns was used for ablating the target. Deposition amount was varied by varying the laser pulse count. While the pulse frequency and the energy density of the laser were maintained at 1 Hz and 1.4 J/cm$^2$, respectively, for all cases. Base pressure inside the PLD chamber was kept at $\sim$ 8.0$\times$10$^{-5}$ mbar. During the growth, 5N argon gas was introduced inside the chamber to maintain a pressure of  $\sim$ 8.0$\times$10$^{-2}$ mbar inside the chamber. 1L-MoS$_2$\xspace layers were coated with Mn for different numbers of laser pulse count ranging from 1 to 10. These samples are labeled by alphabets A to D. In  another sample (sample E), Mn deposition was carried out at the base pressure  ($\sim$ 8.0$\times$10$^{-5}$ mbar) for 10 counts of laser pulses.  Mn was also deposited on a bare \textit{c}-plane sapphire substrate (without MoS$_2$\xspace) for five counts of  laser pulses at 8.0$\times$10$^{-2}$ mbar of Ar-pressure, keeping other parameters constant. This sample is named as sample F. One as-grown 1L-MoS$_2$\xspace sample R, was used as the reference standard for all the measurements. Mn deposition parameters for these samples, are listed in Tab.~\ref{tab1}.

X-ray photoelectron spectroscopy on these samples were performed in Axis Supra spectrometer from Kratos Analytical, UK, where monochromatic Al K${_\alpha}$ x-ray (h$\nu$ = 1486.6~eV) was used as excitation source. Energy axis of the recorded spectra, was calibrated with respect to the binding energy of carbon 1$s$ peak at 284.7~eV. Morphology of the samples were examined by atomic force microscopy (AFM) in tapping mode using  Bruker, NanoScope-IV system. Raman spectra on these monolayers were recorded at room temperature in back scattering geometry with 532~nm diode laser using Horiba Jobin Yvon HR800 confocal Raman spectrometer. A closed-cycle helium refrigerator was used to carry out temperature dependent (from 20~K to 300~K) photoluminescence (PL) measurements, where 532~nm diode laser was used as the excitation source and the spectra were recorded by a 0.55~m focal length monochromator attached with a Peltier cooled CCD camera.

\begin{center}
	\begin{table}
		\caption{Mn deposition parameters for different samples. Substrate temperature, distance between target and substrate, frequency and energy density of the laser are maintained at 300$$^{\small{\text{o}}}$$C, 5~cm, 1~HZ and 1.4 J/cm$^2$, respectively for each deposition.}
		\begin{tabular}{c| c |c |c}

			\hline
			{\small Sample} & {\small Pressure} & {\small Ar flow rate} & {\small Laser Pulse count} \\
			& {\small(mbar)} & {\small(sccm)} &  \\
			\hline
			\hline
			~R~& -- & -- & 0  \\

			~A~& {\small 8.0$\times$10$^{-2}$} & 40 & 1  \\

			~B~& {\small 8.0$\times$10$^{-2}$} & 40 & 5  \\

			~C~& {\small 8.0$\times$10$^{-2}$} & 40 & 7  \\

			~D~& {\small 8.0$\times$10$^{-2}$} & 40 & 10  \\

			~E~& {\small 8.0$\times$10$^{-5}$} & 0 & 10  \\

			~F (Sapphire)~& {\small 8.0$\times$10$^{-2}$} & 40 & 5  \\
			\hline
		\end{tabular}
		\label{tab1}
	\end{table}
\end{center}

\begin{figure}
	\includegraphics[scale=1.8]{fig1.pdf}
	\caption{AFM images for samples (a) R, (b) A, (c) B, (d) C and (e) D. Inset of (a) shows the height distribution along a line drawn across the boundary. Inset of (b) shows the line scan profile over a Mn-deposited island on MoS$_2$\xspace film. (f) Average roughness of MoS$_2$\xspace (black) and sapphire (red) regions as functions of pulse count. Filled symbols represent sample E.}
	\label{fig1}
\end{figure}

Fig.~\ref{fig1} shows AFM images for the MoS$_2$\xspace samples coated with Mn for different counts of laser pulses. In all the images, MoS$_2$\xspace layers can clearly be distinguished from the bare sapphire surface from the height contrast. As shown in Fig.~\ref{fig1}(a), average step-height for the reference sample R has been found to be $\sim$ 0.7~nm, which is expected for a monolayer MoS$_2$\xspace\cite{pkm,tailoring}. In fact, it should be noted that average step-height for all samples before Mn deposition is found to be $\sim$ 0.7~nm confirming monolayer nature of MoS$_2$\xspace in all cases. As shown in Fig.~\ref{fig1}(b), deposition of several flat island type structures with average height of $\sim$ 0.9~nm on the MoS$_2$\xspace regions is evident in sample A, where only one laser pulse is used for Mn deposition. In case of sample B, where Mn is deposited for 5 laser pulses, flat island type structures could not be distinguished clearly on MoS$_2$\xspace regions, instead, several spherical structures are visible on both MoS$_2$\xspace and bare sapphire regions as shown in Fig.~\ref{fig1}(c). Density and size of the spherical droplets clearly increase as one moves from sample C [Fig.~\ref{fig1}(d)]  to D [Fig.~\ref{fig1}(e)], meaning as the laser pulse count for Mn deposition increases (see supplementary Fig.~S1\cite{sup}). Noticeably, in these samples, droplets on bare sapphire are bigger than those on MoS$_2$\xspace regions. It is noteworthy that similar type of droplet formation has also been observed in case of sample E (see supplementary Fig.~S1\cite{sup}). Fig.~\ref{fig1}(f) plots the average value of the roughness of MoS$_2$\xspace (black symbols) and sapphire (red symbols) regions as functions of laser pulse count. Evidently, variation of roughness on MoS$_2$\xspace regions shows a clear increase beyond 5 counts of laser pulse. Note that the average roughness on sapphire regions also shows a sudden jump beyond 5 pulses. These are consistent with the sudden change in the morphology observed in the AFM images and seem to suggest that Mn-deposition for sufficiently less number of laser pulses results in a thin Mn-layer. The strain/surface energy developed in such a thin metal film may not be sufficient to change its morphology when the sample is cooled to room temperature after deposition. It is plausible that in case of the samples, where Mn-deposition is performed for the number of pulses more than a certain value, thickness of the deposited Mn-film crosses a critical mark and the strain/surface energy developed in the layer becomes so high that to minimise energy, the film breaks into Mn-droplets when the sample is brought to the room temperature. This point has been discussed again later (in Fig.~\ref{fig5}).

\begin{figure*}[t]
	\includegraphics[scale=1]{fig2.pdf}
	\caption{XPS core level spectra for Mo $3d$ and S $2s$ levels  in case of samples (a) R, (d) A and (g) D; for S $2p$ levels  in case of the samples (b) R, (e) A and (h) D; for Mn $2p$ levels in case of  the samples (c) F,  (f) A and (i) D. (j) Ratio of the contributions from S-Mn and S-Mo bonds obtained by de-convoluting S~$2s$ (black symbols) and S~$2p_{3/2}$ (red symbols) spectra for these samples versus pulse count. (k) Ratio of the contribution from Mo atoms and Mo-S bonds obtained from the Mo~$3d_{5/2}$ profiles versus pulse count. In (j) and (k), filled symbols represent sample E.}
	\label{fig2}
\end{figure*}

In panels (a-i) of Fig.~\ref{fig2}, XPS spectra for Mo~$3d$, S~$2s$, S~$2p$ and Mn~$2p$ levels for different samples are compared. These spectra are deconvoluted with Voigt functions as mentioned in earlier references\cite{XPS,XPS1}. In all panels, red lines represent the fitted profiles. For the reference sample R [panel (a)], S~$2s$ (orange line) peak  appears (at 226.7~eV) along with Mo~$3d$ doublets. Mo~$3d$ related features can be deconvoluted into two doublets. While, the position of the first doublet, which peaks at 229.3~eV (Mo $3d_{5/2}$) and 232.4~eV (Mo $3d_{3/2}$) [blue line], matches quite well with Mo-S bonds in MoS$_2$\xspace\cite{XPS,XPS1}, the second doublet appearing at  232.6~eV (Mo~$3d_{5/2}$), 235.7~eV (Mo~$3d_{3/2}$) [bright green line] can be assigned to the Mo-O bonds of the unreacted MoO$_3$ particles on the surface\cite{XPS}. Panel (b) shows the S $2p$ features for the reference sample. Deconvolution of the feature results in two peaks centred at  162.2~eV and  163.4~eV, (blue line) which are associated with S~$2p_{3/2}$  and S~$2p_{1/2}$  lines, respectively,  of MoS$_2$\xspace\cite{XPS,XPS1}. Fig.~\ref{fig2}(c) shows Mn $2p$ features for sample F, which is a bare sapphire substrate coated with manganese. The feature can be deconvoluted into two doublets. While the doublet (olive line) appearing at  639.2~eV (Mn~$2p_{3/2}$) and  651.0~eV  (Mn~$2p_{1/2}$) can be attributed to Mn-metal\cite{XPS2,XPS3}, position of the higher energy doublet (purple line) appearing at 644.0~eV (Mn~$2p_{3/2}$) and 655.6~eV (Mn~$2p_{1/2}$) matches quite well with those of Mn-O bonds of manganese oxide (MnO$_x$)\cite{XPS2,XPS5}. Quite evidently, Mo~$3d$, S~$2s$ as well as S~$2p$ features appears much broader for sample A as compared to the reference sample as can be seen in  Fig.~\ref{fig2}(d) and (e). This is a clear indication of the change of the bonding states of both Mo and S ions in this sample. Deconvolution of these features results in the formation of one additional  Mo $3d$ doublet peaking at  227.9~eV (Mo~$3d_{5/2}$)  and  231.0~eV (Mo~$3d_{3/2}$) [magenta line in Fig.~\ref{fig2}(d)], one extra S $2s$ peak (at 225.5~eV) [bottle green line in Fig.~\ref{fig2}(d)] and two extra S $2p$  doublets; one peaking at 161.4~eV (S~$2p_{3/2}$)  and 162.6~eV (S~$2p_{1/2}$) [bottle green line in Fig.~\ref{fig2} (e)] and the other, peaking at 159.9~eV (S~$2p_{3/2}$)  and 161.0~eV (S~$2p_{1/2}$)[brown line in Fig.~\ref{fig2}(e)]. Position of the new Mo~$3d$ doublet [magenta line in Fig.~\ref{fig2} (d)] matches well with neutral Mo atoms\cite{XPS1}. Out of the two extra S~$2p$ doublets, the bottle green line of Fig.~\ref{fig2}(e) can be assigned to S-Mn bond formation and the brown line of Fig.~\ref{fig2}(e) can be attributed to surface adsorbed S atoms\cite{XPS1}. The extra S~$2s$ peak at 225.5~eV [bottle green line in Fig.~\ref{fig2}(d)] can also be attributed S-Mn bonds\cite{XPS4,XPS2}. Interestingly for sample D, Mo~$3d$, S~$2s$ and S~$2p$ profiles are similar to those obtained for the reference sample as shown in Fig.~\ref{fig2}(g) and (h). Deconvolution of these profiles results in a much weaker contributions from the S-Mn bonds and neutral Mo atoms as compared to what has been found in case of sample A. These findings point towards the fact that for sample A, Mn atoms partially substitute Mo ions of MoS$_2$\xspace to form Mn-S bonds and the knocked out Mo ions settle as neutral Mo atoms. While, in sample D, such substitution is very weakly present. It should be noted that like in sample A, significant presence of additional Mo and S related features could also be found in sample B, where Mn is deposited for 5 laser pulses. Similar to the case of sample D, very weak sign of Mo-substitution by Mn-ions could be found for sample C and E, where laser pulse counts of 7 and 10, respectively, are used for Mn deposition. Note that Mo~$3d$, S~$2s$ and S~$2p$ spectra for all the Mn deposited MoS$_2$\xspace samples are shown in supplementary figure S2\cite{sup}. Fig.~\ref{fig2}(f) and (i) show the  Mn $2p$ features for sample A and D, respectively. It is interesting to note that for both the cases, features are very similar to what has been found for sample F [Fig.~\ref{fig2}(c)] showing the contributions only from Mn-metal and Mn-O bonds. In fact, no sign of Mn-S bond formation could be found from Mn-spectra for any of the Mn-coated MoS$_2$\xspace samples. Note that in all these samples, a significant portion of the sapphire surface remains uncovered by 1L-MoS$_2$\xspace. Perhaps, the signal due to Mn-S bond from the Mn-coated MoS$_2$\xspace layers is overwhelmed by the signal arising from the Mn-coated sapphire regions for all these samples. The ratio of the contributions from S-Mn and S-Mo bonds can be obtained in two ways; by deconvoluting (i) S~$2s$ and (ii) S~$2p_{3/2}$ spectra. These results are plotted as a function of the laser pulse count in Fig.~\ref{fig2}(j). Evidently, both the profiles show very similar trend. Similarly, the ratio of the contribution from Mo atoms and Mo-S bonds can be achieved from the Mo~$3d_{5/2}$ peaks, which is plotted versus pulse count in Fig.~\ref{fig2}(k). Interestingly, no signature of knocked out Mo atoms could be found in samples deposited with more than 5 pulse counts. It is quite evident from these plots that substantial amount of Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace layers is achieved only for pulse counts of 1 and 5. Incorporation sharply decreases as the pulse count  goes beyond 5.

\begin{figure}
	\includegraphics[scale=1.8]{fig3.pdf}
	\caption{(a) Room temperature Raman spectra for different Mn-deposited samples in the frequency range containing characteristic first order E$^1_{2g}$ and A$_{1g}$ peaks of MoS$_2$\xspace. (b) Deconvolution of the spectrum with Lorentzian functions for sample D. (c) Peak position and (d) FWHM  of the Lorentzians associated with E$^1_{2g}$ (orange) and A$_{1g}$ (blue) peaks versus pulse count. Filled symbols represent sample E. Solid lines in (C) and (d) are guides to the eyes.}
	\label{fig3}
\end{figure}

Fig.~\ref{fig3}(a) compares room temperature micro-Raman spectra for these samples in the frequency range, where characteristic in-plane E$^1_{2g}$ and out-of-plane A$_{1g}$ vibrational modes for the zone-centre phonons are observed at $\sim$ 385 and 405~cm$^{-1}$, respectively. These spectra are deconvoluted with Lorentzian functions as shown in  the Fig.~\ref{fig3}(b) for sample D as an example. Apart from the  dominant E$^1_{2g}$  (orange line) and A$_{1g}$ (blue line) features, several other phonon modes associated with M and K points of the Brillouin zone (BZ) [magenta lines] can also be seen. Note that these non-zone-centre phonons become visible due to the relaxation of the fundamental Raman selection rule (phonon wave vector $q$ $\approx$ 0) in the presence of disorder in the MoS$_2$ lattice (see supplementary Fig.~S3\cite{sup})\cite{raman1,raman2, raman4,raman5}. There are also a couple of other features (brown lines) in the spectrum, which can be identified as the Raman lines arising from the sapphire substrate (see the supplementary Fig.~S3\cite{sup}). Fig.~\ref{fig3}(c) and (d) plot the peak position and full width at half maximum (FWHM) of the Lorentzians associated with E$^1_{2g}$ (orange) and A$_{1g}$ (blue) peaks as functions of the number of laser pulses used for Mn-deposition in these samples. It is evident from these figures that with the increase of pulse count,  A$_{1g}$  peak shows a tendency of shift towards higher wavenumbers.  On the other hand,  position of E$^1_{2g}$ peak does not vary much with pulse count upto 7, beyond which it reduces. It should be noted that electron concentration in the conduction band has a strong influence on the coupling between electrons and A$_{1g}$ phonons in 1L-MoS$_2$\xspace\cite{raman3}. Increase of electron concentration in the layer is found to result in the softening of frequency as well as broadening of the A$_{1g}$ peak, while E$^1_{2g}$ feature hardly shows any change\cite{raman1,raman2,raman3}. Note that the FWHM of the A$_{1g}$ peak is less than that of the reference sample for samples A and B, where the pulse counts are 1 and 5, respectively. This may suggest a reduction of electron concentration in the two MoS$_2$\xspace samples, which can further imply that Mn-ions incorporate in MoS$_2$\xspace when the number of laser pulses used for Mn deposition is sufficiently less.  This observation is also consistent with the results of the XPS study shown in Fig.~\ref{fig2}. Increasing trend of the line-width for both A$_{1g}$ and E$^1_{2g}$ peaks as well as hardening and softening of A$_{1g}$ and E$^1_{2g}$ modes, respectively, for pulse count beyond 7 may indicate a rise of disorder in the layer with the increase of Mn-deposition time.

\begin{figure}
	\includegraphics[scale=1.8]{fig4.pdf}
	\caption{ (a) Lower energy portion of the room temperature Raman spectra for various Mn-deposited samples. (b) Deconvolution of the structure with Lorentzian functions for sample D. Note that the most dominant peak at $\sim$ 227~cm$^{-1}$  has been assigned to longitudinal acoustic phonon mode at \textit{M} point [LA(\textit{M})]. (c) Intensity ratio between LA(\textit{M}) and A$_{1g}$ peaks [I(LA)/I(A$_{1g}$)] as well as LA(\textit{M}) and E$^1_{2g}$ peaks [I(LA)/I(E$^1_{2g}$)] versus pulse count. Filled symbols represent sample E. Solid lines provide the visual guidance.}
	\label{fig4}
\end{figure}

Fig.~\ref{fig4}(a) highlights the lower energy portion of the Raman spectra for these samples, where certain zone edge phonons give rise to broad structure with multiple features due to the presence of disorder\cite{raman1,raman2,raman4,raman5}. Fig.~\ref{fig4}(b)  shows the Lorentzian deconvolution of the structure for sample D, where the most dominant peak at $\sim$ 227~cm$^{-1}$ (violet line) has previously been assigned to longitudinal acoustic phonon mode at \textit{M} point [LA(\textit{M})]. This peak has also been treated as a measure of disorder in 1L-MoS$_2$\xspace films\cite{raman1} . Please see supplementary Fig.~S3\cite{sup} for identification of other features in the structure. It should be noted that the spectra shown in all the above panels are normalized with respect to the intensity of A$_{1g}$ peak.  From Fig.~\ref{fig4}(a) it is evident that all the disorder related peaks  enhance with the laser pulse count suggesting a gradual increase of disorder with the amount of Mn deposition in these samples. In case of sample E, where Mn-deposition is carried out for 10 pulse-counts in vacuum (switching off the Ar supply), peak-shift and peak-width for both E$^1_{2g}$ and A$_{1g}$ peaks are dramatically higher than the other samples suggesting much higher degree of disorder in this sample [see Fig.~\ref{fig3}(c) and (d)].  Intensity ratios between LA(\textit{M}) and A$_{1g}$ peaks [I(LA)/I(A$_{1g}$)] (blue line) as well as  LA(\textit{M}) and E$^1_{2g}$ peaks [I(LA)/I(E$^1_{2g}$)] (orange line), are plotted in Fig.~\ref{fig4}(f). A gradual rise of disorder with increasing Mn-deposition time is quite evident from the figure. This observation is in line with the results shown in Fig.~\ref{fig3}.

\begin{figure}
	\includegraphics[scale=1.8]{fig5.pdf}
	\caption{(a) Photoluminescence (PL) spectra  recorded at 100~K for various Mn-deposited samples. (b) Log of the integrated intensity associated with A-exciton [$I(A)$] versus inverse of temperature. (c) Thermal energy barrier $E_n$ and $r_{n0}/r$ for non-radiative recombination, which are obtained from the fitting as described in the text, versus pulse count. (d) Logarithm of the intensity of the BL feature as a function of the inverse of temperature for different Mn-deposited samples.}
	\label{fig5}
\end{figure}

Fig.~\ref{fig5}(a) compares photoluminescence (PL) spectra  recorded at 100~K for all samples. Each spectrum is featured by a band-edge exciton/trion (A/A$^-$) peak at $\sim$ 1.9~eV and broad luminescence (BL) peak at $\sim$ 1.7~eV\cite{pkm,tailoring,nihit1}. Deconvolution of these PL spectra by a set of six Gaussian peaks [as shown in the supplementary Fig.~S4(a)\cite{sup}] provides a good fit\cite{tailoring}. Band edge feature is consisting of two Gaussians representing free A-excitons (A) and A-trions (A$^-$), while the BL feature can be fitted with two Gaussian functions centred at 1.71~eV ($D_{hi}$) and 1.6~eV ($D_{lo}$).  $D_{hi}$ and $D_{lo}$ transitions are found to be excitonic in nature and attributed to sulfur vacancy related defects\cite{tailoring}. Note that two narrow Gaussians are used to take into account the sharp feature at $\sim$ 1.78~eV that stems from the sapphire substrate. Temperature dependent PL spectra for all samples are recorded from 20 to 300~K and deconvoluted with six Gaussians as described before. Log of the intensity associated with A-exciton [$I(A)$] is plotted as a function of inverse of temperature in fig.~\ref{fig5}(b), for various samples. In case of the reference sample R (orange), $I(A)$ increases monotonically with temperature. This tendency indicates that the excitons are activated from a bound to free sate as temperature increases. For samples deposited with 1, 5 and 7 pulses of Mn, $I(A)$ shows similar increasing tendency upto a certain temperature, beyond which it shows a reduction with rising temperature. This fall  becomes increasingly steeper  as the pulse count increases from 1 to 7. This result suggests that Mn-deposition might be introducing a non-radiative recombination channel for the A-excitons, which need to overcome a thermal barrier to get access to the non-radiative traps. Above finding further implies that the barrier height must be increasing with the amount of Mn deposited. In sample D, where Mn pulse count is 10, $I(A)$ become almost  temperature independent. In case of sample E, $I(A)$ shows similar trend as sample D [see supplementary Fig.~S4(b)\cite{sup}].  We believe that  the thermal energy barrier for the non-radiative recombination is so high for the two samples that one needs to go to higher than room temperature to see the fall of A-exciton intensity. It should be noted that in samples D and E, the density of Mn-deposition induced disorder in MoS$_2$\xspace is expected to be much more than the rest of the samples as discussed in Fig.~\ref{fig4}. At the steady state condition, density of free excitons ($n_A$)  and excitons bound to $D$ type of defect ($n_D$) should satisfy the following equations.
\begin{eqnarray}
	\frac{dn_A}{dt} &=& G - r n_A - r_n n_A -\beta (N_D - n_D) n_A + \alpha n_D = 0 \nonumber \\
	\frac{dn_D}{dt} &=& \beta (N_D - n_D) n_A -  \alpha n_D - r_D n_D  = 0
	\label{Eq1}
\end{eqnarray}
Where, $G$, $r$ and $r_n$ are , respectively, the rate of generation, radiative and non-radiative recombinations for the free excitons. $N_D$ is the concentration of $D$ type of defect traps and $\beta$ ($\alpha$) is the coefficient of transition of an exciton from free (defect bound) state to defect bound (free) state. $r_D$ is the recombination rate of defect bound excitons. At thermal equilibrium, $\alpha$ = $N_D$ $\beta$ $e^{-E_D/k_B T}$. $r_n$ = $r_{n0}$ $e^{-E_n/k_B T}$, where $E_D$ and $E_n$ are activation energy of excitons from defect bound to free state and the height of the thermal energy barrier for non-radiative recombination, respectively. $k_B$ is the Boltzmann constant and $r_{n0}$ is $r_n$ at $T$ = 0~K. Under the assumptions of $N_D$ $\gg$ $n_D$ and $\beta$$N_D$ $\gg$ $r_D$, the expressions for PL intensities associated with free excitons [$I(A)$= $r$ $n_A$] and the bound excitons[$I(D)$= $r_D$ $n_D$] can be obtained from eq.~\ref{Eq1} as,
\begin{eqnarray}
	I(A) = \frac{G}{1+ \frac{r_D}{r} \exp{(\frac{E_D}{k_BT})} + \frac{r_{n0}}{r} \exp{(-\frac{E_n}{k_BT})} }
	\label{Eq2}\\
	I(D) = \frac{\frac{G r_D}{r} \exp{(\frac{E_D}{k_BT})} }{1+ \frac{r_D}{r} \exp{(\frac{E_D}{k_BT})} + \frac{r_{n0}}{r} \exp{(-\frac{E_n}{k_BT})} }
	\label{Eq3}
\end{eqnarray}

Experimentally obtained temperature variations of $I(A)$ for all samples have been fitted by eq.~\ref{Eq2}. The resultant profiles are the red solid curves shown in fig.~\ref{fig5}(b). The fit for the reference sample R returns $E_D$ = 12.4~meV and $E_n$ = 0. $E_D$ value is kept at 12.4 meV while $E_n$ is varied to fit the data for rest of the Mn-deposited samples. Obtained values for $E_n$ and $r_{n0}/r$ are plotted as functions of pulse count in fig.~\ref{fig5}(c). It can be seen that both the quantities gradually increase with the pulse count. Note that the increase of $r_{n0}/r$ could suggest the enhancement of the density of non-radiative centres in the layer. As discussed before, $E_n$ and $r_{n0}/r$ for sample D and E can not be estimated from this study.  Fig.~\ref{fig5}(d) plots logarithm of combined integrated intensities of $D_{hi}$ and $D_{lo}$ [$I(D_{hi})$ + $I(D_{lo})$], i.e., total intensity of BL feature as a function of the inverse of temperature for sample R, A, D and E. Please refer to the supplementary Fig.~S4(c)\cite{sup} for $ln[I(D_{hi})+I(D_{lo})]$ vs $1/T$ plots for samples B and C. It can be seen that like in the reference sample, in all Mn-deposited samples, intensity of this feature increases and shows a tendency of saturation with the reduction of temperature. This result implies that the Mn disorder induced non-radiative recombination channel, which modifies the variation of intensity of A-exciton with temperature in these samples, does not affect the dynamics of the sulfur deficiency bound excitonic features.  It should be noted that our attempts to simulate the data of Fig.~\ref{fig5}(d) using eq.~\ref{Eq3} with the parameters obtained from the fitting of the data using eq.~\ref{Eq2} as shown in Fig.~\ref{fig5}(b) have failed for every sample. This suggests that BL band has nothing to do with the defect D, which binds  A-excitons in our model (eq.~\ref{Eq1}). This further means that BL transition does not directly involve A-excitons. To find the binding energies associated with $D_{hi}$ and $D_{lo}$ , experimental data of [$I(D_{hi})$ + $I(D_{lo})$] vs $T$ are fitted with the equation below as described in our previous report\cite{tailoring}.
\begin{eqnarray}
	I(T)=\frac{I_0}{1+\sum \gamma_i T^{3/2} \exp(-E_{a_i}/k_B T)}
	\label{Eq4}
\end{eqnarray}
where, $E_{a_i}$ and $\gamma_i$ are the binding energy and a measure of capture cross section of $i$-th type of sulfur-vacancy related defects. Terms under  summation in the denominator take into account the contributions from the of two defect bands, $D_{lo}$ and $D_{hi}$. Temperature variation of the intensity of the defect bound BL feature is fitted with eq.~\ref{Eq4} and the resultant profiles are represented by red solid curves in fig.~\ref{fig5}(d).  Values of $E_{D_{hi}}$ ($\sim$ 100~meV) and $E_{D_{lo}}$ ($\sim$ 5~meV) are found to be almost the same for all the samples. This confirms that Mn deposition does not affect the dynamics of sulfur deficiency related BL transition in these samples.

\begin{figure}
	\includegraphics[scale=1]{fig6.pdf}
	\caption{Schematic depiction of Mn-deposition on MoS$_2$\xspace. Island regime: (a) when Mn-deposition time is sufficiently less, flat-top Mn islands with high area coverage are present on top of MoS$_2$\xspace even after post-growth cooling of the sample to room temperature.  Droplet regime: when Mn-deposition time is sufficiently long, (b) Mn-layer with thickness more than the critical thickness is deposited on MoS$_2$\xspace. (c) However, upon post-growth-cooling of the sample to the room temperature, Mn-droplets are formed. Mo-substitution by Mn as well as the formation of Mo-vacancy (\textit{V$_{Mo}$}) and distorted Mo-S bonds are also presented in these panels.}
	\label{fig6}
\end{figure}

From  Fig.~\ref{fig1} and \ref{fig2}  it appears that in the regime when Mn is deposited on MoS$_2$\xspace forming large flat-top islands (laser pulse count upto 5),  Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace is more than in the regime  when Mn-droplets are formed. It is plausible that since the area coverage in the island regime is more than in the droplet regime, chance of substitution of Mo-ions by Mn-ions is expected to be more in the former case as shown schematically in Fig.~\ref{fig6}.  We believe that even in the droplet regime, Mn-layer covering a large area is deposited on MoS$_2$\xspace at the deposition temperature. A good number of Mo-substitutional incorporation of Mn is thus expected at the deposition temperature in this regime as well. However, during the post-growth-cooling stage droplets are formed in this regime. The droplet formation process give rises to either Mo-vacancies (\textit{V$_{Mo}$}) or distorted Mo-S bonds as Mn-ions are withdrawn from the substitutional sites at this stage (see Fig.~\ref{fig6}). This  can explain why the disorder in MoS$_2$\xspace increases with the pulse count, especially in the droplet regime as observed in Fig.~\ref{fig4}.  Note that reconnection of Mo with S at the Mo-substitutional site upon withdrawal of Mn during droplet formation stage seems to be quite likely in the present case as the relative strength of the knocked out Mo 3$d_{5/2}$ signal with respect to Mo-S signal in Fig.~\ref{fig2}(k) dies off for more than 5 pulse counts of Mn deposition. Results of Fig.~\ref{fig5} show that the non-radiative  recombination centres for the excitons in MoS$_2$\xspace films increase with  Mn-deposition time (pulse count).  Interestingly, the thermal energy barrier for such recombinations $E_n$ has also been found to increase with pulse count.  It is plausible that Mo-related defects act as the non-radiative centres, each one of which offers a particular value of  $E_n$ around it. However, a cluster of such defects may offer an entirely different barrier height.  We believe that the barrier height $E_n$ increases with the cluster size in these samples. Since the probability of formation of bigger clusters  enhances with the concentration of Mo-defects, $E_n$ rises with the pulse count.

The study  shows that Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace is possible when it is deposited below a certain critical amount on the films using PLD technique. In this deposition regime, thin Mn-coating covering large areas can takes place on MoS$_2$\xspace layers. Large area coverage is believed to enhance the chance of substitutional incorporation of Mn. However, dewetting occurs when Mn-deposition is more than the critical amount, which leads to the emergence of Mn-droplets on MoS$_2$\xspace layers. Due to the reduction of the coverage area, substitutional incorporation of Mn is suppressed in this regime. On the other hand,  disorder in the lattice is observed to be enhanced with the deposition time.  Mn-deposition is also found to influence strongly the excitonic recombination dynamics of 1L-MoS$_2$\xspace by introducing non-radiative recombination centres for the excitons. The density of these centres as well as the barrier height for such recombinations  to happen increase rapidly with the amount of Mn-deposition, especially in the droplet-formation regime.  These non-radiative transitions can be attributed to  either Mo-vacancies or distorted Mo-S bonds. A large number of such defects are believed to be produced during the formation of the Mn-droplets as a result of the withdrawal of Mn ions from the Mo-substitutional sites.
\\
\vspace*{-\baselineskip}

\noindent\textbf{Acknowledgment:}
We acknowledge the financial support by Department of Science and Technology (DST) of Government of India under the grant number: CRG/2018/001343. We would like to thank Industrial Research and Consultancy Centre (IRCC), Sophisticated Analytical Instrument Facility (SAIF), and Centre of Excellence in Nanoelectronics (CEN) of IIT Bombay, for providing various experimental facilities.

\bibliographystyle{unsrt}

\section{Introduction}
Monolayer (1L) MoS$_2$\xspace, a two dimensional (2D) transition metal dichalcogenide (TMD), has become one of the most studied materials in the last decade for its unique electronic and optical properties\cite{mak1}, among which spin-valley coupling\cite{mak2,VP,VP1} and direct band gap at the K-points of the Brillouin zone\cite{mak1,bs} are the most important ones. Because of its direct band gap of $\sim$1.9 eV and large exciton binding energy, 1L-MoS$_2$\xspace offers high luminescence yield in the visible range\cite{mak1,bs,VP1}. Atomically thin size and flexible nature make the material a potential candidate for 3D integrated and flexible electronics of the future \cite{3D_flex}. Furthermore, there are multiple theoretical proposals that predict ferromagnetic behaviour even above room temperature in 1L-MoS$_2$\xspace, when Mo is substituted by certain transition metal ions. Such a two dimensional dilute magnetic semiconductor (2D-DMS) has potential for injecting spin polarized excitons and carriers for future 2D spintronic and valleytronic devices\cite{magnet1,magnet2,magnet3,VP2}. Achieving controllable doping is one of the important requirements for any technology application of the material. However, doping of 1L-MoS$_2$\xspace is always onerous because of the strong bonding between Mo and S, which makes it almost inert for any substitutional reaction to take place\cite{XPS1,bs}. This has become even more challenging as the unambiguous detection of doping is extremely difficult in an atomically thin layer. Even though, there are efforts by various groups to incorporate different elements in 1L-MoS$_2$\xspace using varied  means, such as chemical doping\cite{ch1,ch2,ch3},  physisorption\cite{therm,ad2,ad3,ad4}, metal evaporation\cite{evap1,raman2} and ion irradiation\cite{ion1,raman1}, a method for stable and sustainable doping of the layer is yet to be established.  Previously, Zhang \textit{et al.} have reported up to 2 atomic percentage of  Manganese (Mn) doping in 1L-MoS$_2$\xspace by vapor phase deposition technique\cite{XPS2}. Mn incorporation has also been tried using Mn ion irradiation technique\cite{raman1}, which has been found to result in a large amount of structural defects in the lattice. More studies are needed to understand the mechanism of incorporation of Mn and other metals in 1L-MoS$_2$\xspace. It is also important  to investigate how the incorporation can influence the electrical, optical and magnetic properties of the layer.

Here, we present a detailed study of manganese doping in 1L-MoS$_2$\xspace through pulsed laser deposition (PLD) technique and its influence on the optical properties of the films, which are grown on sapphire substrates by microcavity based chemical vapor deposition (CVD) technique\cite{pkm,tailoring}. Our investigation using high resolution X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), Raman and PL spectroscopy reveals two clearly distinct regimes of Mn-incorporation in 1L-MoS$_2$\xspace. When the deposition amount is below a certain critical value, Mn islands covering large areas are found to be formed on MoS$_2$\xspace layers and substitution of Mo ions by Mn are observed. As the deposition amount goes beyond the critical value, dewetting leads to the formation of Mn-droplets on MoS$_2$\xspace layers. In this regime, substitutional incorporation of Mn is decreased, while the disorder in the lattice increases with the deposition time.  PL study shows that Mn-deposition strongly influence the excitonic recombination dynamics by opening up non-radiative recombination channels. Both the density and the barrier height for such recombinations are found to increase  with the amount of Mn-deposited.  These non-radiative transitions are attributed to certain Mo related defects.

Strictly monolayer (1L) MoS$_2$\xspace films were grown on \textit{c}-sapphire substrates by the microcavity based CVD route where molybdenum trioxide (MoO$_3$) and sulfur (S) powders were used as precursors and argon (Ar) was used as the carrier gas.  More details about the growth can be found elsewhere\cite{pkm,tailoring}. Manganese (Mn) was deposited on these layers using pulsed laser deposition (PLD) technique where a polycrystalline Mn pellet was used as target. The pellet was prepared from 99.99\% pure Mn powder by arc melting technique in Ar (5N pure) atmosphere. Substrate (1L-MoS$_2$\xspace/sapphire) temperature was maintained at 300$$^{\small{\text{o}}}$$C during the deposition. Distance between the Mn-target and the sample was kept at 5 cm for all depositions. A KrF excimer laser of wavelength 248 nm and pulse width of 25 ns was used for ablating the target. Deposition amount was varied by varying the laser pulse count. While the pulse frequency and the energy density of the laser were maintained at 1 Hz and 1.4 J/cm$^2$, respectively, for all cases. Base pressure inside the PLD chamber was kept at $\sim$ 8.0$\times$10$^{-5}$ mbar. During the growth, 5N argon gas was introduced inside the chamber to maintain a pressure of  $\sim$ 8.0$\times$10$^{-2}$ mbar inside the chamber. 1L-MoS$_2$\xspace layers were coated with Mn for different numbers of laser pulse count ranging from 1 to 10. These samples are labeled by alphabets A to D. In  another sample (sample E), Mn deposition was carried out at the base pressure  ($\sim$ 8.0$\times$10$^{-5}$ mbar) for 10 counts of laser pulses.  Mn was also deposited on a bare \textit{c}-plane sapphire substrate (without MoS$_2$\xspace) for five counts of  laser pulses at 8.0$\times$10$^{-2}$ mbar of Ar-pressure, keeping other parameters constant. This sample is named as sample F. One as-grown 1L-MoS$_2$\xspace sample R, was used as the reference standard for all the measurements. Mn deposition parameters for these samples, are listed in Tab.~\ref{tab1}.

X-ray photoelectron spectroscopy on these samples were performed in Axis Supra spectrometer from Kratos Analytical, UK, where monochromatic Al K${_\alpha}$ x-ray (h$\nu$ = 1486.6~eV) was used as excitation source. Energy axis of the recorded spectra, was calibrated with respect to the binding energy of carbon 1$s$ peak at 284.7~eV. Morphology of the samples were examined by atomic force microscopy (AFM) in tapping mode using  Bruker, NanoScope-IV system. Raman spectra on these monolayers were recorded at room temperature in back scattering geometry with 532~nm diode laser using Horiba Jobin Yvon HR800 confocal Raman spectrometer. A closed-cycle helium refrigerator was used to carry out temperature dependent (from 20~K to 300~K) photoluminescence (PL) measurements, where 532~nm diode laser was used as the excitation source and the spectra were recorded by a 0.55~m focal length monochromator attached with a Peltier cooled CCD camera.

\begin{center}
	\begin{table}
		\caption{Mn deposition parameters for different samples. Substrate temperature, distance between target and substrate, frequency and energy density of the laser are maintained at 300$$^{\small{\text{o}}}$$C, 5~cm, 1~HZ and 1.4 J/cm$^2$, respectively for each deposition.}
		\begin{tabular}{c| c |c |c}

			\hline
			{\small Sample} & {\small Pressure} & {\small Ar flow rate} & {\small Laser Pulse count} \\
			& {\small(mbar)} & {\small(sccm)} &  \\
			\hline
			\hline
			~R~& -- & -- & 0  \\

			~A~& {\small 8.0$\times$10$^{-2}$} & 40 & 1  \\

			~B~& {\small 8.0$\times$10$^{-2}$} & 40 & 5  \\

			~C~& {\small 8.0$\times$10$^{-2}$} & 40 & 7  \\

			~D~& {\small 8.0$\times$10$^{-2}$} & 40 & 10  \\

			~E~& {\small 8.0$\times$10$^{-5}$} & 0 & 10  \\

			~F (Sapphire)~& {\small 8.0$\times$10$^{-2}$} & 40 & 5  \\
			\hline
		\end{tabular}
		\label{tab1}
	\end{table}
\end{center}

\begin{figure}
	\includegraphics[scale=1.8]{fig1.pdf}
	\caption{AFM images for samples (a) R, (b) A, (c) B, (d) C and (e) D. Inset of (a) shows the height distribution along a line drawn across the boundary. Inset of (b) shows the line scan profile over a Mn-deposited island on MoS$_2$\xspace film. (f) Average roughness of MoS$_2$\xspace (black) and sapphire (red) regions as functions of pulse count. Filled symbols represent sample E.}
	\label{fig1}
\end{figure}

Fig.~\ref{fig1} shows AFM images for the MoS$_2$\xspace samples coated with Mn for different counts of laser pulses. In all the images, MoS$_2$\xspace layers can clearly be distinguished from the bare sapphire surface from the height contrast. As shown in Fig.~\ref{fig1}(a), average step-height for the reference sample R has been found to be $\sim$ 0.7~nm, which is expected for a monolayer MoS$_2$\xspace\cite{pkm,tailoring}. In fact, it should be noted that average step-height for all samples before Mn deposition is found to be $\sim$ 0.7~nm confirming monolayer nature of MoS$_2$\xspace in all cases. As shown in Fig.~\ref{fig1}(b), deposition of several flat island type structures with average height of $\sim$ 0.9~nm on the MoS$_2$\xspace regions is evident in sample A, where only one laser pulse is used for Mn deposition. In case of sample B, where Mn is deposited for 5 laser pulses, flat island type structures could not be distinguished clearly on MoS$_2$\xspace regions, instead, several spherical structures are visible on both MoS$_2$\xspace and bare sapphire regions as shown in Fig.~\ref{fig1}(c). Density and size of the spherical droplets clearly increase as one moves from sample C [Fig.~\ref{fig1}(d)]  to D [Fig.~\ref{fig1}(e)], meaning as the laser pulse count for Mn deposition increases (see supplementary Fig.~S1\cite{sup}). Noticeably, in these samples, droplets on bare sapphire are bigger than those on MoS$_2$\xspace regions. It is noteworthy that similar type of droplet formation has also been observed in case of sample E (see supplementary Fig.~S1\cite{sup}). Fig.~\ref{fig1}(f) plots the average value of the roughness of MoS$_2$\xspace (black symbols) and sapphire (red symbols) regions as functions of laser pulse count. Evidently, variation of roughness on MoS$_2$\xspace regions shows a clear increase beyond 5 counts of laser pulse. Note that the average roughness on sapphire regions also shows a sudden jump beyond 5 pulses. These are consistent with the sudden change in the morphology observed in the AFM images and seem to suggest that Mn-deposition for sufficiently less number of laser pulses results in a thin Mn-layer. The strain/surface energy developed in such a thin metal film may not be sufficient to change its morphology when the sample is cooled to room temperature after deposition. It is plausible that in case of the samples, where Mn-deposition is performed for the number of pulses more than a certain value, thickness of the deposited Mn-film crosses a critical mark and the strain/surface energy developed in the layer becomes so high that to minimise energy, the film breaks into Mn-droplets when the sample is brought to the room temperature. This point has been discussed again later (in Fig.~\ref{fig5}).

\begin{figure*}[t]
	\includegraphics[scale=1]{fig2.pdf}
	\caption{XPS core level spectra for Mo $3d$ and S $2s$ levels  in case of samples (a) R, (d) A and (g) D; for S $2p$ levels  in case of the samples (b) R, (e) A and (h) D; for Mn $2p$ levels in case of  the samples (c) F,  (f) A and (i) D. (j) Ratio of the contributions from S-Mn and S-Mo bonds obtained by de-convoluting S~$2s$ (black symbols) and S~$2p_{3/2}$ (red symbols) spectra for these samples versus pulse count. (k) Ratio of the contribution from Mo atoms and Mo-S bonds obtained from the Mo~$3d_{5/2}$ profiles versus pulse count. In (j) and (k), filled symbols represent sample E.}
	\label{fig2}
\end{figure*}

In panels (a-i) of Fig.~\ref{fig2}, XPS spectra for Mo~$3d$, S~$2s$, S~$2p$ and Mn~$2p$ levels for different samples are compared. These spectra are deconvoluted with Voigt functions as mentioned in earlier references\cite{XPS,XPS1}. In all panels, red lines represent the fitted profiles. For the reference sample R [panel (a)], S~$2s$ (orange line) peak  appears (at 226.7~eV) along with Mo~$3d$ doublets. Mo~$3d$ related features can be deconvoluted into two doublets. While, the position of the first doublet, which peaks at 229.3~eV (Mo $3d_{5/2}$) and 232.4~eV (Mo $3d_{3/2}$) [blue line], matches quite well with Mo-S bonds in MoS$_2$\xspace\cite{XPS,XPS1}, the second doublet appearing at  232.6~eV (Mo~$3d_{5/2}$), 235.7~eV (Mo~$3d_{3/2}$) [bright green line] can be assigned to the Mo-O bonds of the unreacted MoO$_3$ particles on the surface\cite{XPS}. Panel (b) shows the S $2p$ features for the reference sample. Deconvolution of the feature results in two peaks centred at  162.2~eV and  163.4~eV, (blue line) which are associated with S~$2p_{3/2}$  and S~$2p_{1/2}$  lines, respectively,  of MoS$_2$\xspace\cite{XPS,XPS1}. Fig.~\ref{fig2}(c) shows Mn $2p$ features for sample F, which is a bare sapphire substrate coated with manganese. The feature can be deconvoluted into two doublets. While the doublet (olive line) appearing at  639.2~eV (Mn~$2p_{3/2}$) and  651.0~eV  (Mn~$2p_{1/2}$) can be attributed to Mn-metal\cite{XPS2,XPS3}, position of the higher energy doublet (purple line) appearing at 644.0~eV (Mn~$2p_{3/2}$) and 655.6~eV (Mn~$2p_{1/2}$) matches quite well with those of Mn-O bonds of manganese oxide (MnO$_x$)\cite{XPS2,XPS5}. Quite evidently, Mo~$3d$, S~$2s$ as well as S~$2p$ features appears much broader for sample A as compared to the reference sample as can be seen in  Fig.~\ref{fig2}(d) and (e). This is a clear indication of the change of the bonding states of both Mo and S ions in this sample. Deconvolution of these features results in the formation of one additional  Mo $3d$ doublet peaking at  227.9~eV (Mo~$3d_{5/2}$)  and  231.0~eV (Mo~$3d_{3/2}$) [magenta line in Fig.~\ref{fig2}(d)], one extra S $2s$ peak (at 225.5~eV) [bottle green line in Fig.~\ref{fig2}(d)] and two extra S $2p$  doublets; one peaking at 161.4~eV (S~$2p_{3/2}$)  and 162.6~eV (S~$2p_{1/2}$) [bottle green line in Fig.~\ref{fig2} (e)] and the other, peaking at 159.9~eV (S~$2p_{3/2}$)  and 161.0~eV (S~$2p_{1/2}$)[brown line in Fig.~\ref{fig2}(e)]. Position of the new Mo~$3d$ doublet [magenta line in Fig.~\ref{fig2} (d)] matches well with neutral Mo atoms\cite{XPS1}. Out of the two extra S~$2p$ doublets, the bottle green line of Fig.~\ref{fig2}(e) can be assigned to S-Mn bond formation and the brown line of Fig.~\ref{fig2}(e) can be attributed to surface adsorbed S atoms\cite{XPS1}. The extra S~$2s$ peak at 225.5~eV [bottle green line in Fig.~\ref{fig2}(d)] can also be attributed S-Mn bonds\cite{XPS4,XPS2}. Interestingly for sample D, Mo~$3d$, S~$2s$ and S~$2p$ profiles are similar to those obtained for the reference sample as shown in Fig.~\ref{fig2}(g) and (h). Deconvolution of these profiles results in a much weaker contributions from the S-Mn bonds and neutral Mo atoms as compared to what has been found in case of sample A. These findings point towards the fact that for sample A, Mn atoms partially substitute Mo ions of MoS$_2$\xspace to form Mn-S bonds and the knocked out Mo ions settle as neutral Mo atoms. While, in sample D, such substitution is very weakly present. It should be noted that like in sample A, significant presence of additional Mo and S related features could also be found in sample B, where Mn is deposited for 5 laser pulses. Similar to the case of sample D, very weak sign of Mo-substitution by Mn-ions could be found for sample C and E, where laser pulse counts of 7 and 10, respectively, are used for Mn deposition. Note that Mo~$3d$, S~$2s$ and S~$2p$ spectra for all the Mn deposited MoS$_2$\xspace samples are shown in supplementary figure S2\cite{sup}. Fig.~\ref{fig2}(f) and (i) show the  Mn $2p$ features for sample A and D, respectively. It is interesting to note that for both the cases, features are very similar to what has been found for sample F [Fig.~\ref{fig2}(c)] showing the contributions only from Mn-metal and Mn-O bonds. In fact, no sign of Mn-S bond formation could be found from Mn-spectra for any of the Mn-coated MoS$_2$\xspace samples. Note that in all these samples, a significant portion of the sapphire surface remains uncovered by 1L-MoS$_2$\xspace. Perhaps, the signal due to Mn-S bond from the Mn-coated MoS$_2$\xspace layers is overwhelmed by the signal arising from the Mn-coated sapphire regions for all these samples. The ratio of the contributions from S-Mn and S-Mo bonds can be obtained in two ways; by deconvoluting (i) S~$2s$ and (ii) S~$2p_{3/2}$ spectra. These results are plotted as a function of the laser pulse count in Fig.~\ref{fig2}(j). Evidently, both the profiles show very similar trend. Similarly, the ratio of the contribution from Mo atoms and Mo-S bonds can be achieved from the Mo~$3d_{5/2}$ peaks, which is plotted versus pulse count in Fig.~\ref{fig2}(k). Interestingly, no signature of knocked out Mo atoms could be found in samples deposited with more than 5 pulse counts. It is quite evident from these plots that substantial amount of Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace layers is achieved only for pulse counts of 1 and 5. Incorporation sharply decreases as the pulse count  goes beyond 5.

\begin{figure}
	\includegraphics[scale=1.8]{fig3.pdf}
	\caption{(a) Room temperature Raman spectra for different Mn-deposited samples in the frequency range containing characteristic first order E$^1_{2g}$ and A$_{1g}$ peaks of MoS$_2$\xspace. (b) Deconvolution of the spectrum with Lorentzian functions for sample D. (c) Peak position and (d) FWHM  of the Lorentzians associated with E$^1_{2g}$ (orange) and A$_{1g}$ (blue) peaks versus pulse count. Filled symbols represent sample E. Solid lines in (C) and (d) are guides to the eyes.}
	\label{fig3}
\end{figure}

Fig.~\ref{fig3}(a) compares room temperature micro-Raman spectra for these samples in the frequency range, where characteristic in-plane E$^1_{2g}$ and out-of-plane A$_{1g}$ vibrational modes for the zone-centre phonons are observed at $\sim$ 385 and 405~cm$^{-1}$, respectively. These spectra are deconvoluted with Lorentzian functions as shown in  the Fig.~\ref{fig3}(b) for sample D as an example. Apart from the  dominant E$^1_{2g}$  (orange line) and A$_{1g}$ (blue line) features, several other phonon modes associated with M and K points of the Brillouin zone (BZ) [magenta lines] can also be seen. Note that these non-zone-centre phonons become visible due to the relaxation of the fundamental Raman selection rule (phonon wave vector $q$ $\approx$ 0) in the presence of disorder in the MoS$_2$ lattice (see supplementary Fig.~S3\cite{sup})\cite{raman1,raman2, raman4,raman5}. There are also a couple of other features (brown lines) in the spectrum, which can be identified as the Raman lines arising from the sapphire substrate (see the supplementary Fig.~S3\cite{sup}). Fig.~\ref{fig3}(c) and (d) plot the peak position and full width at half maximum (FWHM) of the Lorentzians associated with E$^1_{2g}$ (orange) and A$_{1g}$ (blue) peaks as functions of the number of laser pulses used for Mn-deposition in these samples. It is evident from these figures that with the increase of pulse count,  A$_{1g}$  peak shows a tendency of shift towards higher wavenumbers.  On the other hand,  position of E$^1_{2g}$ peak does not vary much with pulse count upto 7, beyond which it reduces. It should be noted that electron concentration in the conduction band has a strong influence on the coupling between electrons and A$_{1g}$ phonons in 1L-MoS$_2$\xspace\cite{raman3}. Increase of electron concentration in the layer is found to result in the softening of frequency as well as broadening of the A$_{1g}$ peak, while E$^1_{2g}$ feature hardly shows any change\cite{raman1,raman2,raman3}. Note that the FWHM of the A$_{1g}$ peak is less than that of the reference sample for samples A and B, where the pulse counts are 1 and 5, respectively. This may suggest a reduction of electron concentration in the two MoS$_2$\xspace samples, which can further imply that Mn-ions incorporate in MoS$_2$\xspace when the number of laser pulses used for Mn deposition is sufficiently less.  This observation is also consistent with the results of the XPS study shown in Fig.~\ref{fig2}. Increasing trend of the line-width for both A$_{1g}$ and E$^1_{2g}$ peaks as well as hardening and softening of A$_{1g}$ and E$^1_{2g}$ modes, respectively, for pulse count beyond 7 may indicate a rise of disorder in the layer with the increase of Mn-deposition time.

\begin{figure}
	\includegraphics[scale=1.8]{fig4.pdf}
	\caption{ (a) Lower energy portion of the room temperature Raman spectra for various Mn-deposited samples. (b) Deconvolution of the structure with Lorentzian functions for sample D. Note that the most dominant peak at $\sim$ 227~cm$^{-1}$  has been assigned to longitudinal acoustic phonon mode at \textit{M} point [LA(\textit{M})]. (c) Intensity ratio between LA(\textit{M}) and A$_{1g}$ peaks [I(LA)/I(A$_{1g}$)] as well as LA(\textit{M}) and E$^1_{2g}$ peaks [I(LA)/I(E$^1_{2g}$)] versus pulse count. Filled symbols represent sample E. Solid lines provide the visual guidance.}
	\label{fig4}
\end{figure}

Fig.~\ref{fig4}(a) highlights the lower energy portion of the Raman spectra for these samples, where certain zone edge phonons give rise to broad structure with multiple features due to the presence of disorder\cite{raman1,raman2,raman4,raman5}. Fig.~\ref{fig4}(b)  shows the Lorentzian deconvolution of the structure for sample D, where the most dominant peak at $\sim$ 227~cm$^{-1}$ (violet line) has previously been assigned to longitudinal acoustic phonon mode at \textit{M} point [LA(\textit{M})]. This peak has also been treated as a measure of disorder in 1L-MoS$_2$\xspace films\cite{raman1} . Please see supplementary Fig.~S3\cite{sup} for identification of other features in the structure. It should be noted that the spectra shown in all the above panels are normalized with respect to the intensity of A$_{1g}$ peak.  From Fig.~\ref{fig4}(a) it is evident that all the disorder related peaks  enhance with the laser pulse count suggesting a gradual increase of disorder with the amount of Mn deposition in these samples. In case of sample E, where Mn-deposition is carried out for 10 pulse-counts in vacuum (switching off the Ar supply), peak-shift and peak-width for both E$^1_{2g}$ and A$_{1g}$ peaks are dramatically higher than the other samples suggesting much higher degree of disorder in this sample [see Fig.~\ref{fig3}(c) and (d)].  Intensity ratios between LA(\textit{M}) and A$_{1g}$ peaks [I(LA)/I(A$_{1g}$)] (blue line) as well as  LA(\textit{M}) and E$^1_{2g}$ peaks [I(LA)/I(E$^1_{2g}$)] (orange line), are plotted in Fig.~\ref{fig4}(f). A gradual rise of disorder with increasing Mn-deposition time is quite evident from the figure. This observation is in line with the results shown in Fig.~\ref{fig3}.

\begin{figure}
	\includegraphics[scale=1.8]{fig5.pdf}
	\caption{(a) Photoluminescence (PL) spectra  recorded at 100~K for various Mn-deposited samples. (b) Log of the integrated intensity associated with A-exciton [$I(A)$] versus inverse of temperature. (c) Thermal energy barrier $E_n$ and $r_{n0}/r$ for non-radiative recombination, which are obtained from the fitting as described in the text, versus pulse count. (d) Logarithm of the intensity of the BL feature as a function of the inverse of temperature for different Mn-deposited samples.}
	\label{fig5}
\end{figure}

Fig.~\ref{fig5}(a) compares photoluminescence (PL) spectra  recorded at 100~K for all samples. Each spectrum is featured by a band-edge exciton/trion (A/A$^-$) peak at $\sim$ 1.9~eV and broad luminescence (BL) peak at $\sim$ 1.7~eV\cite{pkm,tailoring,nihit1}. Deconvolution of these PL spectra by a set of six Gaussian peaks [as shown in the supplementary Fig.~S4(a)\cite{sup}] provides a good fit\cite{tailoring}. Band edge feature is consisting of two Gaussians representing free A-excitons (A) and A-trions (A$^-$), while the BL feature can be fitted with two Gaussian functions centred at 1.71~eV ($D_{hi}$) and 1.6~eV ($D_{lo}$).  $D_{hi}$ and $D_{lo}$ transitions are found to be excitonic in nature and attributed to sulfur vacancy related defects\cite{tailoring}. Note that two narrow Gaussians are used to take into account the sharp feature at $\sim$ 1.78~eV that stems from the sapphire substrate. Temperature dependent PL spectra for all samples are recorded from 20 to 300~K and deconvoluted with six Gaussians as described before. Log of the intensity associated with A-exciton [$I(A)$] is plotted as a function of inverse of temperature in fig.~\ref{fig5}(b), for various samples. In case of the reference sample R (orange), $I(A)$ increases monotonically with temperature. This tendency indicates that the excitons are activated from a bound to free sate as temperature increases. For samples deposited with 1, 5 and 7 pulses of Mn, $I(A)$ shows similar increasing tendency upto a certain temperature, beyond which it shows a reduction with rising temperature. This fall  becomes increasingly steeper  as the pulse count increases from 1 to 7. This result suggests that Mn-deposition might be introducing a non-radiative recombination channel for the A-excitons, which need to overcome a thermal barrier to get access to the non-radiative traps. Above finding further implies that the barrier height must be increasing with the amount of Mn deposited. In sample D, where Mn pulse count is 10, $I(A)$ become almost  temperature independent. In case of sample E, $I(A)$ shows similar trend as sample D [see supplementary Fig.~S4(b)\cite{sup}].  We believe that  the thermal energy barrier for the non-radiative recombination is so high for the two samples that one needs to go to higher than room temperature to see the fall of A-exciton intensity. It should be noted that in samples D and E, the density of Mn-deposition induced disorder in MoS$_2$\xspace is expected to be much more than the rest of the samples as discussed in Fig.~\ref{fig4}. At the steady state condition, density of free excitons ($n_A$)  and excitons bound to $D$ type of defect ($n_D$) should satisfy the following equations.
\begin{eqnarray}
	\frac{dn_A}{dt} &=& G - r n_A - r_n n_A -\beta (N_D - n_D) n_A + \alpha n_D = 0 \nonumber \\
	\frac{dn_D}{dt} &=& \beta (N_D - n_D) n_A -  \alpha n_D - r_D n_D  = 0
	\label{Eq1}
\end{eqnarray}
Where, $G$, $r$ and $r_n$ are , respectively, the rate of generation, radiative and non-radiative recombinations for the free excitons. $N_D$ is the concentration of $D$ type of defect traps and $\beta$ ($\alpha$) is the coefficient of transition of an exciton from free (defect bound) state to defect bound (free) state. $r_D$ is the recombination rate of defect bound excitons. At thermal equilibrium, $\alpha$ = $N_D$ $\beta$ $e^{-E_D/k_B T}$. $r_n$ = $r_{n0}$ $e^{-E_n/k_B T}$, where $E_D$ and $E_n$ are activation energy of excitons from defect bound to free state and the height of the thermal energy barrier for non-radiative recombination, respectively. $k_B$ is the Boltzmann constant and $r_{n0}$ is $r_n$ at $T$ = 0~K. Under the assumptions of $N_D$ $\gg$ $n_D$ and $\beta$$N_D$ $\gg$ $r_D$, the expressions for PL intensities associated with free excitons [$I(A)$= $r$ $n_A$] and the bound excitons[$I(D)$= $r_D$ $n_D$] can be obtained from eq.~\ref{Eq1} as,
\begin{eqnarray}
	I(A) = \frac{G}{1+ \frac{r_D}{r} \exp{(\frac{E_D}{k_BT})} + \frac{r_{n0}}{r} \exp{(-\frac{E_n}{k_BT})} }
	\label{Eq2}\\
	I(D) = \frac{\frac{G r_D}{r} \exp{(\frac{E_D}{k_BT})} }{1+ \frac{r_D}{r} \exp{(\frac{E_D}{k_BT})} + \frac{r_{n0}}{r} \exp{(-\frac{E_n}{k_BT})} }
	\label{Eq3}
\end{eqnarray}

Experimentally obtained temperature variations of $I(A)$ for all samples have been fitted by eq.~\ref{Eq2}. The resultant profiles are the red solid curves shown in fig.~\ref{fig5}(b). The fit for the reference sample R returns $E_D$ = 12.4~meV and $E_n$ = 0. $E_D$ value is kept at 12.4 meV while $E_n$ is varied to fit the data for rest of the Mn-deposited samples. Obtained values for $E_n$ and $r_{n0}/r$ are plotted as functions of pulse count in fig.~\ref{fig5}(c). It can be seen that both the quantities gradually increase with the pulse count. Note that the increase of $r_{n0}/r$ could suggest the enhancement of the density of non-radiative centres in the layer. As discussed before, $E_n$ and $r_{n0}/r$ for sample D and E can not be estimated from this study.  Fig.~\ref{fig5}(d) plots logarithm of combined integrated intensities of $D_{hi}$ and $D_{lo}$ [$I(D_{hi})$ + $I(D_{lo})$], i.e., total intensity of BL feature as a function of the inverse of temperature for sample R, A, D and E. Please refer to the supplementary Fig.~S4(c)\cite{sup} for $ln[I(D_{hi})+I(D_{lo})]$ vs $1/T$ plots for samples B and C. It can be seen that like in the reference sample, in all Mn-deposited samples, intensity of this feature increases and shows a tendency of saturation with the reduction of temperature. This result implies that the Mn disorder induced non-radiative recombination channel, which modifies the variation of intensity of A-exciton with temperature in these samples, does not affect the dynamics of the sulfur deficiency bound excitonic features.  It should be noted that our attempts to simulate the data of Fig.~\ref{fig5}(d) using eq.~\ref{Eq3} with the parameters obtained from the fitting of the data using eq.~\ref{Eq2} as shown in Fig.~\ref{fig5}(b) have failed for every sample. This suggests that BL band has nothing to do with the defect D, which binds  A-excitons in our model (eq.~\ref{Eq1}). This further means that BL transition does not directly involve A-excitons. To find the binding energies associated with $D_{hi}$ and $D_{lo}$ , experimental data of [$I(D_{hi})$ + $I(D_{lo})$] vs $T$ are fitted with the equation below as described in our previous report\cite{tailoring}.
\begin{eqnarray}
	I(T)=\frac{I_0}{1+\sum \gamma_i T^{3/2} \exp(-E_{a_i}/k_B T)}
	\label{Eq4}
\end{eqnarray}
where, $E_{a_i}$ and $\gamma_i$ are the binding energy and a measure of capture cross section of $i$-th type of sulfur-vacancy related defects. Terms under  summation in the denominator take into account the contributions from the of two defect bands, $D_{lo}$ and $D_{hi}$. Temperature variation of the intensity of the defect bound BL feature is fitted with eq.~\ref{Eq4} and the resultant profiles are represented by red solid curves in fig.~\ref{fig5}(d).  Values of $E_{D_{hi}}$ ($\sim$ 100~meV) and $E_{D_{lo}}$ ($\sim$ 5~meV) are found to be almost the same for all the samples. This confirms that Mn deposition does not affect the dynamics of sulfur deficiency related BL transition in these samples.

\begin{figure}
	\includegraphics[scale=1]{fig6.pdf}
	\caption{Schematic depiction of Mn-deposition on MoS$_2$\xspace. Island regime: (a) when Mn-deposition time is sufficiently less, flat-top Mn islands with high area coverage are present on top of MoS$_2$\xspace even after post-growth cooling of the sample to room temperature.  Droplet regime: when Mn-deposition time is sufficiently long, (b) Mn-layer with thickness more than the critical thickness is deposited on MoS$_2$\xspace. (c) However, upon post-growth-cooling of the sample to the room temperature, Mn-droplets are formed. Mo-substitution by Mn as well as the formation of Mo-vacancy (\textit{V$_{Mo}$}) and distorted Mo-S bonds are also presented in these panels.}
	\label{fig6}
\end{figure}

From  Fig.~\ref{fig1} and \ref{fig2}  it appears that in the regime when Mn is deposited on MoS$_2$\xspace forming large flat-top islands (laser pulse count upto 5),  Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace is more than in the regime  when Mn-droplets are formed. It is plausible that since the area coverage in the island regime is more than in the droplet regime, chance of substitution of Mo-ions by Mn-ions is expected to be more in the former case as shown schematically in Fig.~\ref{fig6}.  We believe that even in the droplet regime, Mn-layer covering a large area is deposited on MoS$_2$\xspace at the deposition temperature. A good number of Mo-substitutional incorporation of Mn is thus expected at the deposition temperature in this regime as well. However, during the post-growth-cooling stage droplets are formed in this regime. The droplet formation process give rises to either Mo-vacancies (\textit{V$_{Mo}$}) or distorted Mo-S bonds as Mn-ions are withdrawn from the substitutional sites at this stage (see Fig.~\ref{fig6}). This  can explain why the disorder in MoS$_2$\xspace increases with the pulse count, especially in the droplet regime as observed in Fig.~\ref{fig4}.  Note that reconnection of Mo with S at the Mo-substitutional site upon withdrawal of Mn during droplet formation stage seems to be quite likely in the present case as the relative strength of the knocked out Mo 3$d_{5/2}$ signal with respect to Mo-S signal in Fig.~\ref{fig2}(k) dies off for more than 5 pulse counts of Mn deposition. Results of Fig.~\ref{fig5} show that the non-radiative  recombination centres for the excitons in MoS$_2$\xspace films increase with  Mn-deposition time (pulse count).  Interestingly, the thermal energy barrier for such recombinations $E_n$ has also been found to increase with pulse count.  It is plausible that Mo-related defects act as the non-radiative centres, each one of which offers a particular value of  $E_n$ around it. However, a cluster of such defects may offer an entirely different barrier height.  We believe that the barrier height $E_n$ increases with the cluster size in these samples. Since the probability of formation of bigger clusters  enhances with the concentration of Mo-defects, $E_n$ rises with the pulse count.

The study  shows that Mo-substitutional incorporation of Mn in 1L-MoS$_2$\xspace is possible when it is deposited below a certain critical amount on the films using PLD technique. In this deposition regime, thin Mn-coating covering large areas can takes place on MoS$_2$\xspace layers. Large area coverage is believed to enhance the chance of substitutional incorporation of Mn. However, dewetting occurs when Mn-deposition is more than the critical amount, which leads to the emergence of Mn-droplets on MoS$_2$\xspace layers. Due to the reduction of the coverage area, substitutional incorporation of Mn is suppressed in this regime. On the other hand,  disorder in the lattice is observed to be enhanced with the deposition time.  Mn-deposition is also found to influence strongly the excitonic recombination dynamics of 1L-MoS$_2$\xspace by introducing non-radiative recombination centres for the excitons. The density of these centres as well as the barrier height for such recombinations  to happen increase rapidly with the amount of Mn-deposition, especially in the droplet-formation regime.  These non-radiative transitions can be attributed to  either Mo-vacancies or distorted Mo-S bonds. A large number of such defects are believed to be produced during the formation of the Mn-droplets as a result of the withdrawal of Mn ions from the Mo-substitutional sites.
\\
\vspace*{-\baselineskip}

\noindent\textbf{Acknowledgment:}
We acknowledge the financial support by Department of Science and Technology (DST) of Government of India under the grant number: CRG/2018/001343. We would like to thank Industrial Research and Consultancy Centre (IRCC), Sophisticated Analytical Instrument Facility (SAIF), and Centre of Excellence in Nanoelectronics (CEN) of IIT Bombay, for providing various experimental facilities.

\bibliographystyle{unsrt}

redpajama_set_name: RedPajamaArXiv


Bateman's offers a curated selection of Used Bikes. We received these bikes as trade-ins from customers and determined they are suitable for re-sale. Used Bikes are on display at 149 Dupont St.
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text:
Glad to Care
As we are all aware, caring is one of the most essential jobs both in the community, hospitals and care homes. During "Glad To Care" week, we interviewed some of our care staff at Oaklands alongside some of the residents, to see what caring meant to them and why our carers are Glad To Care.
Pat Cousins is one of Oaklands longest serving carers with an amazing 23 years under her belt. We asked Pat how long she had been caring for and in total she has worked a staggering 47 years in this sector. We asked Pat why she liked caring so much, she told us "it's in my nature, I like to know people are being looked after properly" and when we wanted Pat's views on what it takes to be a good carer, she said "I really believe it has to be part of your nature, you really need to care about people and respect them, you want to treat them as you would want to be treated yourself".
Carlos Silva worked 3 years as a carer before taking a break to go back to his previous profession as a chef. However, Carlos missed caring and felt that's where his true calling was. Carlos has now been working with us at Oaklands for two years and enjoys his job. We asked Carlos why he wanted to be a carer, Carlos told us "I like helping people and making them happy, I like to help them achieve that happiness, I want our residents to know they are still important people, to be treated with respect and care". We asked Carlos what he thought made a good carer, Carlos replied "It's in the nature of every single human being to be able to care, some have a greater need to care than others, these are the people that make good carers".
Jo Richardson has been caring for 7½ years, 5 years were spent looking after her dad who had dementia. When we asked Jo why she wanted to carry on with care as a career, Jo felt she wanted to help as many people suffering with dementia as she could and to give a little back. Jo joined the team at Oaklands 2 ½ years ago, she told us "you get so much out of it, to see a smile on a residents face and know you are the one who put it there is very rewarding". And when we asked Jo what makes a good carer, she said "really you need to be a good allrounder, you need to be able to relate to people, to have an understanding of what they are going through and a sense of humour helps, but above all you must want to do it, I love my job".
Nandor Feher started his caring career as an agency worker often visiting Oaklands. Nandor enjoyed working at Oaklands and got to know all the residents so when an opportunity arose for Nandor to join the team on a permanent basis, he was happy to be able to fill that role and has now been with us for 1 ½ years. We asked Nandor why he had chosen care as a career, and he told us that he simply liked helping people. And when we asked Nandor what attributes you needed to be a good carer, he replied "Patience, understanding, the want to help people and a good sense of humour comes in handy".
And on the other side of the spectrum, we wanted to get the view of some of our residents about being in care. One resident told us "I have reached the age where I need a little help in certain areas, the care at Oaklands is very good, the staff are very nice and I enjoy the company".
All our carers do a wonderful job, and each one is a vital part of our team at Oaklands.

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text:
Franz Moritz Graf von Lacy (; , tr. ; 21 October 1725 – 24 November 1801) was a Baltic German-born Austrian military leader, he was the son of Count Peter von Lacy and was a famous Austrian field marshal. He served during the reign of Maria Theresa and was a close friend to Joseph II, Holy Roman Emperor, becoming one of the latter's advisers. Lacy was made a Count of the Holy Roman Empire.

Early life

He was born as the son of Count Peter von Lacy and his second wife, Margareta Filippina von Funcken-Lösern (1685–1759), widowed Countess Frölich. His father was a Russian Field Marshal who belonged to the Irish branch of the de Lacy family and had followed James II into exile. Franz Moritz was born in St Petersburg, and was educated in Germany for a military career, and entered the Austrian service. He served in Italy, Bohemia, Silesia and the Netherlands during the War of the Austrian Succession, was twice wounded, and by the end of the war was a lieutenant-colonel. At the age of twenty-five he became full colonel and chief of an infantry regiment.

Seven Years' War
In 1756 with the opening of the Seven Years' War he was again on active service, and in the first battle (Lobositz) he distinguished himself so much that he was at once promoted major-general. He received his third wound on this occasion and his fourth at the battle of Prague in 1757. Later in 1757 Lacy bore a conspicuous part in the great victory of Breslau, and at Leuthen, where he received his fifth wound, he covered the retreat of the defeated army. Soon after this began his association with Field-Marshal Daun, the new generalissimo of the empress's forces, and these two commanders, powerfully assisted later by the genius of Laudon, made headway against Frederick the Great for the remainder of the war.

Lacy, a lieutenant field-marshal at thirty-two, was made chief of staff (quartermaster-general) to Daun in the newly created general staff. They were cautious to the point of timidity since Leuthen and many other bitter defeats had taught the Austrians to respect Frederick; but, having resolved to wear out the enemy by Fabian methods, they were strong enough to persist to the end. Thus for some years the life of Lacy, as of Daun and Laudon, is the story of the war against Prussia. After Hochkirch (October 15, 1758) Lacy received the Grand Cross of the Order of Maria Theresa. In 1759 both Daun and Lacy fell into disfavour for failing to win victories, and Lacy owed his promotion to Feldzeugmeister only to the fact that Laudon had just received this rank for the brilliant conduct of his detachment at Kunersdorf. He shadowed Frederick the Great's Prussian army during the failed Siege of Dresden in 1760. His responsibilities told heavily on Lacy in the ensuing campaigns, and his capacity for supreme command was doubted even by Daun, who refused to give him the command when he himself was wounded at the Battle of Torgau.

Later career

After the Treaty of Hubertusburg a new sphere of activity was opened, in which Lacy's special gifts had the greatest scope. Maria Theresa having placed her son, Emperor Joseph II, at the head of Austrian military affairs, Lacy was made a field marshal, and given the task of reforming and administering the army (1766). He framed new regulations for each arm, a new code of military law, a good supply system. As the result of his work the Austrian army was more numerous, far better equipped, and cheaper than it had ever been before. Joseph soon became very intimate with his military adviser, but this did not prevent his mother, after she became estranged from the young emperor, from giving Lacy her full confidence. His activities were not confined to the army. He was in sympathy with Joseph's innovations, and was regarded by Maria Theresa as a prime mover in the scheme for the partition of Poland. But his self-imposed work broke down Lacy's health, and in 1773, in spite of the remonstrances of Maria Theresa and of the emperor, he laid down all his offices and went to southern France. On returning he was still unable to resume office, though as an unofficial adviser in political and military matters he was far from idle.

In the brief and uneventful War of the Bavarian Succession, Lacy and Laudon were the chief Austrian commanders against the King of Prussia, and when Joseph II at Maria Theresa's death, became the sovereign of the Austrian dominions as well as emperor, Lacy remained his most trusted friend. More serious than the War of the Bavarian Succession was the Turkish war which presently broke out. Lacy was now old and worn out, and his tenure of command therein was not marked by any greater measure of success than in the case of the other Austrian generals. His active career was at an end, although he continued his effective interest in the affairs of the state and the army throughout the reign of Joseph's successor, Leopold II. His last years were spent in retirement at his castle of Neuwaldegg near Vienna. He is buried in a mausoleum in the castle grounds.

Ancestry
His father was the Irish-born Russian field marshal and general, commander of Saint Petersburg and Veliky Novgorod as well as governor of Riga, Count Pierce Edmond de Lacy (German: Peter Graf von Lacy) (1678–1751).

See also

 Irish military diaspora
 Irish regiments

References

Further reading
 "Franz Moritz von Lacy", Allgemeine Deutsche Biographie (ADB) (in German), Volume 17, Leipzig: Duncker & Humblot, 1883,  pp. 487–499; (accessible online).
 .
 Edward de Lacy-Bellingari (1928), The roll of the house of Lacy: pedigrees, military memoirs and synoptical history of the ancient and illustrious family of De Lacy, from the earliest times, in all its branches, to the present day. Full notices on allied families and a memoir of the Brownes (Camas). Baltimore, MD: Waverly Press, viii & 409 pp. (accessible online).
 Stefan Michael Newerkla (2019), "Die irischen Reichsgrafen von Browne-Camus in russischen und österreichischen Diensten. Vom Vertrag von Limerick (1691) bis zum Tod ihres Hausfreunds Ludwig van Beethoven (1827)" [= The Irish counts of Browne-Camus in Russian and Austrian service. From the Treaty of Limerick (1691) to the death of their friend Ludwig van Beethoven (1827)]. In: Lazar Fleishman – Stefan Michael Newerkla – Michael Wachtel (eds.): Скрещения судеб. Literarische und kulturelle Beziehungen zwischen Russland und dem Westen. A Festschrift for Fedor B. Poljakov (= Stanford Slavic Studies, Volume 49). Berlin: Peter Lang, pp. 43–68.
 Stefan Michael Newerkla (2020), "Das irische Geschlecht O'Reilly und seine Verbindungen zu Österreich und Russland" [The Irish O'Reilly family and their connections to Austria and Russia]. In: Diachronie – Ethnos – Tradition: Studien zur slawischen Sprachgeschichte [Diachrony – Ethnos – Tradition: Studies in Slavic Language History]. Eds. Jasmina Grković-Major, Natalia B. Korina, Stefan M. Newerkla, Fedor B. Poljakov, Svetlana M. Tolstaja. Brno: Tribun EU, pp. 259–279 (accessible online), here pp. 259–261.

1725 births
1801 deaths
Military personnel from Saint Petersburg
People from Sankt-Peterburgsky Uyezd
Russian nobility
Emigrants from the Russian Empire to the Austrian Empire
Counts of Austria
Austrian people of Irish descent
Field marshals of Austria
Generals of the Holy Roman Empire
18th-century Austrian people
Austrian military personnel of the War of the Austrian Succession
Austrian military personnel of the Seven Years' War
Austrian military personnel of the War of the Bavarian Succession
Imperial military personnel of the War of Bavarian Succession
Recipients of the Order of St. George of the Third Degree
People of the Silesian Wars
Grand Crosses of the Military Order of Maria Theresa

redpajama_set_name: RedPajamaWikipedia


text:
Forgiveness can make American great again
The political season may not be when people want to hear this. The same could be said about the immediate aftermath of the Orlando massacre. But we all could benefit from being more forgiving.
In fact, this may be the exact time when we need this message the most. Few things destroy health quicker than stress, and the world seems engulfed in stress right now. Having a forgiving heart would wipe that out.
"It's almost entirely erased — it's statistically zero," Loren Toussaint, associate professor of psychology at Iowa's Luther College, recently told Time magazine. He was referring to the almost magical effect a forgiving heart has on stress.
Toussaint is the author of a new study published in the Journal of Health Psychology. Researchers asked 148 young people to fill out surveys assessing how much stress they felt in life, as well as the state of their physical and mental health. Also, the survey asked them about their tendency to forgive people who had wronged them.
The more stress people reported feeling, the worse their physical and mental health. But if people were forgiving — of themselves as well as others — stress had virtually no effect on health.
"If you don't have forgiving tendencies, you feel the raw effects of stress in an unmitigated way," Toussaint told Time. "You don't have a buffer against that stress."
Examples abound of people who have made this work in their lives, despite being victims of awful crimes. More than a decade ago, I first wrote about Vicky Ruvolo, a woman who was driving down a Long Island highway on a frosty night when a teenager in an oncoming car thought it would be fun to toss a stolen 20-pound turkey at her windshield.
Ruvolo nearly died. Her cheeks, jaw and an eye socket were fractured. Her esophagus was damaged and she suffered some brain damage. Yet when the case came to court several months later, she convinced the judge to go easy on her assailant. Then she embraced the young man in court and forgave him.
When I spoke with her years later, Ruvolo had a remarkable sense of humor and seemed genuinely happy. She wrote a book about her experiences. Perhaps most significantly, her assailant, Ryan Cushing, wrote a tribute in it that said, in part, "Your ability to forgive me has had a profound effect on me. It has already made a positive change in my life."
We may read that story and marvel, not only at Ruvolo's capacity to set aside her own pain, but at how her decision affected everything that followed, both in her life and in Cushing's. And yet many people refuse to let go of slights far less significant in their own lives.
A few years ago I interviewed Martin Doblemeier, a documentary filmmaker who was exploring this theme. Among other things, he went around New York City asking people whether a part of the former Twin Towers site, destroyed in the 9/11 attacks, should be dedicated to a "forgiveness garden."
The results were as illuminating as they were varied and emotional. They also remain relevant today. There is little evidence to suggest the health affects of forgiveness are confined only to the perceived harms done to us personally. How we feel about national events, politicians, crimes, groups or circumstances beyond out control are important, as well.
Many people, it appears, get hung up on what forgiveness means. For Ruvolo, it never meant excusing what had been done to her. It wasn't a sign of weakness or capitulation. It was empowering. Doblemeier's reporting found that for many it was a liberating gift people could give themselves.
He and others also have found forgiveness is a constant process, not a one-time event. It can be a difficult attitude to maintain.
But consider this: Time also published a graphic on what stress can do to your health. Among other things, it raises blood pressure, increases cholesterol levels, heightens the risk of type-2 diabetes, causes ulcers and leads to headaches, shoulder and neck pain.
Given all that, a nationwide emphasis on learning to forgive may be just what we all need to regain a sense of sanity.


text:
A university student has the usual dreams and aspirations. It turns out to be a pipe dream and so she murders her own flesh and blood. The usual media heads seek an answer in order to sell copy. A psychologist is writing a book about the character and incident. Her research for the subject matter turn the psychologist inward and things come out of the shadows that perhaps are better left buried deep in the shadows of the psyche.
Director: Yukihiko Tsutsumi
Actors: Hôshi Ishida, Itsuji Itao, Keiko Kitagawa, Kyôko Yoshine, Saki Takaoka, Shô Kiyohara, Tomoya Nakamura, Yoshino Kimura, Yôsuke Kubozuka
Keywords:First Love
In this fifth installment, Death is just as omnipresent as ever, and is unleashed after one man's premonition saves a group of coworkers from a terrifying suspension bridge collapse. But…
Genre: Fantasy, Horror, Mystery, Thriller
The Great Muppet Caper
Kermit and Fozzie are newspaper reporters sent to London to interview Lady Holiday, a wealthy fashion designer whose priceless diamond necklace is stolen. Kermit meets and falls in love with…
Genre: Adventure, Comedy, Crime, Family, Music, Musical, Mystery
Scooby-Doo! Abracadabra-Doo
The gang goes on a trip to check on Velma's younger sister, Madelyn. She's been studying stage magic at the Whirlen Merlin Magic Academy, where apparently there have been sightings…
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Two strangers become dangerously close after witnessing a deadly accident. On a beautiful cloudless day a young couple celebrate their reunion with a picnic. Joe has planned a postcard-perfect afternoon…
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A married couple, struggling to have a child, hires a young woman to be their surrogate, but soon discovers she has a bizarre and deadly agenda.
Trailer: First Love


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Q: How to create excel file with multiple sheets from DataSet using C# How to create excel file with multiple sheets  from DataSet using C#?
I have successfully created an excel file with single sheet. But I am not able to do that for multiple sheets.

A: Here is a simple C# class that programatically creates an Excel WorkBook and adds two sheets to it, and then populates both sheets.  Finally, it saves the WorkBook to a file in the application root directory so that you can inspect the results...
public class Tyburn1
{
    object missing = Type.Missing;
    public Tyburn1()
    {
        Excel.Application oXL = new Excel.Application();
        oXL.Visible = false;
        Excel.Workbook oWB = oXL.Workbooks.Add(missing);
        Excel.Worksheet oSheet = oWB.ActiveSheet as Excel.Worksheet;
        oSheet.Name = "The first sheet";
        oSheet.Cells[1, 1] = "Something";
        Excel.Worksheet oSheet2 = oWB.Sheets.Add(missing, missing, 1, missing)
                        as Excel.Worksheet;
        oSheet2.Name = "The second sheet";
        oSheet2.Cells[1, 1] = "Something completely different";
        string fileName = Path.GetDirectoryName(Assembly.GetExecutingAssembly().Location)
                                + "\\SoSample.xlsx";
        oWB.SaveAs(fileName, Excel.XlFileFormat.xlOpenXMLWorkbook,
            missing, missing, missing, missing,
            Excel.XlSaveAsAccessMode.xlNoChange,
            missing, missing, missing, missing, missing);
        oWB.Close(missing, missing, missing);
        oXL.UserControl = true;
        oXL.Quit();
    }
}

To do this, you would need to add a reference to Microsoft.Office.Interop.Excel to your project (you may have done this already since you are creating one sheet).
The statement that adds the second sheet is...
Excel.Worksheet oSheet2 = oWB.Sheets.Add(missing, missing, 1, missing)
                            as Excel.Worksheet;

the '1' argument specifies a single sheet, and it can be more if you want to add several sheets at once.
Final note:  the statement  oXL.Visible = false; tells Excel to start in silent mode.

A: var groupedSheetList = UserData
    .GroupBy (u => u.date)
    .Select (grp => grp.ToList ())
    .ToList ();

You can try this
using (var package = new ExcelPackage ())
    {
        foreach (var item in groupedSheetList) {
            var workSheet = package.Workbook.Worksheets.Add (item[0].date);
    }
}

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Home/Basketball/College Basketball/Watch Stetson vs Florida Live College Basketball Game Info, TV channel, Sunday Dec 4, 2022
Watch Stetson vs Florida Live College Basketball Game Info, TV channel, Sunday Dec 4, 2022
This Sunday, December 4th, 2022, the Stetson Hatters and the Florida Gators will go head to head in a highly anticipated basketball matchup. The game will be aired on the SEC Network and will start at 2 p.m. ET. The Hatters are hoping to build on their 7-1 record and upset the Gators who are 8-0 so far this season. A win by the Hatters would be an incredible upset, as they have been overlooked in the NCAA rankings this season.
College Basketball Game Info
Who: Stetson vs Florida
When: Sunday at 2 p.m. ET, December 4, 2022
Where: Exactech Arena at Stephen C. O'Connell Center, Gainesville, FL
TV channel: SEC Network
This Sunday, Dec 4th, 2022, the Stetson Hatters will face off against the Florida Gators in a college basketball game on the SEC Network. Both teams' lineups feature several talented players, each with unique skill sets. The Hatters come into the game having won five of their last six games and looking to extend their win streak against the Gators. Florida has been an SEC powerhouse in recent years, winning the conference three times over the last five years.
The Florida Gators and the Stetson Hatters men's basketball teams will face off this Sunday, December 4th in a game sure to excite. The game can be viewed on the SEC Network, which is available to viewers nationwide. This game marks the second matchup between the two teams this season, with the Gators taking the first meeting in Gainesville. With Stetson's impressive season so far, Gators fans are expecting another close game that could go either way.
Sunday, December 4, 2022 marks the date of the highly anticipated Stetson vs Florida basketball game. It will air on SEC Network at 2 p.m. ET and will be the first game of the season. This matchup will pit two of the top teams in the conference against each other. Both teams will be looking to start their season off with a win and make a statement early in the season.
This Sunday, December 4, 2022, Stetson will face off against Florida College in a basketball matchup that is sure to be exciting. The game will be broadcast on the SEC Network, so fans across the country can join in on the fun. This is a matchup between two teams who have an intense rivalry and each team is looking to get the win. Both teams are loaded with talent and have tremendous records for the season, making for an intense and competitive game. More Basketball Info
Watch UCF vs Samford Live College Basketball Game Info, TV channel, Sunday Dec 4, 2022
Watch Cal Poly vs. Portland State Live NCAAB Game Info, TV Channel, Saturday, Dec 10, 2022
Watch Gonzaga Bulldogs vs Xavier Musketeers Live College Basketball info, Start Time, November 27, 2022
The Greatest Houston Cougars Basketball Players Of All Time
[UpDaTeS] UCLA Bruins vs Bellarmine Knights Live College Basketball info, Start Time, November 27, 2022


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11 September 2021 / SF News / Matt Charnock
Dixie and Caldor Fires Each Reach 60% Containment
Saturday morning, a milestone in fighting two of Northern California's most pressing wildfires — the Dixie Fire and Caldor Fire — was achieved: successfully containing at least 60% of each blaze.
This wildfire season has proven especially ominous in how it's mirrored last year's history-making one. We've already seen the largest single-source wildfire in California's history (which could soon reach "gigafire" status as it nears a million acres in size), as well as record-shattering drought levels. And we've even seen the return of orange skies to some parts of Northern California.
Today, more than 13,490 personnel remain assigned to 13 active large wildfires. To date, more than 2.2 million acres have burned statewide. Get the latest on these incidents at: https://t.co/jBh7Rim5k6 pic.twitter.com/RBms4LafDq
— CAL FIRE (@CAL_FIRE) September 11, 2021
But in a situationally positive turn of events, CAL Fire reported earlier today that the Dixie and Caldor Fires are each over 60% contained. More specifically: The Dixie Fire is currently 62%, while the Caldor Fire is about 60% contained.
This news is also a good time to dive into some firefighting lingo.
"Contained fires" are those that have been fully enclosed by control lines; these are around any related spot fires. Though a contained fire might still burn, it's not likely to spread. So when we hear that fire is 60% contained, but not designated as a contained fire, it means that control lines — constructed or natural barriers used to halt wildfires from spreading — surround at least 60% of the blaze.
"Controlled fires," however, are not just contained — but have been extinguished outside control lines and along their perimmerts, as well. Any hot spots which could pose a threat along these lines have been cooled down. Having a controlled fire is the end goal for any fire suppression strategy.
The Caldor Fire still appears to be slowing down... something that can't be said about the Dixie Fire. Before the weekend, on-site firefighters contended with strong winds which whipped up fifty-foot flames and made for even worse smoke conditions. It's scenes like that which leave little doubt that the Dixie Fire could well become the Calfirona's largest wildfire to date in the coming week (or weeks).
If California's increasingly more severe wildfires have taught us anything as of late, it's that we need all understand that the climate crisis is, in fact, real — and will only get worse, should we remain complacent. And we should all, too, have our emergency kits well-stocked and at the ready.
Photo: A chimney remains a home that was destroyed by the Caldor Fire on September 1, 2021 in Twin Bridges, California. The Caldor Fire has burned over 200,000 acres, destroyed hundreds of structures and is currently 20 percent contained. (Photo by Justin Sullivan/Getty Images)
Yep... San Francisco's Cable Cars Are Still Out of Service
Because of an "electrical issue" earlier in the week, San Francisco's iconic cable cars are currently corralled inside SFMTA's "Barn" — with still no operational date in sight.
Bay Area Sports
San Francisco 49ers Season Preview: What To Expect This Year
With the NFL season having officially begun last night with the Buccaneers defeating the Cowboys 31-29 in a thrilling opener, the San Francisco 49ers get set to kick-off their own campaign Sunday in Detroit against the Lions. And the excitement running through the franchise is palpable.